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Effect of steel Alloying Elements 材料基础-钢材合金元素图文简易教材 Descriptive approach

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Page 1 of 180.

Chapter One: Theory of Strengthening

The Mechanism of Strengthening.

Elastic and Plastic Deformation. Slide show on plastic deformation: Mechanical Testing

Understanding Dislocation. The Nature of Dislocation.

"Atomistic" of edge and screw dislocation motions. Origins of Strength from Dislocation Dynamics. Deformation and strengthening Mechanisms.

Crystal Defects. Point Defects Linear Defects Planar Defects. Bulk Defects. Pinning Overview.

Methods of Strengthening. Works or Strain Hardening.

Strain Hardening and Annealing Dislocation and plastic deformation

Grain Boundary Strengthening. Effect of Microstructure Refinement on the Strength and Toughness of low alloy

martensitic steel Effect of microstructure and notches on the fracture toughness of medium CS The influence of austenite grain size on hot ductility of steel

Dispersion Strengthening. Mechanical Alloying.

Solid Solution Strengthening. Precipitation Hardening.

Chapter Two: Effects of Alloying Elements on Iron Carbon Alloy

Introduction Contributing factors to strength of steel. Ways of improving strength of alloy steels. Steels for special applications EF420 Lecture 7: Cast Irons EF420 Lecture 9: Ferrous Alloys EF420 Lecture 12:

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Alloying

Effects of alloying elements. Relative effects in combinations Characteristics of alloying element Summary table. Effects of alloying elements. Effects of Alloying Elements on the Microstructures and Mechanical on ductile cast irons

Carbon. Manganese. Nickel. Chromium. Nickel & Chromium Molybdenum. Vanadium Tungsten Silicon Copper Phosphorous Sulphur

Lead Hydrogen Selenium Tantalum Tin Calcium Cerium Nitrogen Boron Aluminium Zirconium Niobium

Titanium Cobalt Tellurium

Chapter Three: Articles on alloying elements in steel

Carbon contents, steel classification and alloy steels. Carbon steels Stainless steels

Strength in steels. Alloying and its effects on the critical temperature and tensile strength. Control of HSLA properties. Influence of alloying elements on steel microstructures Martensite in austenitic stainless steel welds. Alloying effects on martensite, pearlite and bainite formation. Steel alloys Carbon steels to austenitic steels. Selection of age-hardenable superalloys. Metallurgy of Molybdenum in alloy steel and iron. Maraging steels. Periodic table and atomic radii. 17-4PH precipitates hardening martensitic stainless steel. Thermal Processing of Metal

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Thermal Processing of Metals More reading. Steel Making. 冶金视频 Additional reading:

Slag inclusion formation during solidification of Steel alloys and in cast iron. Clean Steels. Fundamental steps in thermomechanical processing of steels. Strengthening of Alloy Steel by High Temperature Thermomechanical Treatment. Ion-nitriding of Maraging steel. Microstructures and properties of low-alloy fire resistant steel. Some fundamental steps in thermomechanical processing of steel. Strengthening of Alloy Steel by High Temperature Thermomechanical Treatment. High Performance Copper-Precipitation-Hardened Steel. Schaeffler diagram. Comprehensive Characterization of Ageing Behavior in M250 Maraging Steel using Multi-NDE

Techniques. Effect of Aging on Microstructure and Mechanical Property of 1900 MPa Grade Maraging Stainless

Steel. Effect of aging conditions on Maraging Steels. The influence of retained austenite on precipitation hardening of maraging steel. 18Ni 马氏体时效钢时效机理的研究. 1900MPa 级超高强度不锈钢的研制. 固溶处理对新型马氏体时效不锈钢力学性能的影响. 两种马氏体时效不锈钢的相变和力学性能. 马氏体时效不锈钢的发展. 马氏体时效不锈钢时效析出相及位向关系的研究. 无钴马氏体时效钢的研究现状. 无钴马氏体时效钢的研究与应用. 新型 CrCoNiMo 系马氏体时效不锈钢的强韧化机理. 预先冷轧变形对马氏体时效钢强化的影响. Control of High Strength Low Alloy Steel HSLA Properties by alloying Brittle fracture 0f A105 flanges – EN10204 3.1B and shortfall

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Introduction to Dislocations, Fourth Edition by Derek Hull, D J Bacon http://www.arab-eng.org/vb/t93373.html The Science and Engineering of Materials by Donald R. Askeland http://university.arabsbook.com/forum23/thread24766.html

Greek Alphabets.

1. Α α alpha 2. Β β beta 3. Γ γ gamma 4. Δ δ delta 5. Ε ε epsilon 6. Ζ ζ zeta 7. Η η eta 8. Θ θ theta 9. Ι ι iota 10. Κ κ kappa 11. Λ λ lamda 12. Μ μ mu 13. Ν ν nu 14. Ξ ξ xi 15. Ο ο omicron 16. Π π pi 17. Ρ ρ rho 18. Σ σ ς sigma 19. Τ τ tau

http://www.jxcad.com.cn/?u=74032


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Chapter One: The Mechanism of strengthening.

Elastic/Plastic Deformation

When a sufficient load is applied to a metal or other structural material, it will cause the material to change shape. This change in shape is called deformation. A temporary shape change that is self-reversing after the force is removed, so that the object returns to its original shape, is called elastic deformation. In other words, elastic deformation is a change in shape of a material at low stress that is recoverable after the stress is removed. This type of deformation involves stretching of the bonds, but the atoms do not slip past each other.

When the stress is sufficient to permanently deform the metal, it is called plastic deformation. As discussed in the section on crystal defects, plastic deformation involves the breaking of a limited number of atomic bonds by the movement of dislocations. Recall that the force needed to break the bonds of all the atoms in a crystal plane all at once is very great. However, the movement of dislocations allows atoms in crystal planes to slip past one another at a much lower stress levels. Since the energy required to move is lowest along the densest planes of atoms, dislocations have a preferred direction of travel within a grain of the material. This results in slip that occurs along parallel planes within the grain. These parallel slip planes group together to form slip bands, which can be seen with an optical microscope. A slip band appears as a single line under the microscope, but it is in fact made up of closely spaced parallel slip planes as shown in the image.

Plastic deformation occurs when large numbers of dislocations move and multiply so as to result in macroscopic deformation. In other words, it is the movement of dislocations in the material which allows for deformation. If we want to enhance a material's mechanical properties (i.e. increase the yield and tensile strength), we simply need to introduce a mechanism which prohibits the mobility of these dislocations. Whatever the mechanism may be, (work hardening, grain size reduction, etc) they all hinder dislocation motion and render the material stronger than previously. The stress required to cause dislocation motion is orders of magnitude lower than the theoretical stress required to shift an entire plane of atoms, so this mode of stress relief is energetically favorable. Hence, the hardness and strength (both yield and tensile) critically depend on the ease with which dislocations move. Pinning points, or locations in the crystal that oppose the motion of dislocations, can be introduced into the


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lattice to reduce dislocation mobility, thereby increasing mechanical strength. Dislocations may be pinned due to stress field interactions with other dislocations and solute particles, or physical barriers from grain boundaries and second phase precipitates. There are several strengthening mechanisms for metals, however the key concept to remember about strengthening of metallic materials is that it is all about preventing dislocation motion and propagation; you are making it energetically unfavorable for the dislocation to move or propagate. For a material that has been strengthened, by some processing method, the amount of force required to start irreversible (plastic) deformation is greater than it was for the original material. In amorphous materials such as polymers, amorphous ceramics (glass), and amorphous metals, the lack of long range order leads to yielding via mechanisms such as brittle fracture, crazing, and shear band formation. In these systems, strengthening mechanisms do not involve dislocations, but rather consist of modifications to the chemical structure and processing of the constituent material. Unfortunately, strength of materials cannot infinitely increase. Each of the mechanisms elaborated below involves some trade off by which other material properties are compromised in the process of strengthening.

Steel can be strengthened by several basic mechanisms, the most important of which are:

1. Work hardening or strain hardening. 2. Solid solution strengthening by interstitial atoms. 3. Solid solution strengthening by substitutional atoms. 4. Refinement of grain size. 5. Effects of heat treatment on microstructures. 6. Precipitation strengthening. 7. Grain boundary strengthening. 8. Dispersion strengthening, including lamellar and random dispersed structures.

The most distinctive aspect of strengthening of iron and steel is the role of the interstitial solutes carbon and nitrogen. These elements also play a vital part in interacting with dislocations, and in combining preferentially with some of the metallic alloying elements used in steels.


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Plastic Deformation: Mechanical Testing:

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Understanding dislocations Introduction to dislocations http://www.msm.cam.ac.uk/doitpoms/tlplib/dislocations/printall.php

The Nature of Dislocations

Plastid deformation is a measure of material strength, plastic deformation is irreversible. Therefore, the configuration of the atoms must be changed during plastic deformation, for otherwise they would return to their original position on unloading. If we consider shearing a single crystal as an example, it can be deformed plastically by sliding whole layers of atoms against each other as shown in figure above for this sliding to happen, the bonds between the atoms have to be stretched elastically until they can switch to the next atom. The stress required for this process can be estimated and is of the order of one fifth of the shear modulus of the crystal. The yield strength predicted this way for metallic single crystals is thus between 1GPa and 25GPa. If we measure the strength of single crystals of pure metals, the values found are several orders of magnitudes below this theoretical value and even lie below that of engineering alloys. Typical values are in the range of a few mega Pascal. As single crystals always contain lattice defects, one possible explanation could be that these are responsible for the reduced strength. If, however, the number of defects is reduced further, for instance by a heat treatment, the yield strength becomes even smaller. Only an absolutely perfect single crystal without any defects would possess a yield strength agreeing with the theoretical prediction. This can only be nearly realized in so-called whiskers, which, however, are extremely small. The reason for this spectacular failure of the theoretical prediction is that plastic deformation does not occur by sliding of complete layers of atoms. Instead, it proceeds by a mechanism that is based on a special type of lattice defect, the dislocations. To understand plastic deformation of metals thus requires an understanding of dislocations.


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Dislocation densities Dislocation is a lattice imperfection in a crystal structure which exerts a profound effect on a structure sensitive properties such as strength, hardness, ductility and toughness. There are two types, edge and screw or combination of both, all of which are characterized by a Burgers vector which represents the amount and direction of slip when the dislocation moves. Click on the web links provided to read further. Interaction of Dislocations

Transmission Electron Micrograph of Dislocations

Transmission Electron Micrograph of Dislocations

More reading:

http://en.wikipedia.org/wiki/Dislocation http://www-sgrgroup.materials.ox.ac.uk/lectures/microplasticity.html


Page 10 of 180.

"Atomistics" of edge dislocation motion and Asymmetry of screw and edge dislocation motion in Mo Origins of Strength from Dislocation Dynamics

http://www.scribd.com/doc/7833526/12142005-Workshop

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Deformation and strengthening mechanisms

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Crystal Defects

A perfect crystal, with every atom of the same type in the correct position, does not exist. All crystals have some defects. Defects contribute to the mechanical properties of metals. In fact, using the term “defect” is sort of a misnomer since these features are commonly intentionally used to manipulate the mechanical properties of a material. Adding alloying elements to a metal is one way of introducing a crystal defect. Nevertheless, the term “defect” will be used, just keep in mind that crystalline defects are not always bad. There are basic classes of crystal defects: point defects, which are places where an atom is missing or irregularly placed in the lattice structure.

Point defects include lattice vacancies, self-interstitial atoms, substitution impurity atoms, and interstitial impurity atoms

linear defects, which are groups of atoms in irregular positions. Linear defects are commonly called dislocations.

planar defects, which are interfaces between homogeneous regions of the material. Planar defects include grain boundaries, stacking faults and external surfaces.

It is important to note at this point that plastic deformation in a material occurs due to the movement of dislocations (linear defects). Millions of dislocations result for plastic forming operations such as rolling and extruding. It is also important to note that any defect in the regular lattice structure disrupts the motion of dislocation, which makes slip or plastic deformation more difficult. These defects not only include the point and planer defects mentioned above, and also other dislocations. Dislocation movement produces additional dislocations, and when dislocations run into each other it often impedes movement of the dislocations. This drives up the force needed to move the dislocation or, in other words, strengthens the material. Each of the crystal defects will be discussed in more detail in the following pages.

Point Defects

Point defects are where an atom is missing or is in an irregular place in the lattice structure. Point defects include self interstitial atoms, interstitial impurity atoms, substitutional atoms and vacancies. A self interstitial atom is an extra atom that has crowded its way into an interstitial void in the crystal structure. Self interstitial atoms occur only in low concentrations in metals because they distort and highly stress the tightly packed lattice structure.

A substitutional impurity atom is an atom of a different type than the bulk atoms, which has replaced one of the bulk atoms in the lattice. Substitutional impurity atoms are usually close in size (within approximately 15%) to the bulk atom. An example of substitutional impurity atoms is the zinc atoms in brass. In brass, zinc atoms with a radius of 0.133 nm have replaced some of the copper atoms, which have a radius of 0.128 nm. Interstitial impurity atoms are much smaller than the atoms in the bulk matrix. Interstitial impurity atoms fit into the open space between the bulk atoms of the lattice structure. An example of interstitial impurity atoms is the carbon atoms that are added to iron to make steel. Carbon atoms, with a radius of 0.071 nm, fit nicely in the open spaces between the larger (0.124 nm) iron atoms.


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Vacancies are empty spaces where an atom should be, but is missing. They are common, especially at high temperatures when atoms are frequently and randomly change their positions leaving behind empty lattice sites. In most cases diffusion (mass transport by atomic motion) can only occur because of vacancies.

Linear Defects - Dislocations

Dislocations are another type of defect in crystals. Dislocations are areas were the atoms are out of position in the crystal structure. Dislocations are generated and move when a stress is applied. The motion of dislocations allows slip – plastic deformation to occur.

Before the discovery of the dislocation by Taylor, Orowan and Polyani in 1934, no one could figure out how the plastic deformation properties of a metal could be greatly changed by solely by forming (without changing the chemical composition). This became even bigger mystery when in the early 1900’s scientists estimated that metals undergo plastic deformation at forces much smaller than the theoretical strength of the forces that are holding the metal atoms together. Many metallurgists remained skeptical of the dislocation theory until the development of the transmission electron microscope in the late 1950’s. The TEM allowed experimental evidence to be collected that showed that the strength and ductility of metals are controlled by dislocations. There are two basic types of dislocations, the edge dislocation and the screw dislocation. Actually, edge and screw dislocations are just extreme forms of the possible dislocation structures that can occur. Most dislocations are probably a hybrid of the edge and screw forms but this discussion will be limited to these two types.


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Edge Dislocations The edge defect can be easily visualized as an extra half-plane of atoms in a lattice. The dislocation is called a line defect because the locus of defective points produced in the lattice by the dislocation lie along a line. This line runs along the top of the extra half-plane. The inter-atomic bonds are significantly distorted only in the immediate vicinity of the dislocation line.

Understanding the movement of a dislocation is key to understanding why dislocations allow deformation to occur at much lower stress than in a perfect crystal. Dislocation motion is analogous to movement of a caterpillar. The caterpillar would have to exert a large force to move its entire body at once. Instead it moves the rear portion of its body forward a small amount and creates a hump. The hump then moves forward and eventual moves all of the body forward by a small amount.

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As shown in the set of images above, the dislocation moves similarly moves a small amount at a time. The dislocation in the top half of the crystal is slipping one plane at a time as it moves to the right from its position in image (a) to its position in image (b) and finally image (c). In the process of slipping one plane at a time the dislocation propagates across the crystal. The movement of the dislocation across the plane eventually causes the top half of the crystal to move with respect to the bottom half. However, only a small fraction of the bonds are broken at any given time. Movement in this manner requires a much smaller force than breaking all the bonds across the middle plane simultaneously.

Screw Dislocations

There is a second basic type of dislocation, called screw dislocation. The screw dislocation is slightly more difficult to visualize. The motion of a screw dislocation is also a result of shear stress, but the defect line movement is perpendicular to direction of the stress and the atom displacement, rather than parallel. To visualize a screw dislocation, imagine a block of metal with a shear stress applied across one end so that the metal begins to rip. This is shown in the upper right image. The lower right image shows the plane of atoms just above the rip. The atoms represented by the blue circles have not yet moved from their original position. The atoms represented by the red circles have moved to their new position in the lattice and have reestablished metallic bonds. The atoms represented by the green circles are in the process of moving. It can be seen that only a portion of the bonds are broke at any given time. As was the case with the edge dislocation, movement in this manner requires a much smaller force than breaking all the bonds across the middle plane simultaneously.

If the shear force is increased, the atoms will continue to slip to the right. A row of the green atoms will find there way back into a proper spot in the lattice (and become red) and a row of the blue atoms will slip out of position (and become green). In this way, the screw dislocation will move upward in the image, which is perpendicular to direction of the stress. Recall that the edge dislocation moves parallel to the direction of stress. As shown in the image below, the net plastic deformation of both edge and screw dislocations are the same, however.

The dislocations move along the densest planes of atoms in a material, because the stress needed to move the dislocation increases with the spacing between the planes. FCC and BCC metals have many dense planes, so dislocations move relatively easy and these materials have high ductility. Metals are strengthened by making it more difficult for dislocations to move. This may involve the introduction of obstacles, such as interstitial atoms or grain boundaries, to “pin” the dislocations. Also, as a material plastically deforms, more


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dislocations are produced and they will get into each others way and impede movement. This is why strain or work hardening occurs.

In ionically bonded materials, the ion must move past an area with a repulsive charge in order to get to the next location of the same charge. Therefore, slip is difficult and the materials are brittle. Likewise, the low density packing of covalent materials makes them generally more brittle than metals.

Planar Defects

Stacking Faults and Twin Boundaries A disruption of the long-range stacking sequence can produce two other common types of crystal defects: 1) a stacking fault and 2) a twin region. A change in the stacking sequence over a few atomic spacing produces a stacking fault whereas a change over many atomic spacing produces a twin region. A stacking fault is a one or two layer interruption in the stacking sequence of atom planes. Stacking faults occur in a number of crystal structures, but it is easiest to see how they occur in close packed structures. For example, it is know from a previous discussion that face centered cubic (fcc) structures differ from hexagonal close packed (hcp) structures only in their stacking order. For hcp and fcc structures, the first two layers arrange themselves identically, and are said to have an AB arrangement. If the third layer is placed so that its atoms are directly above those of the first (A) layer, the stacking will be ABA. This is the hcp structure, and it continues ABABABAB. However it is possible for the third layer atoms to arrange themselves so that they are in line with the first layer to produce an ABC arrangement which is that of the fcc structure. So, if the hcp structure is going along as ABABAB and suddenly switches to ABABABCABAB, there is a stacking fault present.


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Alternately, in the fcc arrangement the pattern is ABCABCABC. A stacking fault in an fcc structure would appear as one of the C planes missing. In other words the pattern would become ABCABCAB_ABCABC.

If a stacking fault does not corrects itself immediately but continues over some number of atomic spacing, it will produce a second stacking fault that is the twin of the first one. For example if the stacking pattern is ABABABAB but switches to ABCABCABC for a period of time before switching back to ABABABAB, a pair of twin stacking faults is produced. The red region in the stacking sequence that goes ABCABCACBACBABCABC is the twin plane and the twin boundaries are the A planes on each end of the highlighted region.

Grain Boundaries in Polycrystals Another type of planer defect is the grain boundary. Up to this point, the discussion has focused on defects of single crystals. However, solids generally consist of a number of crystallites or grains. Grains can range in size from nanometers to millimeters across and their orientations are usually rotated with respect to neighboring grains. Where one grain stops and another begins is know as a grain boundary. Grain boundaries limit the lengths and motions of dislocations. Therefore, having smaller grains (more grain boundary surface area) strengthens a material. The size of the grains can be controlled by the cooling rate when the material cast or heat treated. Generally, rapid cooling produces smaller grains whereas slow cooling result in larger grains. For more information, refer to the discussion on solidification.

Bulk Defects

Bulk defects occur on a much bigger scale than the rest of the crystal defects discussed in this section. However, for the sake of completeness and since they do affect the movement of dislocations, a few of the more common bulk defects will be mentioned. Voids are regions where there are a large number of atoms missing from the lattice. The image to the right is a void in a piece of metal The image was acquired using a Scanning Electron Microscope (SEM). Voids can occur for a number of reasons. When voids occur due to air bubbles becoming trapped when a material solidifies, it is commonly called porosity. When a void occurs due to the shrinkage of a material as it solidifies, it is called cavitation.

Another type of bulk defect occurs when impurity atoms cluster together to form small regions of a different phase. The term ‘phase’ refers to that region of space occupied by a physically homogeneous material. These regions are often called precipitates. Phases and precipitates will be discussed in more detail latter.


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Pinning points - overview.

In a crystalline material, a dislocation is capable of traveling throughout the lattice when relatively small stresses are applied. This movement of dislocations results in the material plastically deforming. Pinning points in the material act to halt a dislocation's movement, requiring a greater amount of force to be applied to overcome the barrier. This results in an overall strengthening of materials.

Types of pinning points

Point defects

Point defects (as well as stationary dislocations, jogs, and kinks) present in a material create stress fields within a material that disallow traveling dislocations to come into direct contact. Much like two particles of the same electric charge feel a repulsion to one another when brought together, the dislocation is pushed away from the already present stress field.

Alloying elements

The introduction of atom1 into a crystal of atom2 creates a pinning point for multiple reasons. An alloying atom is by nature a point defect, thus it must create a stress field when placed into a foreign crystallographic position, which could block the passage of a dislocation. However, it is possible that the allowing material is approximately the same size as the atom that is replaced, and thus its presence would not stress the lattice (as occurs in cobalt alloyed nickel). The different atom would, though, have a different elastic modulus, which would create a different terrain for the moving dislocation. A higher modulus would look like an energy barrier, and a lower like an energy trough – both of which would stop its movement.

Second phase precipitates

The precipitation of a second phase within the lattice of a material creates physical blockades through which a dislocation cannot pass. The result is that the dislocation must bend (which requires greater energy, or a greater stress to be applied) around the precipitates, which inevitably leaves residual dislocation loops encircling the second phase material and shortens the original dislocation. This is a schematic shows how a dislocation interacts with solid phase precipitates. The dislocation moves from left to right in each frame.


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Dislocation slip mechanism

Grain boundaries

Dislocations require proper lattice ordering to move through a material. At grain boundaries, there is a lattice mismatch, and every atom that lies on the boundary is uncoordinated. This stops dislocations that encounter the boundary from moving.

Dislocations in motion


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Methods of strengthening

Following are brief description s on the methods of strengthening.

Work or Strain hardening

The reason for strain hardening is that the dislocation density increases with plastic deformation (cold work) due to multiplication. The average distance between dislocations then decreases and dislocations start blocking the motion of dislocations.

The primary species responsible for work hardening are dislocations. Dislocations interact with each other by generating stress fields in the material. The interaction between the stress fields of dislocations can impede dislocation motion by repulsive or attractive interactions. Additionally, if two dislocations cross, dislocation line entanglement occurs, causing the formation of a jog which opposes dislocation motion. These entanglements and jogs act as pinning points, which oppose dislocation motion. As both of these processes are more likely to occur when more dislocations are present, there is a correlation between dislocation density and yield strength,

Where G is the shear modulus, b is the Burgers vector, and is the dislocation density. Increasing the dislocation density increases the yield strength which results in a higher shear stress required to move the dislocations. This process is easily observed while working a material. Theoretically, the strength of a material with no dislocations will be extremely high (τ=G/2) because plastic deformation would require the breaking of many bonds simultaneously. However, at moderate dislocation density values of around 107-109 dislocations/m2, the material will exhibit a significantly lower mechanical strength. Analogously, it is easier to move a rubber rug across a surface by propagating a small ripple through it than by dragging the whole rug. At dislocation densities of 1014 dislocations/m2 or higher, the strength of the material becomes high once again. It should be noted that the dislocation density can't be infinitely high because then the material would lose its crystalline structure. Work hardening is an important strengthening process in steel, particularly in obtaining high strength levels in rod and wire, both in plain carbon and alloy steels. For example, the tensile strength of a 0.05% C steel subjected to 95% reduction in area by wire drawing, is raised by no less than 550 MPa while higher carbon steels are strengthened by up to twice this amount. Indeed, without the addition of special alloying elements, plain carbon steels can be raised to strength levels above 1500 MPa simply by the phenomenon of work hardening.

Basic work on the deformation of iron has largely concentrated on the other end of the strength spectrum, namely pure single crystals and polycrystals subjected to small controlled deformations. The diversity of slip


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planes leads to rather irregular wavy slip bands in deformed crystals, as the dislocations can readily move from one type of plane to another by cross slip, provided they share a common slip direction.

The yield stress of iron single crystals are very sensitive to both temperature and strain rate and a similar dependence has been found for less pure polycrystalline iron. Therefore, the temperature sensitivity cannot be attributed to interstitial impurities. It is explained by the effect of temperature on the stress needed to move free dislocations in the crystal, the Peierls-Nabarro stress.

Case Study 1: The effect of cyclic torsion on the dislocation structure of drawn mild steel

1. Introduction Cold forming of metals usually causes their work hardening. The magnitude of this hardening depends on the area reduction, on the temperature and strain rate associated with the processing, and on the way the strain is imposed on the metal. Keeping all other variables constant, the work hardening of a metal submitted to a sequential straining under varying directions or of different natures is different from that resulting from monotonic straining. Changes in the way the material is deformed can alter the hardening rates and even cause strain softening of the metal1-16. Recent research results17-19 show that cyclic straining influences in various ways the mechanical behavior of annealed and drawn metal bars. Annealed Aluminum submitted to cyclic torsion displays higher flow stresses than the annealed material. On the other hand, cyclic torsion softens previously drawn Aluminum. Cyclic torsion also softens steel bars previously drawn in one or two passes and hardens the initially annealed material. Experimental results indicate that the stress-strain curve and the work hardening coefficient (n) of steel drawn in two passes and submitted to cyclic torsion are similar to those for the material submitted to only one drawing pass. This is similar to the case of the Aluminum alloy 6063, where the cyclic torsion after two drawing passes eliminates the hardening associated with the second drawing pass. It is also observed for both materials that their Ultimate Tensile Strength (UTS) tends to remain unaltered by cyclic torsion, in the case of initially annealed material, whereas their Yield Strength (YS) is considerably increased by cyclic torsion. The YS and UTS of both previously drawn materials are decreased by cyclic torsion, with the exception of the YS of Aluminum drawn in a single pass. The decrease in these properties is more pronounced after two drawing passes than after a single drawing pass. Finally, cyclic torsion increases the Tensile Elongation to Fracture of drawn material and decreases this property for initially annealed material. The present research analyzes the relationship between the mechanical effects described above and the dislocation structures in Low Carbon steel. 2. Materials and Experimental Methods The material was an AISI 1010 steel with the following chemical composition: 0.12%C, 0.47%Mn, 0.07%Si, 0.003%Sn, 0.01%Mo, 0.016%P e 0.013%S, received as cylindrical bars 6.4 mm in diameter. The bars were initially annealed and some of them were drawn in one or two passes. A fraction of these bars were then submitted to cyclic torsion. The effects of the strain path were analyzed by Transmission Electron Microscopy (TEM). Annealing was performed under vacuum, at 850 °C for 2400 seconds, leading to an average hardness of 122.6 HV. Drawing was performed in a hydraulic draw bench, using Tungsten Carbide dies with semi-angle of 8?and abundant lubrication with a Molybdenum Disulfide paste. Different dies were employed, guaranteeing a


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fixed reduction of area of 20% in each pass. Cyclic torsion (11.2% plastic strain per cycle, total of 10 cycles) was performed in an especially adapted lathe, where the chuck was manually actuated. All experiments were performed at room temperature, at a strain rate of 0.002 / s5. All deformed samples were stored at temperatures below 0 °C in order to avoid static strain aging effects. TEM was performed in a JEOL-JEM microscope, operated at 200 kV. Analyses were performed in 3 mm samples taken from the cross-section of the bars. Sample preparation involved initial mechanical polishing, followed by electrolytic thinning with a perchloric acid and ethanol solution at room temperature. 3. Results and Discussion The dislocation structures of the annealed and of the drawn material (in one or two passes) are shown in Figures 1 and The effect of cyclic torsion on the dislocation structure of drawn mild steel2 respectively. The dislocation density is quite low for the annealed material, whereas the deformed material displays a much higher dislocation density. Drawn material shows an aligned cell structure, with irregular cell sizes and cell wall thickness. Dense dislocation networks can be observed inside the cells. Higher drawing strains lead to a smaller cell size, as expected.

Figure 1.

Figure 2a.

Figure 2b. TEM of annealed and drawn steel

(8% and 20% per pass) a: 1 pass b: 2 passes

.

The dislocation structure of the material submitted only to cyclic torsion is broadly similar to that resulting from drawing (see Figure 3), but the cell size is higher and the tendency to cell alignment is less pronounced than in


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drawing. It is important to realize that the total strain caused by cyclic torsion is much higher than in drawing, but leads to essentially similar dislocation structures. The analysis of Figure 4 indicates that the dislocation structure of the material after one drawing pass is altered by subsequent cyclic torsion. There is an increase in the cell size, a decrease in their alignment and in the dislocation density within the cells, and an overall evolution of alignment in only one direction to a "checkerboard" appearance, which is typical of the development of two sets of aligned cells, corresponding to the two directions of twisting. This is also the situation for the material initially annealed, drawn in 2 passes and cyclically twisted (Figure 5). Under these circumstances, the dislocation cells tend to be larger and the dislocation density inside the cells lower than for one drawing pass followed by cyclic torsion. Cyclic torsion promotes dynamic recovery of the material, involving the annihilation of cell walls and the decrease of dislocation density inside the cells. This is similar to results from the analysis of the Bauschinger effect20. where such dislocation annihilation stems from dislocation movements in opposing directions.


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Considering the widely established relationship between the material flow stress G, and the corresponding dislocation density r20:

one should expect higher flow stresses in the drawn or cyclic twisted material than in the annealed material. On the other hand, the recovery promoted by cyclic torsion of the previously drawn material should lead to their softening and consequent lower flow stresses. 4. Conclusions Drawing of low carbon steel leads to the formation of an aligned dislocation structure, displaying irregular cell sizes and cell wall thickness as well as dense networks of dislocations inside the cells. Cyclic torsion promotes the softening of material previously strained by drawing. This softening is associated with a restructuring of the previous dislocation arrangement, involving an increase in cell size, a decrease in the dislocation density inside the cells and a "checkerboard" dislocation wall structure.

Strain Hardening and Annealing Slide Show: Full download

http://www.slideshare.net/charliechong/askeland-phule-notes-ch12-printable-presentation

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http://materialteknologi.hig.no/Materiallare/D.R.%20Askeland/

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Dislocation and Plastic Deformation:

http://www.slideshare.net/charliechong/chapter05-1-presentation

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Page 26 of 180.

Grain Boundary Strengthening

In grain boundary strengthening the grain boundaries act as pinning points impeding further dislocation propagation. Since the lattice structure of adjacent grains differs in orientation, it requires more energy for a dislocation to change directions and move into the adjacent grain. The grain boundary is also much more disordered than inside the grain, which also prevents the dislocations from moving in a continuous slip plane. Impeding this dislocation movement will hinder the onset of plasticity and hence increase the yield strength of the material. Grain boundaries act as an impediment to dislocation motion for the following two reasons:

• Dislocation must change its direction of motion due to the differing orientation of grains. • Discontinuity of slip planes from grain 1 to grain

Under an applied stress, existing dislocations and dislocations generated by Frank-Read Sources will move through a crystalline lattice until encountering a grain boundary, where the large atomic mismatch between different grains creates a repulsive stress field to oppose continued dislocation motion. As more dislocations propagate to this boundary, dislocation 'pile up' occurs as a cluster of dislocations are unable to move past the boundary. As dislocations generate repulsive stress fields, each successive dislocation will apply a repulsive force to the dislocation incident with the grain boundary. These repulsive forces act as a driving force to reduce the energetic barrier for diffusion across the boundary, such that additional pile up causes dislocation diffusion across the grain boundary, allowing further deformation in the material. Decreasing grain boundary size decreases the amount of possible pile up at the boundary, increasing the amount of applied stress necessary to move a dislocation across a grain boundary. The higher the applied stress to move the dislocation, the higher the yield strength. Thus, there is then an inverse relationship between grain boundary size and yield strength, as demonstrated by the Hall-Petch equation. A lower number of dislocations per grain results in a lower dislocation 'pressure' building up at grain boundaries. This makes it more difficult for dislocations to move into adjacent grains. This relationship can be mathematically described as follows:

,

Where k is a constant, d is the average grain diameter and σy,0 is the original yield stress. However, when there is a large direction change in the orientation of the two adjacent grains, the dislocation may not necessarily move from one grain to the other but instead create a new source of dislocation in the adjacent grain. The theory remains the same that more grain boundaries create more opposition to dislocation movement and in turn strengthens the material.


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Figure 1: Hall-Petch Strengthening is limited by the size of dislocations. Once the grain size reaches about 10 nm, grain boundaries start to slide. Obviously, there is a limit to this mode of strengthening, as infinitely strong materials do not exist. Grain boundary sizes can range from about 100 μm (large grains) to 1 μm (small grains). Lower than this, the size of dislocations begins to approach the size of the grains. At a grain size of about 10 nm, only one or two dislocations can fit inside of a grain (see Figure 1 above). This scheme prohibits dislocation pile-up and never results in grain boundary diffusion. The lattice resolves the applied stress by grain boundary sliding, resulting in a decrease in the material's yield strength; A phenomenon known as grain-boundary sliding. To understand the mechanism of grain boundary strengthening one must understand the nature of dislocation-dislocation interactions. Dislocations create a stress field around them given by:


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,

Where G is the material's shear modulus, and b is the Burgers vector. If the dislocations are in the right alignment with respect to each other, the local stress fields they create will repel each other. This helps dislocation movement along grains and across grain boundaries. Hence, the more dislocations are present in a grain, the greater the stress field felt by a dislocation near a grain boundary:

This is a schematic roughly illustrating the concept of dislocation pile up and how it effects the strength of the material. A material with larger grain size is able to have more dislocation to pile up leading to a bigger driving force for dislocations to move from one grain to another. Thus you will have to apply less force to move a dislocation from a larger than from a smaller grain, leading materials with smaller grains to exhibit higher yield stress. In a polycrystalline metal, grain size has a tremendous influence on the mechanical properties. Because grains usually have varying crystallographic orientations, grain boundaries arise. While an undergoing deformation, slip motion will take place.

The refinement of the grain size of ferrite provides one of the most important strengthening routes in the heal treatment of steels. The grain size effect on the yield stress can therefore be explained by assuming that a dislocation source operates within a crystal causing dislocations to move and eventually to pile up at the grain boundary. The pile-up causes a stress to be generated in the adjacent grain, which, when it reaches a critical value, operates a new source in that grain. In this way, the yielding process is propagated from grain to grain. The grain size determines the distance dislocations have to move to form grain boundary pile-ups, and thus the number of dislocations involved. With large grain sizes, the pile-ups will contain larger numbers of dislocations, which will in turn cause higher stress concentrations in neighboring grains. In practical terms, the finer the grain size, the higher the resulting yield stress and, as a result, in modern steel working much attention is paid to the final ferrite grain size. While a coarse grain size of d-1/2 = 2, i.e. d = 0.25


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mm, gives a yield stress in mild steels of around 100 MPa, grain refinement to d-1/2 = 20. i.e. d = 0.0025 mm, raises the yield stress to over 500 MPa, so that achieving grain sizes in the range 2-10 μm is extremely worthwhile.

The effect of grain size on yield strength, 　y, is given by the Hall Petch equation for structural steels:

where 　o is the lattice resistance, i.e. friction stress which opposes dislocation motion k is a constant, sometimes called the dislocation locking term d is the ferrite grain size

14 µm

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More reading: Effect of Microstructure on Mechanical Properties of High Strength Steel Weld Metals

http://www.msm.cam.ac.uk/phase-trans/2005/Enda/Thesis.pdf

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Effects of Tempering on the Microstructure and Mechanical Properties of Low Carbon, Low Alloy Martensitic Steel

http://www.insinet.net/jasr/2007/1719-1723.pdf

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Effect of Microstructural Variation on Weld Metal Cold Cracking of HSLA-100 Steel

http://www.journalarchive.jst.go.jp/jnlpdf.php?cdjournal=isijinternational1989&cdvol=43&noissue=5&startpage=706&lang=en&from=jnlabstract


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Effect of Microstructure Refinement on the Strength and Toughness of low alloy martensitic steel

http://www.scribd.com/doc/7824864/Effect-of-Micro-Structure-Refinement-on-the-Strength-and-Toughness-of-Low-Alloy-Martensitic-Steel

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Page 31 of 180.

Effect of microstructure and notches on the fracture toughness of medium carbon steel

http://www.scribd.com/doc/7825015/Effect-of-Micro-Structure-and-Notches-on-the-Fracture-Toughness-of-Medium-Carbon-Steel

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Page 32 of 180.

The influence of austenite grain size on hot ductility of steel

http://www.scribd.com/doc/7825091/The-Influence-of-Austenite-Grain-Size-on-Hot-Ductility-of-Steel

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Page 33 of 180.

Dispersion Strengthening

(Mechanical Alloying) Dispersion strengthening is about the interaction between dislocations and finely dispersed particles in the metal matrix. Traditionally, dispersion particles are obtained as precipitates from phase transformations during traditional metallurgical processes (melting, solidification, and heat treating). This is the case of many high-strength commercial alloys. However, maximum service temperatures are limited by the limited thermal stability of particles or precipitates obtained from thermal processes.

A: Lattice distortion due to the presence of coherent precipitate: B: Noncohereny precipitates produce no lattice distortion

Particles, which may not be metallurgical compatible with a given metal, can be introduced in a metal matrix by violently deforming mixtures of different powders. Such technique is called mechanical alloying . An example is high-energy ball milling. Using this technique, oxides and other highly stable chemical species can be introduced uniformly into the metal microstructure. A further refinement of this technique allows oxidation and other chemical reactions to take place during the mixing process, a technique that allegedly promotes the formation of ultrafine dispersoids as well as allowing control of particle composition and distribution. After mixing, the powder mixture can then be consolidated and compacted using a number of powder metallurgy techniques to produce a solid with a very fine grain structure.

Mechanical alloying methods permit the manufacturing of metallic alloys with a number of interesting properties. Some examples of materials include nanocrystalline and amorphous materials, metastable phases, and alloys with extended solubility limits. Mechanical alloying was originally developed as a means of raising the maximum service temperatures of nickel-based superalloys for aircraft gas-turbine applications. As turbine operating temperatures rose, so too did the demand for materials with increased high-temperature strength and oxidation resistance. Mechanical alloying avoids many of the problems associated with conventional melting and solidification processes, and is now used to make a variety of oxide-dispersion-strengthened (ODS) Fe-Cr, Ni-Cr, and Ni-Cr-gamma superalloys for turbine-engine industrial applications and aluminum alloys for aircraft structural components. Mechanical alloying may play a key role in the development of future aerospace systems by enabling the production of even higher performance materials that are difficult or impossible to make by other methods.

The effectiveness of dispersion strengthening largely depends on (a) mechanical and geometrical characteristics of the dispersoids, such as hardness, continuity, size, and shape, and (b) their density and


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distribution (dispersion factor) in the metal matrix. By controlling process parameters, such as characteristics of the raw powder materials, pressing pressures, atmospheres, cycle time, and temperatures, the microstructure (e.g., grain size, dispersion size, volume fraction, and distribution) can be tailored for specific performance requirements. New research initiatives in this field include the synthesis of refractory carbide nanoparticles and assessment of their use as dispersion strengthening agents, as well as alternative methods to inoculate the metal matrix.

In all steels there is normally more than one phase present, and indeed it is often the case that several phases can be recognized in the microstructure. The matrix, which is usually ferrite (bcc structure) or austenite (fcc structure) strengthened by grain size refinement and by solid solution additions, is further strengthened, often to a considerable degree, by controlling the dispersions of the other phases in the microstructure. The commonest other phases are carbides formed as a result of the low solubility of carbon in α-iron. In plain carbon steels this carbide is normally Fe3C (cementite), which can occur, in a wide range of structures from coarse lamellar form (pearlite), to fine rod or spheroidal precipitates (tempered steels). In alloy steels, the same range of structures is encountered, except that in many cases iron carbide is replaced by other carbides, which are thermodynamically more stable. Other dispersed phases which are encountered include nitrides, intermetallic compounds, and, in cast irons, graphite. Most dispersions lead to strengthening, but often they can have adverse effects on ductility and toughness. In fine dispersions (where ideally small spheres are randomly dispersed in a matrix) are well-defined relationships between the yield stress, or initial flow stress, and the parameters of the dispersion. These relationships can be applied to simple dispersions sometimes found in steels, particularly after tempering, when, in plain carbon steels, the structure consists of spheroidal cementite particles in a ferritic matrix. However, they can provide approximations in less ideal cases, which are the rule in steels, where the dispersions vary over the range from fine rods and plates to irregular polyhedral. Perhaps the most familiar structure in steels is that of the eutectoid pearlite, usually a lamellar mixture of ferrite and cementite. This can be considered as an extreme form of dispersion of one phase in another, and undoubtedly provides a useful contribution to strengthening.


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General behaviour of the powder particles during mechanical alloying processing


Page 36 of 180.

http://www.scribd.com/doc/7823923/Mechanical-Alloy

Case Study:

Effect of mechanical alloying and Ti addition on solution and ageing treatment of an AA7050 aluminium alloy http://www.scielo.br/scielo.php?pid=S1516-14392007000200017&script=sci_arttext Mechanical Alloys & Milling http://www.scribd.com/doc/3629131/Mechanical-alloying-and-milling-Suryanarayana

http://www.scribd.com/doc/3629131/Mechanical-alloying-and-milling-Suryanarayana


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Solid solution strengthening

This is a schematic illustrating how the lattice is strained by the addition of substitutional and interstitial solute. Notice the strain in the lattice that the solute atoms cause. The interstitial solute could be carbon in iron for example. The carbon atoms in the interstitial sites of the lattice create a stress field that impedes dislocation movement.

http://en.wikipedia.org/wiki/Strengthening_mechanisms_of_materials

A grain boundary in a 2D lattice is the interface between two regions of crystalline order. Each region or 'grain' has a different orientation with respect to some arbitrary axis perpendicular to the plane of the lattice.

Grain boundaries A vacancy is a point defect that arises when an atom is 'missing' from the ideal crystal structure.

A vacancy


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A solute atom in a crystal structure is an atomic species that is different from the majority of atoms that form the structure. Solute atoms of similar size to those in the host lattice may substitute for host atoms - these are known as substitutional solutes. Solute atoms that are much smaller than the host atoms may exist within normally empty regions (interstices) in the host lattice, where they are called interstitial solutes.

Substitutional and interstitial solutes. Note that some distortion of the host lattice occurs around the solutes. A dislocation in a 2D close-packed plane can be described as an extra 'half-row' of atoms in the structure. Dislocations can be characterised by the Burgers vector which gives information about the orientation and magnitude of the dislocation.

Dislocation


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Metal Crystal Structure Atomic Radius (nm)

Aluminum FCC 0.1431

Cadmium HCP 0.1490

Chromium BCC 0.1249

Cobalt HCP 0.1253

Copper FCC 0.1278

Gold FCC 0.1442

Iron (Alpha) BCC 0.1241

Lead FCC 0.1750

Magnesium HCP 0.1599

Molybdenum BCC 0.1363

Nickel FCC 0.1246

Platinum FCC 0.1387

Silver FCC 0.1445

Tantalum BCC 0.1430

Titanium (Alpha) HCP 0.1445

Tungsten BCC 0.1371

Zinc HCP 0.1332

A nanometer (nm) equals 10-9 meter or 10 Angstrom units.


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Solid solution strengthening by interstitials

The formation of interstitial atmospheres at dislocations requires diffusion of the solute. As both carbon and nitrogen diffuse much more rapidly in iron than substitutional solutes, it is not surprising that strain ageing can take place readily in the range from 20°C to 150°C. Consequently the atmosphere condenses to form rows of interstitial atoms along the cores of the dislocations. These arise because the temperature is high enough to allow interstitial atoms to diffuse during deformation, and to form atmospheres around dislocations generated throughout the stress-strain curve. Steels tested under these conditions also show low ductility, due partly to the high dislocation density and partly to the nucleation of carbide particles on the dislocations where the carbon concentration is high. The phenomenon is often referred to as blue brittleness, blue being the interference color of the steel surface when oxidized in this temperature range. The break away of dislocations from their carbon atmospheres as a cause of the sharp yield point became a controversial aspect of the theory because it was found that the provision of free dislocations, for example, by scratching the surface of a specimen, did not eliminate the sharp yield point. An alternative theory was developed which assumed that, once condensed carbon atmospheres are formed in iron, the dislocations remain locked, and the yield phenomena arise from the generation and movement of newly formed dislocations. To summarize, the occurrence of a sharp yield point depends on the occurrence of a sudden increase in the number of mobile dislocations. However, the precise mechanism by which this takes place will depend on the effectiveness of the locking of the pre-existing dislocations. If the pinning is weak, then the yield point can arise as a result of unpinning. However, if the dislocations are strongly locked, either by interstitial atmospheres or precipitates, the yield point will result from the rapid generation of new dislocations. Under conditions of dynamic strain ageing, where atmospheres of carbon atoms form continuously on newly-generated dislocations, it would be expected that a higher density of dislocations would be needed to complete the deformation, if it is assumed that most dislocations which attract carbon atmospheres are permanently locked in position.

Strengthening at high interstitial concentrations

Austenite can take into solid solution up to 10% carbon, which can be retained in solid solution by rapid quenching. However, in these circumstances the phase transformation takes place, not to ferrite but to a tetragonal structure referred to as martensite. This phase forms as a result of diffusion less shear transformation leading to characteristic laths or plates. If the quench is sufficiently rapid, the martensite is essentially a supersaturated solid solution of carbon in a tetragonal iron matrix, and as the carbon concentration can be greatly in excess of the equilibrium concentration in ferrite, the strength is raised very substantially. High carbon martensites are normally very hard but brittle, the yield strength reaching as much as 1500 MPa; much of this increase can be directly attributed to increased interstitial solid solution hardening, but there is also a contribution from the high dislocation density, which is characteristic of martensitic transformations in iron-carbon alloys.


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Substitutional solid solution strengthening of iron

Many metallic elements form solid solutions in γ- and α-iron. These are invariably substitutional solid solutions, but for a constant atomic concentration of alloying elements there are large variations in strength. Using single crystal data for several metals, Fig. 1 shows that an element such as vanadium has a weak strengthening effect on α-iron at low concentrations (< 2%), while silicon and molybdenum are much more effective strengthened. Other data indicates that phosphorus; manganese, nickel and copper are also effective strengtheners. However, it should be noted that the relative strengthening might alter with the temperature of testing, and with the concentrations of interstitial solutes present in the steels.

Figure 1. Solid solution strengthening of iron crystals by substitutional solutes. Ratio of the critical resolved shear stress τ0

to shear modulus μ as a function of atomic concentration.

The strengthening achieved by substitutional solute atoms is, in general, greater the larger the difference in atomic size of the solute from that of iron, applying the Hume-Rothery size effect. However, from the work of Fleischer and Takeuchi it is apparent that differences in the elastic behavior of solute and solvent atoms are also important in determining the overall strengthening achieved. In practical terms, the contribution to strength from solid solution effects is superimposed on hardening from other sources, e.g. grain size and dispersions. Also it is a strengthening increment, like that due to grain size,


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which need not adversely affect ductility. In industrial steels, solid solution strengthening is a far from negligible factor in the overall strength, where it is achieved by a number of familiar alloying elements, e.g. manganese, silicon, nickel, molybdenum, several of which are frequently present in a particular steel and are additive in their effect. These alloying elements arc usually added for other reasons, e.g. Si to achieve deoxidation, Mn to combine with sulphur or Mo to promote hardenability. Therefore, the solid solution hardening contribution can be viewed as a useful bonus.

Effect of solid solution strengthening

The graph below shows how different alloying additions affect the yield strength of a ferrite + pearlite structural steel.

Quiz：

1. C and N are interstitial elements rather than substitutional. 2. C and N are much bigger atoms hence introduce more strain energy. 3. C and N have a much smaller atomic mass number therefore for the same weight % addition you are

adding a higher atomic %. 4. Mn and Mo have much lower diffusivities in steel therefore cannot pin the moving dislocations.

1. and 3. - correct, interstitial elements tend to produce a greater strengthening effect because they cause more strain in the lattice than substitutional elements in steels. The fact that C and N have a smaller atomic


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mass number does mean that for the same weight % addition there are more C or N atoms introduced into the steel and hence more lattice strain

Iron and Its Interstitial Solid Solutions

DESCRIPTION Iron, a silvery white solid metal, appears in Group VIII of the periodic table as a transition element. Its atomic number is 26, and its atomic weight is 55.847. Iron is notable among the elements in the abundance of its ores and the vast number of useful alloys that can be formulated with iron as the major constituent. Iron is also biologically important. In its pure form, iron is rather soft and is malleable and ductile at room temperature. It melts at 1,535 deg C and boils at 3,000 deg C. Pure iron can exist in two structural types, or allotropic forms. At room temperature the iron atoms are arranged in a body-centered cubic lattice called the a-form, which is transformed at 910 deg C into a cubic close-packed structure called the gamma-form. At 1,390 deg C iron returns to a body-centered cubic structure, called the delta-form. APPLICATION

Iron is abundant and easily obtainable from its ores. Its desirable mechanical and magnetic properties, as well as its resistance to corrosion, may be improved by mixing iron with other elements, frequently metals, to form alloys. Perhaps the most important alloy of iron is steel, which contains up to approximately 2% carbon. Steels that contain about 0.25% carbon are called mild steels; those with about 0.45% carbon are medium steels; and those with 0.60% to 2% carbon are high-carbon steels. Within this range, the greater the carbon content, the greater the tensile strength of the steel. The hardness of steel may be substantially increased by heating the metal until it is red hot and then quickly cooling it, a process known as quench hardening. An important component of many steels is cementite, a carbon-iron compound. Mild steels are ductile and are fabricated into sheets, wire, or pipe; the harder medium steels are used to make structural steel. High-carbon steels, which are extremely hard and brittle, are used in tools and cutting instruments. The addition of other materials in alloys (for example, manganese or silicon) also increases the hardness of steel. The inclusion of tungsten permits high-speed drills and cutting tools to remain hard even when used at high temperatures. The inclusion of chromium and nickel improves the corrosion resistance of the steel and, within certain limits of composition, is called stainless steel. A common stainless steel contains 0.15% carbon, 18% chromium, and 8% nickel and is used in cooking utensils and food-processing equipment. The inclusion of silicon, ranging from 1 to 5%, results in an alloy that is hard and highly magnetic. An alloy with cobalt is used for permanent magnets.


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THE SIMULATION

Computer simulations of defect energies were run for both vacancies in BCC and FCC iron as well as free surface energy for BCC iron. Parameters for the simulations included the following facts about iron: FCC: lattice parameter = 3.515Å, cohesive energy = 4.196eV BCC: lattice parameter = 2.87Å, cohesive energy = 4.28eV

THE RESULTS

Figure 1 : Vacancy in a BCC crystal lattice. For visualization purposes, the lattice atoms (arbitrarily designated so) are red and connected while the body centered atoms are yellow. The vacancy is "located" within the inner red cube. The BCC iron vacancy simulation resulted in a defect energy of 5.988eV.


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Figure 2 : Vacancy in a FCC crystal lattice.

Figure 3 : Vacancy in a FCC crystal lattice.


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Figure 4 : Vacancy in a FCC crystal lattice. The above three figures show the simulation of a defect in FCC iron. Figure 2 is included for the benefit of the reader. It is not "spiffed up," but rather shows the size and complexity of the simulation. The vacancy is the irregularity toward the center of the figure. For visualization purposes, both Figure 3 (ortho projection) and Figure 4 (perspective projection) show diagonals of all lattice faces. These are the "nearest neighbor" bonds. The lattice face that "contains" the vacancy (a "face" atom) is emphasized by the yellow color of its face's corners. The FCC iron vacancy simulation resulted in a defect energy of 5.857eV.

The study of steels is important because steels represent by far the most widely used metallic materials, primarily due to the fact that they can be manufactured relatively cheaply in large quantities to very precise specifications. They also provide an extensive range of mechanical properties from moderate strength levels (200-300MPa) with excellent ductility and toughness, to very high strengths (2000 MPa) with adequate ductility. It is, therefore, not surprising that irons and steels comprise well over 80% by weight of the alloys in general industrial use. Steels form perhaps the most complex group of alloys in common use. Therefore, in studying them it is useful to consider the behavior of pure iron first, then iron-carbon alloys, and finally examine the many complexities which arise when further alloying additions are made. Pure iron is not an easy material to produce. However, it has recently been made with a total impurity content not exceeding 60 ppm (parts per million), of which 10 ppm is accounted for by non-metallic impurities such as


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carbon, oxygen, sulphur, phosphorus, while 50 ppm represents the metallic impurities. Iron of this purity is extremely weak: the resolved shear stress of a single crystal at room temperature can be as low as 10 MPa, while the yield stress of a polycrystalline sample at the same temperature can be well below 150 MPa. The phase transformation: α- and γ- iron

BCC α-iron


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BCC α- iron

FCC γ- iron

Pure iron exists in two crystal forms, one body-centred cubic (bcc) (α-iron, ferrite) which remains stable from low temperatures up to 910°C (the A3 point), when it transforms to a face-centred cubic (fcc) form (γ-iron, austenite). The γ-iron on remains stable until 1390°C, the A4 point, when it reverts to bcc form, (now δ-iron) which remains stable up to the melting point of 1536°C. The detailed geometry of unit cells of α- and γ-iron crystals is particularly relevant to, for example, the solubility in the two phases of non-metallic elements such as carbon and nitrogen, the diffusivity of alloying elements at elevated temperatures, and the general behavior on plastic deformation. The bcc structure of α-iron is more loosely packed than that of fcc γ-iron. The largest cavities in the bcc structure are the tetrahedral holes existing between two edge and two central atoms in the structure, which together form a tetrahedron. It is interesting that the fcc structure, although more closely-packed, has larger holes than the bcc-structure. These holes are at the centers of the cube edges, and are surrounded by six atoms in the form of an octagon, so they are referred to as octahedral holes.


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Location of octahedral holes (left) and tetrahedral holes (right) relative to lattice spheres in face-centered cubic unit cell.

Note: The holes are gray and the lattice spheres are light blue in the unit cells above.

The α↔γ transformation in pure iron occurs very rapidly, so it is impossible to retain the high-temperature fcc form at room temperature. Rapid quenching can substantially alter the morphology of the resulting α-iron, but it still retains its bcc structure. Carbon and nitrogen in solution in α- and γ- iron The addition of carbon to iron is sufficient to form a steel. However, steel is a generic term which covers a very large range of complex compositions. The presence of even a small concentration of carbon, e.g. 0.1-0.2 weight per cent (wt%); approximately 0.5-1.0 atomic per cent, has a great strengthening effect on iron, a fact known to smiths over 2500 years ago since iron heated in a charcoal fire can readily absorb carbon by solid state diffusion. However, the detailed processes by which the absorption of carbon into iron converts a relatively soft metal into a very strong and often tough alloy have only recently been fully explored. The atomic sizes of carbon and nitrogen are sufficiently small relative to that of iron to allow these elements to enter the α- iron and &gamma- iron lattices as interstitial solute atoms. In contrast, the metallic alloying elements such as manganese, nickel and chromium have much larger atoms, i.e. nearer in size to those of iron, and consequently they enter into substitutional solid solution. However, comparison of the atomic sizes of C and N with the sizes of the available interstices makes it clear that some lattice distortion must take place when these atoms enter the iron lattice. Indeed, it is found that C and N in α-iron occupy not the larger tetrahedral holes, but the octahedral interstices which are more favorably placed for the relief of strain, which occurs by movement of two nearest neighbor iron atoms. In the case of tetrahedral interstices, four iron atoms are of nearest-neighbor status and the displacement of these would require more strain energy. Consequently these interstices are not preferred sites for carbon and nitrogen atoms. The solubility of both C and N in austenite should be greater than in ferrite, because of the larger interstices available. It is, therefore, reasonable to expect that during simple heat treatments, excess carbon and nitrogen will be precipitated. This could happen in heat treatments involving quenching from the γ state, or even after


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treatments entirely within the α field, where the solubility of C varies by nearly three orders of magnitude between 720°C and 20°C. Precipitation of carbon and nitrogen from α-iron. α-iron containing about 0.02 wt % C is substantially supersaturated with carbon if, after being held at 700°C, it is quenched to room temperature. This supersaturated solid solution is not stable, even at room temperature, because of the ease with which carbon can diffuse in α-iron. Consequently, in the range 20-300°C, carbon is precipitated as iron carbide. This process has been followed by measurement of changes in physical properties such as electrical resistivity, internal friction, and by direct observation or the structural changes in the electron microscope. The process of ageing is a two-stage one. The first stage takes place at temperatures up to 200°C and involves the formation or a transitional iron carbide phase (ε) with a close-packed hexagonal structure which is often difficult to identify, although its morphology and crystallography have been established. It forms as platelets on {100}α planes, apparently homogenously in the α-iron matrix, but at higher ageing temperatures (150-200°C) nucleation occurs preferentially on dislocations. The composition is between Fe2.4C and Fe3C. Ageing at 200°C and above leads to the second stage of ageing in which orthorhombic cementite Fe3C is formed as platelets on {110}α. Often the platelets grow on several {110} planes from a common centre giving rise to structures which appear dendritic in character. The transition from ε-iron carbide to cementite is difficult to study, but it appears to occur by nucleation of cementite at the ε-carbide/α interlaces, followed by re-solution of the metastable ε-carbide precipitate. The maximum solubility of nitrogen in ferrite is 0.10 wt %, so a greater volume fraction of nitride precipitate can be obtained. The process is again two-stage with a be tetragonal α" phase, Fe16N2, as the intermediate precipitate, forming as discs on {100}α, matrix planes both homogeneously and on dislocations. Above about 200°C, this transitional nitride is replaced by the ordered fcc γ’, Fe4N. The ageing of α-iron quenched from a high temperature in the α-range is usually referred to as quench ageing, and there is substantial evidence to show that the process can cause considerable strengthening, even in relatively pure iron. In commercial low carbon steels, nitrogen is usually combined with aluminium, or present in too low concentration to make a substantial contribution to quench ageing, with the result that the major effect is due to carbon. This behavior should be compared with that of strain ageing. Some practical aspects. The very rapid diffusivity of carbon and nitrogen in iron compared with that of the metallic alloying elements is exploited in the processes of carburizing and nitriding. Carburizing can be carried out by heating a low carbon steel in contact with carbon to the austenitic range, e.g. 1000°C, where the carbon solubility, c1, is substantial. The result is a carbon gradient in the steel, from c1 at the surface in contact with the carbon, to c at a depth. The diffusion coefficient D of carbon in iron actually varies with carbon content, so the above relationship is not rigorously obeyed. Carburizing, whether carried out using carbon, or more efficiently using a carburizing gas (gas carburizing), provides a high carbon surface on a steel, which, after appropriate heat treatment, is strong


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and wear resistant. Nitriding is normally carried out in an atmosphere of ammonia, but at a lower temperature (500-550°C) than carburizing, consequently the reaction occurs in the ferrite phase, in which nitrogen has a substantially higher solubility than carbon. Nitriding steels usually contain chromium (≈1%), aluminum (≈1%), vanadium or molybdenum (≈0.2%), which are nitride-forming elements, and which contribute to the very great hardness of the surface layer produced.

http://www.msm.cam.ac.uk/doitpoms/tlplib/dislocations/dislocations_in_2D.php http://www.ndt-ed.org/EducationResources/CommunityCollege/Materials/Introduction/introduction.htm http://www.msm.cam.ac.uk/doitpoms/tlplib/solid-solutions/index.php


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Precipitation Hardening

Theory

The primary species of precipitation strengthening are second phase particles. These particles impede the movement of dislocations throughout the lattice. You can determine whether or not second phase particles will precipitate into solution from the solidus line on the phase diagram for the particles. Physically, this strengthening effect can be attributed both to size and modulus effects, and to interfacial or surface energy.

The presence of second phase particles often causes lattice distortions. These lattice distortions result when the precipitate particles differ in size from the host atoms. Smaller precipitate particles in a host lattice leads to a tensile stress, whereas larger precipitate particles leads to a compressive stress. Dislocation defects also create a stress field. Above the dislocation there is a compressive stress and below there is a tensile stress. Consequently, there is a negative interaction energy between a dislocation and a precipitate that each respectively cause a compressive and a tensile stress or vice versa. In other words, the dislocation will be attracted to the precipitate. In addition, there is a positive interaction energy between a dislocation and a precipitate that have the same type of stress field. This means that the dislocation will be repulsed by the precipitate.

Precipitate particles also serve by locally changing the stiffness of a material. Dislocations are repulsed by regions of higher stiffness. Conversely, if the precipitate causes the material to be locally more compliant, then the dislocation will be attracted to that region.

Furthermore, a dislocation may cut through a precipitate particle. This interaction causes an increase in the surface area of the particle. The area created is

where, r is the radius of the particle and b is the magnitude of the burgers vector. The resulting increase in surface energy is

where is the surface energy. The dislocation can also bow around a precipitate particle.

Governing Equations

There are two equations to describe the two mechanisms for precipitation hardening:

Dislocations cutting through particles:


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where τ is material strength, r is the second phase particle radius, γ is the surface energy, b is the magnitude of the Burgers vector, and L is the spacing between pinning points. This governing equation shows that the strength is proportional to r, the radius of the precipitate particles. This means that it is easier for dislocations to cut through a material with smaller second phase particles (small r). As the size of the second phase particles increases, the particles impede dislocation movement and it becomes increasingly difficult for the particles to cut through the material. In other words, the strength of a material increases with increasing r.

Dislocations bowing around particle:

where τ is the material strength, G is the shear modulus, b is the magnitude of the Burgers vector, L is the distance between pinning points, and r is the second phase particle radius. This governing equation shows that for dislocation bowing the strength is inversely proportional to the second phase particle radius r. Dislocation bowing is more likely to occur when there are large particles present in the material.

These governing equations show that the precipitation hardening mechanism depends on the size of the precipitate particles. At small r, cutting will be the dominant strengthening mechanism, while at large r, bowing will be the dominant strengthening mechanism.

Looking at the plot of both equations, it is clear that there is a critical radius at which max strengthening occurs. This critical radius is typically 5-30 nm.


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Some precipitation hardening materials

• 2000-series aluminum alloys (important examples: 2024 and 2019) • 6000-series aluminum alloys • 7000-series aluminum alloys (important examples: 7075 and 7475) • 17-4PH stainless steel (UNS S17400) • Maraging steel • Inconel 718 • Alloy X-750 • Rene 41 • Waspaloy

Because precipitation strengthening is generally associated with a reduction in toughness, grain refinement is often used in conjunction with precipitation strengthening to improve toughness.


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In most binary systems, alloying above a concentration given by the phase diagram will cause the formation of a second phase. A second phase can also be created by mechanical or thermal treatments. The particles that compose the second phase precipitates act as pinning points in a similar manner to solutes, though the particles are not necessarily single atoms. The dislocations in a material can interact with the precipitate atoms in one of two ways (see Figure 2). If the precipitate atoms are small, the dislocations would cut through them. As a result, new surfaces (b in Figure 2) of the particle would get exposed to the matrix and the particle/matrix interfacial energy would increase. For larger precipitate particles, looping or bowing of the dislocations would occur which results in dislocations getting longer. Hence, at a critical radius of about 5nm, dislocations will preferably cut across the obstacle while for a radius of 30nm, the dislocations will readily bow or loop to overcome the obstacle.

Most precipitates in steels are relatively large (compared to GP zones in Al-Cu alloys for example) and strong so dislocations have to travel between them rather than cut through them. This is called Orowan bowing. The effectiveness of the precipitates in the steel depend on their composition (as this controls their thermodynamic stability), size, volume fraction and distribution.

Slowly increase the applied force on the dislocation by holding the button down and see how a dislocation bows around precipitates of different sizes and spacing. Plot a graph of precipitate bowing stress v average separation.

Control of the precipitate type, size and distribution can be achieved in plate products using controlled reheating, rolling and cooling schedules and careful alloying additions. The typical alloying additions that are used to generate precipitate strengthening are Ti, V, Nb and Al. These additions can be made individually but more commonly are made in combination depending on the requirements of strength, toughness etc. from the


Page 56 of 180.

customer. The reason that these additions are used is that their precipitates (carbides and nitrides) have high thermodynamic stability at elevated temperatures which means they also give rise to grain refinement during processing in addition to strengthening at room temperature.

Alloying additions of fine precipitate forming elements can be used to refine the grain size of a steel. A fine precipitate distribution in a steel will restrict the growth of austenite grains at high temperature and will retard recrystallisation of deformed austenite grains. If the combination of precipitates and rolling schedule is used then a 'pancake' structure of deformed austenite grains is created during rolling. These deformed grains provide many nucleation sites for subsequent ferrite formation resulting in a fine grained ferritic structure in the final steel product. The choice of alloying elements is important as the precipitates must be stable at the high temperatures of processing in order to pin the grain boundaries. The typical alloying additions used in High Strength Low Alloy steels (HSLA) are Nb, Al, Ti and / or V. They can be added in isolation or in combination with one another.

Using the diagram above which alloying element gives you the most thermodynamically stable precipitate at 1150°C?

Ans: TiN


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An example of how alloying additions made in combination affect the strength (and toughness) is given below for Al and V additions to a C-Mn ferrite + pearlite steel. The formation of AlN particles causes grain refinement on heat treatment and some strengthening and the VC particles provided strengthening. Note that the toughness decreases as the strength increases except for the strengthening effect of grain refinement which is also beneficial to toughness.

Nb is used as a microalloying addition to control grain size and provide some precipitation strengthening. V is used to provide precipitation strengthening. The strengthening that is obtained by V precipitates will

depend upon their size but a guideline is that for every 0.1 wt% V addition an increase of 50 - 60 MPa is seen in strength

Precipitate hardening of stainless steel grades. http://www.smihq.org/public/publications/past_articles/jan06_zubek.pdf 17-4PH precipitate hardened martensitic stainless steel. http://www.sandmeyersteel.com/17-4PH.html History of precipitate hardening. http://nvl.nist.gov/pub/nistpubs/sp958-lide/014-015.pdf


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Chapter Two: The Effects of Alloying Elements on Iron-Carbon Alloys

Introduction:

Contributing factors to strength of steel.

The main contributions to the strength of steel come from:

Effect of grain size. Effect of microstructure. Effect of solid solution strengthening. Effect of precipitates. Effect of dislocations.

These factors are controlled by the chemistry of the steel and the processing route used to produce the final component. All of these aspects need to be taken into account when selecting or designing steel for a given application.

The processing routes include:

Alloying. Thermal heat treatment. Mechanical straining and deformation. Thermo-mechanical in combination e.g. TMCP steels. Age hardening. Mechanical alloying or dispersion hardening.

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Normalizing Heat Treatment Effect on Low Alloy Steel Weld Metals

http://www.scielo.br/pdf/jbsmse/v26n1/a11v26n1.pdf

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Effect of High Heat input on CTOD Property of the Thick Steel Plate for Offshore Engineering

http://www.posco.com/homepage/docs/eng/dn/company/product/7.Effect%20of%20High%20Heat%20input%20on%20CTOD%20Property%20of%20the%20Thick%20Steel%20Plate%20for%20Offshore%20Engineering.pdf

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Heat treatment of Steels

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Steel forming and heat treating handbook

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热处理网： http://www.rechuliwang.cn/html/rechulichangshi/

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Ways of improving the strength of alloy steels are:

The solid-solution hardening of carbon has a major effect on the strength of martensite, but ductility can only be obtained at low carbon levels. Although alloying elements affect hardenability, they have a minor effect on hardness except to reduce it at high carbon levels by causing austenite to be retained.

1. Grain refinement: which increases strength and ductility. This can be developed by severely curtailing the time after the cessation of forging at some low temperature of austenite stability or by rapid heating, coupled with a short austenitising period. Fine grain is produced in 9% Ni steel by tempering fine lath martensite.

2. Precipitation hardening by carbide, nitride or intermetallic compounds:

By secondary hardening, e.g. 12% Cr steel with additions. Age hardening a low carbon Fe-Ni lath martensite supersaturated with substitutional elements, e.g.

maraging. Age hardening of austenite, e.g. stainless steels. Phosphorus and titanium are common additions.

Stacking faults are often associated with fine carbide precipitates, and strength can be raised by increasing the number of stacking faults (i.e. lower fault energy).

Controlled transformation 18/8 austenite steels in which transformation to martensite is induced by refrigeration or by strain.

3. Thermomechanical treatments

which may be classified into three main groups: Deformation of austenite prior to the transformation.

Ausforming consists of steel deforming in a metastable austenitic condition between Ac1 and Ms (e.g. 500癈 called LT) followed by transformation to martensite and light tempering (Fig. 1). This results in increased dislocation density in the martensite and a finer carbon precipitation on tempering. Strengths up to 1800 MPa can be obtained without impairing the ductility (~6 % deformation). Steels must possess a TTT-curve with a large bay of stable austenite, e.g. 826 M40. Typical application is for leaf springs.

Figure 1. Methods of thermomechanical treatment Deformation of stable austenite just above Ac3 before cooling (called HT). The properties are somewhat inferior to those produced by ausforming.


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Deformation induced transformation originally used in Hadfield 13% Mn steel, but can be adapted to metastable austenitic stainless steels. The fully austenitic steel is severely warm-worked above the lowest temperature at which martensite is produced during the straining. The distinctive property is the high rate of straining hardening, which increases ductility.

Deformation of austenite during the transformation

Isoforming is the deformation of a steel (e.g. 1% Cr) during the isothermal transformation to pearlite, which refines the structure and improves fracture toughness (Fig. 1). A somewhat similar thermomechanical process can be used in the bainitic region, producing bainite and martensite. Zerolling consists in forming martensite by deformation at subzero temperatures to strengthen 18/8 austenitic steels. The amount of martensite is influenced by alloy composition and increased with deformation and lowering of the temperature.

Deformation after the transformation of austenite

Marforming consists of deforming the maraging steel in the soft martensitic condition, generally cold. There is a pronounced increase in strength of the subsequent maraged product. With other steels, considerable increases in strength can be obtained by straining martensite (~3 %) either in the untempered or tempered condition. A strengthening effect also occurs on re-tempering, probably due to the resolution and reprecipitation of the carbides in a more finely dispersed form. Strain tempering and dynamic strain ageing. Both processes involve about 5% deformation at the room temperature between two stages of tempering -- strain tempering -- while in dynamic strain ageing deformation is concurrent with tempering.

4. Solid solution Hardening:

For ferrite and pearlite steels relationships between yield and tensile strength and various compositional and microstructural factors have been developed. These are useful in that they show general characteristics, although they are unable to incorporate all factors that influence strength in modern steels, for example precipitation strengthening. Examples of these relationships are given below:

YS (MPa) = 53.9 + 32.3 Mn + 83.2 Si + 354 Nf + 17.4 d-1/2

UTS (MPa) = 294 + 27.7 Mn + 83.2 Si + 3.85 pearlite % + 7.7 d-1/2

Alloying addition in wt%, d is the ferrite grain size in mm, Nf is the free nitrogen content

ASTM Designation No.

Size Range, Inclusive, in.

Minimum Proof

Strength, kpsi

Minimum Tensile

Strength, kpsi

Minimum Yield

Strength, kpsi

Material

A307 .25-1.5 33 60 36 Low Carbon

A325, Type 1 0.5-1.0, 1.125-1.4 85, 74 120, 105 92, 81 Medium Carbon, Q&T

A325, Type 2 0.5-1.0, 1.125-1.5 85, 75 120, 106 92, 82 Low Carbon Martensite, Q&T

A325, Type 3 0.5-1.0, 1.125-1.6 85, 76 120, 107 92, 83 Weathering Steel, Q&T

A354, Alloy Steel, Q&T


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Grade BC

A354, Grade BD 0.25-4.0 120 150 130 Alloy Steel, Q&T

A449 0.25-1.0, 1.125-1.5, 1.75-3.0

85, 74, 55 120, 105, 90 92, 81, 58 Medium Carbon, Q&T

A490, Type 1 0.5-1.5 120 Alloy Steel, Q&T

A490, Type 3 Weathering Steel, Q&T The microstructure of steel has a significant effect on the strength of steel. For steel with a certain composition the microstructure can be altered through varying the processing route used. For example for strip steels the coiling temperature can be controlled to give varying microstructures in the final coil with varying properties.

In this example, different yield strengths can be achieved for a fixed composition by obtaining different microstructures.

Changing the amount of second phase in a predominantly ferritic microstructure has a pronounced effect on the strength of the steel. This microstructural change can be achieved through control of composition and processing as mentioned earlier. For structural steels produced as plate and sections with a ferrite + pearlite microstructure relatively small changes in the amount of pearlite and ferrite at low levels (~5-10%) have a relatively small effect on the strength level, especially the yield strength as the surrounding ferrite yields first anyway.

The presence of dislocation networks produced through work hardening can result in very high levels of strength, however, it reduces toughness and ductility so is used only for certain steel products. Work hardening is not used to any extent in plate and section steel products but is used in some strip and engineering steel products. For example: Strip products Some strip grades that do not require a high level of formability but an increased strength level are supplied with a degree of work hardening. This work hardening is achieved through controlled cold rolling passes after any annealing process (annealing removes the dislocation structure by recovery or recrystallisation). The level of work hardening introduced is determined by the strain during the cold roll passes (i.e. the thickness reduction per pass) and controlled to give the strength levels required. Work hardening, and hence strengthening, can also occur during fabrication, for example during press forming, bending, drawing etc. High carbon steel products For example steel wire is often supplied in a work hardened condition as the dislocation networks provide much of the strengthening. High strength steel wire is used in suspension cables, tyre cords, pre-stressed wire and wires can have yield strength up to 5000 MPa (lower strength values are usually seen for the applications mentioned). The wire is cold drawn from hot rolled to 90% reduction which produces a heavily worked structure Typical 'mild' steels have a small carbon content, usually under 0.2%. Increasing the carbon content


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hardens the steel and its ability to take and hold an edge but at the expense of toughness. Adding other alloying elements can alter these properties, though usually at the expense of increased cost and manufacturability. The main alloying elements and their effects are described in next section. Carbon steels usually contain less than 1 to 2% carbon and small quantities of manganese, copper, silicon, sulfur, and phosphorus. Alloy steels are carbon steel with other metals added specifically to improve the properties of the steel significantly. Stainless steel is considered a separate group. Plain carbon steel is produced with a wide range of mechanical properties with comparatively low cost. To extend the range of properties of steel, alloys have been developed. The benefits of alloying resulting include

• The maximum UTS is increased. • Thick sections steels are available with high hardness throughout the section. • More controllable quenching with minimum risk of shape distortion or cracking. • Improved impact resistance at high temperature range. • Improved corrosion resistance. • Improved high temperature performance.

The principle elements that are used in producing alloy steel include nickel, chromium, molydenenum, manganese, silicon and vanadium. Cobalt, copper and lead are also used as alloying elements.

Steels for special applications EF420 Lecture 7:

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Cast Irons EF420 Lecture 9:

Ferrous Alloys EF420 Lecture 12:

http://www.slideshare.net/charliechong/askeland-phule-notes-ch12-printable-presentation

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Effects of thermomechanical treatment on microstructure and mechanical properties of T91 steel

http://www.cjmr.org/manage/wenzhang/07019.pdf


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Alloying:

Effect of alloying elements

Elements may encourage formation of graphite from the carbide. Only a small proportion of these elements can be added to the steel before graphite forms destroying the properties of the steel, unless elements are added to counteract the effect. Elements which encourage the formation of graphite include silicon, cobalt, aluminium and nickel.

Alloying elements may go into solid solution in the iron, enhancing the strength. Elements which go into solid solution include silicon, molybdenum, chromium, nickel and magnesium.

Hard carbides (cementite) associated with iron and carbon may be formed with alloying elements. Elements which tend to form carbides include chromium, tungsten, titanium, columbium, vanadium, molybdenum and manganese.

Elements which stabilize austenite include manganese, nickel, cobalt and copper. These increase the range over which austenite is stable e.g. by lowering the eutectoid temperature, and this retards the separation or carbides.

Elements which tend to stabilize ferrite include chromium, tungsten, molybdenum, vanadium and silicon. They reduce the amount of carbon soluble in the austenite and thus increase the volume of free carbide in the steel at a given carbon content. The effectively reduce the austenite ( γ ) phase by raising the eutectoid temperature and lowering the peritectic temperature. Intermediate compounds with iron may be formed e.g. FeCr. Alloying elements may adjust the characteristics such as eutectoid content, quenching rate which produces bainite or martensite.

Relative effect of alloying elements

The combined effect of alloying elements results from many complex interactions resulting from the processing history, the number and quantities of constituents, the heat treatments, the section shape etc. Some basic rules can be identified.

Nickel has reduced carbide forming tendency than iron and dissolves in α ferrite. Silicon combines with oxygen to form nonmetallic inclusions or dissolves in the ferrite. Most of the manganese in alloy steels dissolves in the α ferrite. Any manganese that form carbides result

in (Fe,Mn)3C. Chromium spreads between the ferrite and carbide phases the spread depending on the amount of

carbon and other carbide generating elements present. Tungsten and molybdenum form carbides if sufficient carbon is present which has not already formed

carbides with other stronger carbide forming elements. Vanadium, titanium, and Colombian are strong carbide forming elements and are present in steel as

carbides. Aluminium combines with oxygen and nitrogen to form Al2O an AlN


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Characteristics of alloying

Alloying Additions to Steel

Steels are among the most commonly used alloys. The complexity of steel alloys is fairly significant. Not all effects of the varying elements are included. The following text gives an overview of some of the effects of various alloying elements. Additional research should be performed prior to making any design or engineering conclusions.

Element Influence Uses

Carbon

Hardness - Strength - Wear Most important alloying element. Is essential to the formation of cementite and other carbides, bainite and iron-carbon martensite. Within limits increasing the carbon content increases the strength and hardness of a steel while reducing its toughness and ductility.

Added to construction steels to increase strength, hardness and hardenability.

Aluminum Deoxidation - Ease of Nitriding

Boron Hardenability

Copper Corrosion Resistance - Strength

Lead Machinability

Phosphorus Strength

Silicon

Tellurium Machinabilty

Nickel

Toughness - Strength - Hardenability Stabilises gamma phase by raising A4 and lowering A3. Refines grains in steels and some non-ferrous alloys. Strengthens ferrite by solid solution. Unfortunatly is a powerful graphitiser. Can take into solid solution larger proportions of important elements such as chromium, molybdenum and tungsten than can iron.

Used up to help refine grain size. Used in large amounts in stainless and heat-resisting steels. Nickel based alloys can offer corrosion resistance in more aggressive environments and nickel is used as the basis of complex superalloys for high temperature service.

Manganese

Strength - Hardenability - More Response To Heat Treatment Deoxidises the melt. Greatly increases the hadenability of steels. Stabilises gamma phase. Forms stable carbides.

High manganese (Hadfield) steel contains 12.5% Mn and is austenitic but hardens on abrasion.

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Silicon

Carbide former. Hardenability De-oxidizes melt. Helps casting fluidity. Improves oxidation resistance at higher temperatures.

Up to 0.3% in steels for sand casting, up to 1% in heat resisting steels.

Chromium

Corrosion Resistance - Strength Stabilises alpha phase by raising A3 and depressing A4. Forms hard stable carbides. Strengthens ferrite by solid solution. In amounts above 13% it imparts stainless properties. Unfortunately increases grain growth.

Small amounts in constructional and tool steels. About 1.5% in ball and roller bearings. Larger amounts in Stainless and heat-resisting steels.

Molybdenum

High Temperature Strength - Hardenability Strong carbide-stabilising influence. Raises high temperature creep strength of some alloys. Slows tempering response. When added to stainless steels it greatly improves the pitting and crevice corrosion resistance. There are limits to the proportion that can be taken into an iron based matrix. However up to almost 30% can be incorporated into nickel based alloys which provides excellent corrosion resistance in many aqueous environments.

Reduces 'temper brittleness' in nickel-chromium steels. Increases red-hardness of tool steels. Now used to replace some tungsten in high-speed steels.

Vanadium

Fine Grain - Toughness Strong carbide forming tendency. Stabilises martensite and increases hardenability. Restrains grain growth. Improves resistance to softening at elevated temperatures after hardening.

Used to retain high temperature hardness, e.g. in dies for hot-forging and die casting dies. Increasingly used in high speed steels.

Tungsten

Stabilises alpha phase and forms stable, very hard carbides, which improves creep resistance and renders transformations very sluggish, hence hardened steels resist tempering influences.

Used in high-speed steels and other tool and die steels, particularly those for use at high temperatures. Used in a few stainless steels, in combination with molybdenum. to improve pitting and crevice corrosion resistance. It is also used in some high temperature nickel based alloys and in some high temperature austenitic stainless steels.

Cobalt

Hardness - Wear Has similar corrosion resistance to that of Nickel, but higher cost means that it is not normally used for such applications. Provides matrix - strengthening characteristics to stainless and nickel based alloys designed for high temperature applications.

Used in super high speed steels and maraging steels, permanent magnet steels and alloys.


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Slows the transformation of martensite, hence increases 'red hardness' which is useful in tool steels.

Niobium

Elimination Of Carbide Precipitation In low alloy steels it acts as a carbide former and improves creep resistance. In stainless steels it combines with carbon, stabilising the steel and reducing the susceptibility to intergranular corrosion

Used to stabilize stainless steels.

Titanium Elimination Of Carbide Precipitation In stainless steels combines with excess carbon reducing the risk of intergranular corrosion.

Used in stabilized stainless steels. In nickel based alloys it is used with aluminium to promote age hardening.

Effect of alloying elements. http://www.scribd.com/doc/7833439/Effects-of-Alloying-Elements

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Effects of Alloying Elements on the Microstructures and Mechanical on ductile cast irons

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Carbon

Carbon has a major effect on steel properties. Carbon is the primary hardening element in steel. Hardness and tensile strength increases as carbon content increases up to about 0.85% C as shown in the figure above. Ductility and weldability decrease with increasing carbon.

Carbon is essential in steels which have to be hardened by quenching and for example, in austenitic manganese steel which is required to have high resistance to wear. The maximum hardness obtainable in any carbon steel is a function of the carbon content which may vary up to about 2% according to the purpose for which the steel is to be used. It occurs in varying forms according to the percentage present, and the heat treatment to which the steel has been submitted. (See allotropy and transformation range). Cast irons usually contain from about 1.8% to 4.5% carbon, present either as free carbon (graphite) and/or combined carbon (cementite), the varying distribution of the carbon between these two forms considerably influencing the strength and hardness.

Carbon is a strong austenite former and strongly promotes an austenitic structure. It also substantially increases the mechanical strength. Carbon reduces the resistance to intergranular corrosion. In ferritic stainless steels carbon will strongly reduce both toughness and corrosion resistance. In the martensitic and martensitic-austenitic steels carbon increases hardness and strength. In the martensitic steels an increase in hardness and strength is generally accompanied by a decrease in toughness and in this way carbon reduces the toughness of these steels.

Carbon: Present in all steels, it is the most important hardening element. Also increases the strength of the steel. We usually want knife-grade steel to have >.5% carbon, which makes it "high-carbon" steel.

it has by far the greatest influence of any of the elements. Steel could not exist without carbon. Martensite, along with banite gives steel a microstructure of hard, tough carbide. None of the other elements so dramatically alter the strength and hardness as do small changes in carbon content. Carbon iron crystalline structures have the widest number and variety known to exist in metallurgy. They also combine with other elements to furnish steel with an assortment of iron alloy carbide systems.


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Manganese

Manganese fulfils a variety of functions in steel.

It is used as a deoxidizing agent in nearly all steels. It forms manganese sulphide inclusions which in the ingot are spherical. In the absence of

manganese sulphur forms interdendritic films of iron sulphide causing brittleness at forging temperature (hot shortness).

It effectively increases harden ability and up to 1.5% is added for this purpose. (d) In larger amounts it is used to stabilize austenite, as in 14% manganese steel.

MANGANESE (Mn): Is normally present in all steel and functions as a deoxidizer. It also imparts strength and responsiveness to heat treatment. It is usually present in quantities of 0.5 to 2.0 percent. Range 0.3% to 1.5% always present in steels to reduce the negative effects of impurities carried out forward from the production process e.g. sulphur embrittlement. It promotes the formation of stable carbides in quenched-hardened steels. Alloys containing manganese are pearlitic. Up to 1% acts as hardening agent and from 1% to 2% improves strength and toughness. Alloys containing more than 5% are non-magnetic. Alloys containing large proportions of up to 12.5% manganese have the property that they spontaneously form hard skins when subject to abrasion. (Self-hardening) All commercial steels contain 0.3-0.8% manganese, to reduce oxides and to counteract the harmful influence of iron sulfide. Any manganese in excess of these requirements partially dissolves in the iron and partly forms Mn3C which occurs with the Fe3C. There is a tendency nowadays to increase the manganese content and reduce the carbon content in order to get steel with an equal tensile strength but improved ductility If the manganese is increased above 1,8% the steel tends to become air hardened, with resultant impairing of the ductility. Up to this quantity, manganese has a beneficial effect on the mechanical properties of oil hardened and tempered 0.4% carbon steel. The manganese content is also increased in certain alloy steels, with a reduction or elimination of expensive nickel, in order to reduce costs. Steels with 0.3-0.4% carbon, 1,3-1,6% manganese and 0.3% molybdenum have replaced 3% nickel steel for some purposes. Non-shrinking tool steel contains up to 2% manganese, with 0.8-0.9% carbon. Steels with 5 to 12% manganese are martensitic after slow cooling and have little commercial importance. Hadfield`s manganese steel a specially steel which is austenitic and usually contains approximately 12% Manganese. It is used in mining, earth- moving equipment and in railroad track work. . Hadfield`s manganese steel contains 12 to 14% of manganese and 1,0% of carbon. It is characterized by a great resistance to wear and is therefore used for railway points, rock drills and stone crushers. Austenite is completely retained by quenching the steel from 1000°C, in which soft condition it is used, but abrasion raises the hardness of the surface layer from 200 to 600 VPN (with no magnetic change), while the underlying material remains rough. Annealing embrittles the steel by the formation of carbides at the grain boundaries. Nickel is added to electrodes for welding manganese steel and 2% Mo sometimes added, with a prior carbide dispersion treatment at 600°C, to minimize initial distortion and spreading.


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http://www.arema.org/eseries/scriptcontent/custom/e_arema/library/2005_Conference_Proceedings/00040.pdf

Manganese is generally beneficial to surface quality especially in resulfurized steels. Manganese contributes to strength and hardness, but less than carbon. The increase in strength is dependent upon the carbon content. Increasing the manganese content decreases ductility and weldability, but less than carbon. Manganese has a significant effect on the hardenability of steel.

Manganese aids the grain structure, and contributes to hardenability, strength & wear resistance. Improves the steel (e.g. deoxidizes) during the steel's manufacturing (hot working and rolling). Present in most cutlery steel except for A-2, L-6 and CPM 420V. Manganese slightly increases the strength of ferrite, and also increases the hardness penetration of steel in the quench by decreasing the critical quenching speed. This also makes the steel more stable in the quench. Steels with manganese can be quenched in oil rather than water, and therefore are less susceptible to cracking because of a reduction in the shock of quenching. Manganese is present in most commercially made steels. Manganese is generally used in stainless steels in order to improve hot ductility. Its effect on the ferrite/austenite balance varies with temperature: at low temperature manganese is a austenite stabiliser but at high temperatures it will stabilize ferrite. Manganese increases the solubility of nitrogen and is used to obtain high nitrogen contents in austenitic steels.

Figure 1: Non-metallic inclusion in steel: oxides-dark gray and sulfides-light gray

Manganese is generally used in stainless steels in order to improve hot ductility. Its effect on the ferrite/austenite balance varies with temperature: at low temperature manganese is a austenite stabiliser but at high temperatures it will stabilise ferrite. Manganese increases the solubility of nitrogen and is used to obtain high nitrogen contents in austenitic steels. In general Manganese increases strength and hardness; forms carbide; increases hardenability; lowers the transformation temperature range. When in sufficient quantity produces austenitic steel; always present in a steel to some extent because it is used as a deoxidizer


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Figure 2: Typical duplex oxide–sulfide inclusion (particle A, B and C) and plate-like MnS (particle D) in conventional

continuous casting silicon steel.

SEM of an inclusion.

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Effect of Mn and Ni on the Variation of the Microstructure and Mechanical Properties of Low-carbon Weld Metals

http://www.journalarchive.jst.go.jp/jnlpdf.php?cdjournal=isijinternational1989&cdvol=40&noissue=12&startpage=1237&lang=en&from=jnlabstract

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Nickel

Increase Strength. Improve Toughness. Unable to increase Hardness Ferrite Former.

NICKEL (Ni): Increases strength and toughness but is ineffective in increasing hardness. It is generally added in amounts ranging from 1 percent to 4 percent. In some stainless steels it is sometimes as high as 20 percent. It is used for strength, corrosion resistance, and toughness, nickel increases the strength of ferrite, therefore increasing the strength of the steel. It is used in low alloy steels to increase toughness and hardenability. Nickel also tends to help reduce distortion and cracking during the quenching phase of heat treatment.

Nickel is a ferrite strengthener. Nickel does not form carbides in steel. It remains in solution in ferrite, strengthening and toughening the ferrite phase. Nickel increases the hardenability and impact strength of steels. Range 0.2% to 5% Improves strength, toughness, and hardenability without seriously affecting the ductility. It encourages grain refinement. Nickel and chromium together have opposing properties and are used together to advantage in nickel-chrome steels. The resulting steels have their advantages combined and their undesirable features cancel each other At 5% nickel provides high fatigue resistance. When alloyed at higher proportions significant corrosion resistance results and at 27% a non magnetic stainless steel results.

The addition of nickel, in amounts up to 8% or 10 %, to low carbon steel, increases the tensile strength and considerably raises the impact resistance. 9% nickel steels are useful at very low temperatures. In engineering steels it is widely used, often with chromium and molybdenum. High nickel increases resistance to corrosion, and in combination with chromium, is used in the austenitic corrosion-resisting steels. Certain iron-nickel alloys have unique properties. 25% nickel steel is practically non-magnetic. Alloys with about 36% nickel have very low coefficients of expansion, whilst with 50% to 78.5% nickel; alloys are obtained having very high magnetic permeability in low fields. An alloy containing 29% nickel, 17% cobalt is used for sealing with certain borosilicate glasses.

For stainless steel, the main reason for the nickel addition is to promote an austenitic structure. Nickel generally increases ductility and toughness. It also reduces the corrosion rate and is thus advantageous in acid environments. In precipitation hardening steels nickel is also used to form the intermetallic compounds that are used to increase the strength.

Nickel and manganese are very similar in behavior and both lower the eutectoid temperature. This change point on heating is lowered progressively with increase of nickel (approximately 10°C for 1% of nickel), but the lowering of the change on cooling is greater and irregular. The temperature of this change (Ar1) is plotted for different nickel contents for 0.2% carbon steels in Fig. 1, It will be seen that the curve takes a sudden plunge round about 8% nickel. A steel with 12% nickel begins to transform below 300°C on cooling, but on reheating the reverse change does not occur until about 650°C. Such steels are said to exhibit pronounced lag or hysteresis and are called irreversible steels. This characteristic is made use of in maraging steels and 9% Ni cryogenic steel.


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It Increases strength and toughness but is ineffective in increasing hardness. It is generally added in amounts ranging from 1 percent to 4 percent. In some stainless steels it is sometimes as high as 20 percent.

Maraging steels are a class of high-strength steel with low carbon content and the use of substitutional (as opposed to interstitial) elements to produce hardening from formation of nickel martensites. The name maraging has resulted from the combination of

Martensite + Age hardening Maraging steels contain 18% nickel, along with a amounts of molybdenum, cobalt, and titanium and aluminium, and almost no carbon. These alloys can be strengthened significantly by a precipitation reaction at a relatively low temperature. They can be formed and machined in the solution-annealed condition but not without difficulty. Weldability is excellent. Fracture toughness of the maraging steels is considerably higher than that of the conventional high-strength steels. Maraging steels are hardened by a metallurgical reaction that does not involve carbon. Maraging steels are strengthened by intermetallic compounds such as Ni 3Ti and Ni 3Mo which precipitate at about 500°C. The carbon content provides no real benefit and is kept low as possible in order to minimize the formation of titanium carbide which can adversely affect mechanical properties. Toughness is superior to all low alloy carbon steels of similar strength, particularly the low temperature toughness. These steels are easy to machine and heat treat, so some cost savings result in component production to compensate for the high cost of the steel. A high strength maraging steel (extrusion section MIL-S-46850 grade 300) can have a 0.2% proof stress of 1930MPa and Ultimate Tensile strength of 2068 MPa with an elongation of 4%

The addition of nickel acts similarly to increasing the rate of cooling of a carbon steel. Thus with a constant rate of cooling the 5 to 8% nickel steels become troostitic; at 8 to 10% nickel, where the sudden drop appears, the structure is martensitic, while above 24% nickel the critical point is depressed below room temperature and austenite remains. The lines of demarcation are not so sharp as indicated by Fig. 1, but a gradual transition occurs from one constituent to another. Nickel Strengthens steel; lowers its transformation temperature range; increases hardenability, and improves resistance to fatigue. Strong graphite forming tendency; stabilizes austenite when in sufficient quantity. Creates fine grains and gives good toughness.


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Figure1. Effect of nickel on change points and mechanical properties of 0.2% carbon steels cooled at a constant rate.

The mechanical properties change accordingly as shown in the lower part of Fig. 1. Steels with 0.5% nickel are similar to carbon steel, but are stronger, on account of the finer pearlite formed and the presence of nickel in solution in the ferrite. When 10% nickel is exceeded the steels have a high tensile strength, great hardness, but are brittle, as shown by the Izod and elongation curves. When the nickel is sufficient to produce austenite the steels become non-magnetic, ductile, tough and workable, with a drop in strength and elastic limit. Carbon intensives the action of nickel and the change points shown in Fig. 1 will vary according to the carbon content. The influences of carbon and nickel on the structure are shown in the small inset (Guillet) diagram in Fig. 1, for one rate of cooling. Steels containing 2 to 5% nickel and about 0.1% carbon are used for case hardening; those containing 0.25 to 0.40% carbons are used for crankshafts, axles and connecting rods. The superior properties of low nickel steels are best brought out by quenching and tempering (550-650°C). Since the Ac3 point is lowered, a lower hardening temperature than for carbon steels is permissible and also a wider range of hardening temperatures above Ac3 without excessive grain growth, which is hindered by the slow rate of diffusion of the nickel. Martensitic nickel steels are not utilized and the austenitic alloys cannot compete with similar manganese steels owing to the higher cost. Maraging steels have fulfilled a high tensile requirement in aero and space fields. High nickel alloys are used for special purposes, owing to the marked influence of nickel on the coefficient of expansion of the metal. With 36% nickel, 0.2% carbon, 0.5%


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manganese, the coefficient is practically zero between 0°C and 100°C. This alloy ages with time, but this can be minimized by heating at 100°C for several days. The alloy is called Inver and it is used extensively in clocks, tapes and wire measures, differential expansion regulators, and in aluminum pistons with a split skirt in order to give an expansion approximating to that of cast iron. A carbon-free alloy containing 78.5% nickel and 21.5% iron has a high permeability in small magnetic fields.

Figure 6: Vertical section of Fe-Cr-C diagram for 0.1C wt%.

Figure 7: Schaeffler diagram for weld metals.

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Fun reading on Nickel based superalloy:http://www.msm.cam.ac.uk/phase-trans/2003/Superalloys/superalloys.html

http://www.msm.cam.ac.uk/phase-trans/2003/Superalloys/superalloys.html

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Chromium

Chromium is commonly added to steel to increase corrosion resistance and oxidation resistance, to increase hardenability, or to improve high-temperature strength. As a hardening element, Chromium is frequently used with a toughening element such as nickel to produce superior mechanical properties. At higher temperatures, chromium contributes increased strength. Chromium is strong carbide former. Complex chromium-iron carbides go into solution in austenite slowly; therefore, sufficient heating time must be allowed for prior to quenching.

It is added for wear resistance, hardenability, and (most importantly) for corrosion resistance. As with manganese, chromium has a tendency to increase hardness penetration. When 5 percent chromium or more is used in conjunction with manganese, the critical quenching speed is reduced to the point that the steel becomes air hardening. Chromium can also increase the toughness of steel, as well as the wear resistance. As an alloying element in steel, chromium increases the hardenability and in association with high carbon gives resistance to abrasion and wear. 4%is present in high speed steel and up to 5% is present in hot die steels. In Structural steels it may be present in amounts up to about 3 %. Simple chromium-carbon steels are used for ball bearings having high elastic limit and high uniform hardness due to the uniform distribution of the hard carbide particles, but for most structural purposes chromium is used in conjunction with up to 4 % nickel and small amounts of molybdenum or vanadium. In heat-resisting steels, chromium is present in amounts up to 30%, and it is an important element in many of the highly alloyed heat-resisting materials, whose iron contents are so low that they cannot be regarded as steel. Chromium is also used as an alloying addition to high duty cast irons.

This is the most important alloying element in stainless steels. It is this element that gives the stainless steels their basic corrosion resistance. The corrosion resistance increases with increasing chromium content. It also increases the resistance to oxidation at high temperatures. Chromium promotes a ferritic structure. Chromium is unique in its effect on resistance to corrosion and scaling and is an essential constituent in all stainless steels, e.g., stainless cutlery steels contain 12% to 14% chromium, whilst in steels of the austenitic corrosion-resisting type, 18% chromium is associated with 8% nickel, and small amounts of other elements. Steel with at least 13% chromium is deemed "stainless" steel. Despite the name, all steel can rust if not maintained properly. Chromium can dissolve in either alpha- or gama-iron, but, in the presence of carbon, the carbides formed are cementite (FeCr)3C in which chromium may rise to more than 15%; chromium carbides (CrFe)3C2 (CrFe)7C3 (CrFe)4C, in which chromium may be replaced by a few per cent, by a maximum of 55% and by 25% respectively. Stainless steels contain Cr4C. The pearlitic chromium steels with, say, 2% chromium are extremely sensitive to rate of cooling and temperature of heating before quenching;

It increases the depth penetration of hardening and also the responsiveness to heat treatment. It is usually added with nickel (Ni) for use in stainless steels. Most of the chromium (Cr) bearing alloys contain 0.50 to 1.50 percent chromium; some stainless steels contain as much as 20 percent or more. It can affect forging, causing a tendency in the steel to crack. For example:


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Temperature of Initial Heating, °C Critical Hardening Rate

(Min. to cool from 836°C to 546°C)

836 3.5 S

1010 6.5 S

1200 13 S

The reason is that the chromium carbides are not readily dissolved in the austenite, but the amount increases with increase of temperature. The effect of the dissolved chromium is to raise the critical points on heating (Ac) and also on cooling (Ar) when the rate is slow. Faster rates of cooling quickly depress the Ar points with consequent hardening of the steel. Chromium imparts a characteristic form of the upper portion of the isothermal transformation curve. The percentage of carbon in the pearlite is lowered. Hence the proportion of free cementite (hardest constituent) is increased in high carbon steel and, when the steel is properly heat-treated, it occurs in the spheroidised form which is more suitable when the steel is used for ball bearings. The pearlite is rendered fine. When the chromium exceeds 1.1% in low-carbon steels an inert passive film is formed on the surface which resists attack by oxidizing reagents. Still higher chromium contents are found in heat-resisting steel. Chromium steels are easier to machine than nickel steels of similar tensile strength. The steels of higher chromium contents are susceptible to temper brittleness if slowly cooled from the tempering temperature through the range 550/450°C. These steels are also liable to form surface markings, generally referred to as "chrome lines". The chrome steels are used wherever extreme hardness is required, such as in dies, ball bearings, plates for safes, rolls, files and tools. High chromium content is also found in certain permanent magnets.

Figure 2. Effect of alloying with chromium on the critical temperature of steel and austenite (g -iron) phase transformation

zone on the iron-iron carbide diagram.


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Nickel and chromium

Nickel steels are noted for their strength, ductility and toughness, while chromium steels are characterized by their hardness and resistance to wear. The combination of nickel and chromium produces steels having all these properties, some intensified, without the disadvantages associated with the simple alloys. The depth of hardening is increased, and with 4,5% nickel, 1,25% chromium and 0.35% carbon the steel can be hardened simply by cooling in air. Low nickel-chromium steels with small carbon content are used for casehardening, while for most constructional purposes the carbon content is 0.25-0.35%, and the steels are heat-treated to give the desired properties. Considerable amounts of nickel and chromium are used in steel for resisting corrosion and oxidation at elevated temperatures. Embattlement. The effects of tempering a nickel-chromium steel are shown in Fig. 2, from which it will be noticed that the Izod impact curve No. 1 reaches a dangerous minimum in the range 250-450°C in common with many other steels. This is known as 350°C embattlement. Phosphorus and nitrogen have a significant effect while other impurities (As, Sb, Sn) and manganese in larger quantity may also contribute to the embattlement.

Figure 2. Effect of tempering on the mechanical properties of nickel-chromium steel, C 0.26, Ni 3, Cr 1,2, 29 mm diam, bars hardened in oil from 830°C. Izod (2) for steel with 0.25% molybdenum added Temper brittleness is usually used to describe the notch impact intergranular brittleness (Grain boundaries are revealed in temper brittle samples by etching in 1 gm cetyl trimethyl ammonium bromide; 20 gm picric acid; 100 cc distilled water, 100 cc ether. Shake mixture, allow to stand for 24 hrs; use portion of top layer and return to tube afterwards) induced in some steels by slow cooling after tempering above about 600°C and also from prolonged soaking of tough material between about 400?and 550°C. Temper brittleness seems to be due to grain boundary enrichment with alloying elements-Mn, Cr, Mo-during austenitising which leads to enhanced segregation of embattling elements P, Sn, Sb, As-by chemical


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interaction on slow cooling from 600°C. The return to the tough condition, obtained by rehearing embattled steel to temperatures above 600°C and rapidly cooling, is due to the redistribution and retention in solution of the embattling segregation. Antimony (0-001 %), phosphorus (0-008 %), arsenic, tin, manganese increase, while molybdenum decreases the susceptibility of a steel to embattlement. 0-25 % molybdenum reduces the brittleness as shown by Izod curve No. 2. Table 1 illustrates the effect rate of cooling after tempering and the influence of an addition of 0-45 % molybdenum: Table 1. Steel 0.3% C, 3,5 % Ni, 0.7%, Cr, tempered at 630°C

Steel Cooling Rate

TS MPa

Elongation RA Izod ft lbf

Izod J

Ni-Cr Oil 896 18 60 64 87

Ni-Cr Furnance 880 18 60 19 25

Ni-Cr-Mo Furnance 896 18 61 59 80

Nickel And Chromium Used together for austenitic stainless steels; each element counteracts disadvantages of the other.


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Molybdenum

Molybdenum: A carbide former, prevents brittleness & maintains the steel's strength at high temperatures. Present in many steels, and air-hardening steels (e.g. A-2, ATS-34) always have 1% or more molybdenum -- molybdenum is what gives those steels the ability to harden in air. It is Adds greatly to the penetration of hardness and increases toughness of an alloy. It causes steel to resist softening at high temperatures, which defeats the purpose of forging. If the alloy has below 0.020 percent molybdenum (Mo), you should be able to forge this alloy with little difficulty.

Molybdenum is used very widely because of its powerful effect in increasing hardenability and also because in low alloy steels it reduces susceptibility to temper brittleness. It forms stable carbides, raises the temperature at which softening takes place on tempering and increases resistance to creep. In high speed steel it can be used to replace approximately twice its weight of tungsten. The corrosion resistance of stainless steel is improved by molybdenum additions.

Molybdenum increases the hardenability of steel. Molybdenum may produce secondary hardening during the tempering of quenched steels. It enhances the creep strength of low-alloy steels at elevated temperatures.

Molybdenum substantially increases the resistance to both general and localized corrosion. It increases the mechanical strength somewhat and strongly promotes a ferritic structure. Molybdenum also promotes the formation secondary phases in ferritic, ferritic-austenitic and austenitic steels. In martensitic steels it will increase the hardness at higher tempering temperatures due to its effect on the carbide precipitation.

Molybdenum dissolves in both alpha- and gama-iron and in the presence of carbon forms complex carbides (FeMo)6C, Fe21Mo2C6, Mo2C. Molybdenum is similar to chromium in its effect on the shape of the TTT-curve but up to 0.5% appears to be more effective in retarding pearlite and increasing bainite formation. Additions of 0.5% molybdenum have been made to plain carbon steels to give increased strength at boiler temperatures of 400°C, but the element is mainly used in combination with other alloying elements. Ni-Cr-Mo steels are widely used for ordnance, turbine rotors and other large articles, since molybdenum tends to minimize temper brittleness and reduces mass effect. Molybdenum is also a constituent in some high-speed steels, magnet alloys, heat-resisting and corrosion-resisting steels. Molybdenum Strong carbide forming element, and also improves high temperature creep resistance; reduces temper-brittleness in Ni-Cr steels. Improves corrosion resistance and temper brittleness.


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Vanadium

Ferrite Promoter. Carbide and Nitride Former.

Vanadium acts as a scavenger for oxides, forms vanadium carbide VC, and has a beneficial effect on the mechanical properties of heat-treated steels, especially in the presence of other elements. It slows up tempering in the range of 500-600°C and can induce secondary hardening. Chromium-vanadium (0.15%) steels are used for locomotive forging, automobile axles, coil springs, torsion bars and creep resistance.

Vanadium increases the yield strength and the tensile strength of carbon steel. The addition of small amounts of Vanadium can significantly increase the strength of steels. Vanadium is one of the primary contributors to precipitation strengthening in microalloyed steels. When thermomechanical processing TMCP is properly controlled the ferrite grain size is refined and there is a corresponding increase in toughness. The impact transition temperature also increases when vanadium is added.

All microalloy steels contain small concentrations of one or more strong carbide and nitride forming elements. Vanadium, niobium, and titanium combine preferentially with carbon and/or nitrogen to form a fine dispersion of precipitated particles in the steel matrix.

The presence of vanadium in steel raises the temperature at which grain coarsening sets in and under certain conditions increases the hardenability. It also lessens softening on tempering and confers secondary hardness on high speed and other steels. Vanadium carbide is intensely hard and as much as 5% vanadium may be added to high speed and high chromium tool steel where it improves abrasion resistance. Vanadium is an important constituent in many types of steel, for widely varying applications, e.g., nitriding, heat-resistance, tools, wearing plates and other fully hardened parts. In conjunction with molybdenum, vanadium has a marked effect in enhancing creep resistance.

Vanadium increases the hardness of martensitic steels due to its effect on the type of carbide present. It also increases tempering resistance. Vanadium stabilises ferrite and will, at high contents, promote ferrite in the structure. It is only used in hardenable stainless steels.

It retards grain growth within steel even after long exposures at high temperatures, and helps to control grain structures while heat treating. It is usually present in small quantities of 0.15 to 0.20 percent. Most tool steels which contain this element seem to absorb shock better that those that do not contain vanadium (V). Vanadium contributes to wear resistance and hardenability. A carbide former that helps produce fine-grained steel. A number of steels have vanadium, but M-2, Vascowear, and CPM T440V and 420V (in order of increasing amounts) have high amounts Vanadium Strong carbide forming element; has a scavenging action and produces clean, inclusion free steels. Can cause re-heat cracking when added to chrome molly steels.


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Tungsten

TUNGSTEN (W): Also known as wolfram, is used as an alloying element in tool steels, as it tends to impart a tight, small, and dense grain pattern and keen cutting edges when used in relatively small amounts. It will also cause steel to retain its hardness at higher temperatures and hence will have a detrimental effect upon the steel's forgeability (otherwise known as "red hard") Tungsten dissolves in gama-iron and in alpha-iron. With carbon it forms WC and W2C, but in the presence of iron it forms Fe3W3C or Fe4W2C. A compound with iron, Fe3W2 provides an age-hardening system. Tungsten raises the critical points in steel and the carbides dissolve slowly over a range of temperature. When completely dissolved, the tungsten renders transformation sluggish, especially to tempering, and use is made of this in most hot-working tool ("high speed") and die steels. It Increases wear resistance. When combined properly with chromium or molybdenum, tungsten will make the steel to be a high-speed steel. The high-speed steel M-2 has a high amount of tungsten. Tungsten refines the grain size and produces less tendency to decarburisation during working. Tungsten is also used in magnet, corrosion- and heat-resisting steels.

The effect of the addition of this metal to steel is to increase the strength at normal and elevated temperatures. Owing to the hardness of tungsten carbide and its influence on secondary hardening, tungsten is used as the main alloy addition in high speed tool steels, molybdenum being its only substitute. In addition, tungsten finds considerable application in general tool steels, die and precipitation hardening steels. It has found a useful application in valves and other steels required for use at high temperatures. Tungsten is an essential constituent in the sintered hard metals.

Tungsten Forms hard and stable carbides; raises the transformation temperature range, and tempering temperatures. Hardened tungsten steels resist tempering up to 6000C


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Silicon

Ferrite Former. Encourage Brittleness.

Silicon Si, is one of the principal deoxidizers used in steelmaking. Silicon is less effective than manganese in increasing as-rolled strength and hardness. In low-carbon steels, silicon is generally detrimental to surface quality. Silicon increases the resistance to oxidation, both at high temperatures and in strongly oxidising solutions at lower temperatures. It promotes a ferritic structure.

It has a beneficial effect upon tensile strength and improves hardenability of an alloy. It has a toughening effect when used in combination with certain other elements. Silicon (Si) is usually added to improve electrical conductivity of an alloy. Its average concentration is between 1.5 and 2.5 percent.

Silicon is used as a deoxidizer in the manufacture of steel. It slightly increases the strength of ferrite, and when used in conjunction with other alloys can help increase the toughness and hardness penetration of steel.

Silicon is a powerful deoxidizer, and as such is used in steel making processes in amounts up to about 0·8%. When used as an alloying element, silicon in small percentages will increase the tensile strength and yield point of structural steels. It is used in amounts of 1·5% to 2% in silicon-manganese spring steels and ultra-high tensile steels due to its effect in raising the limit of proportionality and resistance to tempering. Up to 4% in heat resisting steels improves scale resistance owing to the formation of a protective layer (see also Ihrigizing). The higher the silicon, the higher the temperature at which protection against further atmospheric oxidation is given. Water vapour and carbon dioxide, however, attack the layer. Alloys of iron and silicon, containing 15% of the element, are used as acid-resisting materials, but have the properties of cast irons rather than of steels. Carbon-free alloys with up to 4% silicon have a high electrical resistance and low hysteresis loss, and are used as transformer steels. In cast iron, silicon not only serves as a deoxidizer but also has a marked graphitizing effect, thus improving machinability.

Silicon dissolves in the ferrite, of which it is a fairly effective hardener, and raises the Ac change points and the Ar points when slowly cooled and also reduces the gama-alpha volume change. Only three types of silicon steel are in common use-one in conjunction with manganese for springs; the second for electrical purposes, used in sheet form for the construction of transformer cores, and poles of dynamos and motors, that demand high magnetic permeability and electrical resistance; and the third is used for automobile valves.

C Si Mn

1. Silico-manganese 0.5 1,5 0.8

2. Silicon steel 0.07 4,3 0.09

3. Silichrome 0.4 3,5 8


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It contributes oxidation resistance in heat-resisting steels and is a general purpose deoxidizes. In general Silicon strengthens ferrite and raises the transformation temperature temperatures; has a strong graphitizing

tendency. Always present to some extent, because it is used with manganese as a deoxidizer

Copper

Austenite Former. Impair Forging. Use as precipitation hardening alloy e.g17-4PH

Copper dissolves in the ferrite to a limited extent; not more than 3.5% is soluble in steels at normalizing temperatures, while at room temperature the ferrite is saturated at 0.35%. It lowers the critical points, but insufficiently to produce martensite by air cooling. The resistance to atmospheric corrosion is improved and copper steels can be temper hardened.

Copper Cu, in significant amounts is detrimental to hot-working steels. Copper negatively affects forge welding, but does not seriously affect arc or oxyacetylene welding. Copper can be detrimental to surface quality. Copper is beneficial to atmospheric corrosion resistance when present in amounts exceeding 0.20%. Weathering steels are sold having greater than 0.20% Copper.

The addition of about 0.20% copper to low carbon steel may increase its resistance to atmospheric corrosion by as much as 20% to 30%. In amounts of about 0·50% copper appreciably increases the tensile and yield strengths. The addition of increasing amounts of copper leads to defects in rolling. High yield point structural steels containing copper, in association with chromium and appreciable percentages of silicon and phosphorus have been developed. Copper is also added to some stainless steels to improve corrosion resistance.

Copper enhances the corrosion resistance in certain acids and promotes an austenitic structure. In precipitation hardening steels copper is used to form the intermetallic compounds that are used to increase the strength.

Phosphorous

Embrittlement Effects Increase Machinability. Undesirable Element.

Phosphorus P, increases strength and hardness and decreases ductility and notch impact toughness of steel. The adverse effects on ductility and toughness are greater in quenched and tempered higher-carbon steels. Phosphorous levels are normally controlled to low levels. Higher phosphorus is specified in low-carbon free-machining steels to improve machinability.


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Although it has been used to increase the tensile strength of steel and to improve resistance to atmospheric corrosion, phosphorus is usually regarded as an undesirable impurity because of its embrittling effect. In most British specifications the maximum permitted is 0.05 %, but in steel for nitriding it may be restricted to a maximum of 0·02 % since during the nitriding treatment phosphorus has a temper embrittling effect.

Sulphur

Increase Machinability. Undesirable element. Embrittlement Effects. Impair Ductility. Impair Weldability.

SULFUR S, Is usually regarded as an impurity in most alloys and its addition to steel is held to a minimum as it is damaging to the hot forming characteristics of steel. It is, however added to increase machinability. A word of caution, some alloys are offered in different forms, an example is E52100. This particular steel can be had in either a "Bearing Quality" or "Machining Quality" the latter having sulfur added to increase machinability.

Sulfur decreases ductility and notch impact toughness especially in the transverse direction. Weldability decreases with increasing sulfur content. Sulfur is found primarily in the form of sulfide inclusions. Sulfur levels are normally controlled to low levels. The only exception is free-machining steels, where sulfur is added to improve machinability.

A non-metal, which combines with iron to form iron sulphides, in which form its effect is to make the steel red short but combined with manganese its influence is less injurious. In steel the sulphur content is usually specified as less than 0.05 % but it may be added deliberately to improve machinability.

Sulphur is added to certain stainless steels, the free-machining grades, in order to increase the machinability. At the levels present in these grades sulphur will substantially reduce corrosion resistance, ductility and fabrication properties, such as weldability and formability.


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Lead

Improve Machinability. Undesirable Element. Impair Ductility. Impair Toughness. Impair Creep Strength.

Lead Pb, increase the machinability of steel and has no effect upon the other properties of the metal. Lead is virtually insoluble in liquid or solid steel. However, lead is sometimes added to carbon and alloy steels by means of mechanical dispersion during pouring to improve the machinability.

The addition of about 025% lead improves machinability. It also causes a reduction in fatigue strength, ductility and toughness but this only becomes serious in the transverse direction and at high tensile levels. In creep resisting alloys very small amounts of lead may be harmful.

Hydrogen

Hydrogen H, in steel is an undesirable impurity which is introduced from moisture in the atmosphere or the charge during melting. If a large amount of hydrogen is present in the liquid steel, some may be liberated on freezing giving an unsound ingot, evolution of hydrogen subsequently when the solid steel cools may cause hair line cracks. Hydrogen can be reduced to safe proportions by casting in vacuum or by prolonged annealing. It may also be introduced into steel by electrolytic action or by pickling and may then cause brittleness.

Selenium

Selenium Se, A metalloid closely resembling sulphur in its properties. It is sometimes added to steels to the extent of 0·2 % to 0·3 % to improve machinability.


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Tantalum

Tantalum Ta, This metal is associated with niobium and is very similar to it chemically. As an alloying addition to steel, niobium is preferred.

Tellurium is added to steel either alone or together with selenium to promote machinability. It is a powerful carbide stabilizer and has been also added to cast iron where it is said to increase the depth of chill and to prevent shrinkage. It may be added in small amounts to the molten iron or by the use of cores dipped or painted with washes containing tellurium in suspension.

Tin

Tin Sn, Owing to its good resistance to corrosion in many conditions, the major use of tin is in the form of coatings for steel and copper alloys. It is an undesirable impurity in steel giving rise to temper brittleness, but is less harmful than phosphorus.

Calcium

Calcium Ca, This metal in the form of calcium silicide is sometimes added to steel as a deoxidizer and degasefier.

Cerium

Cerium Ce, a metal of the rare earth class which in many respects resemble the alkali metals. The hot working properties of high alloy corrosion- and heat-resistant steels maybe improved by the addition of cerium, whilst in cast iron, cerium acts as a deoxidizer and desulphurizer but when the sulphur content has been reduced to a value of about 0·015%, the cerium enters into solution in the cast iron and functions as a powerful carbide stabilizer. In amounts above 0·02%, cerium is the operative factor in the production of nodular graphite structures in cast iron.

Cerium is one of the rare earth metals (REM) and is added in small amounts to certain heat resistant temperature steels and alloys in order to increase the resistance to oxidation and high temperature corrosion.


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Nitrogen

Nitrogen N, Nitrogen can combine with many metals to form nitrides and is thus applied to the case hardening of steel, the usual source for this purpose being ammonia. The incorporation of nitrogen in austenitic chromium-nickel steels stabilizes the austenite and increases the strength. In carbon steels it has an influence on creep. (See Abnormal Steels)

Nitrogen is a very strong austenite former and strongly promotes an austenitic structure. It also substantially increases the mechanical strength. Nitrogen increases the resistance to localised corrosion, especially in combination with molybdenum. In ferritic stainless steels nitrogen will strongly reduce toughness and corrosion resistance. In the martensitic and martensitic-austenitic steels nitrogen increases both hardness and strength but reduces the toughness.

Boron

Boron B, is added to fully killed steel to improve hardenability. Boron-treated steels are produced to a range of 0.0005 to 0.003%. Whenever boron is substituted in part for other alloys, it should be done only with hardenability in mind because the lowered alloy content may be harmful for some applications.

The addition of about 0.003 % of boron confers increased harden ability to steels in the quenched and tempered condition. Further, it has been found that the addition of 0·003% boron to low carbon, 0·50% molybdenum steel in the normalized condition doubles the yield strength and gives a 30% increase in tensile strength, but the advantage due to boron is very slight when molybdenum is less than 0.35% causes difficulty in forging. As much as 2% may be added to steels used in nuclear engineering.

Boron is a potent alloying element in steel. A very small amount of boron (about 0.001%) has a strong effect on hardenability. Boron steels are generally produced within a range of 0.0005 to 0.003%. Boron is most effective in lower carbon steels.

Boron. In recent years, especially in USA, 0.003-0.005% boron has been added to previously fully killed, fine-grain steel to increase the hardenability of the steel. The yield ratio and impact are definitely improved, provided advantage is taken of the increased hardenability obtained and the steel is fully hardened before tempering. In conjunction with molybdenum boron forms a useful group of high tensile bainitic steels. Boron is used in some hard facing alloys and for nuclear control rods.


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Aluminium

Aluminum Al, is widely used as a deoxidizer. Aluminum can control austenite grain growth in reheated steels and is therefore added to control grain size. Aluminum is the most effective alloy in controlling grain growth prior to quenching. Titanium, zirconium, and vanadium are also valuable grain growth inhibitors, but there carbides are difficult to dissolve into solution in austenite.

As a deoxidizer, up to 0.05% aluminum may be added to steel. For increasing fine grain characteristics or sub-zero impact properties, up to 0·10% may be added. Nitriding steels contain about 1% aluminum for promoting a high surface hardness when heated in ammonia. Still larger additions made to heat resisting steels promote resistance to scaling. Approximately 5% added to chromium steel increases electrical resistivity.

Aluminium improves oxidation resistance, if added in substantial amounts. It is used in certain heat resistant alloys for this purpose. In precipitation hardening steels aluminium is used to form the intermetallic compounds that increase the strength in the aged condition.

Zirconium

Zirconium Zr, can be added to killed high-strength low-alloy steels to achieve improvements in inclusion characteristics. Zirconium causes sulfide inclusions to be globular rather than elongated thus improving toughness and ductility in transverse bending.

Zirconium acts as a deoxidizing element in steel and combines with the sulphur.

Niobium

Niobium Nb (Columbium) increases the yield strength and, to a lesser degree, the tensile strength of carbon steel. The addition of small amounts of Niobium can significantly increase the yield strength of steels. Niobium can also have a moderate precipitation strengthening effect. Its main contributions are to form precipitates above the transformation temperature, and to retard the recrystallization of austenite, thus promoting a fine-grain microstructure having improved strength and toughness.

The metal is also known as columbium. It occurs in association with tantalum, to which it is closely related. Niobium is a strong carbide-forming element and as such is added to certain austenitic corrosion-resistant steels of the 18/8 chromium-nickel type for the prevention of intercrystalline corrosion. Where niobium is used as the stabilizer, it is usually specified that it should be present in an amount at least 8 times that of the carbon content. Further, niobium is often used as a constituent of the electrodes used in the welding of such steels. Niobium is added to heat-resisting steels and


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enhances creep strength. In small amounts, of the order of 0·05%, it increases the yield strength of mild steel.

Niobium is both a strong ferrite and carbide former. As titanium it promotes a ferritic structure. In austenitic steels it is added to improve the resistance to intergranular corrosion but it also enhances mechanical properties at high temperatures. In martensitic steels niobium lowers the hardness and increases the tempering resistance. In U.S. it is also referred to as Columbium (Cb).

Titanium

Titanium Ti, is used to retard grain growth and thus improve toughness. Titanium is also used to achieve improvements in inclusion characteristics. Titanium causes sulfide inclusions to be globular rather than elongated thus improving toughness and ductility in transverse bending.

Titanium is a strong ferrite former and a strong carbide former, thus lowering the effective carbon content and promoting a ferritic structure in two ways. In austenitic steels it is added to increase the resistance to intergranular corrosion but it also increases the mechanical properties at high temperatures. In ferritic stainless steels titanium is added to improve toughness and corrosion resistance by lowering the amount of interstitials in solid solution. In martensitic steels titanium lowers the martensite hardness and increases the tempering resistance. In precipitation hardening steels titanium is used to form the intermetallic compounds that are used to increase the strength.

The principal use of titanium is to stabilize carbon by forming titanium carbide. In austenitic stainless steels it is used in this way to prevent inter crystalline corrosion, the titanium addition being at least four times the carbon content. It is also added to low carbon steels to prevent blistering during vitreous enameling. Titanium carbide is used with tungsten carbide in the manufacture of hard metal tools.

Cobalt

Cobalt Co, Increases strength and hardness, permits quenching at higher temperatures. In some steels used for nuclear engineering cobalt is an undesirable impurity, even in amounts as small as 0·02%. Unlike most other alloying elements cobalt reduces hardenability. It raises the red hardness of steel and this is the reason for adding 5% to 10% cobalt to certain types of high speed steels, developed for the specific purpose of cutting exceptionally hard materials. Heat resisting alloys with high cobalt contents have been developed for use in gas turbines. Cobalt is added to the extent of up to 40 % to magnet steels requiring high coercive force and it is used in electrical- resistance alloys. In the sintered hard metals Cobalt acts as the binding metal.

Cobalt only used as an alloying element in martensitic steels where it increases the hardness and tempering resistance, especially at higher temperatures.

Cobalt has a high solubility in alpha- and gama-iron but a weak carbide-forming tendency. It decreases

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Strong carbide forming element. Not used on its own, but added as a carbide stabiliser to some austenitic stainless steels.


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hardenability but sustains hardness during tempering. It is used in "Stellite" type alloys, gas turbine steel, magnets and as a bond in hard metal.

Tellurium

Tellurium (Te) Tellurium is added to steel either alone or together with selenium to promote machinability. It is a powerful carbide stabilizer and has been also added to cast iron where it is said to increase the depth of chill and to prevent shrinkage. It may be added in small amounts to the molten iron or by the use of cores dipped or painted with washes containing tellurium in suspension.


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Chapter Three: Articles on alloying elements in steel

Carbon content, steel classifications, and alloy steels

Generally, carbon is the most important commercial steel alloy. Increasing carbon content increases hardness and strength and improves hardenability. But carbon also increases brittleness and reduces weldability because of its tendency to form martensite. This means carbon content can be both a blessing and a curse when it comes to commercial steel.

And while there are steels that have up to 2 percent carbon content, they are the exception. Most steel contains less than 0.35 percent carbon. To put this in perspective, keep in mind that’s 35/100 of 1 percent.

Now, any steel in the 0.35 to 1.86 percent carbon content range can be hardened using a heat-quench-temper cycle. Most commercial steels are classified into one of three groups:

1. Plain carbon steels

2. Low-alloy steels

3. High-alloy steels

Plain Carbon Steels

These steels usually are iron with less than 1 percent carbon, plus small amounts of manganese, phosphorus, sulfur, and silicon. The weldability and other characteristics of these steels are primarily a product of carbon content, although the alloying and residual elements do have a minor influence.

Plain carbon steels are further subdivided into four groups:

1. Low

2. Medium

3. High

4. Very high

Low. Often called mild steels, low-carbon steels have less than 0.30 percent carbon and are the most commonly used grades. They machine and weld nicely and are more ductile than higher-carbon steels.

Medium. Medium-carbon steels have from 0.30 to 0.45 percent carbon. Increased carbon means increased hardness and tensile strength, decreased ductility, and more difficult machining.

High. With 0.45 to 0.75 percent carbon, these steels can be challenging to weld. Preheating, postheating (to control cooling rate), and sometimes even heating during welding become necessary to produce acceptable welds and to control the mechanical properties of the steel after welding.

Very High. With up to 1.50 percent carbon content, very high-carbon steels are used for hard steel products such as metal cutting tools and truck springs. Like high-carbon steels, they require heat treating before, during, and after welding to maintain their mechanical properties.

Low-alloy Steels


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When these steels are designed for welded applications, their carbon content is usually below 0.25 percent and often below 0.15 percent. Typical alloys include nickel, chromium, molybdenum, manganese, and silicon, which add strength at room temperatures and increase low-temperature notch toughness.

These alloys can, in the right combination, improve corrosion resistance and influence the steel’s response to heat treatment. But the alloys added can also negatively influence crack susceptibility, so it’s a good idea to use low-hydrogen welding processes with them. Preheating might also prove necessary. This can be determined by using the carbon equivalent formula, which we’ll cover in a later issue.

High-alloy Steels

For the most part, we’re talking about stainless steel here, the most important commercial high-alloy steel. Stainless steels are at least 12 percent chromium and many have high nickel contents. The three basic types of stainless are:

1. Austenitic

2. Ferritic

3. Martensitic

Martensitic stainless steels make up the cutlery grades. They have the least amount of chromium, offer high hardenability, and require both pre- and postheating when welding to prevent cracking in the heat-affected zone (HAZ).

Ferritic stainless steels have 12 to 27 percent chromium with small amounts of austenite-forming alloys.

Austenitic stainless steels offer excellent weldability, but austenite isn’t stable at room temperature. Consequently, specific alloys must be added to stabilize austenite. The most important austenite stabilizer is nickel, and others include carbon, manganese, and nitrogen. Special properties, including corrosion resistance, oxidation resistance, and strength at high temperatures, can be incorporated into austenitic stainless steels by adding certain alloys like chromium, nickel, molybdenum, nitrogen, titanium, and columbium. And while carbon can add strength at high temperatures, it can also reduce corrosion resistance by forming a compound with chromium. It’s important to note that austenitic alloys can’t be hardened by heat treatment. That means they don’t harden in the welding HAZ.

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* Stainless steels always have a high chromium content, often considerable amounts of nickel, and sometimes contain

molybdenum and other elements. Stainless steels are identified by a three-digit number beginning with 2, 3, 4, or 5.

Figure 1

Be sure to check the appropriate AISI and SAE publications for the latest revisions.

Steel Classification Systems

Before we look at a couple of common steel classification systems, let’s consider one more high-carbon metal, cast iron. The carbon content of cast iron is 2.1 percent or more. There are four basic types of cast iron:

1. Gray cast iron, which is relatively soft. It’s easily machined and welded, and you’ll find it used for engine cylinder blocks, pipe, and machine tool structures.

2. White cast iron, which is hard, brittle, and not weldable. It has a compressive strength of more than 200.000 pounds per square inch (PSI), and when it’s annealed, it becomes malleable cast iron.


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3. Malleable cast iron, which is annealed white cast iron. It can be welded, machined, is ductile, and offers good strength and shock resistance.

4. Ductile cast iron, which is sometimes called nodular or spheroidal graphite cast iron. It gets this name because its carbon is in the shape of small spheres, not flakes. This makes it both ductile and malleable. It’s also weldable.

Now let’s take a look at a typical steel classification system (see Figure 1). Both the Society of Automotive Engineers (SAE) and the American Iron and Steel Institute (AISI) use virtually identical systems. Both are based on a four-digit system with the first number usually indicating the basic type of steel and the first two numbers together indicating the series within the basic alloy group.

Keep in mind there may be a number of series within a basic alloy group, depending on the amount of the principal alloying elements. The last two or three numbers refer to the approximate permissible range of carbon content in points (hundredths of a percent).

These classification systems can become fairly complex, and Figure 1 is just a basic representation. Be sure to reference the most recent AISI and SAE publications for the latest revisions.

That’s a look at some basics concerning the iron-carbon-steel relationship and its influences on welding and metal alloys. Next time we’ll look at hardening and ways to make metals stronger. We’ll also consider the influences of some key alloying elements and the effects of welding on metallurgy.

Carbon Steel

Carbon steels and alloy steels are designated by a four digit number, where the first two digits indicate the alloying elements and the last two digits indicate the amount of carbon, in hundredths of a percent by weight. For example, a 1060 steel is a plain carbon steel containing 0.60 wt% C.

designation Type

Carbon steels

10xx Plain carbon (Mn 1.00% max)

11xx Resulfurized

12xx Resulfurized and rephosphorized

15xx Plain carbon (Mn 1.00% to 1.65%)

Manganese steels

13xx Mn 1.75%

Nickel steels

23xx Ni 3.50%

25xx Ni 5.00%

Nickel-chromium steels

31xx Ni 1.25%, Cr 0.65% or 0.80%


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32xx Ni 1.25%, Cr 1.07%

33xx Ni 3.50%, Cr 1.50% or 1.57%

34xx Ni 3.00%, Cr 0.77%

Molybdenum steels

40xx Mo 0.20% or 0.25% or 0.25% Mo & 0.042 S

44xx Mo 0.40% or 0.52%

Chromium-molybdenum (Chromoly) steels

41xx Cr 0.50% or 0.80% or 0.95%, Mo 0.12% or 0.20% or 0.25% or 0.30%

Nickel-chromium-molybdenum steels

43xx Ni 1.82%, Cr 0.50% to 0.80%, Mo 0.25%

43BVxx Ni 1.82%, Cr 0.50%, Mo 0.12% or 0.35%, V 0.03% min

47xx Ni 1.05%, Cr 0.45%, Mo 0.20% or 0.35%

81xx Ni 0.30%, Cr 0.40%, Mo 0.12%

81Bxx Ni 0.30%, Cr 0.45%, Mo 0.12%

86xx Ni 0.55%, Cr 0.50%, Mo 0.20%

87xx Ni 0.55%, Cr 0.50%, Mo 0.25%

88xx Ni 0.55%, Cr 0.50%, Mo 0.35%

93xx Ni 3.25%, Cr 1.20%, Mo 0.12%

94xx Ni 0.45%, Cr 0.40%, Mo 0.12%

97xx Ni 0.55%, Cr 0.20%, Mo 0.20%

98xx Ni 1.00%, Cr 0.80%, Mo 0.25%

Nickel-molybdenum steels

46xx Ni 0.85% or 1.82%, Mo 0.20% or 0.25%

48xx Ni 3.50%, Mo 0.25%

Chromium steels

50xx Cr 0.27% or 0.40% or 0.50% or 0.65%

50xxx Cr 0.50%, C 1.00% min

50Bxx Cr 0.28% or 0.50%

51xx Cr 0.80% or 0.87% or 0.92% or 1.00% or 1.05%

51xxx Cr 1.02%, C 1.00% min

51Bxx Cr 0.80%

52xxx Cr 1.45%, C 1.00% min

Chromium-vanadium steels


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61xx Cr 0.60% or 0.80% or 0.95%, V 0.10% or 0.15% min

Tungsten-chromium steels

72xx W 1.75%, Cr 0.75%

Silicon-manganese steels

92xx Si 1.40% or 2.00%, Mn 0.65% or 0.82% or 0.85%, Cr 0.00% or 0.65%

High-strength low-alloy steels

9xx Various SAE grades

xxBxx Boron steels

xxLxx Leaded steels

Stainless steel

200 Series: austenitic chromium-nickel-manganese alloys 300 Series: austenitic chromium-nickel alloys

Type 301: highly ductile, for formed products. Also hardens rapidly during mechanical working. Type 303: free machining version of 304 via addition of sulfur Type 304: the most common; the classic 18/8 stainless steel. Type 316: the next most common; for food and surgical stainless steel uses; alloy addition of

molybdenum prevents specific forms of corrosion. 316 steel is more resistant to corrosion than 18-8 stainless steels. 316 steel is used in the handling of certain food and pharmaceutical products where it is often required in order to minimize metallic contamination. 316 steel is also known as "marine grade" stainless steel due to its increased ability to resist saltwater corrosion compared to type 304. SS316 is often used for building nuclear reprocessing plants.

400 Series: ferritic and martensitic chromium alloys Type 408: heat-resistant; poor corrosion resistance; 11% chromium, 8% nickel. Type 409: cheapest type; used for automobile exhausts; ferritic (iron/chromium only). Type 410: martensitic (high-strength iron/chromium). Type 416: the most machinable stainless steel; achieved by the addition of extra sulfur which

reduces corrosion resistance. Often used for "stainless" rifle barrels Type 420: "Cutlery grade" martensitic; similar to the Brearley's original "rustless steel". Also

known as "surgical steel". Type 430: decorative, e.g., for automotive trim; ferritic. Type 440: a higher grade of cutlery steel, with more carbon in it, which allows for much better

edge retention when the steel is heat treated properly. 500 Series: heat resisting chromium alloys 600 Series: martensitic precipitation hardening alloys

Type 630: most common PH stainless, better known as 17-4; 17% chromium, 4% nickel


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Stainless steel designations

SAE

designation

UNS

designation % Cr % Ni % C % Mn % Si % P % S % N Other

Austenitic

201 S20100 16–18 3.5–5.5 0.15 5.5–7.5 0.75 0.06 0.03 0.25 -

202 S20200 17–19 4–6 0.15 7.5–10.0 0.75 0.06 0.03 0.25 -

205 S20500 16.5–18 1–1.75 0.12–0.25 14–15.5 0.75 0.06 0.03 0.32–0.40 -

301 S30100 16–18 6–8 0.15 2 0.75 0.045 0.03 - -

302 S30200 17–19 8–10 0.15 2 0.75 0.045 0.03 0.1 -

302B S30215 17–19 8–10 0.15 2 2.0–3.0 0.045 0.03 - -

303 S30300 17–19 8–10 0.15 2 1 0.2 0.15

min -

Mo 0.60

(optional)

303Se S30323 17–19 8–10 0.15 2 1 0.2 0.06 - 0.15 Se min

304 S30400 18–20 8–10.50 0.08 2 0.75 0.045 0.03 0.1 -

304L S30403 18–20 8–12 0.03 2 0.75 0.045 0.03 0.1 -

304Cu S30430 17–19 8–10 0.08 2 0.75 0.045 0.03 - 3–4 Cu

304N S30451 18–20 8–10.50 0.08 2 0.75 0.045 0.03 0.10–0.16 -

305 S30500 17–19 10.50–13 0.12 2 0.75 0.045 0.03 - -

308 S30800 19–21 10–12 0.08 2 1 0.045 0.03 - -

309 S30900 22–24 12–15 0.2 2 1 0.045 0.03 - -

309S S30908 22–24 12–15 0.08 2 1 0.045 0.03 - -

310 S31000 24–26 19–22 0.25 2 1.5 0.045 0.03 - -

310S S31008 24–26 19–22 0.08 2 1.5 0.045 0.03 - -

314 S31400 23–26 19–22 0.25 2 1.5–3.0 0.045 0.03 - -

316 S31600 16–18 10–14 0.08 2 0.75 0.045 0.03 0.10 2.0–3.0 Mo

316L S31603 16–18 10–14 0.03 2 0.75 0.045 0.03 0.10 2.0–3.0 Mo

316F S31620 16–18 10–14 0.08 2 1 0.2 0.10

min - 1.75–2.50 Mo

316N S31651 16–18 10–14 0.08 2 0.75 0.045 0.03 0.10–0.16 2.0–3.0 Mo

317 S31700 18–20 11–15 0.08 2 0.75 0.045 0.03 0.10 max 3.0–4.0 Mo

317L S31703 18–20 11–15 0.03 2 0.75 0.045 0.03 0.10 max 3.0–4.0 Mo

321 S32100 17–19 9–12 0.08 2 0.75 0.045 0.03 0.10 max Ti 5(C+N)

min, 0.70 max

329 S32900 23–28 2.5–5 0.08 2 0.75 0.04 0.03 - 1–2 Mo


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330 N08330 17–20 34–37 0.08 2 0.75–1.50 0.04 0.03 - -

347 S34700 17–19 9–13 0.08 2 0.75 0.045 0.030 - Nb + Ta, 10 x

C min, 1 max

348 S34800 17–19 9–13 0.08 2 0.75 0.045 0.030 -

Nb + Ta, 10 x

C min, 1 max,

but 0.10 Ta

max; 0.20 Ca

384 S38400 15–17 17–19 0.08 2 1 0.045 0.03 - -

Ferritic

405 S40500 11.5–14.5 - 0.08 1 1 0.04 0.03 - 0.1–0.3 Al,

0.60 max

409 S40900 10.5–11.75 0.05 0.08 1 1 0.045 0.03 - Ti 6 x C, but

0.75 max

429 S42900 14–16 0.75 0.12 1 1 0.04 0.03 - -

430 S43000 16–18 0.75 0.12 1 1 0.04 0.03 - -

430F S43020 16–18 - 0.12 1.25 1 0.060.15

min -

0.60 Mo

(optional)

430FSe S43023 16–18 - 0.12 1.25 1 0.06 0.06 - 0.15 Se min

434 S43400 16–18 - 0.12 1 1 0.04 0.03 - 0.75–1.25 Mo

436 S43600 16–18 - 0.12 1 1 0.04 0.03 -

0.75–1.25

Mo; Nb+Ta 5

x C min, 0.70

max

442 S44200 18–23 - 0.2 1 1 0.04 0.03 - -

446 S44600 23–27 0.25 0.2 1.5 1 0.04 0.03 - -

Martensitic

403 S40300 11.5–13.0 0.60 0.15 1 0.5 0.04 0.03 - -

410 S41000 11.5–13.5 0.75 0.15 1 1 0.04 0.03 - -

414 S41400 11.5–13.5 1.25–2.50 0.15 1 1 0.04 0.03 - -

416 S41600 12–14 - 0.15 1.25 1 0.060.15

min -

0.060 Mo

(optional)

416Se S41623 12–14 - 0.15 1.25 1 0.06 0.06 - 0.15 Se min

420 S42000 12–14 - 0.15 min 1 1 0.04 0.03 - -

420F S42020 12–14 - 0.15 min 1.25 1 0.060.15

min -

0.60 Mo max

(optional)

422 S42200 11.0–12.5 0.50–1.0 0.20–0.25 0.5–1.0 0.5 0.025 0.025 - 0.90–1.25

Mo;


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0.20–0.30 V;

0.90–1.25 W

431 S41623 15–17 1.25–2.50 0.2 1 1 0.04 0.03 - -

440A S44002 16–18 - 0.60–0.75 1 1 0.04 0.03 - 0.75 Mo

440B S44003 16–18 - 0.75–0.95 1 1 0.04 0.03 - 0.75 Mo

440C S44004 16–18 - 0.95–1.20 1 1 0.04 0.03 - 0.75 Mo

Heat resisting

501 S50100 4–6 - 0.10 min 1 1 0.04 0.03 - 0.40–0.65 Mo

502 S50200 4–6 - 0.1 1 1 0.04 0.03 - 0.40–0.65 Mo

Unified numbering system

Introduction to the Unified Numbering System of Ferrous Metals and Alloys http://www.key-to-steel.com/Articles/Art111.htm http://en.wikipedia.org/wiki/Unified_numbering_system

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Introduction to steels and cast irons

http://mse.iastate.edu/solidification/Courses/MatE443-Website/instructions/ASM_metallography_guide.pdf

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Page 103 of 180

Strength in Steel

Strength in steels arises from several phenomena, which usually contribute collectively to the observed mechanical properties. The heat treatment of steels is aimed at adjusting these contributions so that the required balance of mechanical properties is achieved. Fortunately the γ/α phase change allows great variations in microstructure to be produced, so that a wide range of mechanical properties can be obtained even in plain carbon steels. The additional use of metallic alloying elements, primarily as a result of their influence on the transformation, provides an even greater control over microstructure, with consequent benefits in the mechanical properties The simplest version of analyzes the effects of alloying elements on iron-carbon alloys would require analysis of a large number of ternary alloy diagrams over a wide temperature range. However, Wever pointed out that iron binary equilibrium systems fall into four main categories (Fig. 1): open and closed γ-field systems, and expanded and contracted γ-field systems. This approach indicates that alloying elements can influence the equilibrium diagram in two ways: by expanding the γ-field, and encouraging the formation of austenite over wider compositional limits. These elements are called γ-stabilizers. by contracting the γ-field, and encouraging the formation of ferrite over wider compositional limits. These elements are called α-stabilizers. The form of the diagram depends to some degree on the electronic structure of the alloying elements which is reflected in their relative positions in the periodic classification. Class 1: open γ-field. To this group belong the important steel alloying elements nickel and manganese, as well as cobalt and the inert metals ruthenium, rhodium, palladium, osmium, iridium and platinum. Both nickel and manganese, if added in sufficiently high concentration, completely eliminate the bcc α-iron phase and replace it, down to room temperature, with the γ-phase. So nickel and manganese depress the phase transformation from γ to α to lower temperatures (Fig. 1a), i.e. both Ac1 and Ac3 are lowered. It is also easier to obtain metastable austenite by quenching from the γ-region to room temperature, consequently nickel and manganese are useful elements in the formulation of austenitic steels. Class 2: expanded γ-field. Carbon and nitrogen are the most important elements in this group. The γ-phase field is expanded, but its range of existence is cut short by compound formation (Fig.1b). Copper, zinc and gold have a similar influence. The expansion of the γ-field by carbon, and nitrogen, underlies the whole of the heat treatment of steels, by allowing formation of a homogeneous solid solution (austenite) containing up to 2.0 wt % of carbon or 2.8 wt % of nitrogen. Class 3: closed γ-field. Many elements restrict the formation of γ-iron, causing the γ-area of the diagram to contract to a small area referred to as the gamma loop (Fig. 1c). This means that the relevant elements are encouraging the formation of bcc iron (ferrite), and one result is that the δ- and γ-phase fields become continuous. Alloys in which this has taken place are, therefore, not amenable to the normal heat treatments involving cooling through the γ/α-phase transformation. Silicon, aluminium, beryllium and phosphorus fall into this category, together with the strong carbide forming elements, titanium, vanadium, molybdenum and


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chromium.

Figure 1. Classification of iron alloy phase diagrams: a. open γ-field; b. expanded γ-field; c. closed γ-field Class 4: contracted y-field. Boron is the most significant element of this group, together with the carbide forming elements tantalum, niobium and zirconium. The γ-loop is strongly contracted, but is accompanied by compound formation (Fig. 1d).


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The distribution of alloying elements in steels. Although only binary systems have been considered so far, when carbon is included to make ternary systems the same general principles usually apply. For a fixed carbon content, as the alloying clement is added the y-field is either expanded or contracted depending on the particular solute. With an element such as silicon the γ-field is restricted and there is a corresponding enlargement of the α-field. If vanadium is added, the γ-field is contracted and there will be vanadium carbide in equilibrium with ferrite over much of the ferrite field. Nickel does not form a carbide and expands the γ-field. Normally elements with opposing tendencies will cancel each other out at the appropriate combinations, but in some cases anomalies occur. For example, chromium added to nickel in a steel in concentrations around 18% helps to stabilize the γ-phase, as shown by 18Cr8Ni austenitic steels. One convenient way of illustrating quantitatively the effect of an alloying element on the γ-phase field of the Fe-C system is to project on to the Fe-C plane of the ternary system the γ-phase field boundaries for increasing concentration of a particular alloying element. For more precise and extensive information, it is necessary to consider series of isothermal sections in true ternary systems Fe-C-X, but even in some of the more familiar systems the full information is not available, partly because the acquisition of accurate data can be a difficult and very time-consuming process. Recently the introduction of computer-based methods has permitted the synthesis of extensive thermochemical and phase equilibria data, and its presentation in the form, for example, of isothermal sections over a wide range of temperatures. If only steels in which the austenite transforms to ferrite and carbide on slow cooling are considered, the alloying elements can be divided into three categories: elements which enter only the ferrite phase elements which form stable carbides and also enter the ferrite phase elements which enter only the carbide phase. In the first category there are elements such as nickel, copper, phosphorus and silicon which, in transformable steels, are normally found in solid solution in the ferrite phase, their solubility in cementite or in alloy carbides being quite low. The majority of alloying elements used in steels fall into the second category, in so far as they are carbide formers and as such, at low concentrations, go into solid solution in cementite, but will also form solid solutions in ferrite. At higher concentrations most will form alloy carbides, which are thermodynamically more stable than cementite. Typical examples are manganese, chromium, molybdenum, vanadium, titanium, tungsten and niobium. Manganese carbide is not found in steels, but instead manganese enters readily into solid solution in Fe3C. The carbide-forming elements are usually present greatly in excess of the amounts needed in the carbide phase, which are determined primarily by the carbon content of the steel. The remainder enters into solid solution in the ferrite with the non-carbide forming elements nickel and silicon. Some of these elements, notably titanium, tungsten, and molybdenum, produce substantial solid solution hardening of ferrite. In the third category there are a few elements which enter predominantly the carbide phase. Nitrogen is the most important element and it forms carbo-nitrides with iron and many alloying elements. However, in the presence of certain very strong nitride forming elements, e.g. titanium and aluminum, separate alloy nitride phases can occur. While ternary phase diagrams, Fe-C-X, can be particularly helpful in understanding the phases which can exist in simple steels, isothermal sections for a number of temperatures are needed before an adequate picture of the equilibrium phases can be built up. For more complex steels the task is formidable and equilibrium diagrams can only give a rough guide to the structures likely to be encountered. It is, however, possible to construct pseudobinary diagrams for groups of steels, which give an overall view of the equilibrium phases likely to be


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encountered at a particular temperature. Structural changes resulting from alloying additions. The addition to iron-carbon alloys of elements such as nickel, silicon, manganese, which do not form carbides in competition with cementite, does not basically alter the microstructures formed after transformation. However, in the case of strong carbide-forming elements such as molybdenum, chromium and tungsten, cementite will be replaced by the appropriate alloy carbides, often at relatively low alloying element concentrations. Still stronger carbide forming elements such as niobium, titanium and vanadium are capable of forming alloy carbides, preferentially at alloying concentrations less than 0.1 wt%. It would, therefore, be expected that the microstructures of steels containing these elements would be radically altered. It has been shown how the difference in solubility of carbon in austenite and ferrite leads to the familiar ferrite/cementite aggregates in plain carbon steels. This means that, because the solubility of cementite in austenite is much greater than in ferrite, it is possible to redistribute the cementite by holding the steel in the austenite region to take it into solution, and then allowing transformation to take place to ferrite and cementite. Examining the possible alloy carbides, and nitrides, in the same way, shows that all the familiar ones are much less soluble in austenite than is cementite. Chromium and molybdenum carbides are not included, but they are substantially more soluble in austenite than the other carbides. Detailed consideration of such data, together with practical knowledge of alloy steel behavior, indicates that, for niobium and titanium, concentrations of greater than about 0.25 wt % will form excess alloy carbides which cannot be dissolved in austenite at the highest solution temperatures. With vanadium the limit is higher at 1-2%, and with molybdenum up to about 5%. Chromium has a much higher limit before complete solution of chromium carbide in austenite becomes difficult. This argument assumes that sufficient carbon is present in the steel to combine with the alloying element. If not, the excess metallic element will go into solid solution both in the austenite and the ferrite. In general, the fibrous morphology represents a closer approach to an equilibrium structure so it is more predominant in steels which have transformed slowly. In contrast, the interphase precipitation and dislocation nucleated structures occur more readily in rapidly transforming steels, where there is a high driving force, for example, in microalloyed steels. The clearest analogy with pearlite is found when the alloy carbide in lath morphology forms nodules in association with ferrite. These pearlitic nodules are often encountered at temperatures just below Ac1, in steels which transform relatively slowly. For example, these structures are obtained in chromium steels with between 4% and 12% chromium and the crystallography is analogous to that of cementitic pearlite. It is, however, different in detail because of the different crystal structures of the possible carbides. The structures observed are relatively coarse, but finer than pearlite formed under equivalent conditions, because of the need for the partition of the alloying element, e.g. chromium between the carbide and the ferrite. To achieve this, the interlamellar spacing must be substantially finer than in the equivalent iron-carbon case. Interphase precipitation. Interphase precipitation has been shown to nucleate periodically at the γ/α interface during the transformation. The precipitate particles form in bands which are closely parallel to the interface, and which follow the general direction of the interface even when it changes direction sharply. A further characteristic is the frequent development of only one of the possible Widmanstätten variants, for example VC plates in a particular region are all only of one variant of the habit, i.e. that in which the plates are most nearly parallel to the interface. The extremely fine scale of this phenomenon in vanadium steels, which also occurs in Ti and Nb steels, is due to the rapid rate at which the γ/α transformation takes place. At the higher transformation temperatures, the slower rate of reaction leads to coarser structures. Similarly, if the reaction is slowed down by addition of further alloying


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elements, e.g. Ni and Mn, the precipitate dispersion coarsens. The scale of the dispersion also varies from steel to steel, being coarsest in chromium, tungsten and molybdenum steels where the reaction is relatively slow, and much finer in steels in which vanadium, niobium and titanium are the dominant alloying elements and the transformation is rapid. Transformation diagrams for alloy steels. The transformation of austenite below the eutectoid temperature can best be presented in an isothermal transformation diagram, in which the beginning and end of transformation is plotted as a function of temperature and time. Such curves are known as time-temperature-transformation, or TTT curves, and form one of the important sources of quantitative information for the heat treatment of steels. In the simple case of a eutectoid plain carbon steel, the curve is roughly C-shaped with the pearlite reaction occurring down to the nose of the curve and a little beyond. At lower temperatures bainite and martensite are formed. The diagrams become more complex for hypo- and hyper-eutectoid alloys as the ferrite or cementite reactions have also to be represented by additional lines


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Alloying and Its Effects on the Critical Temperature, Hardness

and Tensile Strength

Alloying elements have significant effect on the iron-iron carbide equilibrium diagram. The addition of some of these alloying elements will widen the temperature range through which austenite (gamma -iron) is stable while other elements will constrict the temperature range. What this means is that some elements will raise and some elements will lower the critical temperature of steel.

Manganese, cobalt, and nickel increase the temperature range through which austenite is stable. This also means that the lower critical temperature of steel will be lowered by these alloying elements. Other alloying elements that lower the critical temperature of steel are carbon, copper and zinc. The alloying elements that are used to reduce the critical temperature are highly soluble in the gamma iron (austenite). Figure 1 shows the effect of manganese on the critical temperature of steel.

Figure 1. The effect of alloying with manganese on the critical temperature of steel and austenite (γ-iron) phase transformation zone on the iron-iron carbide diagram..

Alloys such as aluminum, chromium, molybdenum, phosphorus, silicon, tungsten tend to form solid solutions with alpha iron (ferrite). This constricts the temperature region through which gamma iron (austenite) is stable. As shown in Figure 2, chromium at different percentages constricts the critical temperature range which results in a marked reduction of the region where austenite is stable.


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Figure 2. Effect of alloying with chromium on the critical temperature of steel and austenite (g -iron) phase transformation zone on the iron-iron carbide diagram.

The elements shown in Figure 3 have the greatest solubility in ferrite and also influence the hardenability of iron when in the presence of carbon. With a slight increase in the carbon content, they respond markedly to heat treating, because carbon acts as a ferrite strengthener. As indicated in Figure 3, Phosphorus will improve the hardness of the ferrite significantly by adding only a very small percentage of Phosphorus, while Chromium will not strengthen the ferrite that well even at very high percentage of Chromium addition to the steel


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Figure 3. The effect of various alloying elements on the hardness of steel.

Figure 4 shows the effect of furnace cooling vs. air cooling on the tensile strength of steel for three different percentages of carbon in the presence of chromium. As this figure indicates, furnace cooling has very little effect on the tensile strength of the material. The addition of chromium does not change the tensile strength properties when the steel is cooled in the furnace. If the same steels are air cooled at the same rate, the slope of the curves increases significantly which means that a slight increase in the chromium content increases the strength drastically when air cooling is applied.

Figure 4. Effect of different percentages of carbon on the tensile strength of steel in the presence of chromium.


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Control of High Strength Low Alloy (HSLA) Steel Properties

Most HSLA steels are furnished in the as-hot-rolled condition with ferritic-pearlitic microstructure. The exceptions are the controlled-rolled steels with an acicular ferrite microstructure and the dual-phase steels with martensite dispersed in a matrix of polygonal ferrite. These two types of HSLA steels use the formation of eutectoid structures for strengthening, while the ferritic-pearlitic HSLA steels generally require strengthening of the ferrite. Pearlite is generally an undesirable strengthening agent in structural steels because it reduces impact toughness and requires higher carbon contents. Moreover, yield strength is largely unaffected by a higher pearlite content. Strengthening Mechanisms in Ferrite The ferrite in HSLA steels is typically strengthened by grain refinement, precipitation hardening, and, to a lesser extent, solid-solution strengthening. Grain refinement is the most desirable strengthening mechanism because it improves not only strength but also toughness. Grain refinement is influenced by the complex effects of alloy design and processing methods. For example, the various methods of grain refinement used in the three different stages of hot rolling (that is, reheating, hot rolling, and cooling) include: The addition of titanium or aluminum to retard austenite grain growth when the steel is reheated for hot deformation or subsequent heat treatment The controlled rolling of microalloyed steels to condition the austenite so that it transforms into fine-grain ferrite The use of alloy additions and/or faster cooling rates to lower the austenite-to-ferrite transformation temperature. The use of higher cooling rates for grain refinement may require consideration of its effect on precipitation strengthening and the possibility of undesirable transformation products. Precipitation strengthening occurs from the formation of finely dispersed carbonitrides developed during heating and cooling. Because precipitation strengthening is generally associated with a reduction in toughness, grain refinement is often used in conjunction with precipitation strengthening to improve toughness. Precipitation strengthening is influenced by the type of carbonitride, its grain size, and, of course, the number of carbonitrides precipitated. The formation of MC is the most effective metal carbide in the precipitation strengthening of microalloyed niobium, vanadium, and/or titanium steels. The number of fine MC particles formed during heating and cooling depends on the solubility of the carbides in austenite and on cooling rates. Steelmaking Precise steelmaking operations are also essential in controlling the properties and chemistry of HSLA steels. Optimum property levels depend on such factors as the control of significant alloying elements and the reduction of impurities and nonmetallic inclusions. Developments in secondary steelmaking such as desulphurization, vacuum degassing, and argon shrouding have enabled better control of steel chemistry and the effective use of microalloyed elements. Compositional limits for HSLA steel grades described in ASTM specifications the use of vacuum degassing equipment allows the production of interstitial-free (IF) steels. The IF steels exhibit excellent formability, high elongation, and good deep draw/ability. Compositions and Alloying Elements Chemical compositions for the HSLA steels are specified by ASTM standards. The principal function of alloying elements in these ferrite-pearlite HSLA steels, other than corrosion resistance, is strengthening of the ferrite by grain refinement, precipitation strengthening, and solid-solution strengthening. Solid-solution strengthening is


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closely related to alloy contents, while grain refinement and precipitation strengthening depend on the complex effects of alloy design and thermo-mechanical treatment. Alloying elements are also selected to influence transformation temperatures so that the transformation of austenite to ferrite and pearlite occurs at a lower temperature during air cooling. This lowering of the transformation temperature produces a finer-grain transformation product, which is a major source of strengthening. At the low carbon levels typical of HSLA steels, elements such as silicon, copper, nickel, and phosphorus are particularly effective for producing fine pearlite. Element such as, manganese and chromium, which are present in both the cementite and ferrite, also strengthen the ferrite by solid-solution strengthening in proportion to the amount, dissolved in the ferrite. In the presence of alloying elements, the practical maximum carbon content at which HSLA steels can be used in the as-cooled condition is approximately 0.20%. Higher levels of carbon tend to form martensite or bainite in the microstructure of as-rolled steels, although some of the higher-strength low-alloy steels have carbon contents that approach 0.30%. The required strength is developed by the combined effect of: Fine grain size developed during controlled hot roiling and enhanced by microalloyed elements (especially niobium) Precipitation strengthening caused by the presence of vanadium, niobium, and titanium in the composition. Nitrogen additions to high-strength steels containing vanadium are limited to 0.005% and have become commercially important because such additions enhance precipitation hardening. The precipitation of vanadium nitride in vanadium-nitrogen steels also improves grain refinement because it has a lower solubility in austenite than vanadium carbide. Manganese is the principal strengthening element in plain carbon high-strength structural steels. It functions mainly as a mild solid-solution strengthener in ferrite, but it also provides a marked decrease in the austenite-to-ferrite transformation temperature. In addition, manganese can enhance the precipitation strengthening of vanadium steels and. to a lesser extent, niobium steels. One of the most important applications of silicon is its use as a deoxidizer in molten steel. Silicon has a strengthening effect in low-alloy structural steels. In larger amounts, it increases resistance to scaling at elevated temperatures. Silicon has a significant effect on yield strength enhancement by solid-solution strengthening and is widely used in HSLA steels for riveted or bolted structures. Copper in levels in excess of 0.50% also increases the strength of both low- and medium-carbon steels by virtue of ferrite strengthening, which is accompanied by only slight decreases in ductility. Copper can be retained in solid solution even at the slow rate of cooling obtained when large sections are normalized, but it is precipitated out when the steel is reheated to about 510 to 605°C (950 to 1125°F). At about 1% copper, the yield strength is increased by about 70 to 140 MPa regardless of the effects of other alloying elements. Copper in amounts up to 0.75% is considered to have only minor adverse effects on notch toughness or weldability. Copper precipitation hardening gives the steel the ability to be formed extensively and then precipitation hardened as a complex shape or welded assembly. The atmospheric-corrosion resistance of steel is increased appreciably by the addition of phosphorus, and when small amounts of copper are present in the steel, the effect of the phosphorus is greatly enhanced. When both phosphorus and copper are present, there is a greater beneficial effect on corrosion resistance than the sum of the effects of the individual elements. Chromium is often, added with copper to obtain improved atmospheric-corrosion resistance. Nickel is often added to copper-bearing steels to minimize hot shortness. Molybdenum in hot-rolled HSLA steels is used primarily to improve hardenability when transformation products


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other than ferrite-pearlite are desired. Molybdenum (0.15 to 0.30%) in microalloyed steels also increases the solubility of niobium in austenite, thereby enhancing the precipitation of NbC(N) in the ferrite. This increases the precipitation-strengthening effect of NbC(N). Aluminum is widely used as a deoxidizer and was the first element used to control austenite grain growth during reheating. During controlled rolling, niobium and titanium are more effective grain refiners than aluminum. Vanadium strengthens HSLA steels by both precipitation hardening the ferrite and refining the ferrite grain size. The precipitation of vanadium carbonitride in ferrite can develop a significant increase in strength that depends not only on the rolling process used, but also on the base composition. Carbon contents above 0.13 to 0.15% and manganese content of 1% or more enhances the precipitation hardening, particularly when the nitrogen content is at least 0.01%. Titanium is unique among common alloying elements in that it provides both precipitation strengthening and sulfide shape control. Small amounts of titanium (<0.025%) are also useful in limiting austenite grain growth. However, it is useful only in fully killed (aluminum deoxidized) steels because of its strong deoxidizing effects, the versatility of titanium is limited because variations in oxygen, nitrogen, and sulfur affect the contribution of titanium as carbide strengthened. Zirconium can also be added to killed high-strength low-alloy steels to improve inclusion characteristics, particularly in the case of sulfide inclusions, for which changes in inclusion shape improve ductility in transverse bending. Boron has no effect on the strength of normal hot-rolled steel but can considerably improve hardenability when transformation products such as acicular ferrite are desired in low-carbon hot-rolled plate. Treatment with calcium is preferred for sulfide inclusion shape control. Controlled Rolling The hot-rolling process has gradually become a much more closely controlled operation, and controlled rolling is now being increasingly applied to microalloyed steels with compositions carefully chosen to provide optimum mechanical properties at room temperature. Controlled rolling is a procedure whereby the various stages of rolling are temperature controlled, with the amount of reduction in each pass predetermined and the finishing temperature precisely defined. This processing is widely used to obtain reliable mechanical properties in steels for pipelines, bridges, offshore platforms, and many other engineering applications. The use of controlled rolling has resulted in improved combinations of strength and toughness and further reductions in the carbon content of microalloyed HSLA steels


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Influence of Alloying Elements on Steel Microstructure

It is a long-standing tradition to discuss the various alloying elements in terms of the properties they confer on steel. For example, the rule was that Chromium (Cr) makes steel hard whereas Nickel (Ni) and Manganese (Mn) make it tough. In saying this, one had certain types of steel in mind and transferred the properties of particular steel to the alloying element that was thought to have the greatest influence on the steel under consideration. This method of reasoning can give false impressions and the following examples will illustrate this point. When we say that Cr makes steel hard and wear-resisting we probably associate this with the 2% C, 12% Cr tool steel grade, which on hardening does in fact become very hard and hard-wearing. But if, on the other hand, we choose a steel containing 0.10% C and 12% Cr, the hardness obtained on hardening is very modest. It is quite true that Mn increases steel toughness if we have in mind the 13% manganese steel, so-called Hadfield steel. In concentrations between l% and 5%, however, Mn can produce a variable effect on the properties of the steel it is alloyed with. The toughness may either increase or decrease. A property of great importance is the ability of alloying elements to promote the formation of a certain phase or to stabilize it. These elements are grouped as austenite-forming, ferrite-forming, carbide-forming and nitride-forming elements. Austenite-forming elements The elements C, Ni and Mn are the most important ones in this group. Sufficiently large amounts of Ni or Mn render a steel austenitic even at room temperature. An example of this is the so-called Hadfield steel which contains 13% Mn, 1,2% Cr and l% C. In this steel both the Mn and C take part in stabilizing the austenite. Another example is austenitic stainless steel containing 18% Cr and 8% Ni. The equilibrium diagram for iron-nickel, Figure 1, shows how the range of stability of austenite increases with increasing Ni-content.


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Figure 1. Fe-Ni equilibrium diagram

More diagrams: http://www.calphad.com/phase_diagrams.html An alloy containing 10% Ni becomes wholly austenitic if heated to 700°C. On cooling, transformation from g to a takes place in the temperature range 700-300°C. Ferrite-forming elements The most important elements in this group are Cr, Si, Mo, W and Al. The range of stability of ferrite in iron-chromium alloys is shown in Figure 2. Fe-Cr alloys in the solid state containing more than 13% Cr are ferritic at all temperatures up to incipient melting. Another instance of ferritic steel is one that is used as transformer sheet material. This is a low-carbon steel containing about 3% Si.


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Figure 2. Cr-Fe equilibrium diagram Multi-alloyed steels The great majority of steels contain at least three components. The constitution of such steels can be deduced from ternary phase diagrams (3 components). The interpretation of these diagrams is relatively difficult and they are of limited value to people dealing with practical heat treatment since they represent equilibrium conditions only. Furthermore, since most alloys contain more than three components it is necessary to look for other ways of assessing the effect produced by the alloying elements on the structural transformations occurring during heat treatment. One approach that is quite good is the use of Schaeffler diagrams (see Figure 3). Here the austenite formers are set out along the ordinate and the ferrite formers along the abscissa. The original diagram contained only Ni and Cr but the modified diagram includes other elements and gives them coefficients that reduce them to the equivalents of Ni or Cr respectively. The diagram holds good for the rates of cooling which result from welding.


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Figure 3. Modified Schaeffler diagram A 12% Cr steel containing 0.3% C is martensitic, the 0.3% C gives the steel a nickel equivalent of 9. An 18/8 steel


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(18% Cr, 8% Ni) is austenitic if it contains 0-0.5% C and 2% Mn. The Ni content of such steels is usually kept between 9% and 10%. Hadfield steel with 13% Mn (mentioned above) is austenitic due to its high carbon content. Should this be reduced to about 0.20% the steel becomes martensitic. Carbide-forming elements Several ferrite formers also function as carbide formers. The majority of carbide formers are also ferrite formers with respect to Fe. The affinity of the elements in the line below for carbon increases from left to right. Cr, W, Mo, V, Ti, Nb, Ta, Zr. Some carbides may be referred to as special carbides, i.e. non-iron-containing carbides, such as Cr7C3 W2C, VC, Mo2C. Double or complex carbides contain both Fe and a carbide-forming element, for example Fe4W2C. High-speed and hot-work tool steels normally contain three types of carbides, which are usually designated M6C, M23C6 and MC. The letter M represents collectively all the metal atoms. Thus M6C represents Fe4W2C or Fe4Mo2C; M23C6 represents Cr23C6 and MC represents VC or V4C3. Carbide stabilizers The stability of the carbides is dependent on the presence of other elements in the steel. How stable the carbides are depends on how the element is partitioned between the cementite and the matrix. The ratio of the percentage, by weight, of the element contained in each of the two phases is called the partition coefficient K. The following values are given for K:

Al Cu P Si Co Ni W Mo Mn Cr Ti Nb Ta

0 0 0 0 0.2 0.3 2 8 11,4 28 Increasing

Note that Mn, which by itself is a very weak carbide former, is a relatively potent carbide stabilizer. In practice, Cr is the alloying element most commonly used as a carbide stabilizer. Malleable cast iron (i.e. white cast iron that is rendered soft by a graphitizing heat treatment called malleablizing) must not contain any Cr. Steel containing only Si or Ni is susceptible to graphitization, but this is most simply prevented by alloying with Cr. Nitride-forming elements All carbide formers are also nitride formers. Nitrogen may be introduced into the surface of the steel by nitriding. By measuring the hardness of various nitrided alloy steels it is possible to investigate the tendency of the different alloying elements to form hard nitrides or to increase the hardness of the steel by a mechanism known as precipitation hardening. The results obtained by such investigations are shown in Figure 4, from which it can be seen that very high hardnesses result from alloying a steel with Al or Ti in amounts of about 1,5%.


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Figure 4. Effect of alloying element additions on hardness after nitriding Base composition: 0.25% C, 0.30% Si, 0.70% Mn On nitriding the base material in Figure 4, hardness of about 400 HV is obtained and according to the diagram the hardness is unchanged if the steel is alloyed with Ni since this element is not a nitride former and hence does not contribute to any hardness increase


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Martensite in Austenitic Stainless Steel Welds

Background

Martensite is a crystal structure that forms in steels during rapid cooling. Cooling rates are dependent on the particular chemistry of the steel. Certain conditions can be met that will cause martensite formation in austenitic stainless steels.

Long, needle-like clusters of crystals in the metal characterize martensite. The martensite crystals have a highly stressed body centered tetragonal structure. In austenitic stainless steels and the austenite phase of a magnetic (austenitic/ferritic) stainless, when martensite is present in a weld, don't expect it to pass the bend test. The weld will not have the toughness and ductility that we normally expect from a stainless steel, and in some circumstances there may be unanticipated corrosion.

Fillet Composition

With only one exception, adding alloying elements to steel allows martensite to form at a slower cooling rate. That exception is cobalt, which works the other way. With cobalt, martensite will form at a faster cooling rate.

We recognize that the actual composition of the weld fillet will vary along a line running across or bisecting the weld fillet. On the centerline, we expect the fillet composition to be closer to the wire composition. Approaching each base metal, the composition shifts from the filler metal composition towards each base metal. In welding metallurgy, we mark on the phase diagram the chemistry of one base metal and the chemistry of the filler wire. We find that the actual fillet compositions lie on the line drawn on the phase diagram. When the base metals are different, it takes different lines on the phase diagram to represent the compositions approaching each base metal.

Predicting Martensite

The results and microstructural consequences of this sort of exercise in physical metallurgy appeared in 1949 as the "Schaeffler Diagram". Our technology did not stand still and the Diagram iterated through several updates. The 1994 Winter Addendum to the ASME Code brought us the Welding Research Council's "WRC-1992 Diagram" which continues to be extensively used. Still, there has been a problem associated with manganese, which brought yet another modification into use. The analysis behind the modification appears, for example, in D.J. Kotecki's, "A martensite boundary on the WRC-1992 diagram" (Welding Journal , Vol. 78, No. 5, pp 180-192).

In lots of cases the 1% manganese line satisfactorily predicts martensite or not martensite. However, we sometimes encounter steels having higher manganese and quite often encounter much lower. High side examples include 1.0-1.5% manganese in a 309L filler wire, 4% in a 307, and 6% in the European 18 8 Mn filler


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wire. On the low side, if you are joining stainless to a modern carbon steel, the manganese can be quite low, perhaps even 0.3%.

(The 1949 Schaeffler diagram - click image to enlarge)

To enter the modified Diagram, we need to calculate two numbers roughly based on chromium and nickel content. We also have to apply a lot of welding "know-how" as to the mixing of the metals.

Calculate the

Nickel Ni Equivalent = %Ni + 35%C + 20%Mn + 0.25%Cu Chrome Cr Equivalent = %Cr + Mo + 1.5%Si + 0.7%Nb

The Amount of Nitrogen

One of the problems we have to also confront is the amount of nitrogen. The best shielding practice with a wire electrode might not introduce nitrogen. Flux cored electrodes tend to add nitrogen. For example, the metal in a flux cored electrode might analyze to 0.05% nitrogen while it deposits as 0.075% nitrogen. Most of us deal with the nitrogen by first plotting the points with only the known nitrogen and then we plot the nearby point based on


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our estimate of the actual nitrogen. The compositions at risk are those on the Diagram and below the indicated manganese brands. Within the bands the diagram is known to be imprecise.

In practice, you have to cope with the range of compositions between one base metal and the filler metal, and between the other base metal and the filler metal. You could put all of your predicted nickel and chromium equivalent compositions on the graph. However, examining the diagram we see that the compositions more at risk for forming martensite are those with low nickel and chromium equivalent numbers, coupled with low manganese. After a little practice with the diagram we know pretty well which compositions are likely to be at risk and we tend to plot only those numbers.

A warning: Notice that all of the alloying elements in the calculation contribute to hardenability. Should either base metal or the filler metal contain any unlisted elements which are known to contribute to hardenability then the Diagram doesn't apply. After decades of absence, tungsten is showing up mostly as a substitute for molybdenum and mostly in Russian and Chinese metals. In doing the calculation the practice is to lump the tungsten with the moly on a 1:1 basis. When forced to do it, welding metallurgists may tweak the calculations a little to reflect other unlisted elements, however, there isn't much of a research foundation on which to base such adjustments.

Ferrite Number

The 300 series of stainless steels is austenitic (non-magnetic) while the duplex stainless steels are mixed austenite and ferrite. In the field, the amount of ferrite is measured through its magnetic response. The portable meter is calibrated for % ferrite and it is called the "ferrite number". Martensite gives a magnetic response, but not as strongly as ferrite, so when it is present it contributes to the "ferrite number".

Recognize also that perfectly good austenitic stainless steel that has been heavily cold worked can become slightly magnetic and give a ferrite number even though ferrite nor martensite are present.

In the upper right side of the diagram there is a cluster of lines of constant ferrite number. The upper left line is 1% ferrite number which is nominally 99% austenite and 1% ferrite. The last line along the lower right side of the cluster is nominally 98% ferrite and 2% austenite.

Learning by Doing (and "cya")

Good record keeping means photocopying a bunch of diagrams and every time you confront the martensite issue you make a record of how you calculated the nickel and chromium equivalents. You plot these on the graph, draw the lines between and estimate an adjustment for nitrogen. You note the ferrite number and whether the bend test is pass or fail.

You don't just do the exercise for samples that fail the bend test. Otherwise you won't learn the limits for welds that pass every time! If testing the weld for its ferrite number is worthwhile then making the record is also worthwhile. The day will come when you note the manganese is low and you make yourself into a hero when you predict the need for a little preheat.


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(The WRC-1992 Diagram modified to reflect experience with manganese. Martensite is predicted to form with composition equivalents below and to the left of the manganese bands. Reprinted from Advanced Materials&Process, June 1000, p. 75 - click image to enlarge)

There will be discrepancies. When it happens it most likely means that you misjudged the effect of cooling rate, or you didn't get the chemistry right, or the basic diagram is imperfect!


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Effects on the martensite, pearlite and bainite formation

Effect on the temperature of martensite formation

All alloying elements with the possible exception of Co, lower Ms the temperature of the start of the martensite formation, as well as Mf, the finish of the martensite formation, i.e. at 100% martensite. For the majority of steels containing more than 0.50% C, Mf lies below room temperature.

This implies that after hardening these steels practically always contain some residual austenite. Ms may be calculated from the equation given below, by inserting the percentage concentration of each alloying element in the appropriate term. The equation is valid only if all the alloying elements are completely dissolved in the austenite.

Ms = 561 - 474C - 33Mn - 17Ni - 17Cr - 21Mo

For high-alloy and medium-alloy steels Stuhlmann has suggested the following equation:

Ms(°C) = 550 - 350C - 40Mn - 20Cr - 10Mo - 17Ni - 8W - 35V - 10Cu + 15Co + 30Al

It can be noted that carbon has the strongest influence on the Ms temperature. Figures 1 and 2 show diagrams with an example of experimental results of the effect of Mn and Ni on the Ms temperature of various types of steel.

Figure 1. Effect of Mn on the Ms - temperature (after Russel and McGuire, Payson and Savage, Zyuzin, Grange and Stewart)


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Figure 2. Effect of Ni on the Ms - temperature (after Russel and McGuire,

Payson and Savage, Zyuzin, Grange and Stewart)

Effect on the formation of pearlite and bainite during the isothermal transformation

All alloying elements except Co delay the formation of ferrite and cementite. It is very difficult to formulate any general rules regarding the influence exerted by the various alloying elements. However, it has definitely been found that some elements affect the bainite transformation more than the pearlite transformation, while other elements act in the opposite manner. Certain elements will, paradoxically, accelerate the transformations if their concentration increases beyond a certain limiting value, this limit been affected by other alloying elements present. For case-hardening and tool steels the time taken to initiate the pearlite-bainite transformation is reduced as the carbon content exceeds about 1%. For tool steels and constructional steels Si-concentrations of 1,5% and above have been found to promote pearlite formation. As a general principle it may be stated that by increasing the concentration of one alloying element by some few percent and the basic carbon content being kept about 0.50%, only a relatively small retardation of the transformation rates is noticed. For plain carbon steels a successive increase in C from 0.30% to 1% produces but a negligible effect. It is only in conjunction with several alloying elements that a more noticeable effect is produced. The diagram in Figure 3, applicable to steel W 1 (l% C) will serve as a basis for this discussion. The shortest transformation time for this steel is less than 1/8th second. Note that the time scale is logarithmic; hence there is no zero time. As has been mentioned previously, both pearlite and bainite form simultaneously in this steel at about 550°C. Since the curves overlap it is customary to draw only one curve. With increasing contents of certain alloying elements, however, the noses of the pearlite and bainite curves will separate. The structures shown in Figure 3 are obtained by austenitizing samples of steel W 1 at 780°C for 10 min and quenching in a salt bath at various temperatures. After holding them for predetermined times at various temperatures they are finally quenched in water. Before the salt-bath quenching the steel contains undissolved carbides but in view of the composition of the austenite the steel may be regarded as an eutectoid one. The


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diagram should be studied with the aid of the explanatory text below.

Figure 3. TTT diagram for isothermal transformation of steel W 1 (1% C)

A = austenite, B = bainite,

Ms = start of martensite transformation,

M50 = 50% M, P = pearlite

1. Quenching in a liquid bath at 700°C; holding time 4 min. During this interval the C has separated out, partly as pearlite lamellae and partly as spheroidized cementite. Hardness 225 HV.

2. Quenching to 575°C, holding time 4 s. A very fine, closely spaced pearlite as well as some bainite has formed. Note that the amount of spheroidized cementite is much less than in the preceding case. Hardness 380 HV.

3. Quenching to 450°C, holding time 60 s. The structure consists mainly of bainite. Hardness 410 HV. 4. Quenching to 20°C (room temperature). The matrix consists of, roughly, 93% martensite and 7%

retained austenite. There is some 5% cementite as well which has not been included in the matrix figure. Hardness 850 HV.


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Steel Alloys

Below is a list of some SAE-AISI designations for Steel (the xx in the last two digits indicate the carbon content in hundredths of a percent)

Carbon Steels 10xx Plain Carbon 11xx Resulfurized 12xx Resulfurized and rephosphorized Manganese steels 13xx Mn 1.75 Nickel steels 23xx Ni 3.5 25xx Ni 5.0 Nickel Chromium Steels 31xx Ni 1.25 Cr 0.65-0.80 32xx Ni 1.75 Cr 1.07 33xx Ni 3.50 Cr 1.50-1.57 34xx Ni 3.00 Cr 0.77 Chromium Molybdenum steels

41xx Cr 0.50-0.95 Mo 0.12-0.30 Nickel Chromium Molybdenum steels

43xx Ni 1.82 Cr 0.50-0.80 Mo 0.25 47xx Ni 1.05 Cr 0.45 Mo 0.20 – 0.35 86xx Ni 0.55 Cr 0.50 Mo 0.20 Nickel Molybdenum steels

46xx Ni 0.85-1.82 Mo 0.20 48xx Ni 3.50 Mo 0.25 Chromium steels 50xx Cr 0.27- 0.65 51xx Cr 0.80 – 1.05

Illustration of effect of Carbon content on Steel Hardness


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Carbon Steel to Austenitic Steel

When a weld is made using a filler wire or consumable, there is a mixture in the weld consisting of approximately 20% parent metal and 80% filler metal alloy ( percentage depends on welding process, type of joint and welding parameters). Any reduction in alloy content of 304 / 316 type austenitic is likely to cause the formation of martensite on cooling. This could lead to cracking problems and poor ductility. To avoid this problem an over alloyed filler metal is used, such as a 309, which should still form austenite on cooling providing dilution is not excessive. The Shaeffler diagram can be used to determine the type of microstructure that can be expected when a filler metal and parent metal of differing compositions are mixed together in a weld. The Shaeffler Diagram

The Nickel and other elements that form Austenite, are plotted against Chrome and other elements that form ferrite, using the following formula:- Nickel Equivalent = %Ni + 30%C + 0.5%Mn Chrome Equivalent = %Cr + Mo + 1.5%Si + 0.5%Nb Example, a typical 304L = 18.2%Cr, 10.1%Ni, 1.2%Mn, 0.4%Si, 0.02%C Ni Equiv = 10.1 + 30 x 0.02 + 0.5 x 1.2 = 11.3 Cr Equiv = 18.2 + 0 + 1.5 x 0.4 + 0 = 18.8 A typical 309L welding consumable Ni Equiv = 14.35, Cr Equiv = 24.9


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The main disadvantage with this diagram is that it does not represent Nitrogen, which is a very strong Austenite former. Ferrite Number The ferrite number uses magnetic attraction as a means of measuring the proportion of delta ferrite present. The ferrite number is plotted on a modified Shaeffler diagram, the Delong Diagram. The Chrome and Nickel equivalent is the same as that used for the Shaeffler diagram, except that the Nickel equivalent includes the addition of 30 times the Nitrogen content.

Examples

The Shaeffler diagram above illustrates a carbon steel C.S , welded with 304L filler. Point A represents the anticipated composition of the weld metal, if it consists of a mixture of filler metal and 25% parent metal. This


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diluted weld, according to the diagram, will contain martensite. This problem can be overcome if a higher alloyed filler is used, such as a 309L, which has a higher nickel and chrome equivalent that will tend to pull point A into the austenite region. If the welds molten pool spans two different metals the process becomes more complicated. First plot both parent metals on the shaeffler diagram and connect them with a line. If both parent metals are diluted by the same amount, plot a false point B on the diagram midway between them. (Point B represents the microstructure of the weld if no filler metal was applied.)

Next, plot the consumable on the diagram, which for this example is a 309L. Draw a line from this point to false point B and mark a point A along its length equivalent to the total weld dilution. This point will give the approximate microstructure of the weld metal. The diagram below illustrates 25% total weld dilution at point A, which predicts a good microstructure of Austenite with a little ferrite.


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The presence of martensite can be detected by subjecting a macro section to a hardness survey, high hardness levels indicate martensite. Alternatively the weld can be subjected to a bend test ( a side bend is required by the ASME code for corrosion resistant overlays), any martensite present will tend to cause the test piece to break rather than bend. However the presence of martensite is unlikely to cause hydrogen cracking, as any hydrogen evolved during the welding process will be absorbed by the austenitic filler metal.

More reading: http://nhml.com/search.cfm?zoom_query=stainless+steel&zoom_per_page=10&zoom_and=0&zoom_sort=0


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Selection of Age-Hardenable Superalloys

By Richard B. Frank High-Temperature R&D Carpenter Technology Corp., Reading, PA, USA

Superalloys are high-performance materials designed to provide high mechanical strength and resistance to surface degradation at high temperatures of 1200°F (650°C) or above. They combine high tensile, creep-rupture, and fatigue strength; good ductility and toughness, with excellent resistance to oxidation and hot corrosion. Furthermore, superalloys are designed to retain these properties during long-term exposures at the elevated temperatures.

This article focuses on the wrought age-hardenable alloys, which are the most commonly used superalloys. Wrought materials can be formed using hot and cold working operations. Not discussed here are the cast, powder (P/M), and oxide dispersion strengthened (ODS) superalloys that can also offer enhanced properties.

The first age-hardenable, high-temperature alloy dates back to about 1929 when various developers added titanium and aluminum to the standard 80% nickel/20% chromium resistance wire alloy. This was a precursor to the 80A nickel-base superalloy, developed in 1940-1944, but still in use today.

Little was done to advance the original age-hardenable alloys until the time period of 1935-1944 when World War II spurred demand for improved alloys that could be used in the early aircraft gas turbine engines. Alloy development activity exploded in the 1950’s and 1960’s to keep pace with the demands of the gas turbine engine industry. Progress in superalloy development not only made the jet engine possible, but allowed for constantly increasing thrust-to-weight ratios over the last 60 years.

Applications

The primary application for superalloys is still in hot sections of aircraft gas turbine engines, accounting for over 50% of the weight of advanced engines. However, the excellent performance of these materials at elevated temperatures has expanded their application far beyond one industry.

In addition to the aerospace industry, these alloys are used in turbine engines for marine, industrial, land-based power generation, and vehicular applications. Specific engine parts using superalloys include turbine discs, blades, compressor wheels, shafts, combustor cans, afterburner parts and engine bolts.

Beyond the gas turbine engine industries, superalloys are commonly used for applications in rocket engines, space, petrochemical/energy production, internal combustion engines, metal forming (hot-working tools and dies), heat-treating equipment, nuclear power reactors, and coal conversion.

While these alloys are primarily used for service at elevated temperatures above 1000°F (540°C), the characteristics of high strength and excellent environmental resistance have made some superalloys an


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excellent choice for lower-temperature applications. Examples are prosthetic devices in the medical industry and components for deep sour gas wells in the oil/gas exploration industry.

Chemical Composition

Table 1 contains the nominal compositions of the most common wrought age-hardenable superalloys. These alloys contain various combinations of nickel, iron, cobalt, and chromium with lesser amounts of other elements including molybdenum, niobium, titanium, and aluminum. With minor additions of beneficial elements such as boron and zirconium, these alloys may contain up to 12 intentional additions. All these additions help to impart and maintain the desired properties at elevated temperatures.

Many other elements such as silicon, phosphorus, sulfur, oxygen, nitrogen and a larger number of tramp elements (like lead, bismuth, selenium) must be tightly controlled in superalloys to avoid detrimental effects on high-temperature properties. These minor and tramp elements are controlled during raw material selection prior to melting, as well as during the melting/remelting processes.

Superalloys can be classified into nickel-base, iron-base, and cobalt-based groups. Nickel-based superalloys (>50% Ni) are the most common group. About half of the alloys in Table 1 are considered nickel-base alloys and


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the others contain large additions of nickel. The nickel base has a high tolerance for alloy additions that might otherwise cause phase instability leading to loss of strength, ductility, and/or environmental resistance.

Iron-based superalloys are less costly, but are less tolerant of alloying additions and typically have lower mechanical properties and maximum temperature limitations. Examples are Pyromet® Alloy A-286 and NCF 3015 (Ni-30) alloy. These alloys contain an austenitic stainless steel base with additions of nickel, titanium, and aluminum to promote age hardening. Pyromet Alloy 706 and Pyromet Alloy 901 have similar amounts of nickel and iron and can be considered nickel-iron-base superalloys. The higher nickel levels of 901 and 706 alloys allow for larger additions of strengthening elements without undesirable effects.

Although there are some cobalt-base superalloys, they are significantly higher in cost and typically cannot be age hardened to high strength levels. However, cobalt is an important alloying addition to nickel-based alloys because it extends the maximum temperature for usage by reducing the solubility of the age-hardening phase. Waspaloy and Pyromet Alloy 41 and Pyromet Alloy 720 are nickel-base alloys with 10-15% cobalt additions. These alloys have the highest temperature capability of the common wrought age-hardenable superalloys.

Chromium, usually in the range of 14 to 23 weight percent, is a critical alloying addition to nearly all superalloys. As in stainless steels, chromium forms a tightly-adherent, protective oxide film (Cr2O3) on the alloy surface to resist oxidation and corrosion at high temperatures as well as corrosion at lower temperatures. This surface layer protects the alloy from the harmful effects of the elements oxygen, nitrogen, and sulfur.

Although most superalloys contain at least 14% chromium, in some applications, it is critical to minimize thermal expansion. Pyromet CTX-909 and Thermo-Span® alloys are considered low-expansion superalloys that have low chromium contents to minimize expansion of the nickel-cobalt-iron base. Resistance to oxidation and hot corrosion are reduced so high-temperature coatings are often applied prior to service. Of the two alloys, 909 alloy provides the lowest expansion coefficient while Thermo-Span alloy (5.5% chromium) provides improved environmental resistance.

Refractory elements like molybdenum, tungsten, and niobium, with their large atomic diameters, increase high temperature strength and stiffness by straining the nickel/iron base matrix. Alloys 901 and 41 contain larger additions of molybdenum to increase this solid solution strengthening effect. Other alloying additions such as chromium and aluminum also contribute to solid solution strengthening but to a lesser extent.

The elements titanium, aluminum, and niobium are added to the nickel or nickel-iron matrix to form an intermetallic Ni3 (Al, Ti, Nb) phase during age-hardening heat treatments. The resultant gamma prime or gamma double prime phases are the primary strengthening agents in superalloys. This will be discussed in more detail in the next section on age-hardening.

Although elements such as boron, zirconium, and magnesium may be added at levels less than 0.1 weight percent, the beneficial effects can be very potent. These elements segregate to and stabilize grain boundaries, which significantly improves hot workability, high temperature strength and ductility. Small additions of carbon also may be added to form carbides that restrict grain growth and grain boundary sliding during high temperature exposure.

Age-Hardening


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The major strengthening method in superalloys is age-hardening. Yield strength of nickel alloys is typically increased by a factor of two or three by precipitation of the gamma prime and/or gamma double prime, Ni3 (Al, Ti, Nb) hardening phase. Although the phase is based on the nickel aluminide (Ni3Al) intermetallic, up to 60% of the aluminum can be replaced by titanium or niobium, which actually increases strength of the alloy.

The gamma prime phase is rather unique in that its strength actually increases with temperature up to 1200°F (650°C) and it is relatively ductile and resistant to oxidation. Gamma prime precipitates as very fine spheroidal or cuboidal particles in the nickel-iron matrix during aging.

While most of the superalloys employ the titanium-rich gamma prime phase for age hardening, a niobium-rich variant called gamma double prime is the primary strengthening phase in some superalloys such as Pyromet Alloy 706 and Pyromet Alloy 718. The niobium-rich phase provides higher strength up to 1200°F (650°C) but is unstable above 1200°F. Thus, 706 and 718 alloys have a lower temperature limit than the alloys strengthened with the titanium-rich gamma prime phase. Since the gamma double prime reaction is more sluggish, these alloys also tend to have better hot workability and weldability.

Heat Treatment

Proper heat treatment is critical to achieving the desired level of properties in age-hardenable superalloys. Typical heat treatments for these alloys are listed in the mechanical property Tables 2 and 3. The initial solution heat treatment typically dissolves all precipitated phases except for some primary carbide and nitride phases. The typical range for the wrought age-hardenable superalloys is 1650-2100°F (900-1150°C) for 1 to 4 hours followed by a rapid air cool or a quench in water, polymer or oil.


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The selection of solution treatment time and temperature varies with the alloy and its phase solvus temperatures, and also depends on the specific properties that are most important for the intended application. Alloys with higher hardener contents (Ti, Al, Nb) require higher temperatures to solution any hardener phase that may have precipitated during hot working or cooling. Best tensile and fatigue properties are typically obtained with lower solution temperatures that result in a finer grain size. In contrast, better long-term stress-rupture and creep properties are generally obtained with higher-temperature solution treatments that result in coarser grain size and lower tensile yield strength. For these reasons, it is common to specify two or more preferred heat treatments for superalloys.

In some cases, another objective of the solution treatment is to form a desirable distribution of a second phase such as carbide in Pyromet 41 alloy and delta phase (Ni3Nb) in Pyromet 718 alloy. After solution treatment, one or more aging treatments are applied to precipitate the hardening phase and possibly other phases in the desired amount and distribution. As with solution treatment, the selection of aging temperatures is dependent on the alloy and the combination of properties desired.


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The aging range for age-hardenable superalloys is 1150-1600°F (620-870°C). Aging times range from 4 hours to 24 hours. Double-aging treatments are quite common to maximize strength and to develop the best combination of short-term tensile and long-term creep-rupture properties. The primary aging treatment precipitates a coarser distribution of the hardener phase and may also improve the type and distribution of carbides on grain boundaries.

The secondary age is typically about 200°F below the primary aging temperature, precipitating a finer dispersion of the gamma prime phase. For some higher-strength applications, the alloy is direct aged after hot, warm, or cold working without an intermediate solution treatment. The strain from working is used to further enhance tensile and fatigue properties with some sacrifice in creep-rupture properties.

Mechanical Properties

For the design engineer or materials specifier, a review of terms defining applicable mechanical properties may be helpful:

Tensile Properties – The design of load-bearing structures is often based on yield strength or, in some cases, the ultimate tensile strength of the material. Yield strength is a measure of the maximum stress a material can withstand before it permanently deforms. Tensile strength is a measure of the maximum stress a material can withstand before it fractures. Elevated temperature tensile properties are most applicable to short-time exposures at higher temperatures. Creep and stress-rupture properties are more applicable for longer exposures.

Creep and Rupture Properties – Creep and rupture strengths become important when the material must withstand the combined effects of high temperature and stress for long periods of time. At elevated temperatures, metals will stretch or "creep" at stresses well below the yield strength. Superalloys are more resistant to creep than low-alloy or stainless steels, but creep will still occur above about 1000°F (540°C). Creep properties are a measure of the alloy’s resistance to stretching under a constant load. Stress-rupture or creep-rupture properties are a measure of resistance to fracture under a constant load (creep test taken to fracture). Both properties are expressed as stress or strength values that will cause a given amount of creep (0.1%-1%) or rupture in a given amount of time (100 to 100,000 hours).

Tables 2 and 3 list typical tensile (yield) and stress-rupture strength properties of the age-hardenable superalloys at temperatures of 1200-1600°F (650-870°C). Yield strengths at room temperature are also listed in Table 2. These properties are shown graphically in Figures 1 and 2. It should be noted that the data represents approximate nominal strength values for specific heat treatments. Actual values can vary by up to 35% due to differences in composition, hot/cold working practices, and heat treatment.


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For example, superalloys like Pyromet 718 and Waspaloy may contain several different aim compositions within the broader industry ranges to optimize properties for specific applications. Higher levels of the age-hardening elements titanium, aluminum, and niobium result in higher strength. Hot or cold working an alloy to obtain a finer grain size typically increases tensile yield strength but decreases stress-rupture strength.

As discussed previously, properties of all age-hardenable superalloys are dependent on heat treatment. Alloys like Pyromet Alloy X-750 and Waspaloy have two or more preferred heat treatments (see Table 2) depending on whether the application requires better short-time tensile and fatigue properties or long-time creep and stress-rupture properties. Examples of alternative heat treatments have been shown for Waspaloy and X-750 alloys but the reader should refer to manufacturers’ datasheets for a more complete listing of alternative heat treatments for the other superalloys.

Other Properties – While tensile and creep-rupture are the most basic mechanical properties considered for high-temperature applications, design criteria may also consider resistance to fatigue (low- and high-cycle), crack growth, and wear/erosion. Hardness and hot hardness tests are sometimes used as a rough measure of yield strength and wear/erosion.

Alloy Selection


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A simplified method known as the Carpenter Selectaloy® system can help designers and engineers select the most suitable superalloy based on strength and maximum temperature requirements. Figures 3 and 4 contain Selectaloy diagrams for the 15 superalloys discussed in this article. Yield strength (Figure 3) or stress-rupture strength (Figure 4) increases vertically on the Selectaloy diagram, and temperature increases from left to right. The alloys are shown multiple times on the diagrams since the alloys are useful over a range of temperatures. The diagram can be used to estimate how the strength of an alloy decreases with temperature, but also how the strength of different alloys compare at different temperatures.


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It should be noted that the alloys were positioned on the Selectaloy diagrams based on average strength values representative of compositions and heat treatments commonly used for each alloy. An alloy’s relative position could move up or down, left or right, with relatively minor modifications of composition, processing and heat treatment. Temperature limits should be considered approximate. Therefore, while the Selectaloy diagrams are useful tools to screen candidate alloys, they are not a substitute for a more detailed evaluation of the critical properties required for an intended application.

Pyromet A-286 alloy is the most basic age-hardenable superalloy in terms of properties and cost. A-286 provides the lowest strength levels, but still higher by a factor of two than other non-age-hardenable stainless alloys. When increased strength or temperature resistance is required, higher nickel alloys are typically preferred. Alloys with the highest levels of strength and temperature resistance typically contain the highest alloy contents and significant levels of cobalt. Relative cost of these alloys will be discussed in the next section.

The Selectaloy diagrams presented in this article provide a method to compare basic strength properties and temperature limitations of common wrought age-hardenable superalloys. However, alloy selection will undoubtedly depend on many other considerations, including other physical and mechanical properties as well as environmental resistance and cost. For example, Thermo-Span and Pyromet CTX-909 alloys provide a benefit of much lower expansion during heating but at the expense of oxidation and corrosion resistance in the


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uncoated condition. Pyromet 31V and Pyromet 751 alloys provide similar strength and temperature resistance, but the higher chromium content of 31V alloy results in much improved resistance to sulfidation and other forms of hot corrosion.

Alloy Cost

From the user’s standpoint, alloy selection must be based on expected cost effectiveness.

In today’s competitive global environment, overdesign is less common. The trend is to select the lowest-cost material to meet design requirements for the application. However, a higher-cost alloy may be justified to minimize overall life cycle cost or for longer service of certain components in a system that is critical or too expensive to be shut down for maintenance. Surely, knowledge of alloy capabilities is critical in making the best decision.

As temperature and strength requirements increase, so does the necessary alloy content. Figure 5 compares the relative alloying costs of the 14 alloys using Pyromet A-286 alloy as a base (cost factor of 1.0). The cost factors are based on 10-year averages of the intrinsic alloying element costs at market prices. Higher temperature strength and resistance typically require higher nickel and cobalt contents. Nickel and cobalt prices have historically been volatile, with high and low prices varying by a factor of 4 to 5. More recently, the price of molybdenum, a potent solid solution strengthener, has increased in price by a factor of nearly ten over the last two years.


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As discussed above, the cost factors in Figure 5 are based only on raw material elemental costs (10-year averages) that fluctuate significantly with time. Differences in melting, working, and other processing costs, which can be substantial, are not included in these factors. Processing yields and specific end user requirements (grain size, ultrasonic testing, etc.) significantly impact product cost. However, the cost comparisons are useful because alloying costs typically represent a large portion of superalloy product cost. Since superalloys are designed for high temperature strength and resistance to deformation, processing difficulty and cost also increase with hot strength and maximum temperature capability.

Figure 6 shows the relationships of stress-rupture strength at 1200-1500°F (650-815°C) with raw material cost factor. It is apparent that the alloys that provide higher levels of strength, temperature resistance, and/or specialized properties also cost more, which reinforces the importance of the alloy selection process.

More Technical Article: http://crswnew.cartech.com/wnew/techarticles/TechLibrarySelector.html


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Metallurgy of Mo in Alloy Steel & Iron

Spherical Agglomerates of Mo powder (lacy appearance), with solid spheres of the Ni-Cr alloy binder for the Mo powder

Molybdenum metal is usually produced by powder metallurgy techniques in which Mo powder is hydrostratically compacted and sintered at about 2100°C. Hot working is done in the 870-1260°C range. Moly forms a volatile oxide when heated in air above about 600°C and therefore high temperature applications are limited to non-oxidizing or vacuum environments.

Moly alloys have excellent strength and mechanical stability at high temperatures (up to 1900°C). Their high ductility and toughness provide a greater tolerance for imperfections and brittle fracture than ceramics.

High temperature furnace with molybdenum hot zone (Courtesy of PLANSEE AG, Austria)

The unique properties of molybdenum alloys are utilised in many applications:

• High temperature heating elements, radiation shields, extrusions, forging dies, etc; • Rotating X-ray anodes used in clinical diagnostics; • Glass melting furnace electrodes and components that are resistant to molten glass; • Heat sinks with thermal expansivity matching silicon for semiconductor chip mounts; • Sputtered layers, only Ångstroms (10-7 mm) thick, for gates and interconnects on integrated circuit

chips; • Sprayed coatings on automotive piston rings and machine components to reduce friction and

improve wear.


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For specialised applications, Mo is alloyed with many other metals:

• Mo-tungsten alloys are noted for exceptional resistance to molten zinc; • Mo is clad with copper to provide low expansion and high conductivity electronic circuit boards; • Mo-25% rhenium alloys are used for rocket engine components and liquid metal heat exchangers

which must be ductile at room temperature.

Hardening

The purpose of quenching steel after heating is hardening, i.e. to produce a hardened microstructure over the full cross section of the workpiece.

A round steel bar quenched in water from a temperature up to 900°C will cool faster near the surface than in the center (Fig 2 below).

Fig 2: Quenching round bar

Fig 3: Simulating cooling of steel sections

On a laboratory scale this is simulated by the Jominy end quench test (Fig 3 above).

The standard sample is heated and then water quenched on one end. The cooling rate in the sample decreases from the water sprayed end where it is highest to the opposite end where it is lowest. When the sample is cool, the surface is ground and the hardness profile is taken. The change of hardness along the sample reflect variations of the microstructure brought about by the different cooling rates.


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The curves in Fig 4 compare the hardness profile of steels with different alloy contents. The carbon manganese steel without molybdenum has only been hardened near the quenched end of the sample and the hardness drops quickly when moving away from the quenched end. With increasing molybdenum content the hard microstructure is maintained at increasing distances from the quenched end. That means that it is possible to harden a steel with higher molybdenum content with slower cooling rates: the hardenability is improved.

Fig 4: Jominy hardenability curves: Hardenability improves with increasing Mo content (after W.W. Cias1)

The hardenability indicates the depth, to which a steel grade can be hardened.

In standard Quenched and Tempered Steels a combination of alloying elements is usually used, including manganese, chromium, molybdenum, nickel and silicon.


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Fig 5: The Hardenability Multiplying Factor shows the rate at which the hardening depth is increased with the percentage of

the alloying element (after Honeycombe2)

The basis of steel hardening lies in the fact, that iron exists in two crystal structures:

Below 912°C and from 1394°C to its melting point iron is body centered cubic – bcc – called ferrite. In the lower temperature range ferrite is also referred to as alpha iron, in the higher temperature range as delta iron.

At a temperatures from 912°C to 1394°C iron is in the face centered cubic crystal structure – fcc – called gamma iron or austenite. Heating pure iron above 912°C transforms the structure from Ferrite into Austenite. Cooling the iron from the austenitizing area below 912°C results in the original bcc iron structure, no matter what cooling rate is applied.

Pure iron can not be hardened.

The addition of carbon converts iron into hardenable steel. (Alloying elements such as manganese, molybdenum and chromium enhance the hardenability).

Carbon is present in iron both in solid solution and in the form of carbides. It is significant that the sides of the face centered cubes of the austenite are about 25% larger than the sides of the body centered cube of the ferrite. The solubility for carbon is therefore much greater in austenite than in ferrite.

When a steel with say 0.4% carbon is heated above the ferrite.austenite (alpha-gamma) transformation point, carbon and the other alloying elements can go into solid solution in the spacious austenitic fcc structure. Subsequent cooling through the gamma – alpha transformation point leads into the narrow ferrite structure. There is not enough space in this structure to keep carbon in solid solution.


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So, if the cooling rate is low, carbide is formed in connection with the transformation process. As a result the microstructure at room temperature consists of ferrite and carbide. (The fine lamellar structure of ferrite and iron carbide is called pearlite - see Fig 6).

Fig 6: Ferrite – Pearlite microstructure – soft and ductile

Fig 7: Martensite microstructure –hard and brittle

The critical factor is, that there is enough time available for the carbon atoms to move through the lattices to form

carbides, which results in the soft microstructure of ferrite and pearlite.

Increasing the cooling rate progressively reduces the carbide formation. A very high cooling rate is achieved with water quenching, which completely supresses the carbide formation. In that case carbon is uncomfortably forced into narrow spaces in the ferrite structure. The microstructure which is generated that way is called martensite. This is the hardest and most brittle form of steel. (See Fig 7)

In plain carbon steels the high cooling rates required for the formation of martensite are only achieved near the quenched surface. Inside the work piece the structure remains soft. Water quenching larger sections also involves the risk of quench cracking.

This is where molybdenum and the other alloying elements enter the scene. Alloying elements slow down the diffusion of carbon atoms through the iron lattice., which delays the transformation from austenite into ferrite. The hardenability of the steel is thus improved since martensite can be produced at slower cooling rates. As shown in Fig 5, molybdenum is very effective in that respect.


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Also, in larger cross sections at intermediate cooling rates a structure called Bainite is formed particularly in Mo alloyed steel. In that case, some nucleation of carbides has taken place during cooling before the austenite- ferrite transformation.

In practice, the microstructure of quenched and tempered steel components consists of a combination of martensite and bainite.

Tempering

Hardened steel must be tempered, i.e. reheated for two reasons:

• Different cooling rates between edge and core of components result in internal stresses, which must be relieved.

• High hardness martensite and bainite must be softened to avoid brittle cracking.

The improvement of ductility is inevitably accompanied by softening, i.e. a loss of strength. This is very pronounced with carbon steel. An important function of alloying elements is to delay temper softening.

Through its capability of forming carbides, molybdenum, carefully combined with chromium and vanadium, is very efficient in delaying the loss of strength during tempering while improving fracture toughness. The resulting structure, tempered martensite, is very strong with an acceptable level of toughness.

Fig 8 shows the effect of molybdenum on the hardness after tempering of a 0.35% carbon steel. It significantly delays softening of the steel. At sufficiently high Mo contents the hardness curve may even increase with increasing tempering temperature. This is known as secondary hardening.

The effect of secondary hardening on tempering is an important function of molybdenum in high speed steels and in some tool- and die steels.


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Fig 8: Influence of Mo content on temper softening (after E.C. Bain3)

Temper brittleness

Temper embrittlement may occur when steels are slowly cooled after tempering through the temperature range between 450 and 550°C. This is due to the segregation of impurities such as phosphorus, arsenic, antimony and tin on the grain boundaries. The molybdenum atom is very large relative to other alloying elements and impurities. It effectively impedes the migration of those elements and thereby provides resistance to temper embrittlement.

Fig 9 shows the ductile-to-brittle transition temperature for two steels. This temperature is an indication for the lower limit of the service temperature without the risk of brittle failure.

If the steels are water quenched after tempering, both steels one without molybdenum and one with 0.15% molybdenum, have essentially the same ductile to brittle behaviour (transition at approx. -50°C). However, if the steels are slowly cooled in the furnace after tempering the picture changes. The transition to brittle fracture has shifted to +25°C for the Cr-steel, while it remained at -50°C for the Cr-Mo steel. The slow cooling rate has not embrittled the molybdenum containing steel, it is, therefore, less susceptible to temper embrittlement.


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Fig 9: Ductile to brittle transition for two tempered steels, as a function of cooling rate after tempering (after Dunn et al4)

Hydrogen embrittlement and sulphide stress cracking

As outlined above, the strength levels obtained in quenched and tempered steels are based mainly on the high strength of martensite, a microstructure characterized by a high density of dislocations and high internal stresses.

Unfortunately, exactly these conditions enhance the diffusion of hydrogen into steel and cause hydrogen embrittlement. Tempering reduces the internal stresses and the dislocation density of martensite, hence reduces hydrogen diffusion. However, the strength can be lowered to insufficient levels. Molybdenum is efficient in mittigating this effect in two ways: through solid solution strengthening and the formation of complex carbides together with other elements such as chromium and niobium.

In cases, where hydrogen sulphide is the source of the hydrogen the phenomenon is called sulphide stress cracking (SCC). The capability of molybdenum to provide resistance to sulphide stress cracking has been the key to the development of a broad range of steel grades used for Oil Country Tubular Goods and in chemical and petrochemical plants.


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High temperature hydrogen attack Carbon steel has severe limitations at conditions of hydrogen attack above about 200 C as are common in processes such as petroleum distilling and catalytic reforming. The hydrogen diffusing into the steel combines with carbon present to form methane and other products. The result is first decarburization and subsequently fissuring due to high gas pressure at localized sites.

Fig. 10 below compares the loss of rupture strength of various steels exposed to pressurized hydrogen at 540 C:

• the unalloyed carbon steel has obviously suffered the damage described above, loosing more tha 50 % of it’s original strength after less tha 50 hours exposure,

• additions of 0.5% Mo or 1%Cr-0.5%Mo show a slight improvement, but are not adequate under the given conditions, whereas

• the alloy content of 2.25%Cr plus 1% Mo provide protection to the extent, that after 500 hours exposure the original rupture strength of the steel has not deteriorated at all.

Fig. 10: Effect of composition and exposure time on the strength of steels exposed to hydrogen at 63 bar at 540 C. The

strength of samples exposed in argon is taken as 100%. (After Nelson5)

The positive effect of Cr plus Mo in this context used to be described as carbide forming and is now referred to as elements lowering grain boundary energy. In any event with suitable selection of Mo and Cr contents the steel will be resistant to hydrogen attack in respect to decarburization, fissuring and loss of strength (4; page 65ff).

http://www.imoa.info/moly_uses/moly_grade_alloy_steels_irons/hardening.html


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Maraging Steels马氏体时效钢

Maraging steel

Al/maraging steel fibre composite prepared via vacuum diffusion bonding of

plasma sprayed monolayers

Maraging steels are carbon free iron-nickel alloys with additions of cobalt, molybdenum, titanium and aluminium. The term maraging is derived from the strengthening mechanism, which is transforming the alloy to martensite with subsequent age hardening. Air cooling the alloy to room temperature from 820°C creates a soft iron nickel martensite, which contains molybdenum and cobalt in supersaturated solid solution. Tempering at 480°C to 500°C results in strong hardening due to the precipitation of a number of intermetallic phases, including, nickel-molybdenum, iron-molybdenum and iron-nickel varieties. Age harden carbon free martensitic steel Principle alloying element Nickel. Precipitates hardening Secondary alloying elements: Mo, Al, Cu, Ti, Ni and Co Intermetallic precipitates such as Ni3Ti and Ni3Mo on age hardening at about 500°C.

Introduction:

The 18% Ni-maraging steels, which belong to the family of iron-base alloys, are strengthened by a process of martensitic transformation, followed by age or precipitation hardening. Precipitation hardenable stainless steels are also in this group. Maraging steels work well in electro-mechanical components where ultra-high strength is required, along with good dimensional stability during heat treatment. Several desirable properties of maraging steels are:

The result is steel which:

• Possesses high strength and toughness. • Allows for easy machining with minimal distortion.


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• Has uniform, predictable shrinkage during heat treatment which results in minimum distortion • Can be easily nitrided. • Resists corrosion and crack propagation. • Superior fracture toughness compared to quenched and tempered steel of similar strength level • Low carbon content, which precludes decarburization problems • Section size is an important factor in the hardening process • Easily fabricated • Can be finely polished. • Good weldability.

These factors indicate that maraging steels could be used in applications such as shafts, and substitute for long, thin, carburized or nitrided parts, and components subject to impact fatigue, such as print hammers or clutches. Maraging steels (from Martensitic Aging) are iron alloys which are known for possessing superior strength without losing malleability. The iron base is alloyed principally with a large percentage of nickel to produce a very specific heat-treatment product. Other alloying elements include molybdenum, aluminum, copper and titanium and are added to produce intermetallic precipitates. Cobalt is added in percentages up to 12% to accelerate the precipitation reactions and ensure profuse and uniform precipitates. Maraging steel is essentially free of carbon, which distinguishes it from most other types of steel. These steels are a special class of low carbon ultra-high strength steels which derive their strength not from carbon, but from precipitation of inter-metallic compounds. The principal alloying element is 15 to 25% nickel. Secondary alloying elements are added to produce intermetallic precipitates, which include cobalt, molybdenum, and titanium. Original development was carried out on 20 and 25% Ni Nickel steels to which small additions of aluminium Al, titanium Ti, and niobium Nb were made. The common, non-stainless grades contain 17~19% nickel, 8~12% cobalt, 3~5% molybdenum and 0.2~1.6% titanium. Stainless grades rely on chromium not only to prevent their rusting, but to augment the hardenability of the alloy as their nickel content is substantially reduced. This is to ensure they can transform to martensite when heat-treated, as high chromium, high nickel steels are generally austenitic, and unable to undergo such a transition. Maraging steels are hardened by a metallurgical reaction that does not involve carbon. These steels are strengthened by intermetallic compounds such as Ni3Ti and Ni3Mo that precipitate at about 500°C. These steels typically have very high nickel, cobalt and molybdenum contents while carbon is essentially an impurity and its concentration is kept as low as possible in order to minimize the formation of titanium carbide which can adversely affect mechanical properties. Ultra-high strengths may be obtained with these steels, and weldability is good. Toughness is superior to all low alloy carbon steels of similar strength, particularly the low temperature toughness. Although they are expensive, they are easy to machine and heat treat, so that some economies result in component production. The common, non-stainless grades contain 17~19% nickel, 8~12% cobalt, 3~5% molybdenum, and 0.2~1.6% titanium. Stainless grades rely on chromium not only to prevent their rusting, but to augment the hardenability of the alloy as their nickel content is substantially reduced. This is to ensure they can transform to martensite when heat treated, as high-chromium, high-nickel steels are generally austenitic, and unable to undergo such a transition.


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Properties:

Due to the low carbon content maraging steels have good machinability. Prior to aging, they may also be cold rolled to as much as 80~90% without cracking. Maraging steels offer good weldability, but must be aged afterward to restore the properties of heat affected zone. When heat treated the alloy has very little dimensional change, so it is often machined to its final dimensions. Due to the high alloy content the alloys have a high hardenability. Since ductile FeNi martensites are formed upon cooling, cracks are non-existent or negligible. They can also be nitrided to increase case hardness. They can be polished to a fine surface finish. Non-stainless varieties of maraging steels are moderately corrosion resistant and resist stress corrosion and hydrogen embrittlement. More corrosion protection can be gained by cadmium plating or phosphating.

Surface markings due to the formation of martensite plates in Fe-Ni single crystals.

The width of the figure corresponds to 4mm of the sample.

• Offer the best available combination of ultra-high yield and tensile strength, ductility and fracture toughness of any ferrous materials.

• Can retain strength at least up to 350 °C. • Being a very low-carbon martensite, the structure is soft. Readily machinable. May be cold rolled to as

much as 80 - 90% without cracking before aging treatment. Hot deformation is also possible. • Possible to finish-machine before aging. Very little dimensional change after aging treatment. • Parts uniformly harden throughout the entire section because they have high hardenability. • Exhibit good weldability. Properties of heat affected zone (HAZ) can be restored by a post-weld aging

treatment. • Since ductile FeNi martensites are formed upon cooling, cracks are non-existent or negligible. • Can be surface hardened by nitriding. • Good corrosion, stress corrosion and hydrogen embrittlement characteristics. • Protection can be provided by cadmium plating or phosphating.


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Heat Treatment:

Annealing.

The steel is first annealed, at approximately 820 °C for 15 to 30 minutes for thin sections and for 1 hour per 25 mm thickness for heavy sections, to ensure formation of a fully austenitized structure. This is followed by air cooling to room temperature to form a soft, heavily dislocated iron-nickel lath (untwinned) martensite (hence the "mar-" in mar-aging). Subsequent aging (precipitation hardening) of the more common commercially used alloys for ~3 hours in 480-500 °C range produces a fine dispersion of Ni3(X,Y) intermetallic phases along dislocations left by martensitic transformation, where X and Y are solute elements added for such precipitation, eg; Mo, Ti, Al, Cu, Si. Newer compositions of maraging steels have revealed other intermetallic stoichiometries and crystallographic relationships with the parent martensite, including rhombohedral and massive complex Ni50(X,Y,Z)50 - usually simplified to Ni50M50. Overaging leads to a reduction in stability of the primary, metastable, coherent precipitates, leading to their dissolution and replacement with semi-coherent Laves phases such as Fe2Ni/Fe2Mo. Further excessive heat-treatment brings about the decomposition of the martensite and reversion to austenite.

Tempering:

Maraging steels are carbonless Fe-Ni alloys additionally alloyed with cobalt, molybdenum, titanium and some other elements. A typical example is an iron alloy with 17-19% Ni, 7-9% Co, 4.5-5% Mo and 0.6-0.9% Ti. Alloys of this type are hardened to martensite and then tempered at 480°C~500°C. The tempering results in strong precipitation hardening owing to the precipitation of intermetallic from the martensite, which is supersaturated with the alloying elements. By analogy with the precipitation hardening in aluminum, copper and other non-ferrous alloys, this process has been termed ageing, and since the initial structure is martensite, the steels have been called maraging.

The structure of commercial maraging steels at the stage of maximum hardening can contain partially coherent precipitates of intermediate metastable phases Ni3Mo and Ni3Ti. Of special practical value is the fact that particles of intermediate intermetallic in maraging steels are extremely disperse, which is mainly due to their precipitation at dislocations.

The structure of maraging steels has a high density of dislocations, which appear on martensitic rearrangement of the lattice. In lath (untwined) martensite, the density of dislocations is of an order of 1011-1012 cm-2, i.e. the same as in a strongly strain-hardened metal. In that respect the substructure of maraging steel (as hardened) differs appreciably from that of aluminum, copper and other alloys which can be quenched without polymorphic change. (Polymorphism (materials science), the ability of a solid material to exist in more than one form or crystal structure)

It is assumed that the precipitation of intermediate phases on tempering of maraging steels is preceded with segregation of atoms of alloying elements at dislocations. The atmospheres formed at dislocations serve as centers for the subsequent concentration stratification of the martensite, which is supersaturated with alloying elements.


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In maraging steels the dislocation structure that forms in the course of martensitic transformation, is very stable during the subsequent heating and practically remains unchanged at the optimum temperatures of tempering (480°C -500°C). Such a high density of dislocations during the whole course of tempering may be due to an appreciable extent, to dislocation pinning by disperse precipitates.

A long holding in tempering at a higher temperature (550°C or more) may coarsen the precipitates and increase the inter particle spacing, with the dislocation density being simultaneously reduced. With a long holding time, semi coherent precipitates of intermediate intermetallic are replaced with coarser incoherent precipitates of stable phases such as Fe2Ni or Fe2Mo.

At increased temperatures of tempering (above 500°C), maraging steels may undergo the reverse martensitic transformation, since the as point is very close to the optimum temperatures of tempering. The formation of austenite is then accompanied with the dissolution of the intermetallic that has precipitated from the gamma-phase.

Application:

Maraging steel's strength and malleability in the pre-aged stage allows it to be formed into thinner rocket and missile skins, allowing more room for payload while still possessing sufficient strength for the application. Maraging steels have very stable microstructural properties, and even after overaging due to excessive temperature only soften very sluggishly. These alloys retain their properties at mildly elevated operating temperatures and have maximum service temperatures of over 400 °C. They are suited to engine component applications such as crankshafts and gears, that work at 'warm' temperatures, and the firing pins of automatic weapons that cycle from hot to cool repeatedly while under substantial loads and impacts. Their uniform expansion and easy machinability, carried out before aging makes maraging steel useful in high wear portions of assembly lines, as well as in the manufacture of dies. Other ultra-high strength steels, such as the secondary hardening 'Aermet©' family are not so amenable to processing because of their ever-present carbide particle dispersion.

In the sport of fencing, testing has shown that the blade breakage patterns in carbon steel and maraging steel blades are identical. Maraging blades are required in foil and épée (there is no such requirement for sabre blades) because the crack propagation in maraging steel is 10 times slower than in carbon steel. This results in less blade breakage and fewer injuries.

Stainless maraging steels have been used in golf club heads and in surgical components and hypodermic syringes. They are not suitable for scalpel blades, as their virtually zero carbon content prevents holding a good cutting edge.

Maraging steel is also an item of great interest to those responsible for national security. Maraging steel is used in creating gas centrifuges for uranium enrichment due to its extremely high strength and balance. Very few other materials will work for this task, and maraging steel’s other uses are very specialized. Nations importing maraging steel often find themselves receiving a great deal of attention.


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Physical Properties:

• Density: 8.1 g/cm³ (0.29 lb/in³) • Specific heat, mean for 0–100 °C (32–212 °F): 813 J/(kg·K) (0.108 Btu/(lb·°F)) • Melting point: 2575 °F, 1413 °C, 1686 K • Thermal conductivity: 25.5 W/(m·K) • Mean Coefficient of Thermal Expansion: 11.3×10-6 • Yield tensile strength: typically 1400-2100 MPa (200,000-300,000 PSI) • Ultimate strength: typically 1600-2500 MPa (230,000-360,000 PSI). Grades exist up to 3.5 GPa

(500,000 PSI) • Elongation at break: up to 15% • KIC fracture toughness: up to 175MPam½ • Young's modulus: 195 GPa • Shear modulus: 77 GPa • Bulk modulus: 140 GPa

Variation of Properties:

The dependence of mechanical properties of maraging steels on the temperature of tempering is of the same pattern as that for all precipitation-hardenable alloys, i.e. the strength properties increase to a maximum, after which softening takes place. By analogy with ageing, the stages of hardening and softening tempering may be separated in the process.

The hardening effect is caused by the formation of segregates at dislocations and, what is most important, by the formation of partially coherent precipitates of intermediate phases of the type Ni3Ti or Ni3Mo. The softening is due, in the first place, to replacement of disperse precipitates having greater inter particle spacing and, in the second place, to the reverse martensitic transformation which is accompanied by the dissolution of intermetallic in the austenite.

The ultimate strength of maraging steels increases on tempering roughly by 80% and the yield limit, by 140%, i.e. the relative gain in strength properties is not greater than in typical age-hardening alloys, such as beryllium bronze or aluminum alloy Grade 1915, but the absolute values of ultimate and yield strength on tempering of maraging steels reach record figures among all precipitation hardening alloys. This is mainly due to the fact that maraging steels have a very high strength (Rm = 1100 MPa) in the initial (as-hardened) state.

The high strength of maraging steels on tempering at 480- 500 ‹C for 1-3 hours may be explained by the precipitation of very disperse semi coherent particles of the size and inter particle spacing of an order of 103 nm in the strong matrix, these intermetallic precipitates also possessing a high strength. Thus, with the same disparity of precipitates as that of G. P. zones in precipitation, hardening non-ferrous alloys, maraging steels possess an appreciably higher ultimate strength (Rm = 1800-2000 MPa).


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As compared with martensite-hardenable carbon-containing steels, carbonless maraging steels show, for the same strength, a substantially greater resistance to brittle fracture, which is their most remarkable merit. On tempering to the maximum strength, the ductility indices and impact toughness, though diminish somewhat, still remain rather high. The high ductility of the carbonless matrix and the high disparity of uniformly distributed intermetallic precipitates are responsible for a very high resistance to cracking, which is the most valuable property of modern high-strength structural materials.

The properties of maraging steels clearly indicate that these steels have many potential applications in mechanical components of electro-mechanical data processing machines. Use of these steels in shafts that require good dimensional control following heat treatment should be pursued for two reasons. First, maintaining dimensions should be easier because quenching and tempering are not necessary. Second, wear data indicate that equivalent or better wear resistance is obtained from the maraging steel than from the more commonly used shaft materials.

Impact-fatigue strength of 18% Ni-maraging steels indicates that these steels could be used in repeated impact loading situations. The good fracture toughness, compared to that of quenched and tempered alloy steels at the same strength level, indicates possible use in high-impact low-cycle load applications.

Finally, due to the relatively low temperature of aging, the use of the maraging steels for long, thin parts should be considered. Here, their use as a replacement for some case hardened or nitrided components is indicated that the potential application should be carefully studied.


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Table 1 summarizes the alloy content of the 18% nickel – cobalt - molybdenum family as developed by Inco in the late 1950s.

Maraging Steels % Alloy content

Type Yield Strength

(0,2% proof stress) (MPa) Ni Co Mo Ti Al

18Ni1400 1400 18 8.5 3 0.2 0.1

18Ni1700 1700 18 8 5 0.4 0.1

18Ni1900 1900 18 9 5 0.6 0.1

18Ni2400 2400 17.5 12.5 3.75 1.8 0.15

17Ni1600 (cast) 1600 17 10 4.6 0.3 0.05

Table 1: Summary of the alloy content of the 18% nickel – cobalt - molybdenum family

With yield strength between 1400 and 2400 MPa maraging steels belong to the category of ultra-high-strength materials. The high strength is combined with excellent toughness properties and weldability.

Typical applications areas include:

• aerospace, e.g. undercarriage parts and wing fittings, • tooling & machinery , e.g. extrusion press rams and mandrels in tube production, gears • Ordnance components and fasteners.


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Maraging 250

Maraging 250 is an 18 percent nickel steel that has been strengthened with cobalt. Maraging 250, like all maraging steels, goes through an aging process that forces the metal to cool from its molten state to its solid state over an artificially long time. This process results in a tempered steel that has both high levels of strength and hardness. It will also resist certain stresses and maintain its structure in environments that would cause irreparable changes to many other steels.

The properties that make Maraging 250 particularly appealing to many industries is its workability. This allows Maraging 250 to be more versatile than many other alloys in its class. However, it is still the alloy's strength and resistance to extreme temperatures that make it a truly effective material in a wide range of atmospheres. After Maraging 250 has undergone heat treatment, it demonstrates excellent mechanical properties. It will reach a yield strength of 240 ksi and a fracture toughness of 75 kic. These properties have made Maraging 250 effective in the construction of missile and rocket motor cases, landing and takeoff gear, and high-performance shafting.

Alloy Steels Maraging 250

Related Metals: VascoMax 250(tm) Specifications: DIN 1.6358 Chemistry Data Aluminum Boron Carbon 0.03 max Cobalt Iron Balance Manganese 0.1 max Molybdenum Nickel Phosphorus 0.01 max Silicon 0.1 max Sulphur 0.01 max Titanium Zirconium Physical Data Density (lb / cu. in.) 0.289

Maraging Alloy 250 - Age Hardenable (Maraging) Iron-Nickel Steel from Super Alloys

http://www.azom.com/Details.asp?ArticleID=4368


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Maraging 300

Maraging 300 is an iron-nickel steel alloy that, as with all maraging steels, exhibits high levels of strength and hardness. However, Maraging 300 also possesses an extreme resistance to crack propagation, even in the most extreme environments. Maraging 300 is often used in applications where high fracture toughness is required or where dimensional changes have to remain at a minimal level.

The unique properties of Maraging 300 have made it a integral part of the aircraft and aerospace industries. It is often used in rocket motor casings and the landing gear for certain planes. Maraging 300 is also effective in the design of power shafts and low-temperature cooling systems.

Alloy Steels Maraging 300

Specifications

The following specifications cover Alloy Steels Maraging 300: DIN 1.6354

Property Results Related Metals: VascoMax 300 (tm)

Chemistry Data Carbon : 0.03 max

Iron : Balance

Manganese : 0.1 max

Phosphorus : 0.01 max

Silicon : 0.1 max

Sulphur : 0.01 max

Principal Design Features： Maraging 300 alloy is a very high strength iron base, nickel , molybdenum, cobalt alloy. It may be age hardened to a strength

level of 300 ksi ultimate with a hardness on the order of Rockwell C 53.

Applications： Used for very high strength structural applications such as aircraft land gear components.

Machinability： Machinability is good, similar to that for 304 stainless steel.

Forming： Maraging 300 steel has good ductility and is readily formed by conventional methods.

Welding： The alloy is weldable by conventional methods, much the same as for 304 stainless steel.

Heat Treatment： Maraging steel is not heat treatable for hardening except by aging. See "Aging".

Forging： The alloy may be forged as for stainless 304 alloy.

Hot Working： Hot working of maraging steel may be done at temperatures in the 500 to 200 F range. Prolonged exposure at these

temperatures may result in some age hardening. However a full aging treatment should be given.


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Cold Working： The alloy is readily cold worked by conventional methods.

Annealing： Annealing, if required after cold working, can be done at 1900 F followed by rapid air cooling.

Aging： The high strength and hardness properties are a result of an aging heat treatment at 900 F

Tempering： Not applicable to this alloy.

Hardening： Maraging 300 hardens by cold working and by aging - see "Aging".

Physical Data Density (lb / cu. in.) 0.289

Specific Gravity 8

Specific Heat (Btu/lb/Deg F - [32-212 Deg F]) 0.108

Melting Point (Deg F) 2575

Thermal Conductivity 136

Mean Coeff Thermal Expansion 5.6

Modulus of Elasticity Tension 27.5

More technical information: http://www.suppliersonline.com/buy/exchange/post/ChooseGradeTokens.asp?Tokens=maraging


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Maraging 350

Maraging refers to a crystalline tempered steel, Martensite, that is created through an aging process. When aging is used, steel is forced to cool from its molten state to its solid state over a prolonged period of time. The result is a metal that is harder and stronger than it would be had the steel been allowed to cool at a natural rate.

Maraging 350 is an alloy that has become an integral material in the airplane and aerospace industries. Due to its strength and its ability to withstand extreme conditions including frequent and sudden changes in speed and temperature, Maraging 350 is used in the production of rocket motor cases, takeoff and landing gear, and certain munitions created by defense companies. Maraging 350 also has uses in less drastic applications such as die casting and high-performance shafting.

Maraging 362

Maraging materials are known for their exceptional strength and hardness. Their ability to resist various forms of stress in extreme environments has made maraging steels commonplace in the aerospace and aircraft industries. Each maraging alloy has its own unique qualities, but many of them are used in similar applications.

Maraging 362, like other maraging steels, undergoes an artificial aging process. This process leads to the material's added strength and hardness. The results of the aging procedure has recently caught the eye of golf club designers and manufacturers who have begun to use maraging alloys on the faces of their clubs in the hope that players will see increased power and that clubs will not corrode over time.

Maraging Steel

Maraging steel gives you an elevated level of strength, hardness, and ductility. These steels are created through an aging process that results in the development of a hard, brittle crystalline called martensite. The term "maraging" is, in fact, a simple combination of martensite and aging.

The construction of maraging steel allows it to withstand atmospheres that would quite simply destroy most standard steel. The aging process instills maraging steel with the ability to withstand sudden changes in speed and temperature, even at extreme levels. This quality has made maraging steel alloys an important component of many of the air and spacecraft used today.

Its specific use in military programs and space programs has made maraging steel a rather unique material. It is not generally carried by small-scale metals companies. Instead, it can be found through distributors


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who regularly work with organizations in the aerospace and aeronautical industries. A metals company that has the capacity to supply maraging steel will usually distribute a wide range of hard-to-find alloys.


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The Periodic Table:

The Atomic Radii:

In specifying values for the radius of an atom, one must keep in mind the fact that atoms are not hard spheres, and the electron distribution in the outer part of the atom does not have a sharp cutoff radius. You could characterize the radius of the atom as a limiting radius where a certain percentage of the electron charge will be found. The illustration above is a plot of "covalent radii" (from Ebbing) which are determined by measuring the bond lengths in the molecules of chemical compounds. Another way to determine characteristic radii is to measure ionic radii in crystals using x-ray diffraction. If the crystalline composition is such that the ions can be considered to be in contact with each other, and you can determine the lattice spacing from x-ray diffraction, then


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you can imply the ionic radius. As might be expected, the ionic radius of negative ions is slightly larger than the covalent radius since they have extra electronic charge, and that of positive ions is slightly smaller.

Atoms and Nuclei

http://www.practicalphysics.org/go/Topic_40.html?topic_id=40

Schaeffler diagram.


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Precipitation hardening stainless steels

Precipitation hardening stainless steels like the martensitic types, can be strengthened by heat treatment. The mechanism is metallurgically different to the process in the martensitic types. This means that either martensitic or austenitic precipitation hardening structures can be produced. This family of stainless alloys utilizes a thermal treatment to intentionally precipitate phases, which cause a strengthening of the alloy. The principle of precipitation hardening is that a supercooled solid solution changes its metallurgical structure on aging. The advantage is that products can be fabricated in the annealed condition and then strengthened by a relatively low temperature treatment, minimizing the problems associated with high temperature treatment. Precipitation hardening stainless steel has high strength and relatively good ductility and corrosion resistance at high temperature. They reach these high strengths by precipitation of intermetallic compounds

Alloy 17-4PH

Sandmeyer Steel Company stocks a large inventory of 17-4PH stainless steel plate in the annealed condition in thicknesses from 3/16" through 3", ready to be processed and shipped to your specific requirements.

Alloy 17-4PH is a precipitation hardening martensitic stainless steel with Cu and Nb/Cb additions. The grade combines high strength, hardness (up to 572°F /300°C), and corrosion resistance. Mechanical properties can be optimized with heat treatment. Very high yield strength up to 1100-1300 MPa (160-190 ksi) can be achieved.

17-4PH (UNS S17400) A 17Cr-4Ni-3Cu Precipitation Hardening Martensitic Stainless Steel

General Properties Applications Standards Corrosion Resistance Chemical Analysis Mechanical Properties Physical Properties Heat Treatment Processing Welding Machining


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General Properties Alloy 17-4 PH is a precipitation hardening martensitic stainless steel with Cu and Nb/Cb additions. The grade combines high strength, hardness (up to 572°F / 300°C), and corrosion resistance.

Mechanical properties can be optimized with heat treatment. Very high yield strength up to 1100-1300 MPa (160-190 ksi) can be achieved.

The grade should not be used at temperatures above 572°F (300°C) or at very low temperatures. It has adequate resistance to atmospheric corrosion or in diluted acids or salts where its corrosion resistance is equivalent to Alloy 304 or 430.

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Applications

• Offshore (foils, helicopter deck platforms, etc.) • Food industry • Pulp and paper industry • Aerospace (turbine blades, etc.) • Mechanical components • Nuclear waste casks

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Standards ASTM.....................A693 grade 630 (AMS 5604B) UNS S17400 EURONORM...........1.4542 X5CrNiCuNb 16-4 AFNOR...................Z5 CNU 17-4PH DIN.........................1.4542

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Corrosion Resistance Alloy 17-4 PH withstands corrosive attacks better than any of the standard hardenable stainless steels and is comparable to Alloy 304 in most media.


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If there are potential risks of stress corrosion cracking, the higher aging temperatures then must be selected over 1022°F (550°C), preferably 1094°F (590°C). 1022°F (550°C) is the optimum tempering temperature in chloride media.

1094°F (590°C) is the optimum tempering temperature in H2S media.

The alloy is subject to crevice or pitting attack if exposed to stagnant seawater for any length of time.

It is corrosion resistant in some chemical, petroleum, paper, dairy and food processing industries (equivalent to 304L grade).

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Chemical Analysis Typical values (Weight %)

C Cr Ni Cu Nb/Cb Mn

0.04 16.5 4.5 3.3 0.3 0.7

PREN (%Cr+3.3%Mo+16%N) ≥17

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Mechanical Properties Room temperature properties (longitudinal direction)

Guaranteed values (ASTM A693 hot rolled plates); thickness from 3/16" up to 3".

YS 0.2% N/mm

UTS N/mm

YS 0.2% ksi

UTS ksi

El% Heat treatment

Min. Typ. Min. Typ. Min. Typ. Min. Typ. Min. Typ.

A 1070 1207 1170 1310 155 175 170 190 8 14

B 790 931 965 1034 115 135 140 150 10 17

A: hardening 925°F (496°C) - 4 hours - air cooling B: hardening 1100°F (593°C) - 4 hours - air cooling 2 examples of heat treatments that may be applied.


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Elevated temperature properties Minimum guaranteed values following EN 10088 hot rolled plates. The EN guaranteed values are valid for a thickness from 3/16" up to 3".

Temperature °F Temperature °C

212 100

302 150

392 200

482250

572300

N/mm_ 730 710 690 670 650YS 0.2%

ksi 106 103 100 97 95

Heat treatment : hardening 1094°F (590°C) - 4 hours - air cooling. 1 example of heat treatments that may be applied.

Minimum guaranteed room temperature impact values

Minimum guaranteed values following ASTM A693 hot rolled plates. The ASTM guaranteed values are valid for a thickness from 3/16" up to 3".

KV transverseHeat treatment

J ft.lbf

Hardening 1100°F (593°C) - 4 hours - air cooling

20 15

Minimum guaranteed room temperature hardness values

Minimum guaranteed values following ASTM A693 hot rolled plates. The ASTM guaranteed values are valid for a thickness from 3/16" up to 3".

Hardness Heat treatment

Rockwell Brinell

Hardening 925°F (496°C) - 4 hours - air cooling

C38 375

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Physical Properties Density: 7800 kg/m3 (.28 lbs/in3)

Following physical properties have been obtained after hardening 896°F (480°C) – 1 hour - air cooling.


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Interval Temperature

°C

Thermal expansion

x10　 -6°C-1 °C °F

Thermal conductivity

(w.m-1.K-1)

Young modulus

(GPa)

0-100 10.8 20 68 14 197

0-200 11 100 212 16 193

0-300 11.3 200 392 18.5 186

0-400 11.6 300 572 20 180

0-500 12 400 752 22 175

500 932 23 170

Room temperature properties: Resistivity : 80 µ*.cm Specific heat : 460 J.kg-1.K-1 Tension modulus : 77 GPa The alloy is magnetic.

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Heat Treatment Martensitic transformation Indicative values Ms : 266°F (130°C) Mf : 86°F (30°C)

Solution annealing 1925°F+ /-50°F (1050°C +/-25°C) – 30 min up to 1 hour. air cooling / oil quenching below 76°F (25°C)

Aging / Tempering The highest mechanical properties are obtained with the following heat treatment : 896°F (480°C) – 1 hour – air cooling. Higher ductilities are obtained when using higher aging temperatures up to 1148°F (620°C).


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Hardness/Temperature tempered for 4 hours after austenitizing at 1904°F (1040°C) for 30 min quenched 212°F/sec (100°C/sec).

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Processing Hot Forming Hot forming should be carried out in a temperature range of 1742-2192°F (950-1200°C). A full heat treatment including solution annealing, cooling lower than 76°F (25°C) and aging at the required temperature must be made after hot forming (function of the requested mechanical properties).

Cold Forming Cold forming can be performed only to a limited extent and only on plates in the fully softened condition. Stress corrosion resistance is improved by re-aging at the precipitation hardening temperature after cold working.

The following processes may be performed: rolling, bending, hydroforming, etc. (fully softened conditions).

Cutting Thermal cutting (plasma, thermal sawing, etc.). Due to the HAZ, the grade requires a suited cutting process. After cutting, grinding is necessary to eliminate the oxide formed layer.

Mechanical cutting (shearing, stamping, cold sawing, etc).

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Welding Alloy 17-4 PH can be welded by the following welding processes: SMAW, GTAW, PAW and GMAW. SAW should not be used without preliminary testing (to check freedom of cracks and toughness of the weld metal).


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Due to a ferrite delta primary type of solidification, the hot cracking risk of the weld metal or the HAZ is reduced.

Generally, no preheating must be done and interpass temperature must be limited to 248°F (120°C). The better toughness is obtained in the weld after a complete heat treatment (solution annealing + precipitation hardening).

Due to the martensitic microstructure, a low oxygen content in the weld metal is preferable to increase ductility and toughness. To avoid cold cracking, the introduction of hydrogen in the weld must be limited.

Alloy 17-4 PH can be welded with homogeneous filler metals such as E 630 (AWS A5.4) electrodes and ER 630 (AWS A5.9) wires.

Austenitic filler material can be used when the mechanical properties of 17-4 PH steel are not required in the weld and, in this case, no post-weld heat treatment must be applied.

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Machining Alloy 17-4 PH can be machined in both solution treated and precipitation hardened conditions. Machining condition may vary according to the hardness of the material. High speed steel tools or preferably carbide tools with standard lubrification are normally used. If very stringent tolerances are required, it is necessary to take into account the dimensional changes during heat treatment.


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Thermal Processing of Metal:

http://www.tntech.edu/me/courses/Zhang/ME30103110/Chap11pt3.ppt

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Online reading materials:

Ductile Iron Tutorial http://www.ductile.org/didata/Section2/2intro.htm Alloys and Stainless Steels http://www.roymech.co.uk/Useful_Tables/Matter/Alloy_Steels.html Metallurgical Engineering http://www.sut.ac.th/Engineering/Metal/course.html Matters http://www.matter.org.uk/steelmatter/metallurgy/default.htm Microplasticity: dislocations and strengthening mechanisms http://www-sgrgroup.materials.ox.ac.uk/lectures/microplasticity.html Steel Heat Treatment: Metallurgy and Technologies http://www.ebookee.com.cn/Steel-Heat-Treatment-Metallurgy-and-Technologies_148046.html Metallurgy fundamentals. http://rapidshare.com/files/9027049/Metallurgy_Fundamentals_Warner_Brandt.djvu.html Solid-State Phase Transformations in Steels http://upload.sms.csx.cam.ac.uk/collection/19

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Stainless Steel Data. http://www.sandmeyersteel.com/precipitation-hardening.html Stainless Steel grades data. http://www.bssa.org.uk/topics.php?article=59 Materials and Processes in Manufacturing http://www.tntech.edu/me/courses/Zhang/ME30103110/ Engineering Materials and physical metallurgy. http://www.sut.ac.th/Engineering/Metal/course.html

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http://www.journalarchive.jst.go.jp/english/jnltoc_en.php?cdjournal=isijinternational1989&cdvol=43&noissue=5

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DEFORMATION-MECHANISM MAPS The Plasticity and Creep of Metals and Ceramics

http://engineering.dartmouth.edu/defmech/

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冶金电子书-免费下载

http://www.gougou.com/search?search=%E5%86%B6%E9%87%91%E7%94%B5%E5%AD%90%E4%B9%A6&id=4009360&pattern=40000


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Additional Reading

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Appendix A.

More reading:

http://www.sut.ac.th/Engineering/Metal/course.html

Applied Metallurgy http://www.esdep.org/members/master/wg02/toc.htm

Steels http://www.key-to-steel.com/default.aspx?ID=Articles&LN=EN

Material Property Search

http://www.matweb.com/search/QuickText.aspx?SearchText=4140

http://www.matter.org.uk/default.htm

The Alloying Elements in Steel, by Edgar C. Bain

http://www.msm.cam.ac.uk/phase-trans/2004/Bain.Alloying/ecbain.html

Material Science Course http://info.lu.farmingdale.edu/depts/met/met205/index.html

http://www.cs.slcc.edu/engr/moore/mse/mse2010.html

Now you can perform virtual tensile testing. http://www.uow.edu.au/cedir/progservs/samples/anim_ENGG153/intro.swf

Articles on steel making and inclusions: Steel making: http://www.sigi.ca/engineering/documents/steel_making.pdf Effects of deoxidation practice on the inclusion formation: http://www.ferroforsk.com/ekstern/publications/mikron/PaperStockholm2000.pdf Slag inclusion formation during solidification: http://www.diva-portal.org/diva/getDocument?urn_nbn_se_kth_diva-4371-2__fulltext.pdf Introduction to materials and processes: http://www.ndt-ed.org/EducationResources/CommunityCollege/Materials/cc_mat_index.htm

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Appendix A.

Chapter1: http://www.cs.slcc.edu/engr/moore/mse/lectures/chapter1_files/frame.htm

All Chapters: http://www.cs.slcc.edu/engr/moore/mse/lectures/

http://mc-mjnde.ornl.gov/Babu/Teaching/Brazil/PTFolder/SolidFol/sld001.htm

http://www.cs.slcc.edu/engr/moore/mse/lectures/chapter1_files/frame.htm

http://www.cs.slcc.edu/engr/moore/mse/lectures/chapter1_files/frame.htm

http://www.cs.slcc.edu/engr/moore/mse/lectures/

http://mc-mjnde.ornl.gov/Babu/Teaching/Brazil/PTFolder/SolidFol/sld001.htm


Appendix A.

Ahindra Ghosh, "Secondary Steelmaking: Principles and Applications" CRC; 1 edition (December 13, 2000) | ISBN: 0849302641 | 344 pages | PDF | 17 Mb well as modify harmful nonmetallic inclusions and achieve the optimum casting temperature, content of alloying elements, and homogeneity. These improvements can come only through the diverse refinement processes that together comprise "secondary steelmaking." Secondary Steelmaking: Principles and Applications reviews the scientific fundamentals and

explores the various unit processes associated with secondary steelmaking. Synthesizing the science and its technology, the author examines the relevant reactions and phenomena, presents an integrated picture of "clean steel" manufacture, and provides an overview of the mathematical modeling important to process research. Solved examples, ample references, and summaries of recent technological advances mean that the steelmaking industry finally has a comprehensive reference, in English, for the all-important secondary steelmaking processes. Students and instructors, steelmakers and R & D engineers will welcome the author's readable style, his knowledge, and his expertise, all gleaned from decades of experience in research, academic, and industrial settings.

http://depositfiles.com/files/2287630 MIRROR : http://www.icefile.net/index.php?page=main&id=e95461423&name=Steelmaking.7z

A wealth of data on metals and their extraction is revealed in this comprehensive handbook. The aim of this book is to provide a clear description of how a particular metal is extracted industrially from different raw materials, and on what its important compounds are. The present work is a collection of 58 articles written by over 280 specialists. It supplies thousands of top-quality illustrations, diagrams and charts, and provides hand-picked references ensuring the most up-to-date coverage. A unique feature of this reference work is its structure. The system used here is according to an economic classification, which reflects mainly the uses, occurrence and economic value of metals. First, the ferrous metals, i.e., those used in the production of iron and steel, are outlined. Then, nonferrous metals are subdivided into primary, secondary, light, precious, refractory, scattered, radioactive, rare earth, ferroalloy metals, and, finally, the alkali and the alkaline earth metals are described. The handbook is an essential aid for the practising metallurgist. Mining

engineers, mineralogists, chemical engineers, chemists and geologists will find it a comprehensive desk reference. It is of interest to engineers and scientists in industry seeking an exhaustive sourcebook, and it should be present in every library. http://rapidshare.com/files/13898521/20062007.part1.rar http://rapidshare.com/files/13945593/20062007.part2.rar

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Appendix A.

Steel Heat Treatment: Metallurgy and Technologies (Steel Heat Treatment Handbook) CRC | 848 pages | 2006-09-28 | ISBN: 0849384559 | PDF | 33.3MB One of two self-contained volumes belonging to the newly revised Steel Heat Treatment Handbook, Second Edition, this book examines the behavior and processes involved in modern steel heat treatment applications. Steel Heat Treatment: Metallurgy and Technologies presents the principles that form the basis of heat treatment processes while incorporating detailed descriptions of advances emerging since the 1997 publication of the first edition. Revised, updated, and expanded, this book ensures up-to-date and thorough discussions of how specific heat treatment processes and different alloy elements affect the structure and the classification and mechanisms of steel transformation, distortion of properties of steel alloys. The book includes entirely new chapters on heat-treated components, and the treatment of

tool steels, stainless steels, and powder metallurgy steel components. Steel Heat Treatment: Metallurgy and Technologies provides a focused resource for everyday use by advanced students and practitioners in metallurgy, process design, heat treatment, and mechanical and materials engineering. http://rapidshare.com/files/59830510/Steel_Heat_Treatment_0849384559.rar

Madeleine Durand-Charre , «Microstructure of Steels and Cast Irons» Publisher: Springer | Number Of Pages: 404 | Publication Date: 2004-04-15 | ISBN / ASIN: 3540209638 | PDF | 24 MB The book comprises three parts. Part 1 gives a historical description of the development of ironworking techniques since the earliest times. Part 2 is the core of the book and deals with the metallurgical basis of microstructures, with four main themes: phase diagrams, solidification processes, diffusion, and solid state phase transformations. Part 3 begins by an introduction to steel design principles. It then goes on to consider the different categories of steels, placing emphasis on their specific microstructural features. Finally, a comprehensive reference list includes several hundred pertinent articles and books. The book is the work of a single author, thus ensuring uniformity and concision. It is

intended for scientists, metallurgical engineers and senior technicians in research and development laboratories, design offices and quality departments, as well as for teachers and students in universities, technical colleges and other higher education establishments.

http://rapidshare.com/files/44188131/Microstructure.of.Steels.and.Cast.Irons_muya.rar

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Appendix A.

Metallurgy Fundamentals By Daniel A. Brandt,&nbspJ. C. Warner, * Publisher: Goodheart-Wilcox Publisher * Number Of Pages: 301 * Publication Date: 2004-02 * ISBN / ASIN: 1590703456 An excellent introduction to metal making I found Metallurgy Fundamentals to be an excellent resource for myself. While not overly technical it covers the basic principles and theory of metal making in clear and straightforward language. I recommend it for persons in drafting, quality, design, or purchasing 114 MB ~ Pdf. Good Quality.

Part 1 http://rapidshare.com/files/37614440/Metallurgy_Fundamentals__Brand_Warnr.part1.rar Part 2 http://rapidshare.com/files/37622602/Metallurgy_Fundamentals__Brand_Warnr.part2.rar -------------- http://rapidshare.com/files/37614440...arnr.part1.rar http://rapidshare.com/files/37622602...arnr.part2.rar

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Scanned copy-poor quality, however it is simple, informative, lots of photos and educational

Slag inclusion formation during solidification of Steel alloys and in cast iron

Sofia Adolfi

Licentiate Thesis

Materials Processing Department of Material Science and Engineering

School of Industrial Engineering and Management Royal Institute of Technology SE-10044 Stockholm, Sweden

Akademisk avhandling som med tillstånd av Kungliga Tekniska Högskolan i Stockholm framlägges till offentlig granskning för avläggande av teknologie Licentiatexamen, fredagen den 11 maj 2007, kl. 10:00, Sal B1, Brinellvägen 23,

KTH, Stockholm.

ISSNKTH/MSE-07/10-SE+CER/AVH, ISBN 978-91-7178-624-1

Sofia Adolfi, Slag inclusion formation during solidification of Steel alloys and in cast iron

School of Industrial Engineering and Management

Department of Materials Science and Technology, Materials Processing

Royal Institute of Technology

SE-100 44 Stockholm, Sweden

ISSN KTH/MSE-07/10-SE+CER/AVH

ISBN 978-91-7178-624-1

© Sofia Adolfi March 2007

Slag inclusion formation during Solidification of Steel alloys and in cast iron Sofia Adolfi 2007

Department of Materials Science, Royal Institute of Technology S-100 44 Stockholm, Sweden

Abstract This thesis explores the formation of segregation and inclusions during solidification

of steel and cast iron. A better understanding of the formation mechanism should result in decreasing fraction of defects during solidification of ingot and strand material.

Density driven macrosegregation was studied both experimentally and theoretically to see the effect of channel segregation on the total segregation. Formation of these pencil-like segregations is due to natural convection in the solidifying metal caused by liquid enrichment of elements with lower density compared to the bulk. It is suggested to change the composition to compensate for this density difference.

Inclusion precipitation can be finite by limitations in segregation. Saturated liquid is found in the last solidified areas, often between dendrites. Here the enrichment of the liquid is possible due to microsegregation. Meanwhile crystals form and solidify the elements with low solubility in the solid is pushed out in the remaining liquid. Soon the liquid is saturated to the level where spontaneous formation of inclusions occurs. Microstructure studies by aid of SEM and micro-probe measurements are analysed to find at what point during solidification process the inclusions start to form. In steel making this formation has a detrimental effect on the mechanical properties in contrary to the production of nodular cast iron where the inclusions have a beneficial effect on the graphite formation.

Inoculation of cast iron aims at reaching higher number density of graphite nodules, nodule morphology modification and control of nodule distribution during solidification. Late precipitation of nucleation sites has shown to have a positive impact on preventing chill. To find the most potent inoculation agent different additives were tested. Special effort has been made to analyse the effect of oxides and sulphides as nucleation sites.

Descriptors: solidification, segregation, precipitation, inclusions, inoculation, EPMA

Supplements The thesis includes the following supplements

Supplement 1

Crack formation during continuous casting of tool steel A. Lagerstedt, S. Adolfi and H. Fredriksson Trans. Indian Inst. Met. Vol. 58, No. 4, August 2005, pp. 671 – 678

I performed the hot-tensile tests, evaluation and part of report writing

Supplement 2

Macrosegregation in ingot cast tool steel A. Lagerstedt, J. Sarnet, S. Adolfi and H. Fredriksson ISRN KTH-MG-INR-04:09 SE TRITA-MG 2004:09

I took part in all experiments and part of final evaluation and report writing. In particular I performed the temperature measurements and EPMA analysis

Supplement 3

A thermodynamic analysis of the inoculation process L. Magnusson, S. Adolfi and H. Fredriksson ISRN KTH-MG-INR-06:03 SE TRITA-MG 2006:03

I performed the experimental part and together with Lena Magnusson the theoretical evaluation

Supplement 4

MnS precipitation during solidification in presence of oxide nuclei S. Adolfi, K. Mori and H. Fredriksson Accepted to the 5th Decennial International Conference on Solidification Processing

I performed the experiments and report writing

Contents

1. Introduction 1

2. Experimental methods 3 2.1 Material preparation 3

2.2 Thermal properties 3 2.2.1Differential Thermal Analysis 3 2.2.2 “In-situ” solidified hot tensile tests 4

2.3 Macrosegregation study 5

2.4 Inclusion study 6 2.4.1 Inclusion characteristics 6 2.4.2 Electron Probe Micro-Analysis 6

2.5 Inoculation in nodular cast iron 7

3. Results 8

3.1 Macrosegregation 10

3.2 Inclusion characteristics 10

3.3 Inoculation in nodular cast iron 10

4. Discussion 13

5. Conclusions 15

6. Future work 17

7. References 21

1

1 Introduction The aim of this study was to analyse the solidification process in steel alloys and cast

iron. Focus has been on segregation and defects and their formation mechanism. Both experimental and theoretical method to analyse this has been used. A great part of the work has been to design the proper experiment. The experimental outcomes are used for further theoretical work.

During solidification of steel ingots and strand compositional variations referred to as macrosegregation range in scale from several millimetres to centimetres or even metres. The in-homogenous distribution of alloy elements have a detrimental impact on the following processing behaviour and the material properties and can lead to rejection of cast components or processed products. The formation of these variations is driven by natural convection where hot metal will flow upward inside the centre of the ingot and cooler solidifying liquid will grow heavier in density and flow downward the mould wall. At this stage some crystal with lower liquidus temperature can remelt while in contact with the high-temperature melt leading to channel formation where highly segregated liquid will flow upwards due to its lower density. During solidification crystals will freeze at the solidus temperature and gradually form a solid. This solid will continue to grow meanwhile the remaining liquid will be enriched by elements with lower solubility in solid phase compared to the liquid. This enrichment or commonly referred to as microsegregation will open up for precipitation of other phases, inclusions which benefit from the saturation of alloying elements.

Hot cracks are often the result from the solidification process. The material will go through a brittle to ductile transformation during solidification. The brittle region causes this crack formation. This brittleness is caused by the presence of thin liquid films in the interdendritic region at the crystal boundaries.1 These are often type II MnS inclusions. By means of hot-tensile-tests, DTA and Electron Probe Micro-analysis the proper material data can be found for further theoretical analysis on crack- and inclusion formation. Inclusions in steel are categorized according to type (chemical composition) and morphology (shape). They origin either from before casting or they spontaneously precipitate during solidification. The first case is often a result from secondary metallurgy and can be an effect of slag or mould material being dragged into the liquid steel. This type is often referred to as macro slag. Precipitation of inclusion during solidification is possible when the supersaturation needed to form a certain compound is reached in the liquid and the formation facilitates by presence of small nuclei, mainly of oxide type. One of the most common inclusions is the MnS. Early it was discovered by Sims and Dahle,2 they are of different morphology due to formation process. They classified the globular shaped as type I, enveloped thin sheeted as type II and the faceted ones as type III. Later work made by Fredriksson and Hillert,3 revise this classification by adding the type IV morphology, a lamellar eutectic structured MnS. They suggest that the formation of types I and II are by a monotectic reaction where MnS forms as a liquid phase. Types III and IV are a eutectic reaction where MnS forms a crystalline phase. A monotectic reaction is a process in which a melt gives a solid phase and another liquid phase, L1→α+L2. A eutectic reaction is a solidification process in which a liquid solidifies to

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two solid phases, L→α+β. The type of MnS inclusion depends on many variables. The main factors are: cooling rate, concentrations of Mn and S in the melt, the solubilities of the alloying elements in molten MnS + Fe and the deoxidation process of the melt. Common deoxidants are Al and Si. Another method to reduce O concentration is vacuum degassing.

During solidification inclusions are pushed in front, entrapped or engulfed by the solidifying front. Stefanescu et al.,4 explain there exist a critical velocity of the planar solid – liquid interface below which particles are pushed ahead of the advancing front and above which particle engulfment occurs. Engulfment is used to describe incorporation of a particle (i.e. inclusion) by a planar interface and entrapment is used to explain how particles incorporate between cells and dendrites. Engulfment will normally lead to a uniform particle distribution, while pushing will result in particle segregation. Later studies by Stefanescu et al.5 show that particle interaction with dendritic solid – liquid fronts can be explained with similar controlling parameters as for planar fronts. Based on dendrite tip radius, particle radius, natural convection induced liquid velocity, VL, and the solidification velocity, VSL, engulfment or entrapment occurs. Engulfment can occur if convection is low and surface velocity is high. Slow interface velocity result in particle being entrapped between dendrite arms. However if VL is high but VSL is small the particle will be pushed in front of the interface. These types of defects in steel production should be possible to decrease by better understanding on the formation mechanism.

Precipitation of oxides and sulphides has a beneficial effect on the graphite formation in production of nodular cast iron. The inclusions act as nucleation sites for graphite. By addition of different inoculation agents we can promote nodularity, refine the graphite structure and suppress carbide formation. In earlier work by Skaland,6 it is suggested that different oxides and sulphides may increase the inoculation effect. Oxide and sulphide inclusions which are formed after addition of the inoculants act as nucleation sites for graphite nodules. To find the most potent inoculants we studied the nucleation process of graphite in presence of different inoculation agents. In contrary to inclusions formed during steel making precipitation of oxides and sulphides is most wanted.

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2 Experimental methods

2.1 Material preparation The material used in this thesis work is presented in table 1, supplement 4. Steel A is

represented by one octagonal 12 ton ingot and one rectangular 10 ton ingot produced by uphill casting at Scana Steel Björneborg AB. Steel B comes from a continuous casting experiment carried out in slab caster 2 at SSAB Oxelösund AB. Steel C is a aluminium-killed 6 ton ingot produced by uphill casting at Ovako AB. Ingots were cold sawed to pieces in order to allow for sulphur printing and drilling for chemical composition to obtain the segregation pattern of C and S. The old but well known Sulphur printing technique was re-discovered and prints were made on sawed surfaces in order to obtain macrographs of the ingot structure. To get as detailed information as possible from the prints some surfaces was milled to give good sharpness. Photo paper was soaked in 5% sulphuric acid and rolled on to the metal surface in good contact. The prints are taken at the vertical cross section. Sample drilling was performed over the vertical cross section to obtain the segregation pattern of C and S.

Test with different inoculation additives are made on a base alloy of white nodular vast iron, presented in table 1, supplement 3.

2.2 Thermal properties

2.2.1 Differential Thermal Analysis Differential Thermal Analysis, DTA was used to study the solidification process. The

result is used as in-data for further theoretical calculations. During heating and cooling the sample temperature was measured and any transformation will be shown as a change in temperature compared to a known reference state. The technique is to measure the difference in temperature between the specimen and a known reference which are exposed to the same heating schedule. The reference could be any material with about the same thermal mass as the sample, which undergoes no transformations in the temperature range of interest. When the sample undergoes a transformation it will either absorb or release heat. The thermocouple will detect and indicate if the transformation is “exothermic” on a plot of temperature versus time. Heating rate is an important consideration in this investigation. Slower heating rate will more accurately depict the onset temperature of a transformation. Furthermore two transformations which are very close in temperature range may be mistaken for a single transformation under rapid heating rate. Measurements were performed using a resistant heated tube furnace. A graphite cylinder was place around the alumina crucible both to be used as reference but also to prevent convection in the melt. A constant heating and cooling rate of 10°C/min was used. Argon is used as shield gas. Temperature was measured by PtRh10%-Pt thermocouples. Calibration was performed with pure silver resulting in corrections of about +3°C and the measuring device work with an accuracy of ± 1°C. Sample

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dimension was 7 mm in diameter, 14 mm in height with a 3 mm wide and 9 mm deep centre hole drilled to fit the thermocouple. Data were sampled 20 times every second.

5000 5200 5400 5600 5800 6000 6200 6400 6600 6800 70001300

1350

1400

1450

1500

Time [s]

Tem

pera

ture

[ °C

]

exothermal process

referencesample

Figure 1: Temperature vs. time curve (DTA)

2.2.2 “In-situ” solidified hot tensile tests The high temperature parameter known as the transition temperature from ductile- to

brittle fracture, TDB can be used as indicator if a metal is crack sensitive. In the study on crack formation during solidification of liquid metal this is a technique to find proper material data for further mathematical modelling. This temperature is measured during solidification and cooling of the metal and correspond to the point when the metal start to have a brittle mechanical behaviour. At temperatures below the transition we find plastic behaviour with necking and ability to high elongation and strains leading to ductile fractures. Low transition temperatures, below the metal solidus temperature indicate crack sensitivity.

This specific temperature can be measured by high temperature tensile testing of in-situ solidified samples. The technique is from the beginning used in the work of Rogberg and Fredriksson, and additionally developed by Karin Hansson.7 The concept is to have a mirror furnace, which can be inserted into a tensile testing machine. The mirror furnace gives a limited heating zone by its focus, which makes it possible to melt a small part of the sample and let it solidify and cool to the tensile test temperature. This technique has the advantage that one can melt the sample and let it solidify with a controlled cooling rate which can simulate the real conditions in the casting process with regards to cooling rate, microsegregation etc.. The mirror furnace consists of three gold plated ellipsoidal reflectors with a halogen lamp inserted at the focal point. Using a maximum power of 1020 W and the centre of the test specimen placed in focal point a 5mm long melting zone was formed and held in place by the surface tension. The temperature was measured with thermocouples of type S made of Pt-Pt10%Rh placed in

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the centre of the melting zone. A quartz tube with argon atmosphere is placed around the specimen to protect it from oxidation. The reflectors and tensile rods were water-cooled. Elongation is measured by an extensometer. Force is measured with a load cell of 5kN in a range of ± 2.5kN. Force, position of the piston, elongation and temperature is sampled by the computer controlling the tensile test machine. Tensile specimens of 40 mm length and 4 mm diameter were heated up to TL+5° and then cooled down with a constant rate of 2°C/s. When reaching the tensile test temperature the tensile test is performed during isothermal conditions with a constant pulling rate of 0.5mm/s which corresponds to a strain rate of 0.1/s if we assume the deformation zone to be of the same length as the heated zone i.e. 5mm.

Results from tensile tests are analyzed by the reduction of area (RA), ultimate tensile stress (σb), elongation (lvdt) and strain to fracture (ε). The fracture surface with zero ductility is used as indication of transition between ductile and brittle fracture. The corresponding temperature is the transition temperature, TDB.

%1000

10 ⋅−

=A

AARA (A1 is a mean value over three measurements)

0

max

AF

b =σ

⎟⎟⎠

⎞⎜⎜⎝

⎛ ∆+=

0

0lnl

llε

2.3 Macrosegregation study Macrosegregation occurs during solidification due to relative movement or flow of

segregated liquid and solid. There are numerous causes of fluid flow and solid movement in casting processes. One reason for this movement of segregated liquid may be density differences of the metal due to temperature or variations in composition. The hot liquid metal becomes cooler close to the chill surfaces and its density increase causing downward movement. Liquid being enriched by rejected solutes with higher density compared to the bulk composition will flow downward and the opposite will happen when low mass elements enrich the liquid. During ingot casting, the most common macrosegregations are the positive, negative and channel segregations8,9. Positive segregation means that the concentration of alloying element exceeds the average bulk concentration. Negative segregation is instead a local lack of alloying element. The positive segregation is often found at the top and is the result of segregated liquid flow toward the top,10,11 and the negative zone with more pure material at the bottom of the ingot is explained by

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sedimentation of equiaxed crystals formed in the bulk liquid.12 Approximate at one third from the surface the so-called A-segregates are found. They have the shape of pencil-like channels, filled with alloy which contains high concentrations of the alloying elements. Liquid jet streams melt the dendrite network allowing for channels to form and be filled by enriched liquid flowing upward due to its lower density.

Drilling for chemical composition was made on ingots along the central axis and at several cross sections at different heights. The results are plotted to obtain the segregation pattern over C and S. Sulphur prints which reveal the S as dark marks are used to analyse segregation channel placement and number density. These results are used as in-data for numerical analysis.

2.4 Inclusion study

2.4.1 Inclusion characteristics Inclusions studied in this present work were analysed with emphasis on its

precipitation process and morphology. Our goal was trying to present a model on inclusion formation based on measurements on liquid supersaturation, inclusion characteristics and solidification process. Comparisons were made between different crystal morphologies such as the columnar and equiaxed crystals and metallurgical processing prior casting. Inclusions were characterized by use of scanning electron microscopy (SEM) and element concentrations are measured by electron probe micro-analysis (EPMA). From SEM images we determine type, shape and size distribution. By use of energy dispersive scanning (EDS) we analyse inclusion chemistry. EPMA results on concentration levels are used to study the precipitation process.

2.4.2 Electron Probe Micro-Analysis Inclusions were investigated by Electron Probe Micro-Analysis, EPMA. This

technique has been known since late 1950. The first instrument was in place at The Swedish Institute of Metals Research in the beginning of the 60’s. An electron beam is focused to a 1 µm point by aid of electromagnetic lenses. The beam is used to penetrate the surface (2 µm deep in steel materials) influencing the atoms to enable emitted x-ray to analyse element specific wavelengths. Intensity is proportional to element amount. To enable this type of measurement some modification must be made to the original SEM equipment. Much higher, 500 – 1000 times, currents are used compared to conventional SEM and sliding sample holder render possible surface scan. Calibration for each element is made by mapping reference materials with known spectrums. One sample with low element concentration and one sample with high element concentration are used. One can choose to scan step wise in the range from 1×1 µm to 50×50 µm. The micro distribution of selected elements are traced and transformed into a two dimensional colour picture, each colour representing a certain amount. Each point represents a mean value of the element concentration. All data about references, coordinates and analyses are saved to be used for further investigation such as line-scan,

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mean values of certain structure areas or to better bring out low content of interesting elements.

2.5 Inoculation in nodular cast iron The same model on precipitation process used in the study on inclusions in steel

alloys can be used when evaluating the graphite formation during solidification of cast iron. Precipitation of graphite nodules during solidification in nodular cast iron was studied to better understand inoculation efficiency. Three different additives were used, FeSi – Ca, FeSi – (Ca, Ce) and FeSi – (Ca, Ce, S, O). The same mirror furnace used for hot tensile test was used to melt samples which solidify with a controlled cooling rate. The samples melt by heat conduction while heating the graphite holder to about 1400 °C. All samples were cast in Zr2O-crucibles. A 250 mm long and 30 mm wide quartz tube is placed around the specimen set-up and was continuously filled with argon gas to protect it from oxidation. Quenching was made in water. Inoculation seeds were placed at the bottom of the crucible before melting. A total of six tests from each alloy have been studied. One from each alloy representing the total solidification interval. Five were quenched within the solidification interval (1150 – 1010 °C). Maximum heating- and cooling rate was set to 300°C and 60 °C/min respectively. Samples were kept for about one minute at the maximum temperature. Temperature was sampled every 0.4 s and measured in the centre of each sample. The time from completely liquid to start of cooling is about 100 s for all tests. A maximum temperature of about 1375 – 1360 °C has been recorded in the melt and is regarded as the inoculation temperature.

Samples were polished to a 3 um diamond paste finish and etched in 2-5% Nital, 2% Pikrin or 5%Br – 95%Methanol solution. Studies were concentrated towards microstructure, nodule count, nodule size distribution, inoculants efficiency and inclusion characteristics. Nodule count and structure fractions were made by means of point counting in 5×-magnifications. The inclusions are analysed by EDS in SEM.

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3 Results

3.1 Macrosegregation Figure 3 shows a plot over the segregation ratio, C/C0 of carbon and sulphur in the

rectangular ingot. As to be expected the macrosegregations follow the well known behaviour. In the figure, the concentration along three horizontal lines, representing three height levels, from surface to centre are shown to the left, the centreline segregation is shown in the middle and the position of each drill sample are shown to the right. A sulphur print of the corresponding surface is shown in figure 4. In this ingot, an increase of the segregation ratio is seen toward the top. The horizontal lines show that the composition are even toward the surface at the mid to lower levels but has clearly unstable segregation ratio towards the centre at the higher level, which is close to the hot-top region. This instability coincides with the A-segregates seen in the sulphur print in figure 4. At the bottom of the ingot a somewhat increasing negative segregation is found.

-300 -200 -100 00

1

2

3

C/C

0

Top

-300 -200 -100 00.5

1

1.5Middle

C/C

0

-300 -200 -100 00.5

1

1.5Bottom

C/C

0

0.5 1 1.5 2 2.5

Centreline

C/C0

-500 0 5000

200

400

600

800

1000

1200

1400

1600

1800

Ingot width [mm]

Ingo

t hei

gth

[mm

]

CS

Figure 3: Segregation ratio of S and C. Sample location is shown in the right figure.

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Figure 4: Sulphur print of vertical cross-section of rectangular ingot.

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3.2 Inclusions characteristics Three types of inclusions were found. To the left in figure 5 we see a duplex oxy-

sulphide inclusion where MnS has grown on Al2O3-nucleis, the centre picture is a MnS inclusion and to the right in the figure we see the Al2O3. The inclusions are found in all three materials but are different in type, shape and size. MnS found in the ingots are classified as type III and I since they are sharp edged to spherical. Continuous cast material contains type II MnS which are smooth and stringy. The oxides are all faceted. Inclusion size and number density are found to be a function of crystal morphology. Columnar structure favour precipitation of MnS and duplex oxy-sulphides. Equiaxed crystal zone contains these inclusions too but to less extend together with pure Al2O3, most of which are very small in size.

Figure 5: Duplex oxy-sulphide, pure MnS and pure Al2O3 in high-sulphur tool steel ingot

3.3 Inoculation in nodular cast iron Nodule size measurements are presented in figure 6. The un-inoculated base iron

samples show a low number of large nodules. The number of nodules then increase at a size of 20 – 25µm, and an additional increase is seen at a size of 5 – 10µm. Addition of Ca and S – O treated inoculants lead to a more even distribution in comparison to addition of FeSi – (Ca, Ce). FeSi – (Ca, Ce, S, O) which show a peak at 5 – 20µm is the additive with most small sized nodules.

This is explained by formation of nodules at the beginning and at the end of solidification. A support for this is found in the thermodynamic calculations [Fig 13a and 14a, supplement 3]. These show that MgO and MgS inclusions are formed at the addition of Mg and nodules are nucleated on those. MgO and MgS inclusions are also formed at the end of solidification acting as nucleation sites for graphite nodules. Addition of inoculants containing Ce or other elements with high affinity to oxygen shows a quite different nodule distribution. There is hardly any favourite size and the distribution is more even [Fig. 6, supplement 3]. This might be explained by sulphides and oxides are repeatedly formed during solidification process.

11

0

10

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40

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5-10

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10-1

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0µm

40-4

5µm

45-5

0µm

Nu

mb

er

of

nod

ule

s(Ca)(Ca, Ce)(Ca, Ce, S, O)(pure base iron)

Figure 6: Nodule size distribution

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4 Discussion Results from supplement 1, 2 and 4 are to be considered as one investigation. We

made different types of experiments and analyses in order to investigate the origin of defects such as segregation, crack formation and inclusions in cast ingots and continuous cast slabs. The experiments are made on the same steel quality through out the work of this thesis. Experiments were designed and tried out in order to achieve the correct material data for further mathematical modelling.

The segregation pattern in the work made in supplement 2 follows the general view on segregation behaviour. It is clear that the areas at which segregation channels are found are prone to house high number density of inclusions. This is supported by EPMA analysis over the area. The channel formation with remelting of the dendrite network causing liquid enrichment with alloy elements being pushed from the solid into the liquid with higher solubility enables the needed supersaturation to form these inclusions. Results on macrosegregation and type of inclusions has been further analysed in supplement 4 with emphasis on the inclusion formation process. The investigation shows two different types of inclusions, (1) duplex inclusion where the core constitute an oxide inclusion with a MnS outer shell and (2) pure MnS and Al2O3 inclusions in different morphologies. It also shows that inclusion precipitation has crystal structure dependence which is a function of solidification process. At higher solidification rates the oxides incorporate into the solid and no small nuclei are available for MnS precipitation.

The theoretical model used in supplement 3 and 4, based on well known segregation models, analyses the precipitation behaviour of inclusions during solidification process. The method to find at what point in the solidification process inclusions are expected to precipitate are based on calculations with the Scheil and Lever rule in combination with calculations on solubility products for nucleation. The theoretical results are compared with experimental observations from EPMA and microstructure investigations. The experimental result supports the calculations that Al2O3 inclusions precipitate before MnS and thus acting nucleation sites for MnS. Results from the structure investigation of Steel C also support this since this material contains less MnS and no oxides. This could be an effect of the metallurgical treatment prior to casting.

In the same way as for the duplex inclusions found in the steel ingots the graphite nodules precipitates on oxide nuclei. Pure base iron show uneven graphite nodule size distribution compared to the iron inoculated with S and O treated additives. When adding inoculants with Ce or elements with high affinity to oxygen the nodule size distribution change from being uneven with a low number of large nodules to a more flat distribution with hardly any favourite size. This is explained by repeatedly formed sulphides and oxides during solidification process.

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5 Conclusions Ingot casting, continuous casting and inoculation experiments has been performed

with the intention to study crack formation, macrosegregation, inclusion characteristics and the effect of metallurgy and solidification process. A crack prediction model used to calculate temperature and elastic stresses was presented together with a model to calculate the influence of the A-segregates in large ingots. The formation of inclusions in steel alloys has been analysed by use of a new way to calculate the precipitation as function of supersaturation and solidification process. The same model can also be applied in the case with inoculation of cast iron.

Macrosegregation in the steel experiment in this study is similar to the general view of segregation pattern, a negative segregation zone in the lower parts of the ingot and an increasing composition toward the top. The primary source of the macrosegregation in the ingot cast material is the transport of segregated liquid in the A-segregates. These areas of saturated metal contain inclusions, mainly oxides and sulphides.

Inclusions are formed in the interdendritic areas where the liquid is highly enriched. Sulphides and oxides are the most common types together with a duplex oxy-sulphide where the MnS grow on oxide nuclei. Inclusion characteristics depend on crystal morphology which can be related to casting process. Probably this is an effect of the solidification rate, both in ingot and continuous cast processing. Calculations by aid of homogeneous nucleation theory together with segregation calculations show that precipitation of oxides start at a solid fraction of about 0.7 and the MnS not until a solid fraction of 95% is reached. Results from structure analysis support the idea that oxides act as nucleation sites for MnS precipitation. In the study on nodular cast iron it shows that oxides present in the liquid promote nodule precipitation. During solidification re-nucleation of oxides and graphite nodules result in a flat nodule size distribution with the beneficial effect of less chill i.e. less problem with micro porosity.

Clean Steel: Part One

Abstract:

Steel cleanliness is an important factor of steel quality and the demand for cleaner steels increases every year. The so-called clean steel generally is the steel in which the content of impurity elements, such as phosphorus, sulphur, total oxygen, nitrogen, hydrogen (including carbon sometimes) and inclusions are very low. The improvement of steel cleanliness has therefore become a more and more important subject in the development of ferrous metallurgical technology, and also an important task for the iron and steel producers.

Steel cleanliness is an important factor of steel quality and the demand for cleaner steels increases every year. The so-called clean steel generally is the steel in which the content of impurity elements, such as phosphorus, sulphur, total oxygen, nitrogen, hydrogen (including carbon sometimes) and inclusions are very low. The improvement of steel cleanliness has therefore become a more and more important subject in the development of ferrous metallurgical technology, and also an important task for the iron and steel producers.

The demand for better mechanical properties of steels was urging steel producers to improve cleanliness of their final products. In order to obtain the satisfactory cleanliness of steel it is necessary to control and improve a wide range of operating practices throughout the steelmaking processes like deoxidant- and alloy additions, secondary metallurgy treatments, shrouding systems and casting practice.

Due to the vague nature of the term "clean steel", some authors imply that it is more precise to refer to:

• steels with low levels of solutes as "high purity steels" • steels with low levels of impurities that originate from the re-melting scrap as "low residual steels"

steels with a low frequency of product defects that can be related to the presence oxides as "clean steels".

It has been well known that the individual or combined effect of carbon [C], phosphorus [P], sulphur [S], nitrogen [N], hydrogen [H] and total oxygen (T.O.) in steel can have a remarkable influence on steel properties, such as tensile strength, formability, toughness, weldability, cracking-resistance, corrosion-resistance, fatigue-resistance, etc. Also, clean steel requires control of non-metallic oxide inclusions and controlling their size distribution, morphology and composition.

The control of the elements mentioned above is different for different performance demands. Those impurity elements also vary with different grades of steel. Table 1 lists the influence of common steel impurities on steel mechanical properties which means that some element is harmful to certain steel grades, but may be less harmful or even useful to another steel grades.

For examples for IF steels, the content of carbon, nitrogen, total oxygen and inclusions should be as low as possible in order to get good flexibility, high "r" value, perfect surface quality etc. In other hands the high quality pipeline steel requires ultra low sulphure, low phosphorus, low nitrogen, low total oxygen content and a certain ratio of Ca/S.

Element Form Mechanical Properties Affected

S, O Sulfide and oxide inclusions

Ductility, Charpy impact value, anisotropy 　 Formability (elongation, reduction of area and bendability) 　

Cold fo　 rgeability, drawability Low temperature toughness 　 Fatigue strength 　

Solid solution Solid solubility (enhanced), hardenability 　

Settled dislocation Strain aging (enhanced), ductility and toughness (lowered)

Pearlite and cementite 　 Dispersion (enhanced), ductility and toughness (lowered) C, N

Carbide and nitride precipitates

Precipitation, grain refining (enhanced), toughness 　

(enhanced) Embrittlement by intergranular precipitation 　

P Solid solution Solid solubility (enhance　 d), hardenability (enhanced) Temper brittleness 　 Separation, secondary work embrittlement 　

Table 1: Influence of typical impurities on mechanical properties

As we mentioned before, steel cleanliness depends on the amount, morphology and size distribution of non-metallic inclusions. The inclusions generate many defects and many applications restrict the maximum size of inclusions so the size distribution of inclusions in steel products is also important. For certain applications where stringent mechanical properties are required the internal cleanliness of steel is very important. Table 2 shows the cleanliness requirements for various steel grades.

Steel product Maximum allowed impurity

fraction Maximum allowed inclusion

size

IF steels [C]≤30 ppm, [N]≤40 ppm, T.O.≤40

ppm [C]≤10 ppm, [N]≤50 ppm

Automotive and deep-drawing Sheets

[C]≤30 ppm, [N]≤30 ppm 100 µm

Drawn and Ironed cans [C]≤30 ppm, [N]≤40 ppm, T.O.≤20

ppm 20 µm

Alloy steel for Pressure vessels [P]≤70 ppm

Alloy steel bars [H]≤2 ppm, [N]≤20 ppm, T.O.≤10

ppm

HIC resistant steel sour gas tubes [P]≤50 ppm, [S] ≤10 ppm

Line pipes [S]≤30 ppm, [N]≤50 ppm, T.O.≤30

ppm 100 µm

Sheets for continuous annealing [N]≤20 ppm

Plates for welding [H]≤1.5 ppm

Bearings T.O.≤10 ppm 15 µm

Tire cord [H]≤2 ppm, [N]≤40 ppm, T.O.≤15 10 µm

ppm

Non-grain-orientated Magnetic Sheets

[N]≤30 ppm

Heavy plate steels [H]≤2 ppm, [N]=30-40 ppm, T.O.≤20

ppm Single inclusion 13 µm

Cluster 200 µm

Wires [N]≤60 ppm, T.O.≤30 ppm 20 µm

Table 2: Cleanliness requirements for various steel grades

As Table 2 shows for sheets used for car body, carbon [C], nitrogen [N], and total oxygen (T.O.) are each required to be very low. For sheets for tin plate application, total oxygen is not only needed below 20 ppm, but the size of the non-metallic inclusions in steel has to be less than 20 µm.

For steel cord used in tires, the size of non-metallic inclusions in steel has to be less than 10 μm and even smaller (5 µm) for TV shadow masks. For ball bearings, in order to improve their fatigue-resistance properties, T.O. in steel has to be below 10 ppm and the size of non-metallic inclusions has to be less than 15 µm. For meeting the specification of increasingly improved toughness for petroleum pipeline and of Hydrogen Induced Cracking (HIC) resistance for the transport of sour natural gas, the sulphur [S] content in steel has to be extremely low, less than 10 ppm.

Steel cleanliness is controlled by a wide range operating practices throughout the steelmaking processes. These include the time and location of deoxidant and alloy additions, the extent and sequence of secondary metallurgy treatments, stirring and transfer operations, shrouding systems, tundish geometry and practices, the absorption capacity of the various metallurgical fluxes, and casting practices.

A one of the steelmaking process routes for the production of clean steels is outlined in Figure 1.

Figure 1: The process route for the production of clean steels

Clean Steel: Part Two

Abstract:

Non-metallic inclusions, which are undesirable components of all steels, play an important role with respect to their effect on the steel properties. Controlling inclusions in steel is closely connected with the concept of "clean steel". The improvement in steel properties by control of non-metallic inclusions plays an important part in defending the applications of steel against newer competitive materials.

Non-metallic inclusions, which are undesirable components of all steels, play an important role with respect to their effect on the steel properties. Controlling inclusions in steel is closely connected with the concept of “clean steel”. The improvement in steel properties by control of non-metallic inclusions plays an important part in defending the applications of steel against newer competitive materials. The aims of the metallurgist are to eliminate undesirable inclusions and control the nature and distribution of the remainder to optimize the properties of the final product.

Generally, non-metallic inclusions in steel normally have a negative contribution to the mechanical properties of steel, since they can initiate ductile and brittle facture. Among various types of nonmetallic inclusions, oxide and sulphide inclusions have been thought harmful for common steels.

All steels contain non-metallic inclusions to a greater or less extent. The type and appearance of these non-metallic inclusions depends on factors such as grade of steel, melting process, secondary metallurgy treatments and casting of steel. Because of this, it is of particular significance to determine how pure the steel is. The term steel cleanness is relative one, since even steel with only 1 ppm each of oxygen and sulfide will still contains 109 -1012 non-metallic inclusions per ton. From the viewpoint of “cleanness” all steels are “dirty”.

Non metallic inclusions in steel are the cause for dangerous and serious material defects such as brittleness and a vide variety of crack formations. However, some of these inclusions can also have a beneficial effect on steels properties by nucleating acicular ferrite during the austenite to ferrite phase transformation especially in low carbon steels. According to definition, the non-metallic inclusions are chemical compounds of metal with nonmetal which are present in steel and alloys like separated parts.

Classification of non-metallic inclusions

Non-metallic inclusions are divided by chemical and mineralogical content, by stableness/stability and origin. By chemical content non-metallic inclusions are divided into the following groups:

• Oxides (simple: FeO, MnO, Cr2O3, TiO2, SiO2, Al2O3 etc.; compound: FeOFe2O3, FeOAl2O3, MgOAl2O3, FeOCr2O3 etc.)

• Sulphides (FeS, MnS, CaS, MgS, Al2S3 etc.; compound: FeSFeO, MnSMnO etc.) • Nitrides (simple: TiN, AlN, ZrN, CeN etc.; compound: Nb(C,N), V(C,N) etc, which can be found in

alloyed steels and has strong nitride-generative elements in its content: titanium, aluminum, vanadium,

cerium etc.) • Phosphides (Fe3P, Fe2P etc.)

The majority of inclusions in steels are oxides and sulphides. Among various types of nonmetallic inclusions, oxide and sulphide inclusions have been thought harmful for common steels. Usually, nitrides are present in special steels (stainless steels, tool steels) which have elements with a strong affinity for nitrogen (e.g. chrome, vanadium), which create nitrides.

Figure 1 shows sulfides and oxides of non metallic inclusion in steel.

Figure 1: Non-metallic inclusion in steel: oxides-dark gray and sulfides-light gray

By mineralogical content oxygen inclusions are divided into the following groups:

• Free oxides – FeO, MnO, Cr2O3, SiO2 (quartz), Al2O3 (corundum) etc. • Spinels-compound oxides which are formed by bi- and tri-valent elements as a ferrites, chromites and

aluminates. • Silicates which are presented in steel like a glass formed with pure SiO2 or SiO2 with admixture of

iron, manganese, chromium, aluminum and tungsten oxides and also crystalline silicates.

Depending on the melting temperature, in liquid steel non-metallic inclusions are in solid or liquid condition.

As mentioned above the majority of inclusions in steels are oxides and sulfides. Sulfides in steel have been paid much attention because their treatment is an important problem in the steelmaking process. They affect on the properties of the final products by their deformation during the steel working process; especially their morphology has a significant effect on the steel properties.

According to analysis based on the steel ingots containing 0.01-0.15% S, the morphology of MnS can be classified into three types:

1) Type I is a globular .MnS with a wide range of sizes, and is often duplex with oxides. 2) Type II has a dendritic structure and is often called grain-boundary sulfide because it is distributed as chain-like formation or thin precipitates in primary ingot grain boundaries. 3) Type III is angular sulfide and always forms as monophase inclusion.

Most of the above mentioned sulfides are formed both during the process of secondary metallurgy or the

solidification process. Recently, with the development of steelmaking technology, the sulfur concentration in steel was lowered drastically. Also, the continuous casting technology of steels with higher cooling rate than the ingot casting almost replaced the ingot casting.

So, the sulfides in the modern commercial steel are usually formed on solidification process or in solid steel during the subsequent cooling process. For example, the Widmanstätten plate-like MnS2, is formed in solid steel and Figure 2 shows the common morphology of MnS in conventional continuously casting steel, including the globular duplex oxide–sulfide (particle A, B and C) and the Widmanstätten plate-like MnS (particle D).

Figure 2: Typical duplex oxide–sulfide inclusion (particle A, B and C) and plate-like MnS (particle D) in conventional continuous casting silicon steel.

Numerous examples of the effect of non-metallic inclusions on steel properties show the importance of the behavior of the inclusions as well as of surrounding metal matrix during plastic working of steels. The aims of the metallurgist are to eliminate undesirable inclusions and control the nature and distribution of the remainder to optimize the properties of the final product.

An attempt by using program ABACUS was performed to model the behavior of slag inclusions and their surrounding matrix material during hot rolling and hot forging of hardenable steels. It is shown that it can be helpful for studying the behavior of inclusions, which is difficult or even impossible to obtain from a conventional experiment.

Figure 3 shows the effective strain contour during plastic deformation. Three regions of strain concentration (red) can be seen and a trihedral void (white region) close to the round inclusion is formed. The strain concentrations arise at the inner surface of the matrix. Another interested thing is that two edges of the pore tend to emerge and a bonding is formed. The difference in mechanical properties between the matrix and the inclusion is found to be the primary reason to create a void. The weak bonding at the interface between the

matrix and the inclusion seems to facilitate to open the void.

Figure 4 shows the effect of rolling temperature on the relative plasticity index during hot rolling of steels. The relative plasticity index of inclusion increases while the rolling temperature rises. There exists a transition region, where the relative plasticity index changes rapidly. This trend agrees with the existing experimental results.

Figure 3: Void formation close to the inclusion.

Figure 4: Effects of rolling temperature on the relative plasticity index.

Clean Steel: Part Three

Abstract:

The presence of non-metallic oxide inclusions is a major cause of incompatibility between the attainable and desirable level of cleanliness in many grades of commercial steel. Generally, inclusions degrade the mechanical properties of the steel and thereby reduce the ductility of the cast metal and increase the risk for mechanical and/or corrosion failure of the final product.

The increasing demand in recent years for high-quality steel products has led to the continuous improvement of steelmaking practices. There is a special interest in the control of non-metallic inclusions due to their harmful effect on the subsequent stages and their great influence on the properties of the final product. Through the control of the amount, size and chemical composition of the inclusions it is possible to obtain a final product of good quality. The control of the formation of non-metallic inclusions and the identification of their constituent phases are of extreme importance for the obtaining of clean steels.

The presence of non-metallic oxide inclusions is a major cause of incompatibility between the attainable and desirable level of cleanliness in many grades of commercial steel. Generally, inclusions degrade the mechanical properties of the steel and thereby reduce the ductility of the cast metal and increase the risk for mechanical and/or corrosion failure of the final product.

Oxide inclusions originate from two sources:

• residual products resulting from intentionally added alloying elements to deoxidize the molten steel after oxygen treatment (endogenous or micro inclusions);

• products resulting from reactions between the melt and atmosphere, slag, or refractory (exogenous or macro inclusions).

Among various types of nonmetallic inclusions, oxide and sulphide inclusions have been thought harmful for common steels.

Alumina inclusions occur as deoxidation products in the aluminum-based deoxidation of steel. Pure alumina has a melting point above 2000°C, i.e., these alumina inclusions are present in a solid state in liquid steel. The addition of calcium to steel which contains such inclusions changes the composition of these inclusions from pure alumina to CaO-containing calcium aluminates.

As it can be see from Figure 1, the, melting point of the calcium aluminates will decrease as the CaO content increases, until liquid oxide phases occur at about 22% of CaO, i.e., when the CaO.2Al2O3 compound is first exceeded at 1600°C. The liquid phase content continues to increase as CaO content rises further and is 100% at 35% of CaO. The minimum melting temperature for the liquid calcium aluminates is around 1400°C, i.e., such liquid calcium aluminates may be present in liquid form until, or even after, the steel solidifies.

Most grades of steel are treated with calcium using either a Ca-Si alloy or a Ca-Fe(Ni) mixture, depending on

the silicon specification. This treatment is made after trim additions and argon rinsing.

In most melt shops the cored wire containing Ca-Si or Ca-Fe(Ni) injection system is used in the calcium treatment of steel. The melting and boiling points of calcium are 839°C and 1500°C respectively. During calcium treatment, the alumina and silica inclusions are converted to molten calcium aluminates and silicate which are globular in shape because of the surface tension effect. The change in inclusion composition and shape is known as the inclusion morphology control.

Figure 1: Binary system CaO-Al2O3

The calcium aluminates inclusions retained in liquid steel suppress the formation of MnS stringers during solidification of steel. This change in the composition and mode of precipitation of sulphide inclusion during solidification of steel is known as sulphide morphology or sulphide shape control.

Several metallurgical advantages are brought about with the modification of composition and morphology of oxide and sulphide inclusions by calcium treatment of steel, as for instance:

• To improve steel castability in continuous casting, i.e. minimize nozzle blockage • To minimize inclusion related surface defects in billet, bloom and slab castings • To improve steel machinability at high cutting speeds and prolong the carbide tool life • To minimize the susceptibility of steel to re-heat cracking, as in the heat-affected zones (HAZ) of

welds

• To prevent lamellar tearing in large restrained welded structures • To minimize the susceptibility of high-strength low alloy (HSLA) linepipe steels to hydrogen-induced

cracking (HIC) in sour gas or sour oil environments. The Ca content in the final product can be controlled within the range of 15 to 20 ppm

• To increase both tensile ductility and impact energy in the transverse and through-thickness directions in steels with tensile strengths below 1400 MPa

When calcium is injected deep into the melt, the following series of reactions are expected to occur to varying extents in Al-killed steels containing alumina inclusions:

Ca + O = CaO (1)

Ca + S = CaS (2)

Ca + (x+1/3)Al2O3 = CaO·x Al2O3 + 2/3[Al] (3)

Depending on the steel composition, the manner of calcium adding in steel bath and other process variables, there will be variations in the conversion of alumina inclusions to aluminates inclusions, the smaller inclusions will be converted to molten calcium aluminates more readily than the larger inclusions.

Thermodynamically, if sulfur or oxygen is dissolved in the steel at moderate levels, or if Al2O3 inclusions are present in steel, calcium will react with oxygen or sulfur until the contents of reactants are very low (< 2ppm). One of the critical questions is whether or not calcium added to steel will react with sulfur by reaction (2) and form CaS or modify Al2O3 to liquid calcium aluminates by reaction (3).

The formation of calcium sulfide can occur if calcium and sulfur contents are sufficiently high. Since calcium has higher affinity for oxygen than for sulfur, the addition of calcium initially results in a more or less pronounced conversion of the alumina into calcium aluminates until the formation of calcium sulfides starts as the addition of calcium continues.

Calcium sulfides are solid at steelmaking temperatures and result in nozzle clogging similar to that caused by alumina. As can be observed from the Figure 2, the conversion of alumina into calcium aluminates occurs until all the inclusions in the steel are present only in liquid form.

Figure 2: Change of inclusions composition during calcium additions

To prevent nozzle clogging in continuous casting by solid inclusions, calcium is added to steel to modify inclusions and desulfurize the steel. Calcium will convert solid alumina (Al2O3) inclusions into lower melting point calcium aluminates, which will help prevent the clogging of the casting nozzles. However, when calcium is added to steel, it will also react with oxygen and sulfur and modify the sulfide inclusions. If the sulfur content of the steel is high, calcium will react with sulfur forming solid CaS, which could clog up the continuous casting nozzle.

The Figure 3 shows influence of calcium treatment on the type of inclusions formed and its relationship with nozzle clogging.

Figure 3: Influence of calcium treatment on the type of inclusions formed and its relationship with nozzle clogging

Calcium treatment cannot be applied to all kinds of steel. For those with high requirement on formability, such as automobile sheet, calcium treatment is not suitable, because this treatment causes the formation of calcium aluminates inclusion which is hard. Therefore, for those kinds of steel, the method of improving molten steel´s purity is usually taken to optimize castability. Through controlling carry-over slag from melting furnace, deformation treatment of ladle slag, metallurgy in tundish, protective casting and other measures, purity of steel is guaranteed and total oxygen content in molten steel decrease

Control of Sulfur in Melts

Abstract:

Sulfur has a strong surface activity both in binary (Fe-S) and ternary (Fe-C-S, Fe-Si-S) alloys. It has been concluded from the results of numerous studies that sulfur can exist in two forms in molten iron: in one case it forms an interstitial solutions, and in other partially substitutional solutions. It has been found that the activity of sulfur increase substantially when carbon and silicon are present in the melt. This explains why pig iron can be desulfurized more readily than steel.

Sulfur (S) is a typical metalloid. The radius of sulfur atom is 1.05 Å. It easily acquires two electrons to form an ion S2-. The coefficient of diffusion of sulfur in liquid iron is 0.74x10-4, 13x10-4, 1.9x10-5 cm/s, according to various experimental data. Sulfur has a strong surface activity both in binary (Fe-S) and ternary (Fe-C-S, Fe-Si-S) alloys. It has been concluded from the results of numerous studies that sulfur can exist in two forms in molten iron: in one case it forms an interstitial solution, and in other partially substitutional solutions. It has been found that the activity of sulfur increase substantially when carbon and silicon are present in the melt. This explains why pig iron can be desulfurized more readily than steel.

A good understanding of the desulphurization of hot metal and liquid steel has been developed in terms of slag-metal reactions, based on a number of studies of the partition of sulfur between liquid slag and liquid iron. These results show that a highly basic slag, high temperature and reducing conditions enhance desulphurization via slag-metal reactions.

Sulfur Equilibrium between Liquid Iron and Slag

The desulphurization of liquid iron with slag may be examined on the basis of the following reaction (1), in which the equilibrium constant can be expressed by equation (2).

[S] + (O2-) = (S2-) + [O] .......... (1)

log K1 = as2-. ao / as. ao

2- .......... (2)

where,

• ao, as: the activities of oxygen and sulfur in liquid iron, respectively • ao

2-, as2-: the activities of oxygen and sulfur ions in slag, respectively.

The sulfur partition ratio between metal and slag is given by equation (3) according to reaction (1).

LS = (wt%S) / [wt%S] = K1 • ao2- • fs / ao • fs2- .......... (3)

where,

• fs, fs2- : the activity coefficients of sulfur in liquid iron and slag, respectively.

Since the values of K1, ao2- and fs2- cannot be determined experimentally, the sulfide capacity is defined as

equation (5) on the basis of reaction (4) and is utilized for the examination of desulphurization in iron and steel making processes.

½ S2 + (O2-) = (S2-) + ½ O2 .......... (4)

CS = (wt%S) • (PO2 / Ps)½ = K4 ao2- / fs2- .......... (5)

where,

• PO2 and Ps are the oxygen and sulfur partial pressure in atm; • K4 the equilibrium constant of reaction (4).

In order to calculate the value of CS, the equilibrium constants of reactions (6) and (7) are substituted into equation (5), and equation (8) is obtained.

½ O2 = [O] .......... (6)

½ S2 = [S] .......... (7)

log Cs = log (wt % S) ao / as + 936 / T-1.375 .......... (8)

By the use of the interaction coefficients of sulfur and oxygen in liquid iron, the values of CS can be calculated from equation (8).

As an illustration, the calculation of sulfide capacity (CS) has been given for CaO-MgO-Al2O3-SiO2 ladle slags.

In order to examine the relationship between CS and composition of slag, the following relationship is assumed to hold at certain temperatures.

log CS = α (NCaO + KMgO • NMgO + KAl2O3• NAl2O3 + KSiO2 • NSiO2)+ β .......... (9)

where:

• α, β: the constants • Ni, Ki: the mole fraction and the lime equivalent coefficient of i-component in the slag phase,

respectively.

At 1600°C, the values of KMgO, KAl2O3 and KSiO2 determined by trial and error are 0.1, -0.8 and -1.0, respectively. As shown in Figure 1, there are linear relations between log Cs and (NCaO + 0.1 NMgO – 0.8 NAl2O3 – NSiO2) at 1575, 1600 and 1650°C.

Figure 1: Plot of the log Cs against NCaO + 0.1NMgO – 0.8NAl2O3 – NSiO2 at 1575, 1600 and 1650°C

The three lines obtained by the method of least squares are shown in the figure. The intercepts of the lines were determined by the use of the slope at 1600°C because the most runs were done at this temperature. The following equation was obtained as a function of temperature.

log CS = 3.44 (NCaO +0.1NMgO – 0.8NAl2O3 – NSiO2) - 9894 / T+2.05 .......... (10)

The observed values of log CS with the correlation coefficient (R) of 0.99 and the standard deviation (σ) of 0.044 can well be expressed by the above equation.

On the other hand, log LS may be expressed by the following general equation derived from equations. (3), (8) and (10).

log LS = α ∑ Ki•Ni - log [wt%O] - β .......... (11)

where,

• ∑ Ki•Ni = NCaO +0.1NMgO – 0.8NAl2O3 – NSiO2.

The experimental data were arranged by the method of least squares according to equation (11) and the following relationship was obtained.

log LS = 3.44 ∑ Ki•Ni - log [wt% O] - 10 980 / T+3.50 .......... (12)

Relationship Between Sulfide Capacity and Theoretical Optical Basicity

Duffy and Ingram defined the theoretical optical basicity (Λ) as follows:

Λ = ∑ (xi / fi) .......... (13)

where, xi, fi : the equivalent cationic fraction and the basicity moderating parameter for the constituent cation i,

respectively.

It was verified by them that fi can be expressed by the Pauling's electronegativity xi as follows:

fi =1.36 (xi -0.26) .......... (14)

It was also found by Duffy and Ingram that log CS are linearly correlated to the theoretical optical basicity.

As an illustration of the above mentioned the experimental results performed on ladle CaO-MgO-Al2O3-SiO2 slags has been given. The values of log CS are plotted against the optical basicities at 1575, 1600 and 1650°C, as shown in Figure 2.

Figure 2: Plot of log CS against theoretical optical basicity Λ

The slopes and intercepts of the lines were determined by the method of least squares, and the following linear relationship was obtained as a function of temperature.

log CS = 14.20Λ - 9 894 / T -7.55 .......... (15)

As described above, it can be seen that both the sum of lime equivalent and the optical basicity are able to be used as parameters for the representation of sulfide capacity. From the relationship between these two parameters, the linear equation (16) from regression was obtained.

Λ = 0.24(NCaO + 0.1NMgO – 0.8NAl2O3 – NSiO2) + 0.67 .......... (16)

The correlation coefficient (R) is larger than 0.99, and the standard deviation (σ) is smaller than 0.01. This shows that the two parameters have almost the same character for the representation of sulfide capacity of slag.

Equilibrium of Oxygen Partition between Metal and Slag

It is well known that the oxygen potential exerts a large influence on the sulfur partition between metal and

slag. The equilibrium of reaction (17) was determined by the use of the value of the equilibrium constant (18) reported in literature.

(FeO) = [Fe] + [O] .......... (17)

log K17 = log ao / aFeO = -6 150 / T + 2.604 .......... (18)

Equation (18) may be modified to the following equation.

log [wt%O] = log NFeO + log γFeO - log fo - 6 150 / T + 2.604 .......... (19)

where γFeO is the activity coefficient of FeO on the basis of mole fraction.

Since log γFeO is a function of temperature and slag composition, log [wt%O] can be expressed by an equation including the term of log NFeO, then the following relationship was obtained.

log [wt%O] = 0.905 log NFeO -0.15 ∑ Ki•Ni - 6 340 / T + 3.115 .......... (20)

where, ∑ Ki•Ni = NCaO + 0.1NMgO – 0.8NAl2O3 – NSiO2.

As shown in Figure 3, the observed values of log [wt% O] agree well with the values calculated from the equation (20). By substituting equation (20) into equation (12), equation (21) is derived:

log LS = 3.59 ∑ Ki•Ni -0.905 log NFeO – 4 640 / T + 0.385 .......... (21)

The above equation can conveniently be used for the estimation of the sulfur partition between liquid iron and slag in which the oxygen content in liquid iron is not analyzed. Next, the values of γFeO calculated from the equation (19) were arranged by the way similar to that described previously, and the following relationship was obtained.

Figure 3: Comparison of the observed log [wt%O] with calculated log [wt%O] from the equation (20)

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Some Fundamental Steps in Thermomechanical Processing

of Steels* .

By Imao TAMURA**

Key words: ~hermomechanicalprocetsing; steel; deformation o f austenite; restoration process; phase transformation; microalloying; ferri te grain re- Jinement.

I. Introduction

Thermomechanical processing is a treatment of which plastic deformation is introduced into the heat treatment of steel. We can find many treatments in such processes. But controlled rolling and accelerated cooling, and the ausforming or deformation- quenching are representative and are used in practical production.

I n controlled rolling, the austenite (y) is work- hardened with plastic deformation (rolling) in high strain rate a t elevated temperatures, and then air cooled to transform to ferrite (a). Sometimes, work- hardened austenite is cooled acceleratedly to take fine ferrite grains in steel plate. -

In the ausforming, work-hardened austenite is quenched to room temperature. We can take the ausformed martensite which increases its strength and not decrease its toughness. The ausforming is not so popular in practice but the deformation-quenching is getting to be popular. The martensite in the deformation-quenching looks like as usual martensite. But, this martensite has a little bit of fine facet size because it is formed from smaller austenite grains than usual austenite. This process can abridge to ration- alize the heat treatment process of steel.

As fundamental steps in thermomechanical proc-

True strain

tn 3

Fig. 1. Schematic illustration of two types of true stress- true strain curves accompanying the steady state deformation at elevated temperatures.

(a) dynamic recovery

essing of steels, we can point out that:

1) The deformation and restoration processes of y during and after hot rolling.

2) The kinetics of y to a transformation and grain size of a.

3) The effect of cooling rate on transformation of work-hardened y to a, and grain size of a.

4) Structure of martensite in ferrous alloys.

5) Structure and properties of martensite formed from work-hardened y.

6) Hardenability of work-hardened y and recrystallized y.

I will explain on these processes based on our research results.

(b) dynamic recrystallization

II. Deformation Behaviors of Austenite during and after Hot Working

I . Restoration Process during Hot Deformation of Auste-

I t is well known that the deformation behavior of austenite a t elevated temperatures is a dynamic recrystallization type. Figure 1 shows schematically both types of dynamic recovery type and dynamic recrystallization type.

I n the case of a dynamic recrystallization type, since the rate of softening due to recovery is slow, the work-hardening can not be balanced only by dynamic recovery even at high temperatures, and the dislocation density is still gradually increased with strain, and finally the recrystallization starts to occur during deformation. The true stress is lowered beyond the peak stress by the occurrence of dynamic recrystallization and reaches a certain steady state stress level at which the work-hardening and softening due to recrystallization occurring repeatedly are balanced. The dynamic recrystallization starts at the strain of around 0.78, (E,: the strain showing the peak stress). In Fig. l(b), the strain range up to about 0 . 7 ~ ~ is the work-hardening region, the strain range between about 0.78, and 8, (the strain at which the steady state deformation starts) is the range of partially dynamic recrystallization, and the strain range after E , ~ is the range of steady dynamic recrystallization. Austenite in steels belongs to this type. In a-Fe (ferrite), since the recovery occurs very fast, the restoration is performed mainly by dynamic recovery. The peak stress (a,) and the steady state stress (a,) in the dynamic recrystallization type change with the Zener-

* Based on the Special Lecture delivered by Prof. Tamura in commemoration of his receiving Nishiyama Medal in the 113th ISIJ Meeting, April 2, 1987, a t The University of Tokyo in Tokyo. Manuscript received on May 18, 1987. 0 1987 ISIJ

* * Professor Emeritus, Kyoto University; Technical Advisor, Sumitomo Metal Industries, Ltd., Nishinagasu-hondori, Amagasaki 660.

Special Lecture ( 763 )

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[ 764 ) Transactions ISIJ, Vol. 27, 1987

Hollomon parameter (2). Z is expressed by

where, i: the strain rate (s-l) I? : a gas constant (J/mol)

Q: the activation energy for deformation (J/ mol)

7: temperature (K). a?,, ep, gs and E , are increased with increase in Z (i.e., with increase in i or decrease in 7).

Figure 2 shows the change in optical microstructure with strain in Fe-31%Ni-0.3%C austenitic alloy which was tensile deformed at 1 273 K (1 000°C) at a strain rate of 1.7 x 10-2/s. Figure 2(a) is the austenite structure before deformation. Austenite grains are equiaxed and their boundaries are smooth and planer. Many annealing twins exist in austenite grains. When deformed slightly (E= 0.08), as shown in Fig. 2 (b), the grain boundaries become finely serrated and the bulg- i n g ~ of grain boundary are observed as indicated by the arrows. At the strain slightly after E,, many small grains are newly formed along the austenite boundaries as shown in Fig. 2(c). I t is characteristic that these newly formed grains hardly contain annealing twins. Furthermore, some of the initial austenite grains becomes larger by the strain-enhanced coarsening. Figure 2(d) shows the structure in the range of steady state deformation at which the work-hardening and the dynamic recrystallization softening are balanced. This structure consists of several grains corresponding to various stages of recrystallization, i.e., from small grains which have just started to recrys- tallize to fairly large grains which have been subjected to deformation after the finishing of recrystallization and are the work-hardened state just before the next recrystallization. I n the stage of steady state deformation, the recrystallization occurs repeatedly during deformation. Dynamic recrystallized grains contain a few annealing twins a t low Z but scarcely contain

them when Z is high. I n the same way as the static recry~tallization,~) it

has been observed that there are two main mechanisms for dynamic recrystallization, i.e., the bulging mechanism and the nucleation-growth mechanism.

Dynamic recrystallization becomes easier to take place when the initial grain size (do) of austenite is smaller. Therefore, in the case of fine initial grain size, c p and E, become small even at the deformation under the same Z condition. However, a,, and a , are independent of do. The average size (d) of dynamically recrystallized grains is the only function of Z and is expressed by the following equation,')

where, A, : material's constant. A is larger for C-steel than for alloy steels, especially small for Nb-containing steel and high alloy steel. p is about 0.3-0.4 for all steels. I t should be noted that d is independent of do.

2. Static Restoration Process during Holding Right after Hot Deformation

When the hot deformation is finished a t the work- hardening stage, the work-hardened austenite is softened by the usual static recrystallization during isothermal holding after deformation. However, if the strain is too small, the recrystallization can not occur and only the recovery takes place. On the other hand, when specimens are isothermally held after deformation to the range of the steady state deformation, in which the dynamic recrystallization has occurred, the recrystallization and the grain growth can succeed to occur continuously during holding after hot deformation. Such a recrystallization is called a metadynamic or postdynamic recrystallization. Namely, the dynamically recrystallized structure is further softened by the metadynamic recrystallization. When specimens are deformed to the range of (11) in Fig.

(a) Before testing (E=O) (b) € =0.08 (c) e=0.24 (d) &=0.40

Pig. 2. Optical micrographs showing tlie formation process of dynamically recrystallized austenite in Fc-31Ni-0.3C austenitic alloy deformed at 1 OOO°C, 1.7 x 10-2/s.

Special Lecture

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Transactions ISIJ, Vol. 27, 1987 ( 765 )

l (b) in which the work-hardened austenite and dynamically recrystallized austenite have coexisted, the static recrystallization occurs in work-hardened austenite and the metadynamic recrystallization occurs in dynamically recrystallized austenite at the same time.

I n case of rolling by hot strip mill in which slabs are rolled by several passes, the static restoration process occurring during interval between rolling passes has great influence on the deformation behavior of next rolling pass. If the softening due to the static restoration is not completed during interval between rolling passes, the work-hardening is successively ac- cumulated by the subsequent rolling passes.

Fraction of softening (X) occurring during isothermal holding right after hot deformation can be evaluated by the two-step deformation test from the following e q ~ a t i o n , ~ )

where, cr,,,: the flow stress immediately before unloading

c ~ , , ~ : the yield stress on reloading after holding the specimen for various periods of time (At)

u,,, : the initial yield stress. As an example, the change in the fractional softening (X) with the holding time (At) in Fe-31 %Ni-O.S%C austenitic alloy is shown in Fig. 3. Specimens were tensile deformed a t 1 373 K (1 100°C) a t a strain rate of 4 x 10-2/s to E = 0.18 (work-hardened austenite) and e= 1.0 (completely dynamically recrystallized austenite), respectively, and then isothermally hild for various periods of time. The change in average aus-

tenite grain size with At is also plotted in Fig. 3. Microstructural changes during isothermal holding a t 1 373 K (1 100°C) are shown in Fig. 4 and Fig. 5. As shown in Fig. 4 (~=0.18) , work-hardened austenites are slightly elongated along the tensile direction (Fig. 4(a)). The austenite hardly changes by holding for 1 s (Fig. 4(b)), although the softening occurs to some extent by the static recovery as can be seen in Fig. 3. However, when specimens were isothermally held for 4 s (Fig. 4(c)), fine austenite grains which are formed by static recrystallization are partially formed mainly along the deformed austenite grain boundaries and the static recrystallization is completed by the holding of 15 s (Fig. 4(d)). Annealing twins in static recrystallized austenites are gradually developed with increase in holding time. By holding for 40 s, recrystallized austenite grains are coarsened (Fig. 3). When the austenite was much more heavily deformed under higher Z condition (work-hardened), the static recrystallized austenite grains become much finer as shown in Fig. 6. As shown in Fig. 5 (a=1.0), dynamically recrystallized austenites are characterized by irregular grain boundaries, a low density of annealing twins and the mixed grain size. With increase in holding time, austenite grain boundaries become fairly straight and austenite structure becomes almost equiaxed grains, and furthermore, the density of annealing twins is increased. This is the typical microstructural change occurring during metadynamic recrystallization process. Metadynamic recrystallization is completed by the holding for about 30 s in the case of Fig. 5, and the usual grain coarsening occurs by further holding. The grain size right after the completion of metadynamic recrystallization is smaller, as the grain size of dynamically recrystallized

Hdding time. A t (s)

Fig. 3. Change in fraction of softening (a) and austenite As-deformed (A t= 0 S) (b) At = 1 s (c) At = 4 s (d) ~t = 15 s grain size (b) with the isothermal holding time of Fig. 4. Change in optical microstructure of work-hardened austenite (de- work-hardened austenite and dynamically recrystal- formed at 1 100°C, t = 4.0 x 10-2/s, c ~ 0 . 1 8 ) during isothermal holding lized austcnite in Fe-31Ni-0.3C austenitic alloy. at 1 100°C for At in Fe-31Ni-0.3C austcnitic alloy.

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[ 766) Transactions ISIJ, Vol. 27, 1987

(a) AS-deformed ( ~ t = O s) (b) At = 1 s (c) At=2 s (d) At=30 s (a) Work-hardened austenite (detormed at 900°C,

Fig. 5. Change in optical microstructure of dynamically recrystallized austenite e=3/s, ~ z 1 . 0 ) (deformed at 1 100°C, d = 4.0 x 10-Z/s, E = 1 .O) during isothermal (h) Static recrystallized austenite (isothermally holding a t 1 100°C for At in Fe-31Ni-0.3C austenitic alloy. held a t 900°C for 50 s of structure (a)

Fig. 6. Optical micrographs showing the static recrystal-

austenite is smaller. lization occurring during isothermal holding after heavily work-hardened austenite in Fe-31Ni-0.3C

3. Efect of Alloying Elements on the lietardation of Re- austenitic alloy.

crystallization5) Table 1. The values of n and m in Eq. (4) for varlous

The effect of alloying elements on the restoration transformation conditions.

process during or after hot deformation is very important. Micro-addition of Nb or Ti acts as a re- Mechanism Site saturation Nucleation and

markable retardation effect on recrystallization due to growth -- - - - - - --

the suppression of grain boundary migration. The Site n m n m suppression of grain boundary migration due to micro-

Grain surface 1 1 4 1 alloying is caused from (1) the solute dragging effect Grain edge 2 2 4 2 due to segregation of alloying elements to the bound- Grain corner 3 3 4 3

aries, or (2) the pinning effect due to precipitates of carbonitride of alloying elements at grain boundaries. For example, Mo segregates to the austenite grain x = l - exp [ - k ( ~ ) $ ] ..... . . . . . . . . . . (4) boundaries, resulting in the suppression of grain boundary migration. Even Nb micro-addition, at where, X: the transfbrmed fraction high 1, it can retard the recrystallization by dragging effect. But, usually, Nb and T i are precipitated as fine carbonitride particles and these precipitates suppress the grain boundary migration by the pinning effect.

III. Transformation Kinetics and Grain Size of Ferrite Transformed from Austenite

I . Isothermal Transformation Kinetics of Austenite

When the ferrite or pearlite Sorms from austenite by isothermal holding below the transformation temperature, these nuclei form preferentially a t the prior austenite grain boundaries. We expressed as Eq. (4)'3 that their transformation kinetics include the effect of austenite grain size, by modification of Johnson-Mehl equation.

t : the isothermal holding time d : the austenite grain size

k ( 7 ) : the rate constant which depends only on transformation temperature and directly corresponds to the C-curve of T T T diagram (isothermal transformation diagram).

The values oS the exponent m and n are listed in Table 1.6) These values depend on both the transformation mechanism and type of nucleation site. Here the transformation which progresses with nucleation and growth over almost entire range of transformation is called " nucleation and growth ". The transformation where all the nuclei are nucleated a t the early stage in the reaction and the progress of transformation is controlled mostly by their growth is called " site saturation ". An austenite grain is assumed to

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Transactions ISIJ, Vol. 27, 1987 ( 767)

be a tetrakaidecahedra, and the nucleation sites a t austenite grain boundary can be classified into surfaces, edges and corners as are shown in Fig. 7.697) Of course the actual nucleation sites are not limited to any one of such types, and the value of m would be the weighted average of the various nucleation sites and not always be an integer.

According to Umemoto and Tamura,e) the values of n and m, and transformation modes were experimentally shown as in Table 2. I t is mentioned in Table 2 that the mode of y +a transformation is site saturation, but i t means the large number of nucleation o f a would occur at the beginning of transformation, which would not be real "site saturation ". Sometimes the nucleation site is varied by supercooling (accelerated cooling). I t seems that the nucleation site of a is moved gradually from edge to surhce by decrease in transformation temperature with cooling rate, as described later.

2. Continuous Cooling Transformation Kinetics of Austenite

Scheils) considered the case where the austenite is cooled continuously through its transformation temperature. He assumed that austenite consumes its fractional nucleation time and when the sum of a number of such fractions equals unity, the transformation starts, i.e.,

Fig. 7. Schematic diagram of austenite grain as a tetrakaidecahedra and showing 3 types of nucleation sites offerritc.

'Table 2. The values of r2 and m for various transformation modes.

Transformation n m Nucleation -

Pearlite 4 2 Edge nucleation Nucleation and growth

1:errite 1 1 Surface nucleation near Site saturation

Bainite 4 0 .6 Grain boundary, Inside grain

Now extend this Scheil's additivity rule to the entire range of transformation. I n this case the incu- bation time ( ~ ( 7 ) ) is replaced with the time required for the reaction to reach a certain fractional completion X by isothermal holding at temperature T. If we combine Eq. (5) with Eq. (4)

where, Q ( T ) = -dT/dt: a cooling rate. This equation expresses the transformed fraction at temperature T during cooling from austenite condition with a cooling rate Q(T).6t9) Furthermore, from Eqs. (4) and (6) we obtain

1 Te K(T1)1/7~ t ( 7 ) = - J - - - d T ' = t., ...... (7)

K T ) T Q(T')

This equation tells that the fraction transformed by cooling from 7, to T with cooling rate Q ( T ) is equal to that obtained by the isothermal holding at temperature T for time t (7) . Thus we call t (T) as an equivalent cooling time, t,,. t,, means the isothermal holding time which gives the equivalent transformed fraction with continuous cooling. The curve produced by the successive plot oS such t,, at each temperature during cooling with cooling rate Q ( T ) will be called an equivalent cooling curve.'O) The transformation behavior for a given cooling curve can be predicted from the T T T diagram by plotting such curve on a T T T diagram.

3. Ferrite Grain Size Transformed from Austenite 1. Expression for Ferrite Grain Size (Isothermal)

In general, ferrites nucleate at austenite grain boundaries and grow into austenite grains. The grain size of ferrite formed from austenite is directly related with the process of transformation.") The nominal ferrite grain diameter D, can be expressed with the total number of ferrite grains uncleated throughout transfbrmation per unit volume of austenite, n,, as12)

where, I : nucleation rate per unit volume oS austenite

X: volume Sraction transformed Thus the problem of obtaining the exprcssion for ferrite grain size is the problem of how to express n,>.

When ferrites nucleate a t the austenite grain boundary surfaces, they grow as ellipsoid with the aspect ratio of 3: 1 1 3 ) as shown in Fig. 8. The total number of rerrite grains nucleated per unit area of austenite grain boundary surface is

9 n, = I,J; exp (-Tzaz~stz dt = - - ) 3:2-(':')

Nuclcation and growth

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ferrite g r a i n

g ra in boundory

r=3Cl(t- 2)

Fig. 8. Schematic drawing of a ferrite grain nucleated at austenite grain boundary surface.

where, I , : the nucleation rate per unit area oL austenite grain surface

a : parabolic rate constant for thickness. The ferrite grain size D, is given as

where, Sg .b : the austenite grain surface area per unit volume

and can be expressed by austenite grain diameter, D,, as

Substituting Eqs. (10) and (12) into Eq. ( l l ) , the ferrite grain size and the y / a transfbrmation ratio can be expressed as

and

In the above derivation it was assumed that the nucleation rate and the parabolic rate constant of growth are time independent and one nuclei grows and forms one grain.

Similarly the D, and y / a transformation ratio in the case of homogeneous and grain edge nucleation has been obtained asl1)

where, I h : the homogeneous nucleation rate per unit volume

I ,: nucleation rate per unit length of austenite grain edge.

I t should be noted that the equations introduced above indicate that ferrite grain size and y/a transformation ratio depend on the ratio of square root of the nucleation rate to the parabolic rate constant in all three types of nucleation sites. Furthermore Sor

grain boundary nucleation a small austenite grain leads to a small ferrite grain. The effect of austenite grain size on ferrite grain size depends on the type of nucleation site, and it becomes large in the following ascending order: homogeneous, grain surSaces, and grain edges.

I t should be noted that the grain refinement on austenite turns out to be less effective in ferrite grain refinement. For instance in the case of grain surface nucleation D, is proportional to D',I3 in Eq. (13). This means that when D, is reduced to 118 of the initial, D, would be reduced only into 112.

Figure 914) shows experimental results of isothermal transformation in a steel (0.2 % C, 0.0016 0/, B). The steel was austenitized between 1 000 and 1200°C to change its austenitic grain size from about 53 to 88 pm, and then transformed at 700, 680 and 660°C into Ferrite. The results are recognized as the dependent of y-grain diameter that a nucleates mainly at grain edges at 700°C, at grain edges and surfaces at 680°C and at the surfaces at 660°C. 2. Estimation of Ferrite Grain Size Formed by Continu-

ous Cooling Transformation It can be assumed ferrite nucleates preferentially at

austenite grain surfaces by supercooling. A continuous cooling transformation can be considered as the sum of short time isothermal holdings at successive temperatures as is schematically shown in Fig. 10.

During cooling, ferrite grains are nucleated and grow at each temperature with the corresponding nucleation and growth rates, as is schematically shown in the Fig. 10 (bottom). During ferrite transformation carbon enriches in retained austenite, however, since ferrite nucleation occurs mostly in the early stage of transformation, it can be assumed that both the nucleation and growth rates of ferrite are a function of an instantaneous temperature only and not a function of thermal history.

Based on this assumption, the number of ferrite nuclei (N,) nucleated at temperature I, during cooling, N, is given as Eq. (16).

where, Q(T) (= -dT/dt) : a cooling rate. Ferrite grains once nucleated grow thereafter with a corresponding growth rate at each temperature during cooling. Thus, at temperature T, the radius of a ferrite grain on an austenite grain surface nucleated at temperaturc T,, is given as Eq. (1 7).

The austenite grain boundary area occupied at tcm- pcrature 7 by a ferrite grain nucleated at temperature T, is given as Eq. (18).

I t will be convenient to use the concept of extended area which is the sum of the areas of austenite grain

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Transactions ISIJ, Vol. 27, 1987 ( 769 ]

A u s t e n i t e G r a i n Size

Fig. 9. An example of D,- or D,/D,,-D, rclationsllip for S15C steel.

surface occupied by ferrite grains, assuming that fcr- rite grain never stop growing and that ferrite grains keep nucleating a t the same rate during transformation as well as on untransformed austenite grain surfaces. The extended grain boundary area occupied by all the ferrite grains nucleated at T , during cooling is given as Eq. (1 9).

Thus the grain boundary arca fraction occupied by the ferrite grains nucleated during cooling horn A3(T , ) to T is givcn by Eq. (20).

{J:$?&; dl-' d T , ............ 1 1 (2 0)

The total number of ferrite grains nucleated per unit area of austenite grain boundary surfacc is given by Eq. (21).

Ferrite grain size is given as Eq. (22).

o ~ * ~ ~ ~ * TI T2 T3 Tl,

d ( h ) d ( T 2 ) d (T3 ) ~ ( T L ) I s (T ) Is(T2) Is(T3) Is(TL )

Fig. 10. Diagrams to explain the relation between isothermal and continuous cooling transformation and schematic drawing showing the formation process of ferrite during continuous cooling.

673 773 873 973 1073 1173

TEMPERATURE, K

Fig. 11. Calculated parabolic rate constant and nucleation rate I, as a function of' temperature for the steel studied. (S15C)

where, Sg.,,(=4/(dn D,)): the austenite grain surface area pcr unit volume.

From these equations, ferrite grain s i ~ e can be calculated.

I t is fairly difficult to estimate of I , ( T ) and tt(T) by some kinds of calculation, but it is not impossible. By using several assumptions and h y p o t h e s i ~ , ~ ~ ' ~ - ~ ~ ) I , ( T ) and a ( T ) are roughly calculated as shown in Fig. 1 lZ4) as a function of transformation temperature. In this figure, I , shows a maximum value at 913 K (640°C) and n shows a maximum value at

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( 770 ) Transactions ISIJ, Vol. 27, 1987

873 K (600°C) fbr S15C steel (0.15 % C, 0.4 % Mn). Figure l2Z4) shows a comparison of calculated

values and experimental values for C15C steel. Wc can get a roughly agreement, that is, a-grain size is getting smaller with cooling rate increases. The Sor- mulations are :

Dtv = Q-0.17(D,)1'~cal~ulat~d) ... ......... ...... (23)

Do = 5.7Q-0.2'3(D,)0.46 (experimental) . . . . .. . . .(24)

Figure 1324) shows experimental results.

IV. Ferrite Grain Refinement by Work-hardening of Austenite

I . Ferrite Nucleation in Deformed Austenite (Isothermal)

Figure 14 is typical micrographs showing the effect of austenite defhrmation on the progress of ferrite transformation.Z5) Specimens of an Fe-0.12C-0.04- Nb-0.04V steel were austenitized at 1 200°C for 30 min, rolled 30 or 50 % in reduction by a single pass a t 840°C (in the unrecrystallized temperature range) and held at 680°C for 15, 30 and 180 s. The ferrite transformation is clearly accelerated by deformation

C O O L I N G IRATE , K is

Fig. 1 l'erritc grain s i ~ e as a function of cooling rate lor the specimens of D, = 15 1 pm. (S 15C) Fig. 13. Ferrite grain size as a function of Q-0.2'3Dt4G.

Not Rolled Rolled (30%) Rolled (50%)

(a) Non-deformed, 15 s (b) 30 % rolled, 15 s (c) 50 % rolled, 15 s

(d) Non-deformed, 30 s (e) 30 % rolled, 30 s (f) 50 % rolled, 30 s (g) Non-defbrmed, 3 min (h) 30 % rolled, 3 min (i) 50 0/6 rolled, 3 min

Fig. 14. Optical micrograph showing the efrect of austenite deformation on the proccss of' ferrite transfbrmation a t 680°C in a Fe-O.12C-0.04Nb-0.04V steel.

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ofaustenite. I n the non-deformed specimen austenite grain boundaries are the predominant nucleation sites of ferrite, but in the deformed specimens nucleation occurs not only on austenite grain boundaries but also within grains. Moreover, the nucleation rate on grain boundaries is substantially accelerated.

Figure 15 shows various types of nucleation sites of ferrite observed in deformed specimens.25) Figure 15(a) shows ferrite grains nucleated a t austenite grain boundaries. The ferrite grains grow into both sidcs of austenite grains, and the prior austenite grain boundaries disappear once ferritcs are nucleated. Fig- ure 15(b) shows the ferrite grains nucleated a t the annealing twin boundaries. In contrast to the non- deformed specimens, the annealing twin boundaries in the deformed specimens are the preferential nuclcation site of ferrite. The ferrite grains nucleated at annealing twin boundaries grow mostly into one side, and the traces of prior annealing twin boundaries remain after transformation. Figure 15(c) shows ferrite grains nucleated intragranularly. These ferrites were considered to be nucleated on the dislocations at the cell wall distributed almost uniformly inside the grains. Figure 15(d) shows the ferrite grains nucleated at decormation bands.

The increase in the ferrite nucleation rate per unit volume of austenite by deformation is attributed (1) to the increase in the austenite grain surface by elongation of grains, (2) to the increase in the nucleation rate per se per unit area of grain surface, and (3) to the formation of additional nucleation site such as annealing twin boundaries, deformation bands, etc.

The grain boundary surface area of austenite per unit volume of a specimen increases by deformation. A spherical austenite grain (Fig. 16(a)) becomes an ellipsoid as shown in Fig. 16(b) by applying the rolling with reduction p. The surface area of a grain with an unit radius before rolling is given as

The surface area of the grain after rolling with reduction p is given as

S,.,,(p) = St''-* [(q x J: J 1 -(2p-pvsiFzB do. -1'1-*

l h c ratio of the surface area before to after rolling, q(=S,.b/SO,.,,), is plotted in Fig. 17 as a function of rolling reduction p. Since the volume of the sphere shown in Fig. 16(a) is equal to that of the ellipsoid shown in Fig. 16(b), the ratio q in Fig. 17 is equal to that obtaining to grain surface areas per unit volume before and after rolling. I t is seen that the increase in the grain surface area per unit volume with the increase in p is quite small when p is smaller than 0.5. For instance 50 % reduction by rolling increases the surface area by 25 %. This increase is equivalent to that obtained by only 20 0/, reduction in austenite grain size.

The number of ferrite grains nucleated on austenite grain boundaries is substantially increased by deformation. Figure 1825) shows the ferrite grains nucleated on the austenite grain boundaries: (a) non- deformed, (b) 30 % rolled, and (c) 50 % rolled, respectively. From these pictures the acceleration of ferrite nucleation rate on austenite grain boundaries by deformation can be estimated. The number of ferrite grains intersect per 1 mm of austenite grain

(a) (b)

(a) Kcfore rolling (sphere) (b) After rolling (ellipsoid)

l i g . 16. Tlic shapc change ol' a n austcnite grain by the rolling with reduction p .

Rolled (30V0) 30s Rol led(30%) 15s

(a) W ~ ~ c . l t , ~ ~ t i o ~ ~ ;it gr~ii11 I>o~~n(laric~s

(I,) Nuc.l(.atiun on ; i~ni~.nl i~ig t \ \ . i r ~ I )ou~~cl; ir ic~~

(c.) S~iclcntion i n d c gr;iinr ((1) Suclc:~tion on clt.IL1.1nntion bantls

Fig. 15. Opticxl n>ic.rogl.nplis showi~lg lllr \ . ; ~ r i o ~ ~ i types 01' ~iucleatiun bites 01' IC3rrite obier\,ccl i l l drll1.111rd specimens. Rolled(50%) 15s Rolled(30%) 15s

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strain & (=-ln(1-p)) 0 -0.5 - 1.0

1 7 --

1.6 --

1. 5 --

1.4 --

1.3 --

o 0.1 0.2 0.3 0.4 0.5 0.6 0.7

Rolling Reduction p

Fig. 17. The ratio of austenitc grain surfacc arca bcforc rolling to that of after rolling as a f~~nction of rolling rcduction $.

boundaries werc counted as 4 1, 214 and 330 for 0, 30 and 50 rolled specimens, respectively. The numbers of ferrite grains nucleated per unit area of austenite grain boundary surface are proportional to the square of these numbers. Furthermore the number of ferrite grains nucleated per unit area of austenite grain surface, n,, is proportional to the square root of I, as is shown in Eq. (10). Since a is not much changed by deformation, the ferrite nucleation rate per unit area of austenite grain boundary surface is estimated to be increased by 740 ((214/41)4) and 4 200 ((330/41)4) times faster by 30 and 50 % rolling, respectively.

As is shown above, the nucleation of ferrite on austenite grain boundaries is substantially accelerated by deformation. The mechanism of enhanced nucleation is considered to be as follows. The deformed austenite grain boundaries may have many ledges (or steps) shown schematically in Fig. 19(a). If ferrite nucleates a t the corner or such a ledge, as is shown in Fig. 19(b), the activation energy for nucleation becomes B/n times less than that fbr nucleation on a planar grain boundary (if the surface energy is assumed to be equal for all kinds of interfaces). Thus if a large number of ledges are produced on grain boundaries by deformation, considerable increase in the nucleation rate can be expected. As is shown in Fig. 15(b) the annealing twin boundaries in work-hardened austenite also act as the preferential nucleation site of ferrite. The reason of this can be considered to be as follows. At the coherent twin boundaries before deformation a large number of ledges are produced on the twin boundaries by deformation. Thus deformed twin boundarics have

Not Rolled

680 'C

3min

Rolled (5O0Io)

680 O C

30s

(a) Non-deformed, 3 min at G80°C

(b) 30 0/, rolled, 3 s at 680°C

(c) 30 % rolled, 3 s at 680°C

1 18. Optical micrographs showing the efycct of deformation on the number of ferrite formed on austenite grain boundaries.

$.,,.,j- rcrrite

,A-F

(a) A ledgc on grain boundary

(b) Ferrite nuclcation at a lcdgc

Fig. 19. Scllematic drawing showing ferrite nucleation at a ledge on deformed austcnite grain boundaries.

the similar structure with that of austenite grain boundary surfaces and act as the preferential nucleation site of ferrite.

Similar mechanisms have bcen proposed for the enhanced nucleation on deformed austenitc grain boundaries by several researchcrs. Sandberg and Robertsz7) have suggested that deformation leads to the formation of grain boundary serrations or bulges, and these act as nucleation sites. Amin and Pic- keringZ8) proposcd that deformed austenitc grain boundaries appear to exhibit greater lattice mismatch, possibly because of greater lattice distortion and a higher dislocation density near deformed grain boundaries. These would be the predominant nucleation sites of ferrite and hence accelerate the nucleation.

Deformation structures within austenite grains have bern known to act as the effective nucleation sites of

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ferrite and play an important role in the ferrite grain refinement. The deformation structure is generally called as deformation bands. These regions are characterized by a high density of cells consisting of tangled dislocations. Dislocations are favor nucleation sites through its stress field. Consequently ferrite nucleation appears to occur frequently at the deformation structures.

2. Growth of Ferrite in Deformed Austenite (Isothermal)

When ferrite grains grow into deformed austenite, their growth rate would be accelerated by the following two factors. One is due to the enhancement of diffusivity by dislocations and supersaturated vacancies. Another is due to the additional driving force associated with stored energy. Since carbon, whose diffusivity controlls the growth rate of ferrite in HSLA steels, is an interstitial element, the former effect is considered to be small. The latter [actor is also not so large estimation by some kinds of calculation (acceleration is less than 10 %).z4) We can find that the ferrite growth rate is slightly enhanced by the deformation of austenite. The acceleration of growth rate by deformation is much less than enhancement of nucleation rate for the transformation kinetics.

3. Ferrite Grain Size firmed ,from Dfirmed Austenite (I~othermal)2~)

For simplicity we assume;

(i) that austenite grain boundary surfaces are the nucleation site and the additional nucleation sites produced by deformation are the deformation bands and annealing twin boundaries,

(ii) that the nucleation rates per unit area of deformation bands and annealing twin boundaries are equal to that of the grain boundary surface,

(iii) that the total area of deformation bands and annealing twin boundaries per unit volume of austenite is quadratic with respect to

the strain E ( = -In (1 -P)), i.e.,

S,.,i(p)+Sl.l,(p) = A(- In (1 -p))2 ......... (27)

where, A is a constant, and

(iv) that the ratio JI,(p)/a(p) can be expressed as a function of strain e as,

where, Io , a,: the nucleation rate and parabolic rate constant of growth in the non-deformed specimen, respectively

B : a constant. An exponent of 1.3 came from an experiment for Fig. 18.

As discussed above the ferrite grain size formed from non-deformed austenite in the case of grain surface nucleation can be expressed by Eq. (1 1). From the assumptions made for deformed specimen Sg.b in Eq. (1 1) should be replaced by the effective interfacial area, S,, which is the sum of the areas ofgrain bound-

ary surface, annealing twin boundary and deformation bands ( S , ( P ) = S g . b ( p ) + S d . ~ ( P ) + S t . h ( P ) ) . Fur- thermore, if both the nucleation and the growth rates are a function of strain, the ferrite grain size in the deformed specimen is considered to be expressed as

Thus if d I o / w o and the coefficients A and B are determined from an experiment, the ferrite grain size formed from deformed austenite can be calculated. Figure 20 shows an example of such calculation. Here the calculated ferrite grain size using the commonly observed values of JIi/cuo, A and B is plotted as a function of the rolling reduction, P. From this figure it is seen that the ferrite grain size formed from work-hardened austenite decreases drastically with the increase in a rolling reduction for a small rolling reduction and decreases gradually for a large rolling reduction.

4 . Ferrite Grain Refinement by Accelerated Cooling

Accelerated cooling after hot rolling is currently being realized as a further advanced thermomechanical treatment in hot rolling proce~s.29>3~) This cooling process is characterized by accelerated cooling in a 7-+a transformation range just after controlled rolling. I t has been shown that the accelerated cooling refines the ferrite grain size and thus further improves both the strength and toughness.

The effect of accelerated cooling on ferrite grain refinement can be seen in a non-deformed specimen. Figure 13 and Eq. (24) show the effect of cooling rate on ferrite grain size in an Fe-0.15C

This dependence of ferrite grain size by cooling rate means that when a cooling rate becomes 10 times faster ferrite grain diameter becomes almost half (0.55).

The principle of grain size refinement by accelerated cooling can be understood by considering the effect of transformation temperature on ferrite grain size in an isothermal transformation. As mentioned in the previous section the grain size of isothermally formed ferrite depends on the ratio of the square root

Da(prn)lnDa True Strain & = -ln(l-p) 1.0 2.0

40

30.-

20-3 0 Austenite Grain Size

10..

2 .o

5-

1.0 0 .I .2 .3 .4 .5 .6 .7 .8

Rolling Reduction, p

Fig. 20. Calculated ferrite grain size formed from work- hardened austenite as a function of rolling reduction for various austenite grain sizes.

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of nucleation rate to the parabolic rate constant for various types of nucleation sites. Thus when the ratio d n a increases with the decrease in the transformation temperature, the grain size of an isothermally formed ferrite decreases with the decrease in the transformation temperature.

The continuous cooling transformation can be considered as the sum of short time isothermal holdings at successive temperatures, as described in previous section and derived two equations of (21) and (22) which show n, and D,, respectively.

As a whole, austenite deformation results in a large increase in the ratio of the nucleation rate to the parabolic rate constant and leads to ferrite grain refinement. The accelerated cooling further refines a ferrite grain size.

V. Structures and Properties of Martensite Formed from Work-hardened Austenite

I . Morfhology of Microstructure Composed of Lath Mar- tensite Formed from Work-hardened Austenite

Among various types of ferrous martensite (such as lath, lenticular, etc.) the lath martensite is the most important one because most of the commercial heat- treatable steels produce lath martensite. Lath martensite structure is characterized by packets and blocks within a original austenite grains as shown schematically in Fig. Zl.31) I t is known that the basic microstructural unit for controlling mechanical properties especially, toughness of lath martensitic steel is size of packets or blo~ks,~2-3~) and thus the refinement of packets or blocks improves the mechanical properties of lath martensitic structure. Therefore, we examined the change in lath martensite morphology with the amount of deformation of austenite using an 18 Ni maraging steel (18.3 % Ni, 9.0 % Co, 5.0 % Mo, 0.7 % Ti, 0.015 % C) and Nb bearing low carbon steel (0.12 % C, 0.038 % Nb, 0.039 % V).

An example of results in 18 Ni maraging steel is shown in Fig. 22.35) In this case, specimens were austenitized at 1 523 K for 120 s in Ar gas atmosphere (austenite grain size of 200 pm was obtained by this treatment) and cooled to 773 K (above M d temperature) and deformed in tension to various strains at 773 K at strain rate of 3 x 10-2/s and then air cooled

to room temperature to produce lath martensite. It appears from Fig. 2233) that the number of packets in an austenite grain markedly decreases and hence the packet size becomes large with an increase in the amount of deformation of austenite. The block becomes narrower and finely segmented, and the block boundaries become heavily waved with an increase in the amount of ausforming. Figure 23 shows the change in the mean packet size and block width of lath martensite in 18 Ni maraging steel with the amount of ausforming. These results indicate that the morphology of ausformed lath martensite is characterized by large packet size and less-developed blocks. And also it is characteristic that the block boundaries are waved in the case of ausformed martensite. Similar tendency was observed in low carbon Nb steel (HSLA steel) as shown in Fig. 24, in which the specimen was rolled at 1 123 K by 50 % after austenitizing a t 1 573 K for 1.8 ks and then water quenched to produce lath martensite. These microstructural characteristics might arise from the preferential formation of lath martensites with the same habit plane along the most active slip plane (one of four (11 11, planes) which is macroscopically bent by heavy deformation.

As shown above, the packet size of ausformed martensite is large compared with that of conventionally quenched lath martensite. I t was observed that the ausformed lath martensites form along the curved {111}, planes as shown in Fig. 22(f)35) and Fig. 24 (b).") This indicates that the crystal planes of austenite (and hence, of martensite) are bent. I n such a situation, there would be a possibility that the propagation of cleavage crack is suppressed since the plane of cleavage fracture ({001],,) is heavily waved. This might be one reason for some improvements in toughness by ausforming in spite of the coarsening of packet size.

2. Mechanical Proferties of Ausformed Martensite

In order to account for the good ductility of ausformed steels, the relation between the ductility of deformed austenite and the ductility of ausformed martensite was studied using an Fe-25.4XNi-0.38XC alloy (Ms=236 K).37) In this experiment, specimens were solution treated at 1 423 K for 10.8 ks. Since the M s temperature of this alloy is below room tem-

Fig. 2 1. Schematic illustration showing the morpholog- I

0.2% Carbon Steel

I ical characteristics of lath martensite structure l8Ni Maraging Steel in 0.2 0/, C steel and 18 % Ni maraging steel.

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( a ) ( c ) : Packed structure (electrirally etched with CrO, 10 g+H,O 90 cc)

( d ) ) : Block structure (chemically etched with FeCI, 10 g+HCl 30 cc+H,O 120 cc)

Fig. 22. Optical micrographs showing the change in packed or block structure of lath martensite with amount of deformation oT austenite (ausformed by tennion at 773 K) in 18Ni maraging steel.

1 0 0.1 0.2 0.3 0.4 0.5 0.6

True Strain

Fig. 23. Change in the packet size and block width of lath martensite with the amount of deformation of austenite (ausformed at 773 K ) in 18Ni maraging

steel. D, = 200 pm.

perature, the solution treated specimen was austenite single phase at room temperature. These specimens were deformed at 473 K by rolling to various amounts of reduction. The rolling temperature of 473 K was applied in order to avoid the formation or strain- induced martensite during deformation. Some of specimens of deformed austenite were subzero cooled into liquid nitrogen to produce the ausformed martensite. Tensile properties were measured in both specimens of deformed austenite and ausformed martensite with an Instron-type tensile machine a t 353 K

(a) 1 573 K for 1.8 ks -+ W.Q.

(b) 1 573 K for 1.8 ks --t 50 % rolled at 1 123 K -t w.9_.

Fig. 24. Optical micrographs of non-ausformed (a) and ausformed (b) lath martensite in Nb steel. Etchant: 2 % nital.

which is above Md temperature. Figure 25 shows the results of tensile strength of the

deformed austenite and the ausformed martensite. Both of tensile strengths of austenite and martensite are increased with an increase in the amount of prior-deformation of austenite. Figure 26 shows the relation between the elongation and the amount of prior-deformation of austenite. In the case of deformed austenite, the elongation is markedly decreased (from 60 to 3 %) with an increase in the amount of prior-derormation. On the other hand, it is characteristic that the elongation of ausformed martensite is

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( 776) Transactions ISIJ, Vol. 27, 1987

Test temperature: 353 K P, T: Parallel and transverse specimens against

rolling direction, respectively S: Subzero cooling to 77 K, namely speci-

mens are of martensite

Fig. 25. Relation between tensile strength and the amount of prior-deformation of austenite at 473 K in Fe-25Ni-0.38C alloy.

hardly effected by the amount of ausforming and is almost same as the elongation of non-ausformed martensite. The ausformed martensite exhibits almost constant elongation of about 8- 10 %, irrespective of the amount of ausforming. Therefore, the elongation of deformed austenite becomes smaller than that of ausformed martensite, when the amount of deformation (ausforming) is more than about 30 % in the case of the present alloy. I t must be emphasized that the heavily deformed austenite is rather enhanced in both strength and ductility by martensitic transformation.

These results indicates that the dislocations in austenite produced by deformation are inherited and that the stress concentrated region a t the obstacles in austenite produced by heavy deformation are healed by the martensitic transformation itself. As a reason for the healing effect, it can be considered that the stress concentration in austenite may be released by the formation of preferential variants of martensite at the stress-concentrated areas.

3. Effect of Ausforming Temperatures on the Increase of Strength ff Ausformed Martensite

In the most published papers, the strength of aus-

0 15 30 L5 60

Pre-deformation ( % )

Test temperature: 353 K

Fig. 26. Relation between elongation and the amount of prior-deformation of austenite at 473 K in Fe- 25Ni-0.38C alloy.

formed martensite has been studied mainly by functions of the amount of deformation and the deformation temperature below Al. And it is generally believed that the increase in the strength of ausformed martensite is closely related to the amount of work- hardening of austenite. However, few attempts have been made to relate the strength of ausformed martensite to the flow stress of austenite during ausforming at elevated temperatures.

I n order to make clear this point, the experiment was carried out using austenitic steels (a high Mn steel : 13.9 % Mn, 0.68 % C) and Fe-high Ni alloy (32.4 % Ni 0.005 % C), and martensitic steels (SNC 815: 3.0 % Ni, 0.8 % Cr, 0.16 % C).38) Ausforming was performed with a tensile machine a t various temperatures between 773 and 1 173 K a t strain rate of 1 . 7 ~ 10-21s. After deformation, specimens were rapidly quenched by water spray. These ausformed specimens were tensile tested at room temperature at strain rate of 1 . 7 ~ 10-3/s to measure the yield stress of deformed austenite (austenitic steels) and ausformed martensite (martensitic steels) a t room temperature.

Figure 27 shows the relation between the increase in 0.2 % proof stress at room temperature (Aa;.,) of ausformed austenite and the amount of work-hardening (Aar) of austenite by ausforming at elevated temperatures in austenitic steels. I t appears that the yield stress increase of austenite (Aa;.,) at room temperature is linearly proportional to the amount of work-hardening of austenite (Aur) by ausforming in both steels, and all data lie on the same straight line, irrespective of ausforming conditions such as deformation temperature and the amount of strain. This

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Transactions ISIJ, Vol. 27, 1987 (777)

result indicates that the strength of ausformed austenite at room temperature is uniquely determined only by the amount of work-hardening by ausforming. This probably suggests that the strengths of austenite at both of high and room temperatures are mainly determined by the dislocation density.

Figure 28 shows the relation between the increase of 0.2 % proof stress of ausformed martensite (Au;:~) (i.e., the difference of 0.2 % proof stress at room temperature between ausformed martensite and conventionally quenched martensite) and the amount of work-hardening (Aar) in austenite at elevated temperature in SNC 815 steel. I t is characteristic that the obtained data lie on the different curves by the ausforming temperature. I t should be emphasized that the d a f , is not determined only by the Aar, but functions of both Aar and ausforming temperature. Even at the same Aar, the effect of ausforming on the strengthening of martensite is different with the de-

a t H.T.

E

Fig. 27. Relation between the increase in 0.2 O//o proof stress a t room temperature of ausformed austenite (do:,) and the amount of work-hardening of austenite (Aor) by ausforming in austenitic steels.

formation temperature, i.e., do;:, is increased with an increase in deformation temperature. Furthermore, when the ausforming temperature is fixed, the contribution of A 0 7 on the An;:, is small at smaller Aur

range and becomes large at larger Aar range. Same tendency was also observed in SKD 61 steel. These results strongly suggest that, in addition to the dislocation density in austenite, the dislocation configuration in austenite is also an important factor for increase in strength of ausformed martensite.

Figure 29 shows the examples of transmission electron micrographs of austenite in a high Mn austenitic steel deformed to the same amount of work-hardening (Aor= I0 kg/mm2) at different temperatures. As can be seen, the total dislocation density seems to be almost same, but the dislocation configuration is different from each other. With increase in the deformation temperature, the dislocation configuration is changed from uniformly distributed pile-up to dislocation cells. I n this austenitic steel, despite the difference of dislocation configuration the 0.2 % proof

Deformat ion temperature: 30 - A -993K

-1023K

25 . o -1073K L1

E E 0 b 2 0 - Y 0

e rT 15- C

0 - N 0

Q 0

5 .

0 5 10 15 2 0 bar at H.T. kg.rnrn-=

Fig. 28. Relation between the increase in 0.2 % proof stress of ausformed martensite (dot,;) and the amount of work-hardening of austenite (do?) by ausforming in SNC 813 steel (3Ni-0.8Cr-0.16C).

Fig. 29. Transmission electron micrographs of austenite deformed to the same amount of work-hardening of austenite ( A u ~ = 10 kg/mm2) a t strain rate of 1.7 x 10-2/s a t (a) 1 073 K, (b) 1 023 K and (c) 873 K, respectively, in Fe-14Mn-0.7C austenitic steel.

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stress of austenite at room temperature is the same as shown in Fig. 27, because the Aar is the same. How- ever, as shown in Fig. 28, in the case of ausformed martensite, Aa$, is increased as the ausforming temperature is raised even at the same Aor. Therefore, it can be considered that the dislocation cell structure in prior austenite is very effective in strengthening of ausformed martensite through the martensitic transformation, and pile-up or uniformly distributed dislocations in prior austenite have little effect on strengthening of ausformed martensite.

VI. Hardenability of Work-hardened Austenite and Recrystallized Austenite

As described in the former section, the work-hardened austenite can transform easily by diffusional process, because the rate of nucleation of ferrite is high. It means that the hardenability of steel is decreased with work-hardening of the austenite.

A bearing steel (SUJ2: 1.33 % Cr, 0.9 % C) with spheroidized carbide particles was austenitized at 850°C at which the steel was in two phase region of y and carbide, and deformed by compression of 60 % (s=3x 10°/s), and then immediately or after holding in some intervals of time (At) at the same temperature as the deformation quenched into water. The hardness at the center portion of SUJ2 specimen at room temperature after quenching changed with holding time (At) as shown in Fig. 30.

The austenite of SUJ2 with spheroidized carbide particles can dynamically recrystallized easier than without carbide particles. The work-hardened austenite of SUJ2 has poor hardenability and becomes its hardenability recover to usual SUJ2 with holding time (At) during the static recrystallization is progress- ing as shown in Fig. 30.

Ausforming is quenching method of work-hardened austenite. Such an austenite has poor hardenability. The steel for ausforming should have high hardenability by including several alloying elements of Ni, Cr, Mo, etc.

For direct quenching it seems that the austenite is deformed and recrystallized and then quenched. In this case, we can not expect to increase strength and toughness of martensite as ausforming does.

I I I I W.Q.

6001 100 1 200 300 400

~t Hold~ng Time ( 5 )

Fig. 30. Change in hardness of water-quenched specimen with holding time after deformation at 850°C in SU.J2.

VII. Conclusion

I have explained the abstract of kinetics, structure and mechanical properties of transformation products of ferrite and martensite formed from work-hardened austenite. Because of too much contents, I am afraid that it would be difficult to understand.

Thank you very much for your kind attention.

REFERENCES

T. Maki, K. Akasaka and I. Tamura: Proc. Int. Conf. Thermomechanical Process. Microalloyed Austenite, ed. by A. J. DeArdo et al., TMS-AIME, Pittsburgh, (1981), 217. T. Maki, S. Okaguchi and I. Tamura: Strength of Metals and Alloys (ICSMA-6), ed. by R. G. Gifkins, Pergamon Press, Melbourne, (1982), 529. P. A. Beck and P. R. Sperry: J. Appl. Phy~., 21 (1950), 150. H. J. McQueen and J. .J. Jonas: Plastic Deformation of Materials, ed. by R. T. Arsenault, Academic Press, N.Y., (1975), 393. J.J. Jonas and I. Weiss: Metal Sci., 13 (1979), 238. M. G. Akben, I. Weiss and J . J. Jonas: Acta Metall., 29 (1981), 111. M. Umemoto, N. Komatsubara and I. Tamura: J. Hent Treating (ASM), 1 (1980), No. 3, 57. J. W. Cahn: Acta Metall., 4 (1956), 449. E. Scheil: Arch. Eisenhiittenwes., 12 (1935), 565. M. Umemoto, N. Nishioka and I. Tamura: J. Heat Treat- ing (ASM), 2 (1981), 130; Trans. Iron Steel Inst. Jpn., 22 (1982), 629. M. Umemoto, N. Komatsubara and I. Tamura: Proc. Int. Conf. Solid-Solid Phase Transf., ed. by H. I. Aaronson el al., TMS-AIME, Pittsburgh, (1981), 11 11. M. Umemoto, H. Ohtsuka and I. Tamura: Proc. Int. Conf. High Strength Low Alloy Steels, ed. by D. P. Dunn and T. Chandra, Univ. Wollongong, Wollongong, (1984), 96. G.H. Gulliver: J. Inst. Met.,19 (1981), 145. J. R. Bradley, J. M. Risbee and H. I. Aaronson: Metall. Trans. A., 8A (1977), 323. M. Umemoto, Z. H. Guo and I. Tamura: Proc. Int. Conf. HSLA Steels (HSLA Steels, Metallurgy and application^, HSLA Steel '85), ed. by J. M. Gray et al., ASM-INTER- NATIONAL, Beijing, (1985), 97. C. Zener: J. Appl. Phys., 20 (1949), 950. M. Hillert and L. I. Staffanson: Acta Chem. Scand., 24 (1970), 3618. H. I. Aaronson and H. A. Domian: Trans. AIME, 236 (1966), 781. C. Wells, W. Batz and R. F. Mehl: Trans. AIME, 188 (1950), 553. L. Kaufman, S. V. Radcliffe and M. Cohen: Decomposi- tion of Austenite by Diffusional Process, ed. by V. E. Zackay and H. I. Aaronson, Interscience Publisher, a divi- sion of John Wiley & Sons, N. Y., (1962), 313. C. Wagner: Trans. AIME, 194 (1952), 91. W. F. Lange and H. I. Aarowon: Metall. Tranc., in press (Private communication). M. Hillert: Lecture on the Theory of Phase Transforma- tion, ed. by H. I. Aaronson, TMS-AIME, Pittsburgh (1975), 1-50. C. A. Dube: PhD Thesis, to Carnegie Institute of Tech- nology, (1948). M. Umemoto, Z. H. Guo and I. Tamura: Muter. Sci. Technol., 3 (1987), 249.

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25) M. Umemoto, H. Ohtsuka, H. Kato and I. Tamura: Proc. Int. Conf. High Strength Low Alloy Steels, ed. by D. P. Dunn and T. Chandra, Univ. Wollongong, Wol- longong, (1984), 107.

26) M. Umemoto, H. Ohtsuka and I. Tamura: Trans. Iron Steel Inst. Jpn., 23 (1983), 775.

27) A. Sandberg and W. Roberts: Proc. Int. Conf. Thermo- mechanical Processing of Microalloyed Austenite, ed. by A. J. DeArdo et al., TMS-AIME, Pittsburgh, (1981), 405.

28) R. K. Amin and F. B. Pickering: Proc. Int. Conf. Thermo- mechanical Processing of Microalloyed Austenite, ed. by A. J. DeArdo et al., TMS-AIME, Pittsburgh, (1981), 377.

29) C. Ouchi, J. Tanaka, I . Kozasu and K. Tsukada: " Op- timization of Processing, Properties and Service Performance through Microstructural Control ", ASTM Spec. Tech. Publ. No. 672, (1979), 105.

30) H. Sekine, T. Maruyama, H. Kageyama and Y. Kawa- shima: Proc. Int. Conf. Thermomechanical Processing of Microalloyed Austenite, ed. by A. J. DeArdo et al., TMS- AIME, Pittsburgh, (1981), 141.

31) T. Maki, K. Tsuzaki and I. Tamura: Trans. Iron Steel Znst. Jpn., 20 (1980), 207.

32) S. Matsuda, T. Inoue, H. Miura and Y. Okamoto: Proc. of Int. Symp. on Toward Improved Ductility and Toughness, ISIJ and JIM, Tokyo, Sendai, (1971), 47.

33) Y. Ohmori, H. Ohtani and T. Kunitake: Met. Sci., 8 (1974), 357.

34) A. R. Marder and G. Krauss: Proc. of Int. Conf. Strength of Metals and Alloys (Znd), 111, ASM, Asilomar, (1970), 822.

3.5) I. Tamura, K. Tsuzaki and T. Maki: Proc. of Int. Conf. Martensitic Transformation (ICOMAT-82), ed. by L. Delaey et al., Les Edition de Physique, Paris, Leuven, (1982), c4-551.

36) I. Tamura, T. Maki and H. T. Yui: to be published.

37) Y. Tomota, K. Tanabe, K. Kuroki and I. Tamura: J . J p n Inst. Met., 41 (1977), 313.

38) C. K. Yao, S. Y. Gao, T. Maki and I. Tamura: Proc. 3rd Int. Cong. Heat Treatment of Materials, ed. by T. Bell, Metals Soc., London, Shanghai, (1983), 5-80.

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Journal of Applied Sciences Research, 2(8): 484-485, 2006© 2006, INSInet Publication

Corresponding Author: 2O.E. Olorunniwo, Department of Material Sciences and Engineering,Obafemi Awolowo University, Ile-ife. Osun state, nigeria

484

Strengthening of Alloy Steel by High Temperature Thermomechanical Treatment1J.A. Omotoyinbo, 2O.E. Olorunniwo, 2O. Ogundare and 2O.O. Oluwole

1Federal University of Technology, Akure. Nigeria.2Obafemi Awolowo University, Ile-Ife. Nigeria.

Abstract: Micro-alloyed steels produced by controlled rolling are of the most attractive propositions in manyengineering applications because of their relative low cost, moderate strength and very good toughness whencompared with as-castalloy steel.. This research work has shownthe possible substitution of cheaper, low-alloysteel for the high alloy steel for the same application. The effect of HTMT onthe mechanical properties of lowalloy steel was investigated.

Key words: Strengthening. High Temperature Themomechanical Treatment (HTMT), micro alloyed,controlled rolling.

INTRODUCTION

Steels are weaker than they should be theoreticallybecause of dislocations and the ease with which thesedislocations moveunder applied shear stress[4]. Therefore,the primary aim of strengthening is to retard thedislocationmovements. However, complete barrier can bedangerous because pile-up of dislocations can lead to acatastrophic crack[2].

Strength in steels arises from several phenomenonincluding solid solution strengthening, dispersionstrengthening and ferrite grain refinement which usuallycontribute collectively to the observed mechanicalproperties such as ductility, strength,etc. Alloyingelements’ presence provides greater control overmicrostructure and consequent benefit in mechanicalproperties.

Therefore, to produce steels with improvedstrength-to-ductility ratio, a combined use of mechanicalworking and heat treatment is employed asstrengthening mechanism [8], resulting in structural steelswith improved weldability, cleannessand inclusion shapecontrol[9].

This paper seeks to investigate and establish thepossible benefits derived from strengthening of alloysteels by HTMT. This includes improved mechanicalproperties, substant ial energy savings (throughelimination of subsequent normalizing treatment) andpossible substitution of cheaper, low alloy steel for highalloyed steel for the same application under the sameconditions.

Experimental procedure: The chemical composition oflow alloy steel as-cast specimens (AISI-SAE 1037),presented in table 1, was obtained from the UniversalSteels Limited, Lagos, Nigeria. All mechanical tests(tensile, hardness, impact) were carried out according tothe A 296 ASTM requirement.

MATERIALS AND METHODS

Tensile test samples: To prepare the tensile testspecimen, the hot-rolled samples which were initially ofround shape were prepared as follows. The 10mmdiameter rod was held tightly in position on the benchvice where it was cut into lengthsof 5mm ach with he aidof hack saw. Each of the samples was then mounted andturned consecutivelyon the lathe machine into cylindricalshape. The vernier caliper was used to take the accuratemeasurement. For reproducibility of results, four sampleswere prepared for the 10mm diameter rod. This processwas repeated for the 12mm and 16mm diameter rod, oneafter the other.

The impact test samples: They were prepared bynotching to 4mm depth at 450 with the aid of the lathemachine. Four samples were prepared each for 10mm,12mm and 16mm diameter rod as well as for the as-castsample to be used in izod impact test.

The hardness test sample-This sample is obtained bygrinding the surface with emery paper. Then, with the aidof hacksaw and lathe machine, cut into 25mm x 25mm x25mm square cube for Rockwell test. Four samples wereprepared for 12mm and 16m diameter samples.

Microstructural test: The specimens were ground insuccession with 220,320,400, 600frit emerypaper. Then,they were polished until mirror-like surfaces wasobtained. The surface was etched with Nital solution andexamined under the microscope. For reproducibility ofresults, four samples were prepared for 10mm,12mm and16mm samples.

RESULTS AND DISCUSSIONS

From the result, it has been shown that the rolled(HTMT) samples possess both the yield strength and the

J. Appl. Sci. Res., 2(8): 484-485, 2006

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97.99%Deformation

98.87%Deformation

92.22%Deformation

AS.CAST60

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Bri

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`% red. Fracture UTS Yield strength Hardness ImpactSample % Elongation in area strength N/mm2 N/mm2 N/mm2 HRC (J)10mm diam as- rolled 1.80 51.41 668.11 688.78 255.10 47.58 48.15-----------------------------------------------------------------------------------------------------------------------------------------------------------------------------------12mm diam. As –rolled 4.63 56.55 367.36 377.55 239.8 41.4 45.57----------------------------------------------------------------------------------------------------------------------------------------------------------------------------------16mm diam as-rolled 7.81 61.26 213.64 221.94 147.96 37.45 41.40--------------------------------------------------------------------------------------------------------------------------------------------------------------------------------As-cast 0.80 3.60 171.94 - - 55.78 11.25

Table 1: The low alloy steel specimens used have the followingcomposition in weight percent as below.

C 0.368 Sn 0.014Si 0.226 Al 0.003S 0.33 Zn 0.003P 0.039 Mo 0.021Mn 0/688 Cu 0.227Ni 0.177 As 0.014Cr 0.119 Co 0.013V 0.002 Pb 0.001W 0.003 Fe 98.055

Fig. 1: Average hardness values.

Fig. 2: Impact values

Ultimate Tensile Strength while the as-cast samples donot. Reasons that could be adduced to these are that,during the HTMT, the as-cast sample has its defectivestructure homogenized, blowholes and gas pores weldedshut and atomic segregation reduced and internal stressesrelieved[3]. These, coupled with hot plastic deformationas well as the presence of alloying elements (carbideformers such as nickel, titanium, etc), inhibitingdislocation movement, to effect strengthening withrespect to Ultimate Tensile Strength (UTS) and yieldstress[2,9].

Also, the as rolled (HTMT) samples are much moreductile than the as cast samples for the same reason citedabove[6]. Conversely, the as-cast samples are harder thanthe as rolled due to the presence of inclusions, residualinternal stress, residual coring and non-uniformity duringsolidification, which are virtually elimination byHTMT[1,8].

Conclusion:The application of HTMT on low alloysteelhas led to the reduction in hardnesswhile the strength andtoughness increase considerably, compared to the as-castsamples. Hence, for applications that require adequateme chanica l proper t ies, high temp eraturethemomechanically treated low alloy steel should beemployed

REFERENCES

1. Adeyeye Oluseyi, 1996. Influence of thermomechanical working on the strengthening ofAluminum alloys, Unpublished B.Sc. Thesis,Metallurgical and Material Engineering department,Federal University of Technology Akure, Nigeria,1998

2. Benjamin, W.N., B.D. Alan, and A.W. Richard,1989. Modern manufacturing process EngineeringMcGrawHill, NewYork, pp: .32-33, 61-62 , 77-114.

3. Bolton, 1988. Production Technology-Processes,Materials and planning Heinemann London pp:174-185.

4. Dieter, G.E. 1988 Mechanical metallurgy McGrawHill, London, pp: 309-312.

5. Degarmo, E.P., J.T. Black and A.K. Ronald, 1993.Materials and processes in manufacturing, 7th Edn.Macmillan, London, pp: 222-223.

6. Higgins, R.A., 1993. Engineering Metallurgy part 1:Applied Physical Metallurgy’ 6th Ed. 124: 219-233,285-330.

7. Llewellyn, D.T., 1992. Steels: metallurgy andapplication’ 1st Edn., Butterworth Heinemann,oxford, pp: 64-77.

8. Rajan, T.V, C.P. Sharma and A. Sharma, 1988. Heattreatment-principle and techniques, McGraw Hill,India, pp: 175-181.

9. Novikov, L., 1978. Theory of Heat treatment ofmetals,’ 1st Edn. Moscow, pp: 105-106, 418-425.

10. Honeycombe, RWK, 1981. Microstructure andproperties’ 1st Ed., Edward Arnold, London, pp:167-185.

Ion-nitriding of Maraging steel (250 Grade) for

Aeronautical application

Kishora Shetty1,2

, Subodh Kumar3 and P. Raghothama Rao

2

2 Regional Centre for Military Airworthiness (Foundry and Forge), CEMILAC, DRDO,

Bangalore-560 037, India.

3 Department of Materials Engineering, Indian Institute of Science, Bangalore-560 012, India.

E-mail: [email protected]

Abstract. Ion nitriding is one of the surface modification processes to obtain better wear

resistance of the component. Maraging steel (250 Grade) is used to manufacture a critical

component in the control surface of a combat aircraft. This part requires high strength and

good wear resistance. Maraging steels belong to a new class of high strength steels with the

combination of strength and toughness that are among the highest attainable in general

engineering alloys. Good wear resistance is achieved by ion-nitriding (also called as plasma

nitriding or glow discharge nitriding) process of case nitriding. Ion-nitriding is a method of

surface hardening using glow discharge technology to introduce nascent (elemental) nitrogen

to the surface of a metal part for subsequent diffusion into the material. In the present

investigation, ion-nitriding of Maraging steel (250 grade) is carried out at 450 0C and its effect

on microstructure and various properties is discussed.

1. Introduction Surface Engineering means ‘engineering the surface’ of a material or components to impart

surface properties, which are different from the bulk of base material [1]. The purpose may be to

reduce wear, minimize corrosion, increase fatigue resistance, reduce frictional energy losses,

provide a diffusion barrier, provide thermal or electrical insulation, exclude certain wave lengths

of radiation, promote radiation, electronic interactions, or simply improve the aesthetic

appearance of the surface. Surface engineering processes, which give required properties at

surfaces include flame hardening, induction hardening, laser hardening, carburizing, nitriding,

cyaniding, plasma nitriding, ion implantation, weld overlay, roll cladding, thermal spraying,

plasma spraying, ion plating, CVD, PVD etc. Nitriding is a process for case hardening of alloy

steel in an atmosphere consisting of a mixture of ammonia gas and dissociated ammonia [2]. In

ion-nitriding (also called as Plasma nitriding or Glow discharge nitriding) method glow discharge

technology is used to introduce nascent (elemental) nitrogen to the surface of a metal part for

subsequent diffusion into the material [3-5].

1 To whom any correspondence should be addressed

IVC-17/ICSS-13 and ICN+T2007 IOP PublishingJournal of Physics: Conference Series 100 (2008) 062013 doi:10.1088/1742-6596/100/6/062013

c© 2008 IOP Publishing Ltd 1

The plasma assisted surface modification techniques offer a great flexibility and are capable

of tailoring desirable chemical and structural surface properties independent of the bulk properties

[3]. It has other advantages like nil or very thin white layer after nitriding and there is no

machining or grinding involved for complex parts after the process. The hardened surface layers

become an integral part of the base material without any significant reduction in properties of

base material. It is also known to provide the modified surface without dimensional change or

distortion of the component. Ion-nitriding provides better control of case chemistry and

uniformity [3,6,7]. This method is one of the most effective techniques for increasing wear

resistance, fatigue strength, surface hardness and corrosion resistance of industrial components

[8]. In this process, vacuum environment and high voltage electrical energy is used to form

plasma through which nitrogen ions are accelerated to impinge on work piece. The ion

bombardment heats the work piece and cleans the surface as the active nitrogen diffuses through

it [3,9].

Maraging steels belong to a new class of high strength steels with the combination of strength

and toughness that are among the highest attainable in general engineering alloys [10]. These

steels differ from conventional steels in that they are hardened by a metallurgical reaction that

does not involve carbon. These steels contain very low carbon (<0.03%) and are strengthened by

the precipitation of intermetallic compounds at temperature about 480 0C [11-13]. The term

maraging is derived from martensite age hardening and denotes the age hardening of a low

carbon, iron – nickel lath martensite matrix [12,14]. Different maraging steels are designed to

provide specific levels of yield strength from 1030 to 2420 MPa (150 to 350 ksi). These steels

typically have very high nickel, cobalt and molybdenum and very low carbon content. Carbon is

treated as an impurity in these steels and is kept as low as commercially feasible (<0.03%) in

order to minimize the formation of Titanium carbide (TiC), which can adversely affect strength,

ductility and toughness [11,12]. Nominal composition of 250 Grade is Fe-18Ni-8.5Co-5Mo-

0.4Ti-0.1Al [10].

Good wear resistance of Maraging steel can be achieved by ion-nitriding process of case

nitriding. In conventional gas nitriding process, the nitriding temperature is 500 0C – 550 0C

[2,15], which is above the ageing temperature of maraging steel. Hence ion-nitriding of Maraging

steel (250 grade) at a temperature lower than the aging temperature has been carried out in the

present investigation. Microstructure of the ion-nitrided specimens was examined and properties

like tensile, low cycle fatigue, hardness, case depth and corrosion by salt spray test were

evaluated. These tests were also carried out on un-nitrided specimens for comparison.

2. Experimental

Figure 1 shows a simplified schematic and figure 2 shows the actual ion-nitriding installation.

The work load is supported on a hearth plate inside a double walled, water cooled vacuum

chamber, connected to vacuum pumps and gas supply. The chamber is evacuated to a pressure of

about 2.5 X 10-2

mbar, a pressure low enough for the background level of oxygen to be within

acceptable limits (less than 50 ppm), then filled with a low pressure mixture of nitrogen and

hydrogen. The use of auxiliary A.C. heaters to heat the cathode to 250 °C is desirable to minimize

cycle time. It can also help provide better temperature uniformity of the part in ion-nitriding

treatment. The discharge is ignited using a D.C. power supply, and pressure and temperature are

raised to the desired operating values by controlling gas flow and pressure, applied voltage and

current. The discharge can be monitored by meters and viewed through inspection windows. The

work is cathode and the vessel is anode. The furnace is electrically grounded, cool to the touch,

and quiet in operation. Maraging steel (250 grade) specimens were prepared from solutionised

material and then aged at 450 0C for 3 hrs. These aged specimens were ion-nitrided in the

nitriding furnace under the vacuum. Plasma is obtained by passing the gas mixture of H2 and N2

gases in the ratio 3:1 into the chamber and maintaining the pressure of 5 mbar. Ion-nitriding is

carried out at 450 0C for 10 h.


2

Figure. 1. Schematic arrangement of an ion-

nitriding system [7].

Figure. 2. Ion-nitriding installation.

3. Results 3.1 Chemical Composition

The chemical analysis of the Maraging steel (250 grade) was carried out using Optical emission

spectroscopy to confirm the material specification. The values obtained are given in Table 1.

Table 1. Chemical composition of the alloy (wt.%).

Element C Ni Mo Co Ti Al S P Fe

Composition 0.01 17.29 4.89 7.90 0.41 0.14 0.005 0.006 Balance

3.2 Visual Check

Ion-nitrided samples were checked for any visible defects and discoloration and no defects were

observed.

3.3 Microstuctural studies

The samples for Metallography were prepared from both ion-nitrided and un-nitrided samples.

Microstructural studies were carried out using both optical microscope and scanning electron

microscope (SEM). Optical micrographs of un-nitrided specimen is shown in figure 3 and that of

ion-nitrided specimen is shown in figure 4. Etchant used here is Ferric Chloride. SEM studies

were carried out by etching the specimens using Ferric Chloride. The microphotographs obtained

from un-nitrided and ion-nitrided samples are shown in figures 5 and 6 respectively.

Figure 3. Optical micrograph of un-

nitrided sample showing martensite

structure.

Figure 4. Optical micrograph of ion-nitrided

sample showing martensite structure at the

core and nitrided layer at the case.


3

Figure 5. SEM micrograph of un-nitrided

sample.

Figure 6. SEM micrograph of ion-nitrided

sample.

3.4 Hardness and Case depth measurement

Hardness was measured for both un-nitrided and ion-nitrided samples using Vickers Hardness

Tester. The hardness values obtained on un nitrided specimen is 616 VPN where as for ion-

nitrided specimen it is 900 VPN.

Case depth was determined by hardness measurements. Hardness values are listed in Table

2. Figure 7 shows the variation of hardness with depth from the surface. Case depth here is taken

as distance from the surface to which hardness is 100 VPN more than the core hardness. From the

graph the case depth estimated is 93 µm.

Table 2. Hardness and corresponding distance from the surface.

Distance from the edge (µm) Hardness (VPN)

50 797

75 761

100 702

125 626

150 623

175 623

200 623

Core 616

Figure 7. Graph of case depth vs hardness

3.5 Salt spray test

Salt spray test was carried out on both un-nitrided and ion-nitrided specimens for 144 hrs at 5 %

NaCl solution in a Salt Spray Test chamber. Figure 8 shows the photograph of samples

500 550 600 650 700 750 800 850

0 25 50 75 100 125 150 175 200 225 Depth from the surface (micrometer)

Hardness (VPN)


4

before and after the salt spray test. Corrosion started in un-nitrided specimen after 48 h of

exposure, whereas no corrosion was noticed in ion-nitrided specimen.

(a)

(b)

1 - Ion nitrided 2 - Un nitrided

Figure 8. (a). Photograph of the samples (a) before the salt spray test and (b) after the salt

spray test.

3.6 Tensile Properties

Tensile properties – Ultimate tensile strength (UTS), 0.2 % Proof stress (0.2 % PS), %

Elongation (% El.) and % Reduction in Area (% RA) were measured using TIRA Test 2820S

Universal Testing Machine (UTM). The results obtained are presented in Table 3.

Table 3. Tensile properties.

Property UTS (MPa) 0.2% PS (MPa) % El. % RA

Sample 1 1694 1606 9.32 60

Sample 2 1696 1619 8.84 56 Un nitrided

Sample 3 1702 1622 9.23 55

Sample 1 1847 1757 5.66 25

Sample 2 1864 1810 6.24 28 Ion nitrided

Sample 3 1877 1813 6.70 30

3.7 Impact Strength

Impact strength was measured on Charpy ‘U’ notch specimens using FIE make Charpy Impact

Testing machine. The values obtained are given in Table 4.

Table 4. Impact strength (Joules).

Sample 1 Sample 2 Sample 3

Un nitrided 21 22 22

Ion nitrided 17 19 19

3.8 Low Cycle Fatigue

Low cycle fatigue testing was carried out on smooth specimens in Zwick Roell UTM with

applied stress of 1172 MPa and with stress ratio (R) of -1. The values obtained are presented in

Table 5.

Table 5. Low cycle fatigue properties (No. of cycles to failure)

Sample 1 Sample 2 Sample 3

Un nitrided 3755 3287 5220

Ion nitrided 5500 14400 11200

1

2

1

2

1

2

1

2


5

4. Discussion Chemical composition analysed conforms to the Maraging steel 250 grade.

Both Optical microscopy and SEM examination confirm that no change in core

microstructure of the material after ion-nitriding occurs.

The ion-nitrided specimens exhibit higher surface hardness and good case depth than the un-

nitrided specimens. The introduction of nitrogen to the surface of Maraging steel by ion nitriding

process improves the wear properties by increasing the hardness.

Salt spray test results show that there is an improvement in corrosion resistance after

nitriding. This is due to the formation of an anodic layer after nitriding.

The ion-nitrided samples show higher UTS and 0.2 % PS values than the un-nitrided ones,

whereas % El and % RA are less for ion-nitrided samples compared to the un nitrided samples.

This can be attributed to the surface layer getting hardened because of ion-nitriding. Impact

strength of ion-nitrided samples is lower compared to un-nitrided samples. This is again due to

the surface layer getting hardened because of ion nitriding.

The number of cycles to failure in low cycle fatigue test also increase on ion-nitriding. The

introduction of nitrogen to surface layers increases the fatigue properties of the maraging steel by

introducing the residual compressive stresses.

5. Conclusions

From this ion nitriding study of Maraging steel (250 grade), the following conclusions can be

drawn.

Case depth obtained is sufficient for the design requirement.

Ion-nitriding improves the surface hardness of Maraging steel 250 grade.

An improvement in UTS, 0.2 % PS and LCF properties is observed on ion-nitriding. The %

El, % RA and impact strength decrease but they satisfy the design requirements.

No change in microstructure observed after ion-nitriding.

Ion-nitriding of 250 Grade appears to be suitable for parts which are subjected to constant

wear.

Acknowledgements The authors thank Mr. B. S Vedaprakash, Group Director (Propulsion), and Mr. K. Tamilmani,

Chief Executive (Airworthiness), CEMILAC for their support.

References [1] Kumar B, Upadhyay N C 2005 IIM Metal News 8 5

[2] Sidney Avner H 1997 Introduction to Physical Metallurgy 2nd

Edition 383

[3] Spalvins T Ion Nitriding Conference proceedings ASM International 1

[4] Moller W, Parascandola S, Telbizova T, Gunzel R and Richter E 2001 Surface & Coatings

Technology 136 73

[5] ASM Handbook 1996 Vol. 5 411

[6] Bernd Edenhofer 1976 Metal Progress 181

[7] EPRI Centre for Materials Fabrication Ohio 1994 Ion nitriding 2

[8] Ahangarani Sh., Mahboubi F and Sabour A. R 2006 Vacuum 80 – Surface Engineering, Surface

Instrumentation & Vacuum Technology 1032

[9] Pavel Novak, Dalibor Vojtech and Jan Serak 2006 Surface & Coatings Technology 200 5229

[10] Gupta B, 1996 Aerospace Materials II 695

[11] ASM Handbook 1991Vol. 4 219

[12] ASM Handbook 1990 Vol. 1 793

[13] INCO Databook, 1976 351

[14] Morito S, X-Huang, Furuhara T, Maki T and Hansen N 2006 Acta Materialia 54 5323

[15] David Pye 2005 Practical Nitriding and Ferritic Nitrocarbourizing ASM International 71


6

Bull. Mater. Sci., Vol. 29, No. 1, February 2006, pp. 59–66. © Indian Academy of Sciences.

59

Microstructures and properties of low-alloy fire resistant steel

BIMAL KUMAR PANIGRAHI R&D Centre for Iron and Steel, Steel Authority of India Limited, Ranchi 834 002, India

MS received 10 June 2004; revised 2 November 2005

Abstract. Microstructures and properties of weldable quality low-alloy fire resistant structural steels (YS: 287–415 MPa) and TMT rebar (YS: 624 MPa) have been investigated. The study showed that it is possible to obtain two-thirds of room temperature yield stress at 600°°C with 0⋅⋅20–0⋅⋅25% Mo and 0⋅⋅30–0⋅⋅55% Cr in low carbon hot rolled structural steel. Microalloying the Cr–Mo steel by niobium or vanadium singly or in combination resulted in higher guaranteed elevated temperature yield stress (250–280 MPa). The final rolling temperature should be maintained above austenite recrystallization stop temperature (~ 900°°C) to minimize dislocation hardening. In a quenched and self-tempered 600 MPa class TMT reinforcement bar steel (YS: 624 MPa), low chromium (0⋅⋅55%) addition produced the requisite yield stress at 600°°C. The low-alloy fire resistant steel will have superior thermal conductivity up to 600°°C (> 30 W/m⋅⋅k) compared to more concentrated alloys. Keywords. Fire resistant steel; thermomechanical processing; microstructure; strength; thermal conductivity.

1. Introduction

Steel structures using mild steel in fire sensitive areas are protected from fire by providing fire resistant coating or insulation. This adds to the constructional cost. The problem with unprotected carbon–manganese mild steel is its poor strength at temperatures above ~ 350°C, which can make a structure unsafe after a major fire. In case of short duration fire, it is rather difficult to assess the damage to structures caused by fire and may call for demolition/renovation of the structures. Most of the research efforts relating to construction in fire sensitive areas were directed towards development of steels that can retain adequate strength after prolonged exposure in fire. The building codes of some specifications require the steel to have a minimum of one-half (ASTM 1996) to two-thirds (Fushioni et al 1995; BIS 2002) of room temperature yield strength at 500–600°C compared to the mild steel which retains about one-third of room temperature yield strength at 600°C. In order to achieve this strength level, the steel chemistry and manufacturing process are closely controlled. Previous studies on microstructure and mechanical properties of fire resistant steels carried out by Chijiwa et al (1993) and Assefpour-Dezfully et al (1990) focussed on processing, structure and properties of fire resistant steels with different combinations of alloying elements. In the present investigation, the effect of lower alloying addition, particularly molybdenum, on the elevated temperature properties will be discussed which is important due to increasing alloy cost. Over and above, there is hardly

any information on alloying required to obtain two-thirds of room temperature yield stress at 600°C for thermome-chanically treated (TMT) rebar particularly, the high strength variety with yield strength above 600 MPa. The present work was carried out to ascertain minimum requirement of molybdenum/chromium in steels with and without micro-alloying elements in order to achieve the guaranteed strength at 600°C in plate, structurals and TMT rebar and their effects on structure and properties.

2. Experimental

Six experimental laboratory heats (steels A–F) and one industrial heat (steel G) were melted in 0⋅1 T air induction (IF) and 6 T electric arc furnace (EAF), respectively for proce-ssing to plate and beam. The ingots from IF heats were soaked at 1250°C for 2 h and thermomechanically processed (TMP) in an experimental rolling mill to 12–14 mm thick plates in nine passes. The ingot from EAF heat was soaked at 1320°C for 4 h and rolled to blooms of size 230 × 160 mm. These blooms were soaked at 1250°C and sub-sequently thermomechanically processed to 200 × 100 mm beam section. The finishing rolling temperatures (FRT) for plates and beam were measured with an infrared pyrome-ter up to an accuracy of ± 5°C and were between 800 and 925°C. The finishing pass reduction was 20–30% for plates and about 10% for beam. The plates and beam were allowed to cool in natural air after rolling. Another heat (steel H) was made in 250 T twin hearth furnace and cast as 9 T ingots for processing to TMT rebar. The ingots were rolled to billets of size 100 × 100 mm after soaking at 1300°C for 4 h. These billets were subsequently reheated at 1250°C

([email protected])

Bimal Kumar Panigrahi

60

and processed to TMT rebar of 32 mm diameter through a Thermex cooling system. Round tensile specimens (dia. 6⋅25 mm) were machined from the quarter width position of the plates (width of plate: 135 mm) in the longitudinal direction as per ASTM A 370 standard and tested at room temperature (30°C) and at elevated temperatures up to 600°C at cross head speed of 2 mm/min in a servohydraulic test machine as per ASTM E 8 and E 21 standards, respectively. Tensile testing of plates was also carried out selectively at room tempera-ture after exposure at 600°C for 3 h under applied stress of 140 MPa. The specimens after exposure were allowed to cool to room temperature inside the furnace. In the elevated temperature test of plates the specimens were heated from room temperature to the test temperature and soaked at the desired temperature for 15 min before tensile stress was applied. The accuracy of the test temperature was con-trolled within ± 1°C. Charpy V-notch impact test specimens of size 10 × 10 × 55 mm from steels A–F were machined across the plate width in the longitudinal direction and tested at RT to – 30°C as per ASTM E 23 standard. The optical microstructure was examined in the longitudinal through thickness direction after polishing and etching by 2% nital. The average grain size was measured by linear intercept method. Transmission electron microscopic (TEM) investigation was carried out using thin foils prepared from mechanically thinned strip of thickness below 0⋅1 mm. The electropolishing was done in a twin jet polisher using a solution of 5% perchloric acid and 95% glacial acetic acid at 20°C and 60V operating voltage. The foils were observed at 200 keV. Flat tensile specimens (gauge length: 5⋅65 √area) were prepared in the longitudinal direction from the web of the beam as per IS 2062 standard (BIS 1998). The ambient and elevated temperature tensile testing was conducted as described for plates. Tensile testing was also done after an aging treatment at 600°C for 3 h under applied tensile stress of 140 MPa. Charpy V-notch specimens (size 7⋅5 × 10 × 55 mm) were prepared from the flange region as described in IS 2062 standard and tested as described for plate in the longitudinal direction. The optical micro-structure was examined in the longitudinal through thickness direction after polishing and etching by 2% nital. The average grain size was measured by linear intercept method. TEM specimens were prepared from the flange region as described for plate and observed at 200 keV. Full scale fire resistance test was carried out on beam section (size 200 × 100 × 3900 mm) as per ASTM E 119 standard (ASTM 1996) at Fire Research Laboratory, CSIR, Roorkee, India. The bare beam was mounted on the specimen frame holder of a gas fired furnace. The furnace was rectangular in shape. The casing was made up of steel reinforced sec-tions. The refractory lining of the furnace consisted of front layer of kyanite insulation bricks backed by hot face insulation. The burner casing was made up of fire bricks. The complete furnace assembly consisted of four walls

and the open roof top for mounting the beam specimen. The flue gas for heating passed through a duct surrounding the sides of the walls and then was disposed off through a chimney. There were ten long flame burners. The capacity of each burner was 35–40 l/h. The furnace was run on positive pressure. The heating of the specimen was done as specified in ASTM E119 standard. The thermocouples were fixed at four different sections of the beam to measure the rise of temperature at different intervals. While plotting the tem-perature vs time curves, average temperature measured by concerned thermocouples in a section was used. Tensile testing of rebar was done using unmachined rebar (gauge length: 5 dia.) at ambient temperature in a 60 T servo hydraulic tensile testing machine at a cross head speed of 10 mm/min. For the elevated temperature test at 600°C, round specimens were used from the core region of rebar and tested as described for plate material. Charpy V-notch impact specimens were also prepared from the location described previously (Panigrahi and Jain 2002) and tested similar to plate. The optical micro-structure was examined in the transverse section after polishing and etching by 3% nital. TEM specimens were prepared from the rim and core regions of the rebar as described for plate and observed at 200 keV.

3. Results and discussion

3.1 Chemical composition

In a structural steel, the composition will be determined by the requirement of strength, toughness and weldabi-lity. The chemical composition of steels is given in table 1. The carbon was 0⋅11–0⋅13% in plate, 0⋅17% in structural beam and 0⋅25% in TMT rebar. The conventional way to improve moderately the strength at elevated temperature (up to ~ 350°C) is by increasing the carbon content (Houdre-mont 1953). Lower carbon in hot-rolled steel improves notch toughness and weldability. In TMT rebar carbon increases the hardenability. Manganese increases the strength by solid solution hardening (Baird and Jamieson 1972) in hot rolled steel. Manganese tends to lower the eutectoid carbon content in steel and promotes segregation particu-larly in hot rolled structural steel with carbon > 0⋅10%. This can adversely affect the mechanical properties. Manga-nese was restricted to 1% in fire resistant hot rolled plate and structural with carbon 0⋅10–0⋅20% and 1⋅5% in TMT rebar where manganese is a potential hardenability en-hancing element. Silicon content in steels A–G varied from 0⋅29–0⋅35%. Silicon is a solid solution strengthener and a deoxidizer. Silicon retards softening at high tem-peratures (Irvine 1962) and also increases the hardenability. However, in steel H, its content was low due to tech-noeconomic reasons. The aluminium (0⋅01–0⋅02%) is required to tie-up free nitrogen to improve toughness. In the steels A, D, E and F, the content of aluminium was

Fire resistant steel

61

Table 1. Chemical composition of steels (wt.%).

Steel A B C D E F G H

C 0⋅11 0⋅12 0⋅12 0⋅13 0⋅13 0⋅12 0⋅17 0⋅25 Mn 1⋅02 1⋅02 1⋅08 0⋅99 1⋅00 0⋅96 0⋅74 1⋅42 Si 0⋅35 0⋅35 0⋅29 0⋅31 0⋅31 0⋅30 0⋅31 0⋅056 S 0⋅022 0⋅032 0⋅029 0⋅025 0⋅026 0⋅025 0⋅034 0⋅024 P 0⋅019 0⋅025 0⋅027 0⋅021 0⋅022 0⋅021 0⋅027 0⋅024 Cr 0⋅31 0⋅29 0⋅40 0⋅38 0⋅39 0⋅38 0⋅54 0⋅55 Mo 0⋅26 0⋅20 0⋅14 0⋅26 0⋅26 0⋅26 0⋅22 – Nb – – – 0⋅02 – 0⋅02 – – V – – – – 0⋅07 0⋅12 – – Al 0⋅0033 0⋅099 0⋅097 0⋅0034 0⋅0039 0⋅0039 0⋅052 0⋅004 N 0⋅0042 0⋅0063 0⋅0067 0⋅0047 0⋅0047 0⋅0070 0⋅0124 0⋅0075

lower than the prescribed range. The rebar steel was semi-killed. Chromium in small percent forms (FeCr)3C (Houdremont 1953) in the cementite of pearlite. It is a ferrite stabilizer and increases the hardenability. However, its potential as a solid solution strengthening element is small due to its high affinity to carbon. The amount of chro-mium was 0⋅29–0⋅40% in plate and 0⋅55% in structural beam and TMT rebar. Molybdenum increases elevated temperature yield strength (Argent et al 1970; Honey-combe 1981) even when present in small percent (~ 0⋅25%) in solid solution in low carbon steel. It also forms carbide (Fe, Mo)3C in the cementite of pearlite resisting softening on prolonged exposure in a fire (Houdremont 1953). Molyb-denum tends to hinder self-diffusion of iron (Houdremont 1953), thus increasing the recrystallization temperature of steel. Microalloying elements, vanadium and niobium, increase the elevated temperature strength (Pickering 1978; Sage 1983). They also increase the recrystallization tempera-ture of steel (Borato et al 1988). Their carbides also offer resistance to softening when the steel is exposed to fire. The chemistry of the present investigation differs from the previous studies (Assefpour-Dezfully et al 1990; Chi-jiwa et al 1993; Fushioni et al 1995) in the following ways: (a) the content of molybdenum and chromium is low, (b) low niobium and vanadium are used in micro-alloyed steels and (c) chromium without molybdenum has been used in TMT rebar.

3.2 Thermomechanical processing (TMP)

Processing plays an important role since it determines the final microstructure and mechanical properties of hot rolled steel (Panigrahi 2004). The TMP schedule for plate and beam is shown in figure 1. TMP of plate and beam involves controlled deformation of austenite at higher temperatures above the recrystallization stop temperature (tnr) (Borato et al 1988) during which static and dynamic restoration processes occur (Tanaka 1981). Table 2 shows data on total reduction, approximate recrystallization stop temperature, finishing temperature and grain size of steels.

As the finish rolling temperatures (FRT) of all steels ex-cept steel F and a few structurals of steel G were above tnr, an equiaxed ferrite microstructure could be obtained with low dislocation density. Steel F was finish rolled at 800°C, quite below tnr (884⋅1°C) but above Ar3 temperature (771°C) (Ouchi et al 1982). Deformation in this region strain hardened the austenite, increasing its dislocation density. Deformation bands are also formed (Tanaka 1981). Since ferrite nucleates both on austenite grain boundaries and deformation bands, the grains are finest (7⋅7 µm). How-ever, the cellular structure retained by rolling below tnr, lowered the yield strength at elevated temperatures. In the processing of TMT rebar (figure 1), initially at higher temperatures (1200–1100°C) the material recrys-tallizes statically. At the intermediate and finishing stages of processing (1000–1050°C), static, dynamic as well as metadynamic recrystallization of austenite can occur (Poliak and Jonas 2003) depending upon strain at different passes producing an equiaxed austenite structure. On leaving the finishing stand at 1000–1050°C, the bar enters a Thermex cooling chamber where the surface of the bar is cooled by pressurized water at ~ 200°C/s forming a thin rim of lath martensite while the core is still austenite. On emergence from the cooling chamber, the martensite rim gets self-tempered by the heat of the core and the core finally transforms to bainite.

3.3 Microstructure

Typical optical microstructures of some steels in as-rolled condition are shown in figure 2. The Cr–Mo (steel A) and Cr–Mo–Nb (steel D) alloyed plates (figures 2a and b) and beam (figure 2c) showed a predominantly polygonal fer-rite–pearlite microstructure. At higher magnification, steel A (figure 2d) also showed upper bainite (arrow mark). The bainite in steel A resembles pearlitic bainite (Bhadeshia 1992) as the colonies tend to have crystallographic facets rather than nicely rounded colonies of conventional pear-lite. The vanadium (steel E) and niobium plus vanadium (steel F) microalloyed steels also showed similar micro-


62

Figure 1. Schematic of TMP schedule for plate, beam and TMT rebar (M, martensite; γ, austenite).

Figure 2. As hot rolled microstructures of a. Cr–Mo plate (steel A), b. Cr–Mo–Nb plate (steel D), c. Cr–Mo beam (steel G) and d. pearlitic bainite of Cr–Mo plate (steel A).

Table 2. Process parameters and microstructural data.

Thickness Total FRT Grain size Steels (mm) reduction (%) tnr (°C) (°C) Structure (µm) Product

A 14 86 814⋅0 900 F + P + B 10⋅4 Plate B 12 88 853⋅5 900 F + P 13⋅5 Plate C 12 88 874⋅3 900 F + P 12⋅5 Plate D 12 88 876⋅9 900 F + P + B 11⋅0 Plate E 12 88 828⋅5 875 F + P + B 12⋅5 Plate F 12 88 884⋅1 800 F + P + B 7⋅7 Plate G – 94* 874 850–925 F + P + B 10⋅3 Beam H – 92 984⋅4 1000 M + B – Rebar

F, Ferrite; P, pearlite; B, bainite; M, tempered martensite; *flange region; tnr = 887 + 464%C + (6645%Nb−664√%Nb) + (732%V−230√%V) + 890% Ti + 363%Al−357%Si.


63

Figure 3. Bright field TEM of a. Cr–Mo–Nb–V plate (steel F), b. and c. Cr–Mo beam (steel G) and d. TMT rebar.

structures. The bainite ‘colonies’ in steel G were coarser than steels A, D, E and F and were more numerous possi-bly due to presence of somewhat higher carbon and chromium in this steel, lower reduction per pass and coarser austenite grains prior to transformation. The bainite was not observed in steels B and C. The finishing rolling in the austenite region above recrystallization stop tempera-ture helped to form polygonal ferrite grain structure in Cr–Mo, Cr–Mo–Nb and Cr–Mo–V plate steels. In this case ferrite nucleates on austenite grain boundaries (Tanaka 1981). Cr–Mo–Nb–V plate (steel F) was finish rolled in the unrecrystallized austenite region producing grain size finer than other steels (table 2) due to nucleation of ferrite on austenite grain boundaries, transgranular twins and deformation bands (DeArdo 1995; Panigrahi 2001) after transformation. The partial cellular structure of steel F with dislocations retained is shown in figure 3a. The beam steel also showed upper bainite (figure 3b) with cementite particles distributed between ferrite platelets (figure 3c). An exposure treatment up to 3 h at 600°C did not show changes in the appearance of microstructure. Unlike martensite in which dissolved carbon in solid solution is high, bainite has little carbon in solid solution and is much less sensitive to tempering at low temperatures (~ 600°C). However, cementite particles can coarsen and a general recovery of dislocation substructure occurs (Bhadeshia 1992). The TMT rebar steel showed a tempered martensite rim. The core structure was bainitic (figure 3d) due to high hardenability of this steel.

3.4 Elevated temperature yield stress

The dependence of yield stress and UTS with temperature for steels A–G is shown in figure 4. Steel F showed a faster drop of yield stress after 500°C. However, except steel C (0⋅14%Mo) whose yield stress was marginally below the desired level at 600°C, all other steels retained two-thirds of their room temperature yield strength at 600°C (table 3). The elevated temperature yield stress de-pends mainly on the stability of microstructure at higher temperatures. Chromium and molybdenum have a stabi-lizing effect on cementite of pearlite (Houdremont 1953). A higher amount of chromium in steel C prevented exce-ssive drop of yield stress at 600°C despite low molybdenum (0⋅14%) in it. Steels D and E that are microalloyed with Nb and V, respectively have highest level of yield stress at 600°C viz. 82% and 85%, respectively. Nb and V have added advantages due to their precipitation hardening potential and the loss of strength at higher temperature is minimal. The rapid drop of yield stress of steel F above 500°C was due to lowering of dislocation density on soaking above 500°C. The UTS of all steels fell sharply above 400°C (figure 4). The elongation of all steels at 600°C was higher than elongation at the ambient tem-perature. Exposure at 600°C for 3 h increased the yield stress by about 50 MPa without significant change in elongation in steels A, D, E, F and G. The core of rebar steel H with bainitic structure also showed two-thirds of room temperature yield stress at 600°C (table 3). Due to presence of


64

Figure 4. Dependence of (a) yield stress and (b) UTS on temperatures of fire resistant steels and mild steel (MS).

Table 3. Tensile properties of the investigated steels.

Steel A B C D E F G H

Ambient temperature (RT) YS (MPa) 287 293 298 338 293 415 320 624 UTS (MPa) 505 486 539 582 530 622 541 819 El. (%) 38⋅0 36⋅2 32⋅0 30⋅0 37⋅8 26⋅8 27⋅3 12⋅5

Elevated temperature (600°C) YS (MPa) 219 191 188 279 250 280 214 419 UTS (MPa) 326 259 247 332 327 338 292 426 El. (%) 42⋅2 41⋅7 49⋅0 40⋅4 36⋅2 46⋅0 33⋅0 30⋅5 % of room temperature YS 76⋅3 65⋅1 63⋅0 82⋅5 85⋅3 67⋅4 66⋅8 67⋅1

After thermal exposure at 600°C under applied stress YS (MPa) 360 – – 408 348 469 369 – UTS (MPa) 432 – – 523 475 584 492 – El. (%) 22⋅5 – – 30⋅8 34⋅0 22⋅8 27⋅0 –

tempered martensite rim, the overall strength of TMT rebar at 600°C should be higher than two-thirds of its room temperature yield strength. In a composite micro-structure comprising tempered martensite rim and bainitic core, the overall yield stress is given by area fraction of rim multiplied by yield stress of rim plus area fraction of core multiplied by yield stress of core. The area fraction of bainite core was 0⋅766 and that of tempered martensite rim was 0⋅234. Previous investigation (Panigrahi 2002) showed that low-alloyed tempered martensite structure had more than two-thirds of room temperature yield stress at 600°C.

3.5 Impact toughness

The ambient temperature impact toughness, 50% energy ductile brittle transition temperature (DBTT) (Dahl 1992) and fracture characteristics are given in table 4. The upper shelf energy of as-rolled plate and beam varied from 65–118 J (figure 5). The steels A, D, E and F showed DBTT at – 4°C to –12°C. The DBTT of steel B was below –30°C and that of steel C was –28°C. Presence of upper bainite in all steels except steels B and C, low aluminium content (steels A, D, E and F) and cellular structure (steel F) were responsible for increase in DBTT. Notwithstanding

Temperature (°C) Temperature (°C)


65

higher carbon content, steel H showed a DBTT of –14°C due to presence of tempered martensite and some acicular ferrite in it. Steel G showed a DBTT of –5°C in as rolled condition and –34°C in aged condition. The DBTT of steel G in as rolled condition was higher than steel A possibly due to presence of relatively coarser ‘colonies’ of pearlitic bainite in its microstructure and a higher carbon content (0⋅17%) (Pickering 1978) compared to steel A (0⋅11%

Figure 5. Charpy energy vs temperature curves of plates and beam.

Figure 6. Typical time–temperature curves for Cr–Mo beam (steel G) and mild steel (T1, T2 and T3 are locations of thermo-couples at section 1).

carbon). On exposure at 600°C, the improved toughness (DBTT: –34°C) of steel G could be due to lowering of dislocation density of bainite and diffusion of embrittling atoms present on the grain boundaries to the dislocations. Improvement of impact transition temperature by diffusion of phosphorus from prior austenite grain boundaries to lattice defects was reported earlier for spring steel (Wettlaufer and Kasper 2000). However, upper shelf energy of aged steel was lower than the as-rolled beam possibly due to microstructural heterogeneity leading to formation of local-ized carbide rich areas that aided lowering of the absorbed energy.

3.6 Fire resistance

The average temperature rise of fire resistant steel (FRS) beam and mild steel (MS) beam (0⋅17C, 0⋅69Mn, 0⋅04Si, 0⋅048S, 0⋅028P, 0⋅007Al) is shown in figure 6 for section Table 4. Charpy toughness data.

Toughness at RT DBTT Fracture Steel (Joule) (°C) surface at RT

A 118 − 12 Dimple B 116 <− 30 Dimple C 100 − 28 Dimple D 94 − 9 Dimple E 82 − 8 Dimple F 73 − 4 Dimple G (as rolled) 87 − 5 Dimple G (after thermal exposure 50 − 34 Dimple at 600°C) H 88 − 14 Quasi cleavage

Note: All values of toughness are for full size specimens.

Figure 7. Dependence of thermal conductivity of dilute steel alloys on temperatures.

Temperature (°C)

Tem

pera

ture

(°C

)

Temperature (°C)


66

1 (inset). The rise of temperatures up to 600°C for other sections was similar to section 1. It was observed that the temperature of the fire resistant steel was at a lower level compared to mild steel. The data points for these curves remain below the ASTM E 119 standard time–temperature curve in a fire test (ASTM 1996). The time to reach the critical temperature (538°C) was 32 min and 21 min for fire resistant steel and mild steel, respectively. The rate of heating is dependent on the thermal conductivity which is a material property governing the flow of heat through a material at steady state. Figure 7 shows the dependence of thermal conductivity with temperature for pure iron and some constructional steels (Riemann 1953; McGannon 1966). The thermal conductivity of pure iron and dilute alloys decreases with rising temperature. The likely be-haviour of 0⋅5Cr–0⋅22Mo fire resistant steel which is a dilute alloy is also shown in this figure. Its conductivity should be higher than 30 W/m⋅k at 600°C. The dominant carrier of the thermal conduction in pure metals and dilute alloys is electrons (Ho et al 1975; Isachenko et al 1980; Touloukian and Ho 1981; Landolt 1991). When small amount of alloying elements (impurities) are introduced in iron, it causes discontinuities in the crystal lattice struc-ture and promotes scattering of free electrons decreasing the room temperature thermal conductivity (Riemann 1953; McGannon 1966). With rising temperature, the scattering of free electrons is intensified (Isachenko et al 1980). This causes a net reduction of thermal conductivity at higher temperature. Even then the thermal conductivity of dilute steel alloy is still high (> 30 W/m⋅k) at 600°C. Since in a fire, temperature gradients are encountered, a high ther-mal conductivity material is desired to avoid degradation by local necking.

4. Conclusions

Small additions of Mo (0⋅20–0⋅25%) and Cr (0⋅30–0⋅55%) in low carbon steel produced hot rolled fire resistant steel with a minimum of two-thirds of room temperature yield stress at 600°C. Microalloying by niobium and vanadium improved the ratio of YS600°C/YSRT. However, final rolling temperature should be maintained in the recrystallized austenite region to ensure a microstructure free from high dislocation density. The steels should be adequately treated by aluminium to tie-up free nitrogen in order to realize lower DBTT. In the quenched and self-tempered TMT rebar steel, addition of ~ 0⋅55% chromium was adequate to ensure the guaranteed strength at 600°C due to transforma-tion to tempered martensite rim and bainitic core. Despite decreasing trend of thermal conductivity with the rise of temperature in dilute alloys, the low alloy fire resistant steel will have higher thermal conductivity up to 600°C (> 30 W/m⋅k) compared to more concentrated alloys and is less expensive.

References

American Society of Testing Materials 1996 Standard test methods for fire tests of building construction and materials, Philadelphia, E119

Argent B B, Niekenk M N and Redfern G A 1970 J. Iron & Steel Inst. 208 830

Assefpour-Dezfully M, Hugas B A and Brownrigg A 1990 Mater. Sci. & Technol. 6 1210

Baird J D and Jamieson A 1972 J. Iron & Steel Inst. 210 847 Bhadeshia H K D H 1992 Bainite in steels (London: Institute of

Materials) Borato F, Barbosa R, Yue S and Jonas J J 1988 Proc. Thermec’88

(ed.) I Tamura (Tokyo: Iron and Steel Inst. Japan) p. 388 Bureau of Indian Standards 1998 Indian Standards IS 2062,

New Delhi Bureau of Indian Standards 2002 Indian Standards IS 15103,

New Delhi Chijiwa R, Tamehiro H, Yoshida Y, Funato K, Uemori R and

Horii Y 1993 Nippon Steel Tech. Report 58 47 Dahl W 1992 Steel (Dusseldorf: Springer Verlag and Verlag

Stahl Eisen) 1 DeArdo A J 1995 Microalloying’95 (Warrendale: Iron and Steel

Society) p. 15 Fushioni M, Chikaraishi H and Keira K 1995 Nippon Steel Tech

Report 66 29 Ho C Y, Powell R W and Liley P E 1975 Thermal conductivity

of the elements: A comprehensive review (NewYork: AIP) Honeycombe R W K 1981 Steel microstructure and properties

(London: Edward Arnold; Ohio : ASM) Houdremont E 1953 Handbook of special steels (Berlin:

Springer Verlag) 1 Irvine K J 1962 J. Iron & Steel Inst. 200 820 Isachenko V P, Osipova V A and Sukomel A S 1980 Heat

transfer (Moscow: Mir Publisher) Landolt B 1991 Thermal conductivity of pure metals and alloys

(eds) O Madelun and G K White (Berlin: Springer Verlag) 15C

McGannon H E (ed.) 1966 Making, shaping and treating of steels (Pittsburgh : USS)

Ouchi C, Sampei T and Kozasu I 1982 Iron & Steel Inst., Japan 22 214

Panigrahi B K 2001 Bull. Mater. Sci. 24 361 Panigrahi B K 2002 Unpublished result Panigrahi B K 2004 Seminar on Structural steel for construction

industry (NIT, Rourkela: The Institution of Engineers) Panigrahi B K and Jain S K 2002 Bull. Mater. Sci. 25 319 Pickering F B 1978 Physical metallurgy and design of steels

(London: Applied Science Pub.) Poliak E I and Jonas J J 2003 Iron & Steel Inst. Japan Int. 43

692 Riemann W 1953 Stahl und Eisen 73 721 Sage A M 1983 Proc. int. conf. steels for line pipe and pipe line

fittings (London: Metals Soc.) p. 39 Tanaka T 1981 Int. Metal. Rev. 26 185 Touloukian Y S and Ho C Y 1981 Properties of selected fer-

rous alloying elements (New York: McGraw Hill Book Co.) III.1

Wettlaufer M and Kasper R 2000 Steel Res. 71 357

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Some Fundamental Steps in Thermomechanical Processing

of Steels* .

By Imao TAMURA**

Key words: ~hermomechanicalprocetsing; steel; deformation o f austenite; restoration process; phase transformation; microalloying; ferri te grain re- Jinement.

I. Introduction

Thermomechanical processing is a treatment of which plastic deformation is introduced into the heat treatment of steel. We can find many treatments in such processes. But controlled rolling and accelerated cooling, and the ausforming or deformation- quenching are representative and are used in practical production.

I n controlled rolling, the austenite (y) is work- hardened with plastic deformation (rolling) in high strain rate a t elevated temperatures, and then air cooled to transform to ferrite (a). Sometimes, work- hardened austenite is cooled acceleratedly to take fine ferrite grains in steel plate. -

In the ausforming, work-hardened austenite is quenched to room temperature. We can take the ausformed martensite which increases its strength and not decrease its toughness. The ausforming is not so popular in practice but the deformation-quenching is getting to be popular. The martensite in the deformation-quenching looks like as usual martensite. But, this martensite has a little bit of fine facet size because it is formed from smaller austenite grains than usual austenite. This process can abridge to ration- alize the heat treatment process of steel.

As fundamental steps in thermomechanical proc-

True strain

tn 3

Fig. 1. Schematic illustration of two types of true stress- true strain curves accompanying the steady state deformation at elevated temperatures.

(a) dynamic recovery

essing of steels, we can point out that:

1) The deformation and restoration processes of y during and after hot rolling.

2) The kinetics of y to a transformation and grain size of a.

3) The effect of cooling rate on transformation of work-hardened y to a, and grain size of a.

4) Structure of martensite in ferrous alloys.

5) Structure and properties of martensite formed from work-hardened y.

6) Hardenability of work-hardened y and recrystallized y.

I will explain on these processes based on our research results.

(b) dynamic recrystallization

II. Deformation Behaviors of Austenite during and after Hot Working

I . Restoration Process during Hot Deformation of Auste-

I t is well known that the deformation behavior of austenite a t elevated temperatures is a dynamic recrystallization type. Figure 1 shows schematically both types of dynamic recovery type and dynamic recrystallization type.

I n the case of a dynamic recrystallization type, since the rate of softening due to recovery is slow, the work-hardening can not be balanced only by dynamic recovery even at high temperatures, and the dislocation density is still gradually increased with strain, and finally the recrystallization starts to occur during deformation. The true stress is lowered beyond the peak stress by the occurrence of dynamic recrystallization and reaches a certain steady state stress level at which the work-hardening and softening due to recrystallization occurring repeatedly are balanced. The dynamic recrystallization starts at the strain of around 0.78, (E,: the strain showing the peak stress). In Fig. l(b), the strain range up to about 0 . 7 ~ ~ is the work-hardening region, the strain range between about 0.78, and 8, (the strain at which the steady state deformation starts) is the range of partially dynamic recrystallization, and the strain range after E , ~ is the range of steady dynamic recrystallization. Austenite in steels belongs to this type. In a-Fe (ferrite), since the recovery occurs very fast, the restoration is performed mainly by dynamic recovery. The peak stress (a,) and the steady state stress (a,) in the dynamic recrystallization type change with the Zener-

* Based on the Special Lecture delivered by Prof. Tamura in commemoration of his receiving Nishiyama Medal in the 113th ISIJ Meeting, April 2, 1987, a t The University of Tokyo in Tokyo. Manuscript received on May 18, 1987. 0 1987 ISIJ

* * Professor Emeritus, Kyoto University; Technical Advisor, Sumitomo Metal Industries, Ltd., Nishinagasu-hondori, Amagasaki 660.

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Hollomon parameter (2). Z is expressed by

where, i: the strain rate (s-l) I? : a gas constant (J/mol)

Q: the activation energy for deformation (J/ mol)

7: temperature (K). a?,, ep, gs and E , are increased with increase in Z (i.e., with increase in i or decrease in 7).

Figure 2 shows the change in optical microstructure with strain in Fe-31%Ni-0.3%C austenitic alloy which was tensile deformed at 1 273 K (1 000°C) at a strain rate of 1.7 x 10-2/s. Figure 2(a) is the austenite structure before deformation. Austenite grains are equiaxed and their boundaries are smooth and planer. Many annealing twins exist in austenite grains. When deformed slightly (E= 0.08), as shown in Fig. 2 (b), the grain boundaries become finely serrated and the bulg- i n g ~ of grain boundary are observed as indicated by the arrows. At the strain slightly after E,, many small grains are newly formed along the austenite boundaries as shown in Fig. 2(c). I t is characteristic that these newly formed grains hardly contain annealing twins. Furthermore, some of the initial austenite grains becomes larger by the strain-enhanced coarsening. Figure 2(d) shows the structure in the range of steady state deformation at which the work-hardening and the dynamic recrystallization softening are balanced. This structure consists of several grains corresponding to various stages of recrystallization, i.e., from small grains which have just started to recrys- tallize to fairly large grains which have been subjected to deformation after the finishing of recrystallization and are the work-hardened state just before the next recrystallization. I n the stage of steady state deformation, the recrystallization occurs repeatedly during deformation. Dynamic recrystallized grains contain a few annealing twins a t low Z but scarcely contain

them when Z is high. I n the same way as the static recry~tallization,~) it

has been observed that there are two main mechanisms for dynamic recrystallization, i.e., the bulging mechanism and the nucleation-growth mechanism.

Dynamic recrystallization becomes easier to take place when the initial grain size (do) of austenite is smaller. Therefore, in the case of fine initial grain size, c p and E, become small even at the deformation under the same Z condition. However, a,, and a , are independent of do. The average size (d) of dynamically recrystallized grains is the only function of Z and is expressed by the following equation,')

where, A, : material's constant. A is larger for C-steel than for alloy steels, especially small for Nb-containing steel and high alloy steel. p is about 0.3-0.4 for all steels. I t should be noted that d is independent of do.

2. Static Restoration Process during Holding Right after Hot Deformation

When the hot deformation is finished a t the work- hardening stage, the work-hardened austenite is softened by the usual static recrystallization during isothermal holding after deformation. However, if the strain is too small, the recrystallization can not occur and only the recovery takes place. On the other hand, when specimens are isothermally held after deformation to the range of the steady state deformation, in which the dynamic recrystallization has occurred, the recrystallization and the grain growth can succeed to occur continuously during holding after hot deformation. Such a recrystallization is called a metadynamic or postdynamic recrystallization. Namely, the dynamically recrystallized structure is further softened by the metadynamic recrystallization. When specimens are deformed to the range of (11) in Fig.

(a) Before testing (E=O) (b) € =0.08 (c) e=0.24 (d) &=0.40

Pig. 2. Optical micrographs showing tlie formation process of dynamically recrystallized austenite in Fc-31Ni-0.3C austenitic alloy deformed at 1 OOO°C, 1.7 x 10-2/s.

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l (b) in which the work-hardened austenite and dynamically recrystallized austenite have coexisted, the static recrystallization occurs in work-hardened austenite and the metadynamic recrystallization occurs in dynamically recrystallized austenite at the same time.

I n case of rolling by hot strip mill in which slabs are rolled by several passes, the static restoration process occurring during interval between rolling passes has great influence on the deformation behavior of next rolling pass. If the softening due to the static restoration is not completed during interval between rolling passes, the work-hardening is successively ac- cumulated by the subsequent rolling passes.

Fraction of softening (X) occurring during isothermal holding right after hot deformation can be evaluated by the two-step deformation test from the following e q ~ a t i o n , ~ )

where, cr,,,: the flow stress immediately before unloading

c ~ , , ~ : the yield stress on reloading after holding the specimen for various periods of time (At)

u,,, : the initial yield stress. As an example, the change in the fractional softening (X) with the holding time (At) in Fe-31 %Ni-O.S%C austenitic alloy is shown in Fig. 3. Specimens were tensile deformed a t 1 373 K (1 100°C) a t a strain rate of 4 x 10-2/s to E = 0.18 (work-hardened austenite) and e= 1.0 (completely dynamically recrystallized austenite), respectively, and then isothermally hild for various periods of time. The change in average aus-

tenite grain size with At is also plotted in Fig. 3. Microstructural changes during isothermal holding a t 1 373 K (1 100°C) are shown in Fig. 4 and Fig. 5. As shown in Fig. 4 (~=0.18) , work-hardened austenites are slightly elongated along the tensile direction (Fig. 4(a)). The austenite hardly changes by holding for 1 s (Fig. 4(b)), although the softening occurs to some extent by the static recovery as can be seen in Fig. 3. However, when specimens were isothermally held for 4 s (Fig. 4(c)), fine austenite grains which are formed by static recrystallization are partially formed mainly along the deformed austenite grain boundaries and the static recrystallization is completed by the holding of 15 s (Fig. 4(d)). Annealing twins in static recrystallized austenites are gradually developed with increase in holding time. By holding for 40 s, recrystallized austenite grains are coarsened (Fig. 3). When the austenite was much more heavily deformed under higher Z condition (work-hardened), the static recrystallized austenite grains become much finer as shown in Fig. 6. As shown in Fig. 5 (a=1.0), dynamically recrystallized austenites are characterized by irregular grain boundaries, a low density of annealing twins and the mixed grain size. With increase in holding time, austenite grain boundaries become fairly straight and austenite structure becomes almost equiaxed grains, and furthermore, the density of annealing twins is increased. This is the typical microstructural change occurring during metadynamic recrystallization process. Metadynamic recrystallization is completed by the holding for about 30 s in the case of Fig. 5, and the usual grain coarsening occurs by further holding. The grain size right after the completion of metadynamic recrystallization is smaller, as the grain size of dynamically recrystallized

Hdding time. A t (s)

Fig. 3. Change in fraction of softening (a) and austenite As-deformed (A t= 0 S) (b) At = 1 s (c) At = 4 s (d) ~t = 15 s grain size (b) with the isothermal holding time of Fig. 4. Change in optical microstructure of work-hardened austenite (de- work-hardened austenite and dynamically recrystal- formed at 1 100°C, t = 4.0 x 10-2/s, c ~ 0 . 1 8 ) during isothermal holding lized austcnite in Fe-31Ni-0.3C austenitic alloy. at 1 100°C for At in Fe-31Ni-0.3C austcnitic alloy.

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(a) AS-deformed ( ~ t = O s) (b) At = 1 s (c) At=2 s (d) At=30 s (a) Work-hardened austenite (detormed at 900°C,

Fig. 5. Change in optical microstructure of dynamically recrystallized austenite e=3/s, ~ z 1 . 0 ) (deformed at 1 100°C, d = 4.0 x 10-Z/s, E = 1 .O) during isothermal (h) Static recrystallized austenite (isothermally holding a t 1 100°C for At in Fe-31Ni-0.3C austenitic alloy. held a t 900°C for 50 s of structure (a)

Fig. 6. Optical micrographs showing the static recrystal-

austenite is smaller. lization occurring during isothermal holding after heavily work-hardened austenite in Fe-31Ni-0.3C

3. Efect of Alloying Elements on the lietardation of Re- austenitic alloy.

crystallization5) Table 1. The values of n and m in Eq. (4) for varlous

The effect of alloying elements on the restoration transformation conditions.

process during or after hot deformation is very important. Micro-addition of Nb or Ti acts as a re- Mechanism Site saturation Nucleation and

markable retardation effect on recrystallization due to growth -- - - - - - --

the suppression of grain boundary migration. The Site n m n m suppression of grain boundary migration due to micro-

Grain surface 1 1 4 1 alloying is caused from (1) the solute dragging effect Grain edge 2 2 4 2 due to segregation of alloying elements to the bound- Grain corner 3 3 4 3

aries, or (2) the pinning effect due to precipitates of carbonitride of alloying elements at grain boundaries. For example, Mo segregates to the austenite grain x = l - exp [ - k ( ~ ) $ ] ..... . . . . . . . . . . (4) boundaries, resulting in the suppression of grain boundary migration. Even Nb micro-addition, at where, X: the transfbrmed fraction high 1, it can retard the recrystallization by dragging effect. But, usually, Nb and T i are precipitated as fine carbonitride particles and these precipitates suppress the grain boundary migration by the pinning effect.

III. Transformation Kinetics and Grain Size of Ferrite Transformed from Austenite

I . Isothermal Transformation Kinetics of Austenite

When the ferrite or pearlite Sorms from austenite by isothermal holding below the transformation temperature, these nuclei form preferentially a t the prior austenite grain boundaries. We expressed as Eq. (4)'3 that their transformation kinetics include the effect of austenite grain size, by modification of Johnson-Mehl equation.

t : the isothermal holding time d : the austenite grain size

k ( 7 ) : the rate constant which depends only on transformation temperature and directly corresponds to the C-curve of T T T diagram (isothermal transformation diagram).

The values oS the exponent m and n are listed in Table 1.6) These values depend on both the transformation mechanism and type of nucleation site. Here the transformation which progresses with nucleation and growth over almost entire range of transformation is called " nucleation and growth ". The transformation where all the nuclei are nucleated a t the early stage in the reaction and the progress of transformation is controlled mostly by their growth is called " site saturation ". An austenite grain is assumed to

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Transactions ISIJ, Vol. 27, 1987 ( 767)

be a tetrakaidecahedra, and the nucleation sites a t austenite grain boundary can be classified into surfaces, edges and corners as are shown in Fig. 7.697) Of course the actual nucleation sites are not limited to any one of such types, and the value of m would be the weighted average of the various nucleation sites and not always be an integer.

According to Umemoto and Tamura,e) the values of n and m, and transformation modes were experimentally shown as in Table 2. I t is mentioned in Table 2 that the mode of y +a transformation is site saturation, but i t means the large number of nucleation o f a would occur at the beginning of transformation, which would not be real "site saturation ". Sometimes the nucleation site is varied by supercooling (accelerated cooling). I t seems that the nucleation site of a is moved gradually from edge to surhce by decrease in transformation temperature with cooling rate, as described later.

2. Continuous Cooling Transformation Kinetics of Austenite

Scheils) considered the case where the austenite is cooled continuously through its transformation temperature. He assumed that austenite consumes its fractional nucleation time and when the sum of a number of such fractions equals unity, the transformation starts, i.e.,

Fig. 7. Schematic diagram of austenite grain as a tetrakaidecahedra and showing 3 types of nucleation sites offerritc.

'Table 2. The values of r2 and m for various transformation modes.

Transformation n m Nucleation -

Pearlite 4 2 Edge nucleation Nucleation and growth

1:errite 1 1 Surface nucleation near Site saturation

Bainite 4 0 .6 Grain boundary, Inside grain

Now extend this Scheil's additivity rule to the entire range of transformation. I n this case the incu- bation time ( ~ ( 7 ) ) is replaced with the time required for the reaction to reach a certain fractional completion X by isothermal holding at temperature T. If we combine Eq. (5) with Eq. (4)

where, Q ( T ) = -dT/dt: a cooling rate. This equation expresses the transformed fraction at temperature T during cooling from austenite condition with a cooling rate Q(T).6t9) Furthermore, from Eqs. (4) and (6) we obtain

1 Te K(T1)1/7~ t ( 7 ) = - J - - - d T ' = t., ...... (7)

K T ) T Q(T')

This equation tells that the fraction transformed by cooling from 7, to T with cooling rate Q ( T ) is equal to that obtained by the isothermal holding at temperature T for time t (7) . Thus we call t (T) as an equivalent cooling time, t,,. t,, means the isothermal holding time which gives the equivalent transformed fraction with continuous cooling. The curve produced by the successive plot oS such t,, at each temperature during cooling with cooling rate Q ( T ) will be called an equivalent cooling curve.'O) The transformation behavior for a given cooling curve can be predicted from the T T T diagram by plotting such curve on a T T T diagram.

3. Ferrite Grain Size Transformed from Austenite 1. Expression for Ferrite Grain Size (Isothermal)

In general, ferrites nucleate at austenite grain boundaries and grow into austenite grains. The grain size of ferrite formed from austenite is directly related with the process of transformation.") The nominal ferrite grain diameter D, can be expressed with the total number of ferrite grains uncleated throughout transfbrmation per unit volume of austenite, n,, as12)

where, I : nucleation rate per unit volume oS austenite

X: volume Sraction transformed Thus the problem of obtaining the exprcssion for ferrite grain size is the problem of how to express n,>.

When ferrites nucleate a t the austenite grain boundary surfaces, they grow as ellipsoid with the aspect ratio of 3: 1 1 3 ) as shown in Fig. 8. The total number of rerrite grains nucleated per unit area of austenite grain boundary surface is

9 n, = I,J; exp (-Tzaz~stz dt = - - ) 3:2-(':')

Nuclcation and growth

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ferrite g r a i n

g ra in boundory

r=3Cl(t- 2)

Fig. 8. Schematic drawing of a ferrite grain nucleated at austenite grain boundary surface.

where, I , : the nucleation rate per unit area oL austenite grain surface

a : parabolic rate constant for thickness. The ferrite grain size D, is given as

where, Sg .b : the austenite grain surface area per unit volume

and can be expressed by austenite grain diameter, D,, as

Substituting Eqs. (10) and (12) into Eq. ( l l ) , the ferrite grain size and the y / a transfbrmation ratio can be expressed as

and

In the above derivation it was assumed that the nucleation rate and the parabolic rate constant of growth are time independent and one nuclei grows and forms one grain.

Similarly the D, and y / a transformation ratio in the case of homogeneous and grain edge nucleation has been obtained asl1)

where, I h : the homogeneous nucleation rate per unit volume

I ,: nucleation rate per unit length of austenite grain edge.

I t should be noted that the equations introduced above indicate that ferrite grain size and y/a transformation ratio depend on the ratio of square root of the nucleation rate to the parabolic rate constant in all three types of nucleation sites. Furthermore Sor

grain boundary nucleation a small austenite grain leads to a small ferrite grain. The effect of austenite grain size on ferrite grain size depends on the type of nucleation site, and it becomes large in the following ascending order: homogeneous, grain surSaces, and grain edges.

I t should be noted that the grain refinement on austenite turns out to be less effective in ferrite grain refinement. For instance in the case of grain surface nucleation D, is proportional to D',I3 in Eq. (13). This means that when D, is reduced to 118 of the initial, D, would be reduced only into 112.

Figure 914) shows experimental results of isothermal transformation in a steel (0.2 % C, 0.0016 0/, B). The steel was austenitized between 1 000 and 1200°C to change its austenitic grain size from about 53 to 88 pm, and then transformed at 700, 680 and 660°C into Ferrite. The results are recognized as the dependent of y-grain diameter that a nucleates mainly at grain edges at 700°C, at grain edges and surfaces at 680°C and at the surfaces at 660°C. 2. Estimation of Ferrite Grain Size Formed by Continu-

ous Cooling Transformation It can be assumed ferrite nucleates preferentially at

austenite grain surfaces by supercooling. A continuous cooling transformation can be considered as the sum of short time isothermal holdings at successive temperatures as is schematically shown in Fig. 10.

During cooling, ferrite grains are nucleated and grow at each temperature with the corresponding nucleation and growth rates, as is schematically shown in the Fig. 10 (bottom). During ferrite transformation carbon enriches in retained austenite, however, since ferrite nucleation occurs mostly in the early stage of transformation, it can be assumed that both the nucleation and growth rates of ferrite are a function of an instantaneous temperature only and not a function of thermal history.

Based on this assumption, the number of ferrite nuclei (N,) nucleated at temperature I, during cooling, N, is given as Eq. (16).

where, Q(T) (= -dT/dt) : a cooling rate. Ferrite grains once nucleated grow thereafter with a corresponding growth rate at each temperature during cooling. Thus, at temperature T, the radius of a ferrite grain on an austenite grain surface nucleated at temperaturc T,, is given as Eq. (1 7).

The austenite grain boundary area occupied at tcm- pcrature 7 by a ferrite grain nucleated at temperature T, is given as Eq. (18).

I t will be convenient to use the concept of extended area which is the sum of the areas of austenite grain

Special Lecture

Transactions ISIJ, Vol. 27, 1987 ( 769 ]

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A u s t e n i t e G r a i n Size

Fig. 9. An example of D,- or D,/D,,-D, rclationsllip for S15C steel.

surface occupied by ferrite grains, assuming that fcr- rite grain never stop growing and that ferrite grains keep nucleating a t the same rate during transformation as well as on untransformed austenite grain surfaces. The extended grain boundary area occupied by all the ferrite grains nucleated at T , during cooling is given as Eq. (1 9).

Thus the grain boundary arca fraction occupied by the ferrite grains nucleated during cooling horn A3(T , ) to T is givcn by Eq. (20).

{J:$?&; dl-' d T , ............ 1 1 (2 0)

The total number of ferrite grains nucleated per unit area of austenite grain boundary surfacc is given by Eq. (21).

Ferrite grain size is given as Eq. (22).

o ~ * ~ ~ ~ * TI T2 T3 Tl,

d ( h ) d ( T 2 ) d (T3 ) ~ ( T L ) I s (T ) Is(T2) Is(T3) Is(TL )

Fig. 10. Diagrams to explain the relation between isothermal and continuous cooling transformation and schematic drawing showing the formation process of ferrite during continuous cooling.

673 773 873 973 1073 1173

TEMPERATURE, K

Fig. 11. Calculated parabolic rate constant and nucleation rate I, as a function of' temperature for the steel studied. (S15C)

where, Sg.,,(=4/(dn D,)): the austenite grain surface area pcr unit volume.

From these equations, ferrite grain s i ~ e can be calculated.

I t is fairly difficult to estimate of I , ( T ) and tt(T) by some kinds of calculation, but it is not impossible. By using several assumptions and h y p o t h e s i ~ , ~ ~ ' ~ - ~ ~ ) I , ( T ) and a ( T ) are roughly calculated as shown in Fig. 1 lZ4) as a function of transformation temperature. In this figure, I , shows a maximum value at 913 K (640°C) and n shows a maximum value at

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873 K (600°C) fbr S15C steel (0.15 % C, 0.4 % Mn). Figure l2Z4) shows a comparison of calculated

values and experimental values for C15C steel. Wc can get a roughly agreement, that is, a-grain size is getting smaller with cooling rate increases. The Sor- mulations are :

Dtv = Q-0.17(D,)1'~cal~ulat~d) ... ......... ...... (23)

Do = 5.7Q-0.2'3(D,)0.46 (experimental) . . . . .. . . .(24)

Figure 1324) shows experimental results.

IV. Ferrite Grain Refinement by Work-hardening of Austenite

I . Ferrite Nucleation in Deformed Austenite (Isothermal)

Figure 14 is typical micrographs showing the effect of austenite defhrmation on the progress of ferrite transformation.Z5) Specimens of an Fe-0.12C-0.04- Nb-0.04V steel were austenitized at 1 200°C for 30 min, rolled 30 or 50 % in reduction by a single pass a t 840°C (in the unrecrystallized temperature range) and held at 680°C for 15, 30 and 180 s. The ferrite transformation is clearly accelerated by deformation

C O O L I N G IRATE , K is

Fig. 1 l'erritc grain s i ~ e as a function of cooling rate lor the specimens of D, = 15 1 pm. (S 15C) Fig. 13. Ferrite grain size as a function of Q-0.2'3Dt4G.

Not Rolled Rolled (30%) Rolled (50%)

(a) Non-deformed, 15 s (b) 30 % rolled, 15 s (c) 50 % rolled, 15 s

(d) Non-deformed, 30 s (e) 30 % rolled, 30 s (f) 50 % rolled, 30 s (g) Non-defbrmed, 3 min (h) 30 % rolled, 3 min (i) 50 0/6 rolled, 3 min

Fig. 14. Optical micrograph showing the efrect of austenite deformation on the proccss of' ferrite transfbrmation a t 680°C in a Fe-O.12C-0.04Nb-0.04V steel.

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