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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 decreases.

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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.

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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 forgeability, drawability
# Low temperature toughness
# Fatigue strength
# C, N Solid solution Solid solubility (enhanced), hardenability
# Settled dislocation Strain aging (enhanced), ductility and toughness (lowered)
# Pearlite and cementite Dispersion (enhanced), ductility and toughness (lowered)
# Carbide and nitride precipitates Precipitation, grain refining (enhanced), toughness (enhanced)
# Embrittlement by intergranular precipitation
# P Solid solution Solid solubility (enhanced), 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 ppm

10 µm
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

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Steel is made by the Bessemer, Siemens Open Hearth, basic oxygen furnace, electric arc, electric high-frequency and crucible processes.

Crucible and high-frequency methods

The Huntsman crucible process has been superseded by the high frequency induction furnace in which the heat is generated in the metal itself by eddy currents induced by a magnetic field set up by an alternating current, which passes round water-cooled coils surrounding the crucible. The eddy currents increase with the square of the frequency, and an input current which alternates from 500 to 2000 hertz is necessary. As the frequency increases, the eddy currents tend to travel nearer and nearer the surface of a charge (i.e. shallow penetration). The heat developed in the charge depends on the cross-sectional area which carries current, and large furnaces use frequencies low enough to get adequate current penetration.

Automatic circulation of the melt in a vertical direction, due to eddy currents, promotes uniformity of analysis. Contamination by furnace gases is obviated and charges from 1 to 5 tonnes can be melted with resultant economy. Consequently, these electric furnaces are being used to produce high quality steels, such as ball bearing, stainless, magnet, die and tool steels.

Figure 1.
Furnaces used for making pig iron and steels. RH side of open hearth furnace shows use of oil instead of gas

Acid and basic steels

The remaining methods for making steel do so by removing impurities from pig iron or a mixture of pig iron and steel scrap. The impurities removed, however, depend on whether an acid (siliceous) or basic (limey) slag is used. An acid slag necessitates the use of an acid furnace lining (silica); a basic slag, a basic lining of magnesite or dolomite, with line in the charge. With an acid slag silicon, manganese and carbon only are removed by oxidation, consequently the raw material must not contain phosphorus and sulphur in amounts exceeding those permissible in the finished steel.

In the basic processes, silicon, manganese, carbon, phosphorus and sulphur can be removed from the charge, but normally the raw material contains low silicon and high phosphorus contents. To remove the phosphorus the bath of metal must be oxidised to a greater extent than in the corresponding acid process, and the final quality of the steel depends very largely on the degree of this oxidation, before deoxidisers-ferro-manganese, ferro-silicon, aluminium-remove the soluble iron oxide and form other insoluble oxides, which produce non-metallic inclusions if they are not removed from the melt:

2Al + 3FeO (soluble) « 3Fe + Al2O3 (solid)

In the acid processes, deoxidation can take place in the furnaces, leaving a reasonable time for the inclusions to rise into the slag and so be removed before casting. Whereas in the basic furnaces, deoxidation is rarely carried out in the presence of the slag, otherwise phosphorus would return to the metal. Deoxidation of the metal frequently takes place in the ladle, leaving only a short time for the deoxidation products to be removed. For these reasons acid steel is considered better than basic for certain purposes, such as large forging ingots and ball bearing steel. The introduction of vacuum degassing hastened the decline of the acid processes.

Bessemer steel

In both the Acid Bessemer and Basic Bessemer (or Thomas) processes molten pig iron is refined by blowing air through it in an egg-shaped vessel, known as a converter, of 15-25 tonnes capacity (Fig. 1). The oxidation of the impurities raises the charge to a suitable temperature; which is therefore dependent on the composition of the raw material for its heat: 2% silicon in the acid and 1,5-2% phosphorus in the basic process is normally necessary to supply the heat. The “blowing” of the charge, which causes an intense flame at the mouth of the converter, takes about 25 minutes and such a short interval makes exact control of the process a little difficult.

The Acid Bessemer suffered a decline in favour of the Acid Open Hearth steel process, mainly due to economic factors which in turn has been ousted by the basic electric arc furnace coupled with vacuum degassing.

The Basic Bessemer process is used a great deal on the Continent for making, from a very suitable pig iron, a cheap class of steel, e.g. ship plates, structural sections. For making steel castings a modification known as a Tropenas converter is used, in which the air impinges on the surface of the metal from side tuyeres instead of from the bottom. The raw material is usually melted in a cupola and weighed amounts charged into the converter.

Open-hearth processes

In the Siemens process, both acid and basic, the necessary heat for melting and working the charge is supplied by oil or gas. But the gas and air are preheated by regenerators, two on each side of the furnace, alternatively heated by the waste gases. The regenerators are chambers filled with checker brickwork, brick and space alternating.

The furnaces have a saucer-like hearth, with a capacity which varies from 600 tonnes for fixed, to 200 tonnes for tilting furnaces (Fig. 1). The raw materials consist essentially of pig iron (cold or molten) and scrap, together with lime in the basic process. To promote the oxidation of the impurities iron ore is charged into the melt although increasing use is being made of oxygen lancing. The time for working a charge varies from about 6 to 14 hours, and control is therefore much easier than in the case of the Bessemer process.

The Basic Open Hearth process was used for the bulk of the cheaper grades of steel, but there is a growing tendency to replace the OH furnace by large arc furnaces using a single slag process especially for melting scrap and coupled with vacuum degassing in some cases.

Electric arc process

The heat required in this process is generated by electric arcs struck between carbon electrodes and the metal bath (Fig. 1). Usually, a charge of graded steel scrap is melted under an oxidising basic slag to remove the phosphorus. The impure slag is removed by tilting the furnace. A second limey slag is used to remove sulphur and to deoxidise the metal in the furnace. This results in a high degree of purification and high quality steel can be made, so long as gas absorption due to excessively high temperatures is avoided. This process is used extensively for making highly alloyed steel such as stainless, heat-resisting and high-speed steels.

Oxygen lancing is often used for removing carbon in the presence of chromium and enables scrap stainless steel to be used. The nitrogen content of steels made by the Bessemer and electric arc processes is about 0,01-0,25% compared with about 0,002-0,008% in open hearth steels.

Oxygen processes

The high nitrogen content of Bessemer steel is a disadvantage for certain cold forming applications and continental works have, in recent years, developed modified processes in which oxygen replaces air. In Austria the LID process (Linz-Donawitz) converts low phosphorus pig iron into steel by top blowing with an oxygen lance using a basic lined vessel (Fig. 2b). To avoid excessive heat scrap or ore is added. High quality steel is produced with low hydrogen and nitrogen (0,002%). A further modification of the process is to add lime powder to the oxygen jet (OLP process) when higher phosphorus pig is used.

Figure 2.

The Kaldo (Swedish) process uses top blowing with oxygen together with a basic lined rotating (30 rev/min) furnace to get efficient mixing (Fig. 2a). The use of oxygen allows the simultaneous removal of carbon and phosphorus from the (P, 1,85%) pig iron. Lime and ore are added. The German Rotor process uses a rotary furnace with two oxygen nozzles, one in the metal and one above it (Fig. 2c). The use of oxygen with steam (to reduce the temperature) in the traditional basic Bessemer process is also now widely used to produce low nitrogen steel. These new techniques produce steel with low percentages of N, S, P, which are quite competitive with open hearth quality.

Other processes which are developing are the Fuel-oxygen-scrap, FOS process, and spray steelmaking which consists in pouring iron through a ring, the periphery of which is provided with jets through which oxygen and fluxes are blown in such a way as to “atomise” the iron, the large surface to mass ratio provided in this way giving extremely rapid chemical refining and conversion to steel.

Vacuum degassing is also gaining ground for special alloys. Some 14 processes can be grouped as stream, ladle, mould and circulation (e.g. DH and RH) degassing methods, Fig. 3. The vacuum largely removes hydrogen, atmospheric and volatile impurities (Sn, Cu, Pb, Sb), reduces metal oxides by the C – O reaction and eliminates the oxides from normal deoxidisers and allows control of alloy composition to close limits. The clean metal produced is of a consistent high quality, with good properties in the transverse direction of rolled products. Bearing steels have greatly improved fatigue life and stainless steels can be made to lower carbon contents.

Figure 3. Methods of degassing molten steel

Vacuum melting and ESR. The aircraft designer has continually called for new alloy steels of greater uniformity and reproducibility of properties with lower oxygen and sulphur contents. Complex alloy steels have a greater tendency to macro-segregation, and considerable difficulty exists in minimising the non-metallic inclusions and in accurately controlling the analysis of reactive elements such as Ti, Al, B. This problem led to the use of three processes of melting.

(a) Vacuum induction melting within a tank for producing super alloys (Ni and Co base), in some cases for further remelting for investment casting. Pure materials are used and volatile tramp elements can be removed.
(b) Consumable electrode vacuum arc re-melting process (Fig. 4) originally used for titanium, was found to eliminate hydrogen, the A and V segregates and also the large silicate inclusions. This is due to the mode of solidification. The moving parts in aircraft engines are made by this process, due to the need for high strength cleanness, uniformity of properties, toughness and freedom from hydrogen and tramp elements.
(c) Electroslag refining (ESR) This process, which is a larger form of the original welding process, re-melts a preformed electrode of alloy into a water-cooled crucible, utilising the electrical resistance heating in a molten slag pool for the heat source (Fig. 5). The layer of slag around the ingot maintains vertical unidirectional freezing from the base. Tramp elements are not removed and lead may be picked up from the slag.

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A designation is the specific identification of each grade, type, or class of steel by a number, letter, symbol, name, or suitable combination. Unique to a particular steel grade, type and class are terms used to classify steel products. Within the steel industry, they have very specific uses: grade is used to denote chemical composition; type is used to indicate deoxidation practice; and class is used to describe some other attribute, such as strength level or surface smoothness.

In ASTM specifications, however, these terms are used somewhat interchangeably. In ASTM A 533, for example, type denotes chemical composition, while class indicates strength level. In ASTM A 515, grade identifies strength level; the maximum carbon content permitted by this specification depends on both plate thickness and strength level. In ASTM A 302 grade denotes requirements for both chemical composition and mechanical properties. ASTM A 514 and A 5117 are specifications for high-strength quenched and tempered plate for structural and pressure vessel applications, respectively, each contains several compositions that can provide the required mechanical properties. However, A 514 type A has the identical composition limits as A 517 grade.

Chemical composition is by far the most widely used basis for classification and/or designation of steels. The most commonly used system of designation in the United States is that of the Society of Automotive Engineers (SAE) and the American Iron and Steel Institute (AISI). The Unified Numbering System (UNS) is also being used with increasing frequency.

SAE-AISI Designations
As stated above, the most widely used system for designating carbon and alloy steels is the SAE-AISI system. As a point of technicality, there are two separate systems, but they are nearly identical and have been carefully coordinated by the two groups. It should be noted, however, that AISI has discontinued the practice of designating steels.

The SAE-AISI system is applied to semi-finished forgings, hot-rolled and cold-finished bars, wire rod and seamless tubular goods, structural shapes, plates, sheet, strip, and welded tubing.

Carbon steels contain less than 1.65% Mn, 0.60% Si, and 0.60% Cu; they comprise the lxxx groups in the SAE-AISI system and are subdivided into four distinct series as a result of the difference in certain fundamental properties among them.

Designations for merchant quality steels include the prefix M. A carbon steel designation with the letter B inserted between the second and third digits indicates the steel contains 0.0005 to 0.003% B. Likewise, the letter L inserted between the second and third digits indicates that the steel contains 0.15 to 0.35% Pb for enhanced machinability. Resulfurized carbon steels of the 11xx group and resulfurized and rephosphorized carbon steels of the 12xx group are produced for applications requiring good machinability. Steels that having nominal manganese contents of between 0.9 and 1.5% but no other alloying additions now have 15xx designations in place of the 10xx designations formerly used.

Alloy steels contain manganese, silicon, or copper in quantities greater than those listed for the carbon steels, or they have specified ranges or minimums for one or more of the other alloying elements. In the AISI-SAE system of designations, the major alloying elements are indicated by the first two digits of the designation. The amount of carbon, in hundredths of a percent, is indicated by the last two (or three) digits.

For alloy steels that have specific hardenability requirements, the suffix H is used to distinguish these steels from corresponding grades that have no hardenability requirement. As with carbon steels, the letter B inserted between the second and third digits indicates that the steel contains boron. The prefix E signifies that the steel was produced by the electric furnace process.

HSLA Steels. Several grades of HSLA steel are described in SAE Recommended Practice J410. These steels have been developed as a compromise between the convenient fabrication characteristics and low cost of plain carbon steels and the high strength of heat-treated alloy steels. These steels have excellent strength and ductility as-rolled.

UNS Designations The Unified Numbering System (UNS) has been developed by ASTM and SAE and several other technical societies, trade associations, and United States government agencies.

A UNS number, which is a designation of chemical composition and not a specification, is assigned to each chemical composition of a metallic alloy. The UNS designation of an alloy consists of a letter and five numerals. The letters indicate the broad class of alloys; the numerals define specific alloys within that class. Existing designation system, such as the AISI-SAE system for steels, have been incorporated into UNS designations. UNS is described in greater detail in SAE J1086 and ASTM E 527.
AMS Designation
Aerospace Materials Specifications (AMS), published by SAE, are complete specifications that are generally adequate for procurement purposes. Most of the AMS designations pertain to materials intended for aerospace applications; the specifications may include mechanical property requirements significantly more severe than those for grades of steel having similar compositions but intended for other applications. Processing requirements, such as for consumable electrode remelting, are common in AMS steels.

ASTM (ASME) Specifications The most widely used standard specifications for steel products in the United States are those published by ASTM. These are complete specifications, generally adequate for procurement purposes. Many ASTM specifications apply to specific products, such as A 574 for alloy steel socket head cap screws. These specifications are generally oriented toward performance of the fabricated end product, with considerable latitude in chemical composition of the steel used to make the end product.

ASTM specifications represent a consensus among producers, specifiers, fabricators, and users of steel mill products. In many cases, the dimensions, tolerances, limits, and restrictions in the ASTM specifications are similar to or the same as the corresponding items of the standard practices in the AISI Steel Products Manuals.

Many of the ASTM specifications have been adopted by the American Society of Mechanical Engineers (ASME) with little or no modification; ASME uses the prefix S and the ASTM designation for these specifications. For example, ASME-SA213 and ASTM A 213 are identical.

Steel products can be identified by the number of the ASTM specification to which they are made. The number consists of the letter A (for ferrous materials) and an arbitrary, serially assigned number. Citing the specification number, however, is not always adequate to completely describe a steel product. For example, A 434 is the specification for heat-treated (hardened and tempered) alloy steel bars. To completely describe steel bars indicated by this specification, the grade (SAE-AISI designation in this case) and class (required strength level) must also be indicated. The ASTM specification A 434 also incorporates, by reference, two standards for test methods (A 370 for mechanical testing and E 112 for grain size determination) and A 29, which specifies the general requirements for bar products.

SAE-AISI designations for the compositions of carbon and alloy steels are sometimes incorporated into the ASTM specifications for bars, wires, and billets for forging. Some ASTM specifications for sheet products include SAE-AISI designations for composition. The ASTM specifications for plates and structural shapes generally specify the limits and ranges of chemical composition directly, without the SAE.AISI designations.

General Specifications. Several ASTM specifications, such as A 20 covering steel plate used for pressure vessels, contain the general requirements common to each member of a broad family of steel products. These general specifications are often supplemented by additional specifications describing a different mill form or intermediate fabricated product.

European and Japanese Designation Systems
Below some basics of European and Japanese designation systems are explained. Please refer to articles about corresponding national and international standards for more details.

DIN standards are developed by Deutsches Institut fur Normung in the Federal Republic of Germany. All West German steel specifications are preceded by the uppercase letters DIN followed an alphanumeric or numeric code. The latter method, known as the Werkstoff number, uses numbers only with a decimal point after the first digit.

JIS standards are developed by the Japanese Industrial Standards Committee, which is part of the Ministry of International Trade and Industry in Tokyo. The JIS steel specifications begin with the uppercase letters JIS and are followed by an uppercase letter (G in the case of carbon and low-alloy steels) designating the division (product form) of the standard. This letter is followed by a series of numbers and letters that indicate the specific steel.

British standards (BS) are developed by the British Standards Institute in London, England. Similar to the JIS standards, each British designation includes a product form and an alloy code.

AFNOR standards are developed by the Association Francaise de Normalisation in Paris, France. The correct format for reporting AFNOR standards is as follows. An uppercase NF is placed to the left of the alphanumeric code. This code consists of an uppercase letter followed by a series of digits, which are subsequently followed by an alphanumeric sequence.

UNI standards are developed by the Ente Nazionale Italiano di Unificazione in Milan, Italy. Italian standards are preceded by the uppercase letter UNI followed by a four-digit product form code subsequently followed by an alphanumeric alloy identification.

Swedish standards (SS) are prepared by the Swedish Standards Institution in Stockholm. Designations begin with the letters SS followed by the number 14 (all Swedish carbon and low-alloy steels are covered by SS14). What subsequently follows is a four digit numerical sequence similar to the German Werkstoff number.

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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 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
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.

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 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 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.

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A study of the constitution and structure of all steels and irons must first start with the iron-carbon equilibrium diagram. Many of the basic features of this system (Fig. 1) influence the behavior of even the most complex alloy steels. For example, the phases found in the simple binary Fe-C system persist in complex steels, but it is necessary to examine the effects alloying elements have on the formation and properties of these phases. The iron-carbon diagram provides a valuable foundation on which to build knowledge of both plain carbon and alloy steels in their immense variety.

Fig. 1. The iron-carbon diagram.

It should first be pointed out that the normal equilibrium diagram really represents the metastable equilibrium between iron and iron carbide (cementite). Cementite is metastable, and the true equilibrium should be between iron and graphite. Although graphite occurs extensively in cast irons (2-4 wt % C), it is usually difficult to obtain this equilibrium phase in steels (0.03-1.5 wt %C). Therefore, the metastable equilibrium between iron and iron carbide should be considered, because it is relevant to the behavior of most steels in practice.

The much larger phase field of γ-iron (austenite) compared with that of α-iron (ferrite) reflects the much greater solubility of carbon in γ-iron, with a maximum value of just over 2 wt % at 1147°C (E, Fig.1). This high solubility of carbon in γ-iron is of extreme importance in heat treatment, when solution treatment in the γ-region followed by rapid quenching to room temperature allows a supersaturated solid solution of carbon in iron to be formed.

The α-iron phase field is severely restricted, with a maximum carbon solubility of 0.02 wt% at 723°C (P), so over the carbon range encountered in steels from 0.05 to 1.5 wt%, α-iron is normally associated with iron carbide in one form or another. Similarly, the δ-phase field is very restricted between 1390 and 1534°C and disappears completely when the carbon content reaches 0.5 wt% (B).

There are several temperatures or critical points in the diagram, which are important, both from the basic and from the practical point of view.

* Firstly, there is the A1, temperature at which the eutectoid reaction occurs (P-S-K), which is 723°C in the binary diagram.
* Secondly, there is the A3, temperature when α-iron transforms to γ-iron. For pure iron this occurs at 910°C, but the transformation temperature is progressively lowered along the line GS by the addition of carbon.
* The third point is A4 at which γ-iron transforms to δ-iron, 1390°C in pure iron, hut this is raised as carbon is added. The A2, point is the Curie point when iron changes from the ferro- to the paramagnetic condition. This temperature is 769°C for pure iron, but no change in crystal structure is involved. The A1, A3 and A4 points are easily detected by thermal analysis or dilatometry during cooling or heating cycles, and some hysteresis is observed. Consequently, three values for each point can be obtained. Ac for heating, Ar for cooling and Ae (equilibrium}, but it should be emphasized that the Ac and Ar values will be sensitive to the rates of heating and cooling, as well as to the presence of alloying elements.

The great difference in carbon solubility between γ- and α-iron leads normally to the rejection of carbon as iron carbide at the boundaries of the γ phase field. The transformation of γ to α - iron occurs via a eutectoid reaction, which plays a dominant role in heat treatment.

The eutectoid temperature is 723°C while the eutectoid composition is 0.80% C(s). On cooling alloys containing less than 0,80% C slowly, hypo-eutectoid ferrite is formed from austenite in the range 910-723°C with enrichment of the residual austenite in carbon, until at 723°C the remaining austenite, now containing 0.8% carbon transforms to pearlite, a lamellar mixture of ferrite and iron carbide (cementite). In austenite with 0,80 to 2,06% carbon, on cooling slowly in the temperature interval 1147°C to 723°C, cementite first forms progressively depleting the austenite in carbon, until at 723°C, the austenite contains 0.8% carbon and transforms to pearlite.

Steels with less than about 0.8% carbon are thus hypo-eutectoid alloys with ferrite and pearlite as the prime constituents, the relative volume fractions being determined by the lever rule which states that as the carbon content is increased, the volume percentage of pearlite increases, until it is 100% at the eutectoid composition. Above 0.8% C, cementite becomes the hyper-eutectoid phase, and a similar variation in volume fraction of cementite and pearlite occurs on this side of the eutectoid composition.

The three phases, ferrite, cementite and pearlite are thus the principle constituents of the infrastructure of plain carbon steels, provided they have been subjected to relatively slow cooling rates to avoid the formation of metastable phases.

The austenite- ferrite transformation
Under equilibrium conditions, pro-eutectoid ferrite will form in iron-carbon alloys containing up to 0.8 % carbon. The reaction occurs at 910°C in pure iron, but takes place between 910°C and 723°C in iron-carbon alloys.

However, by quenching from the austenitic state to temperatures below the eutectoid temperature Ae1, ferrite can be formed down to temperatures as low as 600°C. There are pronounced morphological changes as the transformation temperature is lowered, which it should be emphasized apply in general to hypo-and hyper-eutectoid phases, although in each case there will be variations due to the precise crystallography of the phases involved. For example, the same principles apply to the formation of cementite from austenite, but it is not difficult to distinguish ferrite from cementite morphologically.

The austenite-cementite transformation
The Dube classification applies equally well to the various morphologies of cementite formed at progressively lower transformation temperatures. The initial development of grain boundary allotriomorphs is very similar to that of ferrite, and the growth of side plates or Widmanstaten cementite follows the same pattern. The cementite plates are more rigorously crystallographic in form, despite the fact that the orientation relationship with austenite is a more complex one.

As in the case of ferrite, most of the side plates originate from grain boundary allotriomorphs, but in the cementite reaction more side plates nucleate at twin boundaries in austenite.

The austenite-pearlite reaction
Pearlite is probably the most familiar micro structural feature in the whole science of metallography. It was discovered by Sorby over 100 years ago, who correctly assumed it to be a lamellar mixture of iron and iron carbide.

Pearlite is a very common constituent of a wide variety of steels, where it provides a substantial contribution to strength. Lamellar eutectoid structures of this type are widespread in metallurgy, and frequently pearlite is used as a generic term to describe them.

These structures have much in common with the cellular precipitation reactions. Both types of reaction occur by nucleation and growth, and are, therefore, diffusion controlled. Pearlite nuclei occur on austenite grain boundaries, but it is clear that they can also be associated with both pro-eutectoid ferrite and cementite. In commercial steels, pearlite nodules can nucleate on inclusions.

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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.

Figure 1. Classification of iron alloy phase diagrams: a. open γ-field; b. expanded γ-field; c. closed γ-field
(Wever, Archiv, Eisenhüttenwesen, 1928-9, 2, 193)

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 chromium.

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).

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 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 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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