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.

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.

High-strength low-alloy (HSLA) steels, or microalloyed steels, are designed to provide better mechanical properties and/or greater resistance to atmospheric corrosion than conventional carbon steels in the normal sense because they are designed to meet specific mechanical properties rather than a chemical composition.

The HSLA steels have low carbon contents (0.05-0.25% C) in order to produce adequate formability and weldability, and they have manganese contents up to 2.0%. Small quantities of chromium, nickel, molybdenum, copper, nitrogen, vanadium, niobium, titanium and zirconium are used in various combinations.

HSLA Classification:

* Weathering steels, designated to exhibit superior atmospheric corrosion resistance
* Control-rolled steels, hot rolled according to a predetermined rolling schedule, designed to develop a highly deformed austenite structure that will transform to a very fine equiaxed ferrite structure on cooling
* Pearlite-reduced steels, strengthened by very fine-grain ferrite and precipitation hardening but with low carbon content and therefore little or no pearlite in the microstructure
* Microalloyed steels, with very small additions of such elements as niobium, vanadium, and/or titanium for refinement of grain size and/or precipitation hardening
* Acicular ferrite steel, very low carbon steels with sufficient hardenability to transform on cooling to a very fine high-strength acicular ferrite structure rather than the usual polygonal ferrite structure
* Dual-phase steels, processed to a micro-structure of ferrite containing small uniformly distributed regions of high-carbon martensite, resulting in a product with low yield strength and a high rate of work hardening, thus providing a high-strength steel of superior formability.

The various types of HSLA steels may also have small additions of calcium, rare earth elements, or zirconium for sulfide inclusion shape control.
Low-alloy Steels
Low-alloy steels constitute a category of ferrous materials that exhibit mechanical properties superior to plain carbon steels as the result of additions of alloying elements such as nickel, chromium, and molybdenum. Total alloy content can range from 2.07% up to levels just below that of stainless steels, which contain a minimum of 10% Cr.

For many low-alloy steels, the primary function of the alloying elements is to increase hardenability in order to optimize mechanical properties and toughness after heat treatment. In some cases, however, alloy additions are used to reduce environmental degradation under certain specified service conditions.

As with steels in general, low-alloy steels can be classified according to:

* Chemical composition, such as nickel steels, nickel-chromium steels, molybdenum steels, chromium-molybdenum steels
* Heat treatment, such as quenched and tempered, normalized and tempered, annealed.

Because of the wide variety of chemical compositions possible and the fact that some steels are used in more than one heat-treated, condition, some overlap exists among the alloy steel classifications. In this article, four major groups of alloy steels are addressed: (1) low-carbon quenched and tempered (QT) steels, (2) medium-carbon ultrahigh-strength steels, (3) bearing steels, and (4) heat-resistant chromium-molybdenum steels.

Low-carbon quenched and tempered steels combine high yield strength (from 350 to 1035 MPa) and high tensile strength with good notch toughness, ductility, corrosion resistance, or weldability. The various steels have different combinations of these characteristics based on their intended applications. However, a few steels, such as HY-80 and HY-100, are covered by military specifications. The steels listed are used primarily as plate. Some of these steels, as well as other, similar steels, are produced as forgings or castings.

Medium-carbon ultrahigh-strength steels are structural steels with yield strengths that can exceed 1380 MPa. Many of these steels are covered by SAE/AISI designations or are proprietary compositions. Product forms include billet, bar, rod, forgings, sheet, tubing, and welding wire.

Bearing steels used for ball and roller bearing applications are comprised of low carbon (0.10 to 0.20% C) case-hardened steels and high carbon (-1.0% C) through-hardened steels. Many of these steels are covered by SAE/AISI designations.

Chromium-molybdenum heat-resistant steels contain 0.5 to 9% Cr and 0.5 to 1.0% Mo. The carbon content is usually below 0.2%. The chromium provides improved oxidation and corrosion resistance, and the molybdenum increases strength at elevated temperatures. They are generally supplied in the normalized and tempered, quenched and tempered or annealed condition. Chromium-molybdenum steels are widely used in the oil and gas industries and in fossil fuel and nuclear power plants.

The American Iron and Steel Institute (AISI) defines carbon steel as follows:

Steel is considered to be carbon steel when no minimum content is specified or required for chromium, cobalt, columbium [niobium], molybdenum, nickel, titanium, tungsten, vanadium or zirconium, or any other element to be added to obtain a desired alloying effect; when the specified minimum for copper does not exceed 0.40 per cent; or when the maximum content specified for any of the following elements does not exceed the percentages noted: manganese 1.65, silicon 0.60, copper 0.60.

Carbon steel can be classified, according to various deoxidation practices, as rimmed, capped, semi-killed, or killed steel. Deoxidation practice and the steelmaking process will have an effect on the properties of the steel. However, variations in carbon have the greatest effect on mechanical properties, with increasing carbon content leading to increased hardness and strength. As such, carbon steels are generally categorized according to their carbon content. Generally speaking, carbon steels contain up to 2% total alloying elements and can be subdivided into low-carbon steels, medium-carbon steels, high-carbon steels, and ultrahigh-carbon steels; each of these designations is discussed below.

As a group, carbon steels are by far the most frequently used steels. More than 85% of the steel produced and shipped in the United States is carbon steel.

Low-carbon steels contain up to 0.30% C. The largest category of this class of steel is flat-rolled products (sheet or strip), usually in the cold-rolled and annealed condition. The carbon content for these high-formability steels is very low, less than 0.10% C, with up to 0.4% Mn. Typical uses are in automobile body panels, tin plate, and wire products.

For rolled steel structural plates and sections, the carbon content may be increased to approximately 0.30%, with higher manganese content up to 1.5%. These materials may be used for stampings, forgings, seamless tubes, and boiler plate.

Medium-carbon steels are similar to low-carbon steels except that the carbon ranges from 0.30 to 0.60% and the manganese from 0.60 to 1.65%. Increasing the carbon content to approximately 0.5% with an accompanying increase in manganese allows medium carbon steels to be used in the quenched and tempered condition. The uses of medium carbon-manganese steels include shafts, axles, gears, crankshafts, couplings and forgings. Steels in the 0.40 to 0.60% C range are also used for rails, railway wheels and rail axles.

High-carbon steels contain from 0.60 to 1.00% C with manganese contents ranging from 0.30 to 0.90%. High-carbon steels are used for spring materials and high-strength wires.

Ultrahigh-carbon steels are experimental alloys containing 1.25 to 2.0% C. These steels are thermomechanically processed to produce microstructures that consist of ultrafine, equiaxed grains of spherical, discontinuous proeutectoid carbide particles.

Steels can be classified by a variety of different systems depending on:

• The composition, such as carbon, low-alloy or stainless steel.
• The manufacturing methods, such as open hearth, basic oxygen process, or electric furnace methods.
• The finishing method, such as hot rolling or cold rolling
• The product form, such as bar plate, sheet, strip, tubing or structural shape
• The deoxidation practice, such as killed, semi-killed, capped or rimmed steel
• The microstructure, such as ferritic, pearlitic and martensitic
• The required strength level, as specified in ASTM standards
• The heat treatment, such as annealing, quenching and tempering, and thermomechanical processing
• Quality descriptors, such as forging quality and commercial quality.

Stainless steels are iron-based alloys containing at least 10.5% Cr. Few stainless steels contain more than 30% Cr or less than 50% Fe. They achieve their stainless characteristics through the formation of an invisible and adherent chromium-rich oxide surface film. This oxide forms itself in the presence of oxygen.

Other elements added to improve characteristics include nickel, molybdenum, copper, titanium, aluminum, silicon, niobium, nitrogen, sulfur, and selenium. Carbon is normally present in amounts ranging from less than 0.03% to over 1.0% in certain martensitic grades.

The selection of stainless steels may be based on corrosion resistance, fabrication characteristics, availability, mechanical properties in specific temperature ranges and product cost. However, corrosion resistance and mechanical properties are usually the most important factors in selecting a grade for a given application.

Stainless steels are commonly divided into five groups: martensitic stainless steels, ferritic stainless steels, austenitic stainless steels, duplex (ferritic-austenitic) stainless steels, and precipitation-hardening stainless steels.

The development of precipitation-hardenable stainless steels was spearheaded by the successful production of Stainless W by U.S. Steel in 1945. The problem of obtaining raw materials has been a real one, particularly in regard to nickel during 1950s when civil wars raged in Africa and Asia, prime sources of nickel, and Cold War politics played a role because Eastern-bloc nations were also prime sources of the element. This led to the development of a series of alloys (AISI 200 type) in which manganese and nitrogen are partially substituted for nickel. These stainless steels are still produced today.

Over the years, stainless steels have become firmly established as materials for cooking utensils, fasteners, cutlery, flatware, decorative architectural hardware, and equipment for use in chemical plants, dairy and food-processing plants, health and sanitation applications, petroleum and petrochemical plants, textile plants, and the pharmaceutical and transportation industries. Some of these applications involve exposure to either elevated or cryogenic temperatures; austenitic stainless steels are well suited to either type of service.

Modifications in composition are sometimes made to facilitate production. For instance, basic compositions are altered to make it easier to produce stainless steel tubing and casting. Similar modifications are made for the manufacture of stainless steel welding electrodes; here combinations of electrode coating and wire composition are used to produce desired compositions deposited weld metal.

Martensitic stainless steels are essentially alloys of chromium and carbon that possess a distorted body-centered cubic (bcc) crystal structure (martensitic) in the hardened condition. They are ferromagnetic, hardenable by heat treatments, and are generally resistant to corrosion only to relatively mild environments. Chromium content is generally in the range of 10.5 to 18%, and carbon content may exceed 1.2%. The chromium and carbon contents are balanced to ensure a martensitic structure after hardening.

General corrosion is often much less serious than localized forms such as stress corrosion cracking, crevice corrosion in tight spaces or under deposits, pitting attack, and intergranular attack in sensitized material such as weld heat-affected zones (HAZ). Such localized corrosion can cause unexpected and sometimes catastrophic failure while most of the structure remains unaffected, and therefore must be considered carefully in the design and selection of the proper grade of stainless steel.

Corrosive attack can also be increased dramatically by seemingly minor impurities in the medium that may be difficult to anticipate but that can have major effects, even when present in only part-per-million concentrations; by heat transfer through the steel to or from the corrosive medium; by contact trimmed only on the ends.

Stainless steels are available in the form of plate, sheet, strip, foil, bar, wire, semi-finished products, pipes, tubes, and tubing.

Sheet
Sheet is a flat-rolled product in coils or cut lengths at least 610 mm wide and less than 4.76 mm thick. Stainless steel sheet is produced in nearly all types except the free machining and certain martensitic grades. Sheet from the conventional grades is almost exclusively produced on continuous mills. Hand mill production is usually confined to alloys that cannot be produced economically on continuous mills, such as certain high-temperature alloys.

The steel is cast in ingots, and the ingots are rolled on a slabbing mill or a blooming mill into slabs or sheet bars. The slabs or sheet bars are then conditioned prior to being hot rolled on a finishing mill. Alternatively, the steel may be continuous cast directly into slabs that are ready for hot rolling on a finishing mill. The current trend worldwide is toward greater production from continuous cast slabs.

Sheet produced from slabs on continuous rolling mills is coiled directly off the mill. After they are descaled, these hot bands are cold rolled to the required thickness and coils off the cold mill are either annealed and descaled or bright annealed. Belt grinding to remove surface defects is frequently required at hot bands or at an intermediate stage of processing. Full coils or lengths cut from coils may then be lightly cold rolled on either dull or bright rolls to produce the required finish. Sheet may be shipped in coils, or cut sheets may be produced by shearing lengths from a coil and flattening them by roller leveling or stretcher leveling.

Strip
Strip is a flat-rolled product, in coils or cut lengths, less than 610 mm wide and 0.13 to 4.76 mm thick. Cold finished material 0.13 mm thick and less than 610 mm wide fits the definitions of both strip and foil and may be referred to by either term.

Cold-rolled stainless steel strip is manufactured from hot-rolled, annealed, and pickled strip (or from slit sheet) by rolling between polished rolls. Depending on the desired thickness, various numbers of cold rolling passes through the mill are required for effecting the necessary reduction and securing the desired surface characteristics and mechanical properties.

Hot-rolled stainless steel strip is a semi-finished product obtained by hot-rolling slabs or billets and is produced for conversion to finished strip by cold rolling.

Heat Treatment. Strip of all types of stainless steel is usually either annealed or annealed and skin passed, depending on requirements. When severe forming, bending, and drawing operations are involved, it is recommended that such requirements be indicated so that the producer will have all the information necessary to ensure that he supplies the proper type and condition. When stretcher strains are objectionable in ferritic stainless steels such as type 430, they can be minimized by specifying a No 2 finish. Cold-rolled strip in types 410, 414, 416, 420, 431, 440A, 440B, and 440C can be produced in the hardened and tempered condition.

Experience in the use of stainless steels indicates that many factors can affect their corrosion resistance. Some of the more prominent factors are:

* Chemical composition of the corrosive medium including impurities
* Physical state of the medium-liquid, gaseous, solid, or combinations thereof
* Temperature
* Temperature variations
* Aeration of the medium
* Oxygen content of the medium
* Bacteria content of the medium
* Ionization of the medium
* Repeated formation and collapse of bubbles in the medium
* Relative motion of the medium with respect to the steel
* Chemical composition of the metal
* Nature and distribution of microstruc-tural constituents etc.

Surface Finish. Other characteristics in the stainless steel selection checklist are vital for some specialized applications but of little concern for many applications. Among these characteristics, surface finish is important more often than any other except corrosion resistance. Stainless steels are sometimes selected because they are available in a variety of attractive finishes. Surface finish selection may be made on the basis of appearance, frictional characteristics, or sanitation.

Plate
Plate is a flat-rolled or forged product more than 250 mm (10 in.) in width and at least 4.76 mm (0.1875 in.) in thickness. Exceptions include highly alloyed ferritic stainless steels, some of the martensitic stainless steels, and a few of the free-machining grades. Plate is usually produced by hot rolling from slabs that have been directly cast or rolled from ingots and that usually have been conditioned to improve plat surface. Some plate may be produced by direct rolling from ingot.

For strip, edge condition is often more important than it usually is for sheet. Strip can be furnished with various edge specifications:

* Mill edge (as produced, condition unspecified)
* No.1 edge (edge rolled, rounded, or square)
* No.3 edge (as slit)
* No.5 edge (square edge produced by rolling or filing after slitting)

Foil
Foil is a flat-rolled product, in coil form, up to 0.13 mm thick and less than 610 mm wide. Foil is produced in slit widths with edge conditions corresponding to No.3 and No.5 edge conditions for strip. Foil is made from types 201, 202, 301, 302, 304, 304L, 305, 316, 316L, 321, 347, 430, and 442, as well as from certain proprietary alloys.

The finishes, tolerances, and mechanical properties of foil differ from those of strip because of limitations associated with the way in which foil is manufactured. Nomenclature for finishes, and for width and thickness tolerances, vary among producers.

Mechanical Properties. In general, mechanical properties of foil vary with thickness. Tensile strength is increased somewhat, and ductility is lowered, by a decrease in thickness.

Carbon. Most of the microalloyed steels developed for forging have carbon contents ranging from 0.30 to 0.50%, which is high enough to form a large amount of pearlite. The pearlite is responsible for substantial strengthening. This level of carbon also decreases the solubility of the microalloying constituents in austenite.

Niobium, Vanadium, and Titanium. Formation of carbonitride precipitates is the other major strengthening mechanism of microalloyed forging steels. Vanadium, in amounts ranging from 0.05 to 0.2%, is the most common microalloying addition used in forging steels. Niobium and titanium enhance strength and toughness by providing control of austenite grain size. Often niobium is used in combination with vanadium to obtain the benefits of austenite grain size control (from niobium) and carbonitride precipitation (from vanadium).

Manganese is used in relatively large amounts (1.4 to 1.5%) in many microalloyed forging steels. It tends to reduce the cementite plate thickness while maintaining the interlamellar spacing of pearlite developed; thus high manganese levels require lower carbon contents to retain the large amounts of pearlite required for high hardness. Manganese also provides substantial solid solution strengthening, enhances the solubility of vanadium carbonitrides, and lowers the solvus temperature for these phases.

The silicon content of most commercial microalloyed forging steels is about 0.30%; some grades contain up to 0.70%. Higher silicon contents are associated with significantly higher toughness, apparently because of an increased amount of ferrite relative to that formed in ferrite-pearlite steels with lower silicon contents.

Sulfur. Many microalloyed forging steels, particularly those destined for use in automotive forgings in which machinability is critical, have relatively high sulfur contents. The higher sulfur contents contribute to their machinability, which is comparable to that of quenched-and-tempered steels.

Aluminum and Nitrogen. As in hardenable fine-grain steels, aluminum is important for austenite grain size control in microalloyed steels. The mechanism of aluminum grain size control is the formation of aluminum nitride particles. It has been shown that nitrogen is the major interstitial component of vanadium carbonitride. For this reason, moderate to high nitrogen contents are required in vanadium-containing microalloyed steels to promote effective precipitate strengthening.

The driving force behind the development of microalloyed forging steels has been the need to reduce manufacturing costs. This is accomplished in these materials by means of a simplified thermomechanical treatment (that is, a controlled cooling following hot forging) that achieves the desired properties without the separate quenching and tempering treatments required by conventional carbon and alloy steels. Figure 1 shows typical thermal cycles for conventional quench and temper and for microalloy process routes.

Fig. 1 Processing cycles for conventional quenched-and-tempered steels (top) and microalloyed steels (bottom)

Control of Properties

In order to realize the full strengthening potential of microalloying additions, it is necessary to use a soaking temperature prior to forging that is high enough to dissolve all vanadium-bearing precipitates. A soaking temperature above 1100°C (2010°F) is preferred. Rapid induction heating methods for bar and billet to conventional commercial forging temperatures of 1250°C are acceptable and allow sufficient time for the dissolution of the micro alloying constituents.

Tensile strength decreases slightly as the finish forging temperature is reduced, but there is not significant effect on yield strength. Ductility and toughness show a significant increase with a reduction in finishing temperature; this is due to grain refinement of the austenite and increased ferrite content. Forgers are beginning to use this approach to enhance the toughness of as-forged microalloyed steel; however, low finish forging temperatures are often avoided to minimize die wear. The specified properties of microalloyed forging steels can be achieved over a wide range of finishing temperatures.

One of the most important processing factors affecting the properties of as-forged micro alloyed steels is the post forging cooling procedure. Increasing the cooling rate generally increases the yield and tensile strength because it enhances grain refinement and precipitation hardening. At high cooling rates, an optimum can be reached; above this rate the strength reduces due to the suppression of precipitation and the introduction of low-temperature transformation products.

The optimal cooling rate and maximum hardness are significantly influenced by the alloy and residual element content of the steel. Nevertheless, through control of the steel composition it is possible to ensure that the specified mechanical properties are achieved over a wide range of section sizes and cooling conditions.

Effects of Microalloying Elements

The combination of low-carbon and very low alloying content in High-Strength Low-Alloy (HSLA) steels does not affect processing significantly. For example, niobium or niobium-titanium HSLA steels develop high drawability values in interstitial-free steels. Dual-phase steels, which feature hard martensite particles in a soft ductile ferrite matrix are also highly formable. The low-carbon content of HSLA also makes them highly weldable. The machinability of HSLA steels is also comparable to carbon steels. Improved machining characteristics can be achieved by using inclusion-shape-controlled HSLA steels.

The only challenge associated with processing of HSLA steels has been associated with forging. The application of microalloying technology to forging steels has lagged behind that of flat-rolled products because of the different property requirements and thermomechanical processing of forging steels.

Forging steels are commonly used in applications in which high strength, fatigue resistance, and wear resistance are required. These requirements are most often filled by medium-carbon steels. Thus, the development of microalloyed forging steels has centered on grades containing 0.30 to 0.50% C, although steels with carbon contents as low as 0.20% have also been developed.

Bainite frequently occurs in alloy steels during quenching to form martensite. The cooling rate towards the centre of a steel bar is lower than the outside, so in large sections bainite can form in the inner regions with martensite predominating towards the surface.

However, low carbon fully bainitic steels have been developed, as described, using 0.5% Mo and very small concentrations of boron, which allow bainite to form over a wide range of cooling rates. Further control of the reaction is obtained by use of metallic alloying elements such as Ni, Cr, Mn which depress the temperature of maximum rate of formation of bainite. As the transformation temperature is lowered, for a constant cooling rate, the strength of the steel increases substantially. For a series of steels with 0.2% C the tensile strength can be varied between 600 and 1200 MPa. However, this increase in strength is accompanied by a loss of ductility.

The practical advantage of bainitic steels is that relatively high strength levels together with adequate ductility can be obtained without further heat treatment, after the bainite reaction has taken place. The steels are readily weldable, because bainite rather than martensite, will form in the heat-affected zone adjacent to the weld metal, and so the incidence of cracking will be reduced. Furthermore, the steels have a low carbon content, which improves the weldability and reduces stresses arising from transformation.

Reaction Kinetics of Bainite Formation
In plain carbon steels, it is often difficult to separate the bainite reaction from the ferrite and pearlite reactions, because these phases can form under similar continuous to bainitic. For example, the TTT diagram for a 0.8% C steel is a continuous curve although there is both a pearlite and bainite reaction occurring, but it is difficult to disentangle the reactions sufficiently to study their kinetics.

However, the addition of certain alloying elements separates the reactions to the extent that they can be represented as individual curves on the TTT diagram, which then takes on a more complex form than the familiar C-curve.

There are two important features of bainite kinetics which can be shown by a variety of techniques, e.g. dilatometry, electrical resistivity, magnetic measurements and by metallography. First, there is a well defined temperature Bs, above which no bainitic will form, which has been confirmed for a wide range of alloy steels and has been correlated with the structural transition from Widmanstätten ferrite to upper bainitic laths. Second, below Bs temperature and time dependent process take place over a wide temperature range (up to 150°C), which does not go to completion.

The bainitic reaction has several basic features of a nucleation and growth process. It takes place isothermally, starting with an incubation period during which no transformation occurs, followed by an increasing rate of transformation to a maximum and then a gradual slowing down.

Using techniques such as thermionic emission microscopy it has been possible to study directly the progress of the bainite reaction. It has been found that upper bainitic plates lengthen and thicken during transformation by the movement along the plate boundaries of small steps which appear to be diffusion-controlled.

The plates grow at a constant rate edgewise, which leads to a model for the reaction in which the driving force is provided by partition of the carbon from the ferrite to the austenite, the actual growth rate being determined by the rate of diffusion of carbon in austenite away from the γ/α interface.

Role of Alloying Elements
Carbon has a large effect on the range of temperature over which upper and lower bainite occur. The Bs temperature is depressed by many alloying elements but carbon has the greatest influence, as indicated by the following empirical equation:

Bs(°C) = 830 - 270(%C) - 90(%Mn) - 37(%Ni) - 70(%Cr) - 83(%Mo)

Other alloying elements. In plain carbon steels, the bainitic reaction is kinetically shielded by the ferrite and pearlite reactions which commence at higher temperatures and shorter times, so that in continuously cooled samples bainitic structures are difficult to obtain. Even using isothermal transformation difficulties arise if, for example, the ferrite reaction is particularly rapid the addition of metallic alloying elements usually results in retardation of the ferrite and pearlite reactions. In addition, the bainite reaction is depressed to lower temperatures.

This often leads to greater separation of the reactions, and the TTT curves for many alloy steels show much more clearly separate C-curves for the pearlite and bainitic reactions. However, it is still difficult to obtain a fully bainitic structure because of its proximity to the martensite reaction.

A very effective means of isolating the bainite reaction in low carbon steels has been found by adding about 0.002% soluble boron to a 0.5% Mo steel. While the straight molybdenum steel encourages the bainite reaction, the boron markedly retards the ferrite reaction, probably by preferential segregation to the prior austenite boundaries. This permits the bainite reaction to occur at shorter times. Consequently, by use of a range of cooling rates, fully bainitic steels can be obtained.

Lower bainite (temperature range 400-250°C) appears more acicular than upper bainite, with more clearly defined individual plates adopting a lenticular habit. Viewed on a single surface they misleadingly suggest an acicular morphology.

However, two-surface optical microscopy of lower bainite indicates that the ferrite plates are much broader than in upper bainite, and closer in morphology to martensite plates. While these plates nucleate at austenitic grain boundaries, there is also much nucleation within the grains, i.e. intragranular nucleation, and secondary plates form from primary plates away from the grain boundaries.

Electron microscopy shows that the plates have a similar lath substructure to upper bainite, with the ferrite subunits about 0.5 μm wide and slightly disoriented from each other. The plates possess a higher dislocation density than upper bainite, but not as dense as in martensites of similar composition.

The crystallography of the plates seems to depend both on the temperature of transformation, and on the carbon content of the steel. Moreover they showed that the phenomenological theory of martensite could be used for lower bainite to give satisfactory agreement between theory and experiment.

Ohmori and coworkers have found that, in a 0.1% C steel, bainite formed near the Ms has a {011}α habit plane and a <111>α growth direction similar in behavior to low carbon lath martensite. However, on increasing the carbon to 0.6-0.8% C, the habit of the bainitic ferrite plates changes to {122}α//{496}γ, which is not the same as for martensite of the same composition which has a {225}γ habit plane. Because of such variations, it has been suggested that lower bainite is not a true martensitic reaction. However, there is no reason to expect the transformations to be identical, and anyway the inhomogeneous shear would be expected to occur by slip in lower bainite, whereas twinning is the mode adopted in higher carbon martensites.

However, in contrast to tempered martensite the cementite particles in lower bainite exhibit only one variant of the orientation relationship, such that they form parallel arrays at about 60° to the axis of the bainite plate. This feature of the precipitate suggests strongly that it has not precipitated within plates supersaturated with respect to carbon, but that it has nucleated at the γ/α interface and grown as the interface has moved forward. It thus appears that the lower bainite reaction is basically an interface-controlled process leading to cementite precipitation, which then decreases the carbon content of the austenite and enhances the driving force for further transformation.

The morphology of upper bainite (temperature range 550-400°C) bears a close resemblance to Widmanstätten ferrite, as it is composed of long ferrite laths free from internal precipitation.

Two-surface optical micrography decisively reveals that the ferrite component of upper bainite is composed of groups of thin parallel laths with a well-defined crystallographic habit. Like Widmanstätten ferrite, the bainitic ferrite laths exhibit the Kurdjumov-Sachs relationship with the parent austenite, but the relationship is less precise as the transformation temperature is lowered.

A widely-accepted view is that the crystallography of upper bainite is very similar to that of low-carbon lath martensite. However, a detailed examination of the crystallography reveals that there are significant differences, and that upper bainite ferrite formation cannot be understood in terms of the crystallographic theory of martensite-formation.

Electron microscopy shows that upper bainite laths have a fine structure comprising smaller laths about 0.5 μm wide. These laths all possess the same variant of the Kurdjumov-Sachs relationship, so they are only slightly disoriented from each other. The longitudinal boundaries are, therefore, low angle boundaries.

A typical austenite grain will have numerous sheaves of bainitic ferrite exhibiting the several variants of the Kurdjumov-Sachs orientation relationship, so large angle boundaries will occur between sheaves. The dislocation density of the laths increases with decreasing transformation temperature, but even at the highest transformation temperatures the density is greater than that in Widmanstätten ferrite.

The upper bainitic ferrite has a much lower carbon concentration (<0.03% C) than the austenite from which it forms, consequently as the bainitic laths grow, the remaining austenite is enriched in carbon. This is an essential feature of upper bainite which forms in the range 550-400°C when the diffusivity of carbon is still high enough to allow partition between ferrite and austenite. Consequently, carbide precipitation does not occur within the laths, but in the austenite at the lath boundaries when a critical carbon concentration is reached.

The morphology of the cementite formed at the lath boundaries is dependent on the carbon content of the steel. In low carbon steels, the carbide will be present as discontinuous stringers and isolated particles along the lath boundaries, while at higher carbon levels the stringers may become continuous. With some steels, the enriched austenite does not precipitate carbide, but remains as a film of retained austenite. Alternatively, on cooling it may transform to high carbon martensite with an adverse effect on the ductility. This type of bainite is often referred to as granular bainite.

The Bainite Reaction

Examination of the TTT diagram for a eutectoid carbon steel, Fig. 1, bearing in mind the fact that the pearlite reaction is essentially a high temperature one occurring between 550°C and 720°C and that the formation of martensite is a low temperature reaction, reveals that there is a wide range temperature, usually 250-550°C, when neither of these phases forms.

Fig.1: Time-Temperature-Transformation (TTT) diagram for a 0.89 carbon steel
(US Steel Co., Atlas of Isothermal Diagrams)

This is the region in which lath-shaped fine aggregates of ferrite and cementite are formed, which possess some of the properties of the high temperature reactions involving ferrite and pearlite as well as some of the characteristics of the martensite reaction.

The generic term for these intermediate structures is bainite after Edgar Bain who with Davenport first found them during their pioneer systematic studies of the isothermal decomposition of austenite. Bainite also occurs during thermal treatments at cooling rates too fast for pearlite to form, yet not rapid enough to produce martensite. The nature of bainite changes as the transformation temperature is lowered. Two main forms can be identified, upper and lower.

The hot-rolling process has gradually become a much more closely controlled operation, and is being increasingly applied to low alloy steels with compositions carefully chosen to provide optimum mechanical properties when the hot deformation is complete. This process, in which the various stages of rolling are temperature-controlled, the amount of reduction in each pass is predetermined and the finishing temperature is precisely defined, is called controlled rolling and is now of greatest importance in obtaining reliable mechanical properties in steels for pipelines, bridges, and many other engineering applications.

On the other hand, in more highly alloy steels, it is possible to subject the steels to heavy deformations in the metastable austenitic conditions prior to transformation to martensite. This process, ausforming, allows the attainment of very high strength levels combined with good toughness and ductility.

Before World War II, strength in hot-rolled low alloy steels was achieved by the addition of carbon up to 0.4% and manganese up to 1.5%, giving yield stresses of 350-400 MPa. However, such steels are essentially ferrite pearlite aggregates, which do not possess adequate toughness for many modern applications. Indeed, the toughness, as measured by the ductile/brittle transition decreases dramatically with carbon content, i.e. with increasing volume of pearlite in the steel. Furthermore, with the introduction of welding as the main fabrication technique, the high carbon contents led to serious cracking problems, which could only be eliminated by the use of lower carbon steels. The great advantage of producing in these steels a fine ferrite grain size soon became apparent, so controlled rolling in the austenitic condition was gradually introduced to achieve this.

Fine ferrite grain sizes in the finished steel were found to be greatly expedited by the addition of small concentrations (< 0.1 wt %) of grain refining elements such as niobium, titanium, vanadium and aluminum. On adding such elements to steels with 0.03-0.08% C and up to 1.5% Mn, it became possible to produce fine-grained material with yield strengths between 450 and 550 MPa, and with ductile/brittle transition temperatures as low as -70°C. Such steels are now referred to as high strength low alloy steels (HSLA), or micro-alloyed steels. This progress, from the relatively low strength of ordinary mild steel (220-250 MPa) in a period of twenty years represents a major metallurgical development, the importance of which, in engineering applications, cannot be overstated.

The primary grain refinement mechanism in controlled rolling is the recrystallization of austenitic during hot deformation, known as dynamic recrystallization. This process is clearly influenced by the temperature and the degree of deformation, which takes place during each pass through the rolls. However, in austenite devoid of second phase particles, the high temperatures involved in hot rolling lead to marked grain growth, with the result that grain refinement during subsequent working is limited.

Clearly, the control of grain size at high austenitizing temperatures requires as fine a grain boundary precipitate as possible, and one which will not dissolve completely in the austenite, even at the highest working temperatures (1200-1300°C). The best grain refining elements are very strong carbide and nitride formers, such as niobium, titanium and vanadium, also aluminum which forms only a nitride. As both carbon and nitrogen are present in control-rolled steels, and as the nitrides are even more stable than the carbides it is likely that the most effective grain refining compounds are the respective carbo-nitrides, except in the case of aluminum nitride. As a result of the combined use of controlled rolling and fine dispersions of carbo-nitrides in low alloy steels, it has been possible to obtain ferrite grain sizes between 5 and 10 μm in commercial practice.

The solubility data implies that, in micro-alloyed steel, carbides and carbo-nitrides of Nb, Ti and V will precipitate progressively during controlled rolling as the temperature falls. While the primary effect of these fine dispersions is to control grain size, dispersion strengthening will take place. The strengthening arising from this cause will depend both on the particle size and the interparticle spacing which is determined by the volume fraction of precipitate. These parameters will depend primarily on the type of compound which is precipitating, and that is determined by the micro alloying content of the steel. However, the maximum solution temperature reached and the detailed schedule of the controlled rolling operation are also important variables.

It is now known, not only that precipitation takes place in the austenite, but that further precipitation occurs during the transformation to ferrite. The precipitation of niobium, titanium and vanadium carbides has been shown to take place progressively as the interphase boundaries move through the steel As this precipitation is normally on an extremely line scale occurring between 850 and 650°C, it is likely to be the major contribution to the dispersion strengthening. In view of the higher solubility of vanadium carbide in austenite, the effect will be most pronounced in the presence of this element, with titanium and niobium in decreasing order of effectiveness. If the rate of cooling through the transformation is high, leading to the formation of supersaturated acicular ferrite, the carbides will tend to precipitate within the grains, usually on the dislocations, which are numerous in this type of ferrite.

In arriving at optimum compositions of micro-alloyed steels, it should be noted that the maximum volume fraction of precipitate, which can be put into solid solution in austenite at high temperatures, is achieved by use of stoichiometric compositions.

In modern control-rolled micro-alloyed steels, there are at least three strengthening mechanisms, which contribute to the final strength achieved. The relative contribution from each is determined by the composition of the steel and, equally important, the details of the thermomechanical treatment to which the steel is subjected. Firstly, there are the solid solution strengthening increments from manganese, silicon and uncombined nitrogen. Secondly, the grain size contribution to the yield stress is shown as a very substantial component, the magnitude of which is very sensitive to the detailed thermomechanical history. Finally, a typical increment is dispersion strengthening. The total result is a range of yield strengths between about 350 and 500 MPa. In this particular example, the steel was normalized (air cooled) from 900°C, but had it been control rolled down to 800°C or even lower, the strength levels would have been substantially raised.

The effect of the finishing temperature for rolling is important in determining the grain size and, therefore, strength level reached for particular steel. It is now becoming common to roll through the transformation into the completely ferritic condition, and so obtain fine subgrain structures in the ferrite, which provide an additional contribution to strength. Alternatively, the rolling is finished above the γ/α transformation, and the nature of the transformation is altered by increasing the cooling rate. Slow rates of cooling obtained by coiling at a particular temperature will give lower strengths than rapid rates imposed by water spray cooling following rolling.

Dual phase steels are referred to as dual phase low alloy (DPLA) steels. They exhibit continuous yielding, i.e. no sharp yield point, and a relatively low yield stress (300-350 MPa) together with a rapid rate of work hardening and high elongations (≈ 30 %) which gives excellent formability. As a result of the work hardening, the yield stress in the final formed product is as high as in HSLA steels (500-700 MPa). The simplest steels in this category contain 0.08-0.2 % C, 0.5-1.5 % Mn, but steels micro-alloyed with vanadium are also suitable, while small additions of Cr (0.5 %) and Mo (0.2-0.4%) are frequently used.

The simplest way of achieving a duplex structure is to use intercritical annealing in which the steel is heated to the (α + γ) region between Ac1 and Ac3 and held, typically, at 790°C for several minutes to allow small regions of austenite to form in the ferrite. As it is essential to transform these regions to martensite, recooling must be rapid or the austenite must have a high hardenability. This can be achieved by adding between 0.2 and 0.4% molybdenum to the steel already containing 1.5 % manganese. The required structure can then be obtained by air-cooling after annealing.

To eliminate an extra heat treatment step, dual phase steels have now been developed. These steels can be given the required structure during cooling after controlled rolling. Typically, these steels have additions of 0.5% Cr and 0.4% Mo. After completion of hot rolling around 870°C, the steel forms approximately 80% ferrite on the water-cooled run-out table from the mill. The material is then cooled in the metastable region (510-620°C) below the pearlite/ferrite transformation and, on subsequent cooling, the austenite regions transform to martensite.

Micro-alloyed steels produced by controlled rolling are a most attractive proposition in many engineering applications because of their relatively low cost, moderate strength, and very good toughness and fatigue strength, together with their ability to be readily welded. They have, to a considerable degree, eliminated quenched and tempered steels in many applications.

These steels are most frequently available in control rolled sheet, which is then cooled over a range of temperatures between 750 and 550°C. The cooling temperature has an important influence as it represents the final transformation temperature, and this influences the microstructure. The lower this temperature, under the same conditions, the higher the strength achieved.

The normal range of yield strength obtained in these steels varies from about 350 to 550 MPa. The strength is controlled both by the detailed thermomechanical treatment, by varying the manganese content from 0.5 to 1.5 wt%, and by using the microalloying additions in the range 0.03 to above 0.1 wt%. Niobium is used alone, or with vanadium, while titanium can be used in combination with the other two carbide-forming elements. The interactions between these elements are complex, but in general terms niobium precipitates more readily in austenite than does vanadium as carbide or carbo-nitride, so it is relatively more effective as a grain refiner. The greater solubility of vanadium carbide in austenite underlines the superior dispersion strengthening potential of this element shared to a lesser degree with titanium. Titanium also interacts with sulphur and can have a beneficial effect on the shape of sulphide inclusions. Bearing in mind that the total effect of these elements used in conjunction is not a simple sum of their individual influence, the detailed metallurgy of these steels becomes extremely complex.

One of the most extensive applications is in pipelines for the conveyance of natural gas and oil, where the improved weldability due to the overall lower alloying content (lower hardenability) and, particularly, the lower carbon levels is a great advantage.

When the brave and experienced fighting Christian knights trekked the long route to the holy lands to fight the Muslims they were in for a terrible surprise. They were beaten and beaten hard and repeatedly and eventually sent packing back to England and Europe defeated.

What they were not prepared for was the secret weapon of the enemy they came to easily slaughter in the name of God. The Muslim fighters had a superior weapon and that was Damascus steel.

Their swords and daggers were made from this vastly superior steel. It is said, and probably a fanciful bit of exaggeration, that a Damascus sword could slice through a silk scarf falling through the air. True or not, the Damascus sword was amazingly sharp but also much stronger than normal steel.

So the Muslims had the capability to cut a strand of hair off or hack through the armor of the crusader. It is frustrating to fight sword against sword when your opponent’s weapon can break yours in half.

This technology is used in some special jewellery called Damascus Jewellery today. But it is not the strength of the metal that is sought but the look of it. Damascus steel is characterized by the pattern on the metal. “Damas” in Arabic means water and it is sometimes thought that the term Damascus steel originates from this and the pattern on the metal which resembles various forms of windswept, rippling water.

So today we have some stunning jewellery in the form of rings, bangles, chokers and bracelets called Damascus Jewellery. It has this characteristic pattern, mostly a swirling, twisting, overlapping effect.

Ore was smelted in Hyberdad in southern India and put into crucibles and made into solid pieces of metal called billets. These billets were taken to Damascus in Syria and forged into swords and daggers. The metal was folded and melted together, folded again, and this process was continued until it was finally melted and hammered into a sword with great strength and beauty.

For a couple of hundred years the metallurgy technique was lost to the world and but the sword smith techniques have continued so today we can get a Damascus ring made from two different metals hammered and forged together with the original looking patterns of the Damascus steel. In a sense it is wrong to call it Damascus as it is only part of the original process but it doesn’t need the original sharpness or strength either.

Unlike some other kinds of jewellery, bracelets are incredibly diverse. You can make a bracelet from almost any material you can think of, whether it’s metal, fabric, leather, glass, wire or plastic, or natural materials like shells and stones. Bracelets are both cheap to buy and easy to make, so it’s easy to personalise them to your taste and give them as gifts to your friends. Many people cherish bracelets that they have been given, especially little girls, who will collect friendship bracelets to represent each one of their friends.

If you want to make a bracelet yourself, the easiest way to do it is probably to get some beads and a piece of string. Threads woven together can also work well, as can shells if you make tiny holes in them to thread a string through. The best thing to do is to keep the bracelet as a flat string, and then tie it around the hand of the person you’re giving it to so that it fits well.

Another interesting use of bracelets, only invented quite recently, is the charity bracelet. These are usually made from silicone, and bear a slogan such as ‘make poverty history’ or ‘help the aged’. The idea is that if you see someone wearing one, you are likely to mention it, and then they will tell you a little about the cause and why it matters to them – a good way of getting people to talk about charities they care about.

In some parts of the world, bracelets have special meanings. In India, for example, they are often made from glass, and make a musical-sounding noise when the arm is moved. In Latin America, it is believed that bracelets made from gold and coral can scare away evil spirits, protecting the wearer from the jealousy of others.

Although pure iron is a weak material, steels cover a wide range of the strength spectrum from low yield stress levels (around 200 MPa) to very high levels (approaching 2000 MPa). These mechanical properties are usually achieved by the combined use of several strengthening mechanisms, and in such circumstances it is often difficult to quantify the different contributions to the strength. These results should then be helpful in examining the behavior of more complex steels.

Like other metals, iron can be strengthened by several basic mechanisms, the most important of which are:

* Work hardening
* Solid solution strengthening by interstitial atoms
* Solid solution strengthening by substitutional atoms
* Refinement of grain size
* Dispersion strengthening, including lamellar and random dispersed structures.

The most distinctive aspect of strengthening of iron 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.

Work hardening
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 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.

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`s, 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.

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 interparticle spacing and, in the second place, to the reverse ������ martensitic transformation which is accompanied by the dissolution of intermetallics 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 interparticle spacing of an order of 103 nm in the strong matrix, these intermetallic precipitates also possessing a high strength. Thus, with the same dispersity 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).

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 dispersity 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.

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:

* Ultra-high strength at room temperature
* Simple heat treatment, which results in minimum distortion
* 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
* 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.

Tempering of maraging steels
Tempering as an operation of heat treatment has been well known from the Middle Ages. It is used with martensite-quenched alloys. The processes of tempering will be considered here for steels only, sinse steels constitute an overwhelming majority of all marensite-hardenable alloys.

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 480-500��C. The tempering results in strong precipitation hardening owing to the precipitation of intermetallics 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. Ni3Ti phase is similar to hexagonal ��-carbide in carbon steels. Of special practical value is the fact that particles of intermediate intermetallics 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.

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.

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 (480-500��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 interparticle spacing, with the dislocation density being simultaneously reduced. With a long holding time, semi coherent precipitates of intermediate intermetallics 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 intermetallics that have precipitated from the ��-phase.