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1. Classification of Steels

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1.1 Classification of Stainless Steels Abstract: 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. Stainless steels are available in the form of plate, sheet, strip, foil, bar, wire, semifinished products, pipes, tubes, and tubing.

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

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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, semifinished 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

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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 freemachining 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:

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• • • •

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, varies 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.

Bar Bar is a product supplied in straight lengths; it is either hot or cold finished and is available in various shapes, sizes, and surface finishes. This category includes small shapes whose dimensions do not exceed 75 mm and, second, hot-rolled flat stock at least 3.2 mm thick and up to 250 mm wide. Hot-finished bar is commonly produced by hot rolling, forging, or pressing ingots to blooms or billets of intermediate size, which are subsequently hot rolled, forged, or extruded to final dimensions.

1.2. Classification of Carbon and Low-Alloy Steels Abstract: 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.

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

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• • • •

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.

Carbon Steels 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

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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 highstrength 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.

High-Strength Low-Alloy Steels 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

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• • • •

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 chromiummolybdenum 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.

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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) throughhardened 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. Chromiummolybdenum steels are widely used in the oil and gas industries and in fossil fuel and nuclear power plants.

1.3. The Effects of Alloying Elements on Iron-Carbon Alloys Abstract: 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: open and closed γ-field systems, and expanded and contracted γ-field systems. 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.

The simplest version of analyzes the effects of alloying elements on iron-carbon alloys would require analysis of a large number of ternary alloy diagrams over a wide temperature range. However, Wever pointed out that iron binary equilibrium systems fall into four main categories (Fig. 1): open and closed γ-field systems, and expanded and contracted γ-field systems. This approach indicates that alloying elements can influence the equilibrium diagram in two ways: • •

by expanding the γ-field, and encouraging the formation of austenite over wider compositional limits. These elements are called γ-stabilizers. by contracting the γ-field, and encouraging the formation of ferrite over wider compositional limits. These elements are called α-stabilizers.

The form of the diagram depends to some degree on the electronic structure of the alloying elements which is reflected in their relative positions in the periodic classification. Class 1: open γ-field. To this group belong the important steel alloying elements nickel and manganese, as well as cobalt and the inert metals ruthenium, rhodium, palladium, osmium, iridium and platinum. Both nickel and manganese, if added in sufficiently high concentration, completely eliminate the bcc α-iron phase and replace it, down to room temperature, with the γ-phase. So nickel and manganese depress the phase transformation from γ to α to lower temperatures (Fig. 1a), i.e. both Ac1

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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.

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

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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 γloop 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

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• •

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 ironcarbon 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

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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 ironcarbon 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.

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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.

1.4. The Iron-Carbon Equilibrium Diagram Abstract: 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 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.

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. 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.

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Fig. 1. The iron-carbon diagram.

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.

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• • •

Firstly, there is the A1, temperature at which the eutectoid reaction occurs (PS-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

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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.

1.5. Iron and Its Interstitial Solid Solutions Abstract: 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). 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 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-

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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

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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 nearestneighbor 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 εcarbide/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

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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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2. Iron and Carbon Steels

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2.1. Designation of Carbon and Low-Alloy Steels Abstract: 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. This article describes basics of SAE, AISI, UNS, AMS, European and Japanese designation systems.

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 coldfinished 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.

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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.

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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,

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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.

2.2. Cast steel: Microstructure and grain size Abstract: The equilibrium diagram does not tell us what form is taken by the ferrite or cementite ejected from the austenite on cooling. Without going too deeply into the matter, it may be considered that the ferrite has a choice of three different positions, which, in order of degree of supercooling or ease of forming nuclei, are: (1) boundaries of the austenite crystals; (2) certain crystal planes (octahedral); (3) about inclusions

The equilibrium diagram does not tell us what form is taken by the ferrite or cementite ejected from the austenite on cooling. Without going too deeply into the matter, it may be considered that the ferrite has a choice of three different positions, which, in order of degree of supercooling or ease of forming nuclei, are: (1) boundaries of the austenite crystals; (Fig. 1) (2) certain crystal planes (octahedral); (Fig. 2) (3) about inclusions (Fig. 3).

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Figure 1.

Figure 2.

Figure 3. Thus, ferrite starts to form at the grain boundaries, and if sufficient time is allowed for the diffusion phenomena a ferrite network structure is formed, while the pearlite occupies the centre, as in Fig. 1. The size of the austenite grains existing above A3 is thereby betrayed. If the rate of cooling is faster, the complete separation of the ferrite at the boundaries of large austenite grains is not possible, and ejection takes place within the crystal along certain planes, forming a mesh-like arrangement known as a Widmanstätten structure, shown in Fig. 2. In steels containing more than 0,9% carbon, cementite can separate in a similar way and Widmanstätten structures are also found in other alloy systems. Steel with Widmanstätten structures are characterised by (1) low impact value, (2) low percentage elongation since the strong pearlite is isolated in ineffective patches by either weak ferrite or brittle cementite, along which cracks can be readily propagated. This structure is found in overheated steels and cast steel, but the high silicon used in steel castings modifies. It is highly desirable that Widmanstätten and coarse network structures generally be avoided, and as these partly depend upon the size of the original austenite grain, the methods of securing small grains are of importance. Large austenite grains may be refined by (a) hot working, (b) normalising.

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Such refined austenite grains are liable to coarsen when the steel is heated well above the Ac3 temperature, in such operations as welding, forging and carburising unless the grain growth is restrained. This restraint can be brought about by a suitable mode of manufacture of the steel. Controlled grain size It is now possible to produce two steels of practically identical analysis with inherently different grain growth characteristics so that at a given temperature each steel has an "inherent austenite grain size", one being fine relative to the other. The so-called "fine-graine" steel increases its size on heating above Ac3 but the temperature at which the grain size becomes relatively coarse is definitely higher than that at which a "coarse-grained" steel develops a similar size. The fine-grained steels are "killed" with silicon together with a slight excess of aluminium which forms aluminium nitride as submicroscopic particles that obstruct austenite grain growth and is an example of a general phenomenon. At the coarsening temperature the AIN goes into solution rapidly above 1200°C and virtually completely at 1350°C. The austenite grain size is frequently estimated by the following tests: (1) McQuaid-Ehn Test. Micro-sections of structural steels carburised for not less than 8 hours at 925°C and slowly cooled to show cementite networks are photographed at a magnification of 100. Comparison is made with a grain-size chart issued by the American Society for Testing Materials. This test is also valuable in detecting "abnormality" of pearlite. (2) The Quench and Fracture test consists in heating normalised sections of the steel, above Ac3 quenching them at intervals of 30°C. An examination of the fractured surface enables the depth of hardness and grain size to be estimated by comparison with standard fractures.

2.3. Steel-making processes Abstract: Steel is made by the Bessemer, Siemens Open Hearth, basic oxygen furnace, electric arc, electric highfrequency and crucible processes. 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 1525 tonnes capacity. In the Siemens process, both acid and basic, the necessary heat for melting and working the charge is supplied by oil or gas. Both 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 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.

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

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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 watercooled 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

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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

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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.

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Figure 2.

Figure 3. Methods of degassing molten steel 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

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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. 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 macrosegregation, 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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2.4. Structure of plain steel Abstract: The essential difference between ordinary steel and pure iron is the amount of carbon in the former, which reduces the ductility but increases the strength and the susceptibility to hardening when rapidly cooled from elevated temperatures. On account of the various micro-structures which may be obtained by different heat-treatments, it is necessary to emphasise the fact that the following structures are for "normal" steels, i.e. slowly cooled from 760-900°C depending on the carbon contents.

The essential difference between ordinary steel and pure iron is the amount of carbon in the former, which reduces the ductility but increases the strength and the susceptibility to hardening when rapidly cooled from elevated temperatures. On account of the various micro-structures which may be obtained by different heattreatments, it is necessary to emphasise the fact that the following structures are for "normal" steels, i.e. slowly cooled from 760-900°C depending on the carbon contents. The appearance of pure iron is illustrated in Fig. 1. It is only pure in the sense that it contains no carbon, but contains very small quantities of impurities such as phosphorus, silicon, manganese, oxygen, nitrogen, dissolved in the solid metal. In other words, the structure is typical of pure metals and solid solutions in the annealed condition. It is built up of a number of crystals of the same composition, given the name ferrite in metallography (Brinell hardness 80). The addition of carbon to the pure iron results in a considerable difference in the structure (Fig. 2), which now consists of two constituents, the white one being the ferrite, and the dark parts representing the constituent containing the carbon, the amount of which is therefore an index of the quantity of carbon in the steel. Carbon is present as a compound of iron and carbon (6-67 %) called cementite, having the chemical formula Fe3 C. This cementite is hard (Brinell hardness 600 +), brittle and brilliantly white.

x200

x200 Figure 1. Armco iron: ferrite grains

Figure 2. 0,4% carbon steel. Ferrite + pearlite

On examination the dark parts will be seen to consist of two components occurring as wavy or parallel plates alternately dark and light (Fig. 3). These two phases are

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ferrite and cementite which form a eutectic mixture, containing 0,87% carbon and known as pearlite. The appearance of this pearlite depends largely upon the objective employed in the examination and also on the rate of cooling from the elevated temperature.

Figure 3. 0,87% carbon steel Allotropy of iron Certain substances can exist in two or more crystalline forms; for example charcoal, graphite and diamonds are allotropic modifications of carbon. Allotropy is characterized by a change in atomic structure which occurs at a definite transformation temperature. Four changes occur in iron, which give rise to forms known as alpha, beta, gamma and delta. Of these, α, β and δ forms have the same atomic structure (body centred cubic) while γ -iron has a face centred cubic structure. Iron can, therefore, be considered to have two allotropic modifications. The A2 change at 769°C, at which the α-iron loses its magnetism, can be ignored from a heat-treatment point of view. These changes in structure are accompanied by thermal changes, together with discontinuities in other physical properties such as electrical, thermo-electric potential, magnetic, expansion and tenacity. The A3 change from a b.c.c. to an f.c.c. atomic structure at 937°C is accompanied by a marked contraction while the reverse occurs at 1400°C. These changes in structure are accompanied by recrystallisation, followed by grain growth. Critical points The addition of carbon to iron, however, produces another change at 695°C, known as A1 and associated with the formation of pearlite. These structural changes, which occur during cooling, give rise to evolutions of heat, which cause arrests on a cooling curve. The temperatures of these arrests are known as critical points or "A" points. These arrests occur at slightly higher temperatures on heating, as compared with cooling, and this lag effect, increased by rapid cooling, is known as thermal hysteresis.

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To differentiate between the arrests obtained during heating and cooling, the letters c and r respectively are added to the symbol A (from chauffage and refroidissement). In a steel containing about 0,8-0,9% carbon the evolution of the heat at Ar1 is sufficient to cause the material to become visibly hotter and the phenomenon is called "recalescence". Iron-cementite equilibrium diagram The addition of carbon to iron not only gives rise to the A1 point but also influences the critical points in pure iron. The A4 point is raised; and the A3 point lowered until it coincides with A1. The α, β and δ modifications, which may be called ferrite, have only slight solubility for carbon, but up to 1,7% of carbon dissolves in y-iron to form a solid solution called Austenite. These effects are summarised in the iron-Fe3 C equilibrium diagram (Fig. 4), which is of much importance in the study of steels. The iron-iron carbide system is not in true equilibrium, the stable system is irongraphite, but special conditions are necessary to nucleate graphite. Will be seen that the complicated Fe-Fe3C diagram can be divided into several simple diagrams: Peritectic transformation CDB - δ-iron transforms to austenite. Eutectic at E austenite and cementite. Solid solution D to F - primary dendrites of austenite form. Eutectic point at P - formation of pearlite.

Figure 4. Iron-cementite equilibrium diagram

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The ferrite solubility line, A3P, denotes the commencement of precipitation of ferrite from austenite. The cementite solubility line, FP, indicates the primary deposition of cementite from austenite. The pearlite line, A1PG, indicates the formation of the eutectic at a constant temperature. Let us consider the freezing of alloys of various carbon contents. 0,3% carbon Dendrites of δ-iron form, the composition of which is represented eventually by C (0,07 %), and the liquid, enriched in carbon, by B. The solid crystals then react with the liquid to form austenite of composition D. Diffusion of carbon occurs as the solid alloy cools to line A3P. Here α-ferrite commences to be ejected from the austenite, consequently the remaining solid solution is enriched in carbon, until point P is reached at which cementite can be also precipitated. The alternate formation of ferrite and cementite at 695°C gives rise to pearlite. The structure finally consists of masses of pearlite embedded in the ferrite. 0,6% carbon When line BE is reached dendrites of austenite form, and finally the alloy completely freezes as a cored solid solution, which, on cooling through the critical range (750695°C), decomposes into ferrite and pearlite. 1,4% carbon Again, the alloy solidifies as a cored solid solution, but on reaching line FP, cementite starts to be ejected and the residual alloy becomes increasingly poorer in carbon until point P is reached, when both cementite and ferrite form in juxtaposition. The structure now consists of free cementite and pearlite.

2.5. Corrosion of Carbon Steel Abstract: Carbon steel, the most widely used engineering material, accounts for approximately 85%, of the annual steel production worldwide. Despite its relatively limited corrosion resistance, carbon steel is used in large tonnages in marine applications, nuclear power and fossil fuel power plants, transportation, chemical processing, petroleum production and refining, pipelines, mining, construction and metal-processing equipment.

Carbon steel, the most widely used engineering material, accounts for approximately 85%, of the annual steel production worldwide. Despite its relatively limited corrosion resistance, carbon steel is used in large tonnages in marine applications, nuclear power and fossil fuel power plants, transportation, chemical processing, petroleum production and refining, pipelines, mining, construction and metalprocessing equipment. The cost of metallic corrosion to the total economy must be measured in hundreds of millions of dollars (or euros) per year. Because carbon steels represent the largest single class of alloys in use, both in terms of tonnage and total cost, it is easy to understand that the corrosion of carbon steels is a problem of enormous practical

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importance. This is the reason for the existence of entire industries devoted to providing protective systems for irons and steel. Carbon steels are by their nature of limited alloy content, usually less than 2% by weight for total of additions. Unfortunately, these levels of addition do not generally produce any remarkable changes in general corrosion behavior. One possible exception to this statement would be weathering steels, in small additions of copper, chromium, nickel and phosphorus produce significant reduction in corrosion rate in certain environments. Because corrosion is such a multifaceted phenomenon, it is generally useful to attempt to categorize the various types. This is usually done on environmental basis. In this article, atmospheric corrosion, aqueous corrosion and some other corrosion types of interest, such as corrosion in soils, concrete and boilers and heating plants will be addressed.

Atmospheric corrosion Atmospheres are often classified as being rural, industrial or marine in nature. Two decidedly rural environments can differ widely in average yearly temperature and rainfall patterns, mean temperature, and perhaps acid rain, can make extrapolations from past behavior less reliable. The corrosion of carbon steel in the atmosphere and in many aqueous environments is best understood from a film formation and brake down standpoint. It is an inescapable fact that iron in the presence of oxygen and water is thermodynamically unstable with respect to its oxides. Because atmospheric corrosion is an electrolytic process, the presence of an electrolyte is required. This should not be taken to mean that the steel surface must be awash in water; a very thin adsorbed film of water is all that is required. During the actual exposure, the metal spends some portion of the time awash with water because of rain or splashing and a portion of the time covered with a thin adsorbed water film. The portion of time spent covered with the thin water film depends quite strongly on relative humidity at the exposure site. This fact has led many corrosion scientists to investigate the influence of the time of wetness on the corrosion rate. Rusting of iron depends on relative humidity and time of exposure in atmosphere containing 0.01% SO2. The increase in corrosion rate produced by the addition of SO2 is substantial. Oxides of nitrogen in the atmosphere would also exhibit an accelerating effect on the corrosion of steel. Indeed, any gaseous atmospheric constituent capable of strong electrolytic activity should be suspected as being capable of increasing the corrosion rate of steel. Because carbon steels are not very highly alloyed, it is not surprising that most grades do not exhibit large differences in atmospheric-corrosion rate. Nevertheless, alloying can make changes in the atmospheric-corrosion rate of carbon steel. The elements generally found to be most beneficial in this regard are copper, nickel, silicon, chromium and phosphorus. Of these, the most striking example is that of copper, increases from 0.01-0.05%, decrease the corrosion rate by a factor of two to

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three. Additions of the above elements in combination are generally more effective than when added singly, although the effects are not additive.

Aqueous Corrosion Carbon steel pipes and vessels are often required to transport water or are submerged in water to some extent during service. This exposure can be under conditions varying temperature, flow rate, pH, and other factors, all of which can alter the rate of corrosion. The relative acidity of the solution is probably the most important factor to be considered. At low pH the evolution of hydrogen tends to eliminate the possibility of protective film formation so that steel continues to corrode but in alkaline solutions, the formation of protective films greatly reduces the corrosion rate. The greater alkalinity, the slower the rate of attack becomes. In neutral solutions, other factors such as aeration, became determining so that generalization becomes more difficult. The corrosion of steels in aerated seawater is about the same overall as in aerated freshwater, but this is somewhat misleading because the improved electrical conductivity of seawater can lead to increased pitting. The concentration cells can operate over long distance, and this leads to a more nonuniform attack than in fresh water. Alternate cycling through immersion and exposure to air produces more pitting attack than continuous immersion. The effect of various alloying addition and exposure conditions on the corrosion behavior is shown in Table 1. Table 1. Comparison of results under different type of exposure Effects of alloy selection, chemical Sea air composition and alloy additions

Freshwater

Alternately Continuously wet with wet with seawater or seawater Spray and dry

Ferrous alloys

Pockmarked

Vermiform on cleaned bars

Pitting, Pitting, particularly on particularly on bars with scale bars with scale

Wrought iron versus carbon steel

Steel superior to Iron and steel wrought and equal in lowingot irons moor areas

Low-moor iron Low-moor iron superior to superior to carbon steel carbon steel

Sulfur and phosphorus content

Best results when S and P are low

Best results when S and P are low

Apparently little influence

Addition of copper

Beneficial: Beneficial: 0.635% Cu Effect increasing almost as with copper good as content 2.185% Cu

Beneficial: 0.635% and 2.185% Cu much the same

0.635% Cu slightly beneficial: 2.185% Cu somewhat less so

Addition of nickel

3.75% Ni 3.75%Ni superior even to superior even 2% Cu; 36% Ni to 2%Cu;

3.75%Ni beneficial usually more

3.75% Ni slightly beneficial and slightly superior

Best results when S and P are low

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Addition of 13.5% Cr

almost perfect after 15-year exposure

36%Ni excellent resistance

Excellent resistance to corrosion: cold blast metal perfect after 15year exposure: equal to 36% Ni steel

Excellent resistance to corrosion: equal to 36% Ni steel

Excellent resistance to corrosion: cold Undergoes Behavior of cast graphitic blast metal irons superior to hot: corrosion no graphitic corrosion

to Cu: 36% Ni so than Cu: 36%Ni the best the best metal in metal in the the set set Subject to severe localized corrosion that virtually destroys the metal

Subject to severe localized corrosion that virtually destroys the metal

Undergoes graphitic corrosion

Undergoes graphitic corrosion

Interestingly, the corrosion rates of specimens completely immersed in seawater do not appear to depend on the geographical location of the test site; therefore, by inference, the mean temperature does not appear to play an important role. This constancy of the corrosion rate in seawater has been attributed to the more rapid fouling of the exposed steel by marine organisms, such as barnacles and algae, in warmer seas. It is further speculated that this fouling offsets that increases expected from the temperature rise.

Soil Corrosion The response of carbon steel to soil corrosion depends primarily on the nature of the soil and certain other environmental factors, such as the availability to moisture and oxygen. These factors can lead to extreme variations in the rate of the attack. For example, under the worst condition a buried vessel may perforate in less than one year, although archeological digs in arid desert regions have uncovered iron tools that are hundreds of years old. Some general rules can be formulated. Soils with high moisture content, high electrical conductivity, high acidity, and high dissolved salts will be most corrosive. The effect of aeration on soils is somewhat different from the effect of aeration in water because poorly aerated conditions in water can lead to accelerated attack by sulfate-reducing anaerobic bacteria. The effect of low levels of alloying additions on the soil corrosion of carbon steels is modest. Some data seems to show a small benefit of 1%Cu and 2.5% Ni on plain carbon steel. The weight loss and maximum pit depth in soil corrosion can be represented by an equation of the form:

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Z = a·tm Where: Z - either the weight of loss of maximum pit depth T - time of exposure a and m - constants that depend on the specific soil corrosion situation.

2.6. Gray Iron Abstract: Cast irons are alloys of iron, carbon, and silicon in which more carbon is present than can be retained in solid solution in austenite at the eutectic temperature. In gray cast iron, the carbon that exceeds the solubility in austenite precipitates as flake graphite. Gray irons usually contain 2.5 to 4% C, 1 to 3% Si, and additions of manganese, depending on the desired microstructure (as low as 0.1% Mn in ferritic gray irons and as high as 1.2% in pearlitics). Sulphur and phosphorus are also present in small amounts as residual impurities.

Cast irons are alloys of iron, carbon, and silicon in which more carbon is present than can be retained in solid solution in austenite at the eutectic temperature. In gray cast iron, the carbon that exceeds the solubility in austenite precipitates as flake graphite. Gray irons usually contain 2.5 to 4% C, 1 to 3% Si, and additions of manganese, depending on the desired microstructure (as low as 0.1% Mn in ferritic gray irons and as high as 1.2% in pearlitics). Sulphur and phosphorus are also present in small amounts as residual impurities. The composition of gray iron must be selected in such a way to satisfy three basic structural requirements: • • •

The required graphite shape and distribution The carbide-free (chill-free) structure The required matrix

For common cast iron, the main elements of the chemical composition are carbon and silicon. High carbon content increases the amount of graphite or Fe3C. High carbon and silicon contents increase the graphitization potential of the iron as well as its castability. The combined influence of carbon and silicon on the structure is usually taken into account by the carbon equivalent (CE): CE = %C + 0.3x(%Si) + 0.33x(%P) - 0.027x(%Mn) + 0.4x(%S) Although increasing the carbon and silicon contents improves the graphitization potential and therefore decreases the chilling tendency, the strength is adversely affected. This is due to ferrite promotion and the coarsening of pearlite.

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The manganese content varies as a function of the desired matrix. Typically, it can be as low as 0.1% for ferritic irons and as high as 1.2% for pearlitic irons, because manganese is a strong pearlite promoter. The effect of sulfur must be balanced by the effect of manganese. Without manganese in the iron, undesired iron sulfide (FeS) will form at grain boundaries. If the sulfur content is balanced by manganese, manganese sulfide (MnS) will form, which is harmless because it is distributed within the grains. The optimum ratio between manganese and sulfur for a FeS-free structure and maximum amount of ferrite is: %Mn = 1.7x(%S) + 0.15 Other minor elements, such as aluminum, antimony, arsenic, bismuth, lead, magnesium, cerium, and calcium, can significantly alter both the graphite morphology and the microstructure of the matrix. In general, alloying elements can be classified into three categories. Silicon and aluminum increase the graphitization potential for both the eutectic and eutectoid transformations and increase the number of graphite particles. They form colloid solutions in the matrix. Because they increase the ferrite/pearlite ratio, they lower strength and hardness. Nickel, copper, and tin increase the graphitization potential during the eutectic transformation, but decrease it during the eutectoid transformation, thus raising the pearlite/ferrite ratio. This second effect is due to the retardation of carbon diffusion. These elements form solid solution in the matrix. Since they increase the amount of pearlite, they raise strength and hardness. Chromium, molybdenum, tungsten, and vanadium decrease the graphitization potential at both stages. Thus, they increase the amount of carbides and pearlite. They concentrate in principal in the carbides, forming (FeX)nC-type carbides, but also alloy the aFe solid solution. As long as carbide formation does not occur, these elements increase strength and hardness. Above a certain level, any of these elements will determine the solidification of a structure with Fe3C (mottled structure), which will have lower strength but higher hardness. Generally, it can be assumed that the following properties of gray cast irons increase with increasing tensile strength from class 20 to class 60: • • • •

All strengths, including strength at elevated temperature Ability to be machined to a fine finish Modulus of elasticity Wear resistance.

On the other hand, the following properties decrease with increasing tensile strength, so that low-strength irons often perform better than high-strength irons when these properties are important: • •

Machinability Resistance to thermal shock

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• •

Damping capacity Ability to be cast in thin sections.

Successful production of a gray iron casting depends on the fluidity of the molten metal and on the cooling rate, which is influenced by the minimum section thickness and on section thickness variations. Casting design is often described in terms of section sensitivity. This is an attempt to correlate properties in critical sections of the casting with the combined effects of composition and cooling rate. All these factors are interrelated and may be condensed into a single term, castability, which for gray iron may be defined as the minimum section thickness that can be produced in a mold, cavity with given volume/area ratio and mechanical properties consistent with the type of iron being poured. Scrap losses resulting from missruns, cold shuts, and round corners are often attributed to the lack of fluidity of the metal being poured. Mold conditions, pouring rate, and other process variables being equal, the fluidity of commercial gray irons depends primarily on the amount of superheat above the freezing temperature (liquidus). As the total carbon content decreases, the liquidus temperature increases, and the fluidity at a given pouring temperature therefore decreases. Fluidity is commonly measured as the length of flow into a spiral-type fluidity test mold. The significance of the relationships between fluidity, carbon content, and pouring temperature becomes apparent when it is realized that the gradation in strength in the ASTM classification of gray iron is due in large part to differences in carbon content (~3.60 to 3.80% for class 20; ~2.70 to 2.95% for class 60). The fluidity of these irons thus resolves into a measure of the practical limits of maximum pouring temperature as opposed to the liquidus of the iron being poured. The usual microstructure of gray iron is a matrix of pearlite with graphite flakes dispersed throughout. Foundry practice can be varied so that nucleation and growth of graphite flakes occur in a pattern that enhances the desired properties. The amount, size, and distribution of graphite are important. Cooling that is too rapid may produce so-called chilled iron, in which the excess carbon is found in the form of massive carbides. Cooling at intermediate rates can produce mottled iron, in which carbon is present in the form of both primary cementite (iron carbide) and graphite. Flake graphite is one of seven types (shapes or forms) of graphite established in ASTM A 247. Flake graphite is subdivided into five types (patterns), which are designated by the letters A through E. Graphite size is established by comparison with an ASTM size chart, which shows the typical appearances of flakes of eight different sizes at l00x magnification. Type A flake graphite (random orientation) is preferred for most applications. In the intermediate flake sizes, type A flake graphite is superior to other types in certain wear applications such as the cylinders of internal combustion engines.

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Type B flake graphite (rosette pattern) is typical of fairly rapid cooling, such as is common with moderately thin sections (about 10 mm) and along the surfaces of thicker sections, and sometimes results from poor inoculation. The large flakes of type C flake graphite are formed in hypereutectic irons. These large flakes enhance resistance to thermal shock by increasing thermal conductivity and decreasing elastic modulus. On the other hand, large flakes are not conducive to good surface finishes on machined parts or to high strength or good impact resistance. The small, randomly oriented interdendritic flakes in type D flake graphite promote a fine machined finish by minimizing surface pitting, but it is difficult to obtain a pearlitic matrix with this type of graphite. Type D flake graphite may be formed near rapidly cooled surfaces or in thin sections. Frequently, such graphite is surrounded by a ferrite matrix, resulting m soft spots in the casting. Type E flake graphite is an interdendritic form, which has a preferred rather than a random orientation. Unlike type D graphite, type E graphite can be associated with a pearlitic matrix and thus can produce a casting whose wear properties are as good as those of a casting containing only type A graphite in a pearluic matrix. There are, of course, many applications in which flake type has no significance as long as the mechanical property requirements are met.

2.7. Hardenable Carbon Steels Abstract: Carbon steels are produced in greater tonnage and have wider use than any other metal because of their versatility and low cost. There are now almost 50 grades available in the nonresulfurized series 1000 carbon steels and nearly 30 grades in the resulfurized series 1100 and 1200. The versatility of the carbon steel group has also been extended by availability of the various grades with lead additions.

Carbon steels are produced in greater tonnage and have wider use than any other metal because of their versatility and low cost. There were several reasons why carbon steels proved satisfactory on reappraisal: a. their hardenability, though less than that of alloy steels, was adequate for many parts, and for some parts shallower hardening was actually an advantage because of minimized quench cracking; b. refinements in heat treating methods, such as induction hardening, flame hardening, and "shell quenching", made it possible to obtain higher properties from carbon steels than previously; and c. new compositions were added to the carbon steel group, permitting more discriminating selection. There are now almost 50 grades available in the nonresulfurized series 1000 carbon steels and nearly 30 grades in the resulfurized series 1100 and 1200. The versatility of the carbon steel group has also been extended by availability of the various grades with lead additions. Carbon steels can be divided into three arbitrary classifications based on carbon content.

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Steels with 0.10 to 0.25% C. Three principal types of heat treatment are used for this group of steels: a. conditioning treatments, such as process annealing, that prepare the steel for certain fabricating operations, b. case hardening treatments, and c. quenching and tempering to improve mechanical properties. The improvement in mechanical properties that can be gained by straight quenching and tempering of the low-carbon steels is usually not worth the cost. An example of process annealing is in the treatment of low-carbon cold-headed bolts made from cold drawn wire. Sometimes the strains introduced by cold working weaken the heads so much that they break through the most severely worked portion under slight additional strain. Process annealing overcomes this condition. Since the temperatures used are close to the lower transformation temperature, this treatment results in considerable reduction of the normal mechanical properties of the shank of the cold headed bolt. A more suitable treatment is stress relieving at about 1000oF (540oC). This treatment is used in order to retain much of the strength acquired in cold working and to provide ample toughness. A common practice is to combine a stress-relieving treatment with a quench from the upper transformation temperature, or slightly above, producing mechanical properties that approach those of cold drawn stock. A common quenching medium is a water solution of soluble oil, the use of which produces two desirable results: a. the surface of the parts acquires a pleasing black color accepted as a commercial finish, and b. the speed of the quench is slowed to the point where fully quenched hardness is not produced, so it is not necessary to temper the parts. Heat treatments are frequently employed to improve machinability. The generally poor machinability of the low-carbon steels, except those containing sulfur or other special alloying elements, results principally from the fact that the proportion of free ferrite to carbide is high. This situation cannot be changed fundamentally, but the machinability can be improved by putting the carbide in its most voluminous form, pearlite, and dispersing this pearlite evenly throughout the ferrite mass. Normalizing is commonly used with success, but best results are obtained by quenching the steel in oil from 1500 to 1600oF (815-870oC). With the exception of steels 1024 and 1025, no martensite is formed, and the parts do not require tempering. Steels with 0.25 to 0.55% C. Because of their higher carbon content, these steels are usually used in the hardened and tempered condition. By selection of quenching medium and tempering temperature a wide range of mechanical properties can be produced. They are the most versatile of the three groups of carbon steels and are most commonly used for crankshafts, couplings, tie rods and many other machinery parts where the required hardness values are within the range from 229 to 447 HB. This group of steels shows a continuous change from water-hardening to oilhardening types. The hardenability is very sensitive to changes in chemical composition, particularly to the content of manganese, silicon and residual elements, and to grain size; the steels are sensitive to section changes.

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The rate of heating parts for quenching has a marked effect on hardenability under certain conditions. If the structure is non-uniform, as a result of severe banding or lack of proper normalizing or annealing, extremely rapid heating such as may be obtained in liquid baths, will not allow sufficient time for diffusion of carbon and other elements in the austenite. As a result, non-uniform or low hardness will be produced unless the duration of heating is extended. In heating steels that contain free carbide (for example, spheroidized material), sufficient time must be allowed for the solution of the carbides; otherwise the austenite at the time of quenching will have a lower carbon content than is represented by the chemical composition of the steel, and disappointing results may be obtained. These medium-carbon steels should usually be either normalized or annealed before hardening, in order to obtain the best mechanical properties after hardening and tempering. Parts made from bar stock are frequently given no treatment prior to hardening, but it is common practice to normalize or anneal forgings. Most of bar stocks, both, hot finished and cold finished, are machined as received, except the higher-carbon grades and small sizes, which require annealing to reduce the asreceived hardness. Forgings are usually normalized, since this treatment avoids the extreme softening and consequent reduction of machinability that result from annealing. In some instances a "cycle treatment" is used. In this practice the parts are heated as for normalizing, and are then cooled rapidly in the furnace to a temperature somewhat above the nose of the S-curve - that is, within the transformation range that produces pearlite. Then the parts are held at temperature or cooled slowly until the desired amount of transformation has taken place; thereafter they are cooled in any convenient manner. Specially arranged furnaces are usually required. Details of the treatments vary widely and are frequently determined by the furnace equipment available. Cold headed products are commonly made from these steels, especially from the ones containing less than 0.40% C. Process treating before cold working is usually necessary because the higher carbon decreases the workability. For certain uses, these steels are normalized or annealed above the upper transformation temperature, but more frequently a spheroidizing treatment is used. The degree of spheroidization required depends on the application. After shaping operations are finished, the parts are heat treated by quenching and tempering. These medium-carbon steels are widely used for machinery parts for moderate duty. When such parts are to be machined after heat treatment, the maximum hardness is usually held to 321HB, and is frequently much lower. Salt solutions are often successfully used. Salt solutions are not dangerous to operators but their corrosive action on iron or steel parts of equipment is very serious. When the section is light or the properties required after heat treatment are not high, oil quenching is often used. This nearly always eliminates the breakage problem and is very effective in reducing distortion. A wide range in austenitizing temperatures is made necessary in order to meet required conditions. Lower temperatures should be used for the higher-manganese

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steels, light sections, coarse-grained material and water quenching; higher temperatures are required for lower manganese, heavy sections, fine grain and oil quenching. From these steels are made many common hand tools, such as pliers, open-end wrenches, screwdrivers, and a few edged tools - for example, tin snips and brush knives. The cutting tools are necessarily quenched locally on the cutting edges, in water, brine or caustic, and are subsequently given suitable tempering treatments. In some instances the edge is time quenched; then the remainder of the tool is oil quenched for partial strengthening. When made of these grades of steel, pliers, wrenches and screwdrivers are usually quenched in water, either locally or completely, and are then suitably tempered. Steels with 0.55 to 1.00% C. Carbon steels with these higher carbon contents are more restricted in application than the 0.25 to 0.55% C steels since they are more costly to fabricate, because of decreased machinability, poor formability and poor weldability. They are also more brittle in the heat treated condition. Higher-carbon steels such as 1070 to 1095 are especially suitable for springs where resistance to fatigue and permanent set are required. They are also used in the nearly fully hardened condition (Rockwell C 55 and higher) for applications where abrasion resistance is the primary requirement, as for agricultural tillage tools such as plowshares, and knives for cutting hay or grain. Forged parts should be annealed because refinement of the forging structure is important in producing a high-quality hardened product, and because the parts come from the hammer too hard for cold trimming of the flash or for economical machining. Ordinary annealing practice, followed by furnace cooling to 1100oF (590oC), is satisfactory for most parts. Most of the parts made from steels in this group are hardened by conventional quenching. However, special technique is necessary sometimes. Both oil and water quenching are used - water, for heavy sections of the lower-carbon steels and for cutting edges and oil, for general use. Austempering and martempering are often successfully applied; the principal advantages from such treatments are considerably reduced distortion, elimination of breakage, in many instances, and greater toughness at high hardness. For heavy machinery parts, such as shafts, collars and the like, steels 1055 and 1061 may be used, either normalized and tempered for low strength, or quenched and tempered for moderate strength. Other steels in the list may be used, but the combination of carbon and manganese in the two mentioned makes them particularly well adapted for such applications. It must be remembered that even with all hardenability factors favorable, including the use of a drastic quench, these steels are essentially shallow hardening, as compared with alloy steels, because carbon alone, or in combination with manganese in the amounts involved here, does not promote deep hardening to any significant extent. Therefore, the sections for which such steels are suited will be definitely limited. In spite of this limitation the danger of breakage is real and must be carefully guarded against when such parts are being treated, especially whenever changes in section are involved.

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Hand tools made from these steels include open-end wrenches, Stillson wrenches, hammers, mauls, pliers and screw drivers and cutting tools, such as hatchets, axes, mower knives and band knives. The combination of carbon and manganese in the steels used may vary widely for the same type of tool, depending partly on the equipment available for manufacture and partly on personal experience with, or preference for, certain combinations. A manganese content lower than standard will be used in some tools. This is justified when it makes a particular carbon range easier to handle, but it should be understood that for many applications, a combination of lower carbon and higher manganese would serve just as well.

2.8. Cast Carbon Steels Abstract: Carbon steels contain only carbon as the principal alloying element. Other elements are present in small quantities, including those added for deoxidation. Silicon and manganese in cast carbon steels typically range from 0.25 to about 0.80% Si, and 0.50 to about 1.00% Mn. Carbon steels can be classified according to their carbon content into three broad groups:

• • •

Low-carbon steels: < 0.20% C Medium-carbon steels: 0.20 to 0.50% C High-carbon steels: > 0.50% C

Carbon steels contain only carbon as the principal alloying element. Other elements are present in small quantities, including those added for deoxidation. Silicon and manganese in cast carbon steels typically range from 0.25 to about 0.80% Si, and 0.50 to about 1.00% Mn. Carbon steels can be classified according to their carbon content into three broad groups: Low-carbon steels: < 0.20% C Medium-carbon steels: 0.20 to 0.50% C High-carbon steels: > 0.50% C Low-alloy steels contain alloying elements, in addition to carbon, up to a total alloy content of 8%. Cast steels containing more than the following amounts of a single alloying element are considered low-alloy cast steels: Element

Mn

Amount (%)

1.00

Si 0.80

Ni

Cu

Cr

Mo

V

W

0.50

0.50

0.25

0.10

0.05

0.05

For deoxidation of carbon and low-alloy steels, aluminum, titanium, and zirconium are used. Aluminum is more frequently used because of its effectiveness and low cost. Unless otherwise specified, the normal sulphur limit for carbon and low-alloy steels is 0.06%, and the normal phosphorus limit is 0.05%.

Structure and Property Correlations Carbon steel castings are produced to a great variety of properties because composition and heat treatment can be selected to achieve specific combinations of

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properties, including hardness, strength, ductility, fatigue resistance, and toughness. Although selections can be made from a wide range of properties, it is important to recognize the interrelationships among these properties. For example, higher hardness, lower toughness, and lower ductility values are associated with higher strength values. The relationships among these properties and mechanical properties will be discussed in the following text. Property trends among carbon steels are illustrated as a function of the carbon content in Fig. 1.

Figure 1: Properties of cast carbon steels as a function of carbon content and heat treatment. (a) Tensile strength and reduction of area; (b) Yield strength and reduction of area; (c) Brinell hardness; (d) Charpy V-notch impact energy

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Strength and Hardness. Depending on alloy choice and heat treatment, ultimate tensile strength levels from 414 to 1724 MPa can be achieved with cast carbon and low-alloy steels. For carbon steels, the hardness and strength values are largely determined by carbon content and the heat treatment. Strength and Ductility. Ductility depends greatly on the strength, or hardness, of the cast steel (Fig. 2). Actual ductility requirements vary with the strength level and the specification to which steel is ordered. Quenched-and-tempered steels exhibit higher ductility values for a given yield strength level than normalized, normalizedand-tempered, or annealed steels.

Figure 2: Tensile properties of cast carbon steels as a function of Brinell hardness. Strength and Toughness. Several test methods are available for evaluating the toughness of steels or the resistance to sudden or brittle fracture. These include the Charpy V-notch impact test, the drop-weight test, the dynamic tear test, and specialized procedures to determine plane-strain fracture toughness.

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Charpy V-notch impact energy trends at room temperature reveal the distinct effect of strength and heat treatment on toughness. Higher toughness is obtained when steel is quenched and tempered, rather than normalized and tempered. Quenching, followed by tempering, produces superior toughness as indicated by the shift of the impact energy transition curve to lower temperatures. Nil ductility transition temperatures (NDTT) ranging from 38°C to as low as -90°C have been recorded in tests on normalized-and-tempered cast carbon and low-alloy steels in the yield strength range of 207 to 655 MPa. When cast steels are quenched and tempered, the range of strength and of toughness is broadened. Depending on alloy selection, NDTT values of as high as 10°C to as low as -107°C can be obtained in the yield strength range of 345 to 1345 MPa. An approximate relationship exists between the Charpy V-notch impact energy temperature behavior and the NDTT value. The NDTT value frequently coincides with the energy transition temperature determined in Charpy V-notch tests. Plane-strain fracture toughness (KIc) data for a variety of steels reflect the important strength-toughness relationship. Fracture mechanics tests have the advantage over conventional toughness tests of being able to yield material property values that can be used in design equations. Strength and Fatigue. The most basic method of presenting engineering fatigue data is by means of the S-N curve, which relates the dependence of the life of the fatigue specimen in terms of the number of cycles to failure N to the maximum applied stress. Other tests have been used, and the principal findings for constant amplitude tests and fatigue notch sensitivity for cast carbon steels are highlighted below. The endurance ratio (endurance limit divided by the tensile strength) of cast carbon and low-alloy steels as determined by rotating-beam bending fatigue tests is generally taken to be approximately 0.40 to 0.50 for smooth bars. The results indicate that this endurance ratio is largely independent of strength, alloying additions, and heat treatment. The fatigue notch sensitivity factor determined in rotating-beam bending fatigue tests is related to the microstructure of the steel (composition and heat treatment) and the strength. The quenched-and-tempered steels with a martensitic structure are less notch-sensitive than the normalized-and-tempered steels with a ferritepearlite microstructure. Section Size and Mass Effects. Mass effects are common to steels, whether rolled, forged, or cast, because the cooling rate during heat treating varies with section size and because the microstructure constituents, grain size, and nonmetallic inclusions increase in size from surface to center. Mass effects are metallurgical in nature and are distinct from the effect of discontinuities, which are discussed in the following section in this article. The section size or mass effect is of particular importance in steel castings because mechanical properties are typically assessed from test bars machined from standardized coupons having fixed dimensions and are cast separately from or attached to the castings. The removal of test bars from the casting is impractical

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because removal of material for testing would destroy the usefulness of the component. Test specimens removed from a casting will not routinely exhibit the same properties as test specimens machined from the standard test coupon designs for which minimum properties are established in specifications. The mass effect discussed above, shows that the difference in cooling rate between the test coupons and the part being produced, is the fundamental reason for this situation. Several specifications such as ASTM E 208, A 356, and A 757 provide for the mass effect by permitting the testing of coupons that are larger than normal and that have cooling rates more representative of those experienced by the part being produced.

2.9. Cold Rolled Steels Abstract: Cold rolled steels provide excellent press formability, surface finish, and thickness and flatness tolerances. Steel companies manufacture three groups of low- or ultra-low-carbon grades to meet a variety of customer formability requirements: CS Type B, DS Type B, EDDS, and EDDS+. They also produce HSLA steels and structural steel grades for those applications that require specified strength levels.

Cold rolled steels provide excellent press formability, surface finish, and thickness and flatness tolerances. Steel companies manufacture three groups of low- or ultralow-carbon grades to meet a variety of customer formability requirements: CS Type B, DS Type B, EDDS, and EDDS+. They also produce HSLA steels and structural steel grades for those applications that require specified strength levels. Cold rolled steels can also be specified as dent resistant or bake hardenable for applications that require dent resistance after forming and painting. Each grade can be processed with several surface finishes depending on customer requirements. Lubricants can be applied to enhance formability and to avoid at-press lubrication. Cold rolled steels have the following features: • • • •

Excellent Surface Appearance. Cold Rolled Steels have manufacturing controls in place assuring consistent surface quality to satisfy customer requirements. Formability. Cold Rolled Steels can be used to produce parts containing simple bends to parts with extreme deep drawing requirements. Paintability. Due to stringent surface roughness controls, Cold Rolled Steels are readily paintable using essentially any paint system. Weldability. Cold Rolled Steels can be joined using virtually any accepted welding practice.

Standard grades for cold rolled steels are: • • •

Commercial Steel (CS Type B). May be moderately formed; a specimen cut in any direction can be bent flat on itself without cracking. Drawing Steel (DS Type B). DS Type B is made by adding aluminum to the mol steel and may be used in drawing applications. Extra Deep Drawing Steel (EDDS). Interstitial Free (I-F) steels are made Drawing Steel by adding titanium and/or niobium to the molten steel after vacuum degassing and offer excellent drawability.

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•

Extra Deep Drawing Steel Plus (EDDS+). Interstitial Free (I-F) steels are made by adding titanium and/or niobium to the molten steel after vacuum degassing and offer excellent drawability.

Surface Finish Cold rolled steels are manufactured with a matte finish obtained by rolling with specially roughened rolls on the cold mill and the temper mill. This finish helps to maintain effective lubrication during metal forming and improves the appearance of painted surfaces. Non-standard matte finishes can be provided that optimize the opposing effects of surface roughness on painted part appearance and lubrication during press forming.

Surface Protection and Lubrication To prevent rusting in transit and storage, cold rolled steels can be supplied with a rust protective oil film or press forming lubricants. A pre-applied press forming lubricant provides uniform lubrication and eliminates the housekeeping problems. A dry film (acrylic/polymer) lubricant can also be supplied by further processing the cold rolled product through a coil coating facility. These specialty organic coatings are easily removed with a mild alkaline cleaner.

Formability and Mechanical Properties The formability of all steel products is a result of the interaction of many variables, the main ones being the mechanical properties of the steel, the forming system (tooling) used to manufacture parts, and the lubrication used during forming. Tight control over chemical composition, hot rolling parameters, amount of cold reduction, annealing time and temperature, and the amount of temper rolling allow the production of high-quality cold rolled steel products to meet customers requirements. Commercial Steel (CS Type B) should be used for moderate forming or bending applications. CS Type B products are produced from aluminum-killed continuously cast slabs and, unless otherwise specified, have a carbon content of less than 0.15%. To prevent the occurrence of fluting or stretcher strains during forming, CS products are tempered as a normal step in the mill processing. For more severe forming applications, Drawing Steel Type B (DS Type B) should be used. DS Type B has a controlled carbon content (200 m/sec), yielding good coating densities, potentially approaching theoretical density. Plasma spraying results in fine, essentially equiaxed grains, without extensive columnar boundaries, of particular advantage in certain ceramics applications (thermal-barrier coatings, for example). Coatings are chemically homogeneous; there is no (or controllable) change in composition with thickness. It is possible, however, to change from depositing a metal, to a continuously varying metalceramic mixture, to a ceramic-rich mixture, and finally to a completely ceramic outer layer, using programmed automation without intermediate delays in spraying or in part handling. Off-the-shelf plasma-spray equipment offers the capability of high coating-feedstock throughput (3 kg/hr), and special high-power guns can achieve a feedstock (e.g., alumina) throughput of over 25 kg/hr. Aside from normally spraying in air, it is possible, and sometimes essential, to plasma spray in a reduced-pressure environment chamber. Underwater spraying also is possible.

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The plasma flame is maintained by a steady, continuous-arc discharge of flowing inert gas, generally argon plus a small percentage of an enthalpy-enhancing diatomic gas, such as hydrogen. Feedstock powder {10 to 70 μm diameter} is carried by an inert gas into the emerging plasma flame. The particles melt in transit without vaporizing excessively, are accelerated, and impinge on the substrate where they flatten and solidify at cooling rates similar to those achieved in rapid-solidification processes. Much of the heat contained within the particles being deposited, as well as the heat of solidification and the heat of the plasma flame, is removed by conduction through the substrate. Consequently, precautions must be taken to prevent thermal degradation of substrate properties, or to prevent a metal substrate and/or coating from becoming excessively oxidized. Both the substrate and coating contract upon cooling, which can generate high residual stresses if a significant difference in coefficients of thermal expansion exists; these stresses can lead to coating delamination. While there are hundreds of parameters that influence the plasma-arc spraying process, about 12 have been identified as having the strongest influence on coating properties and the survivability of the coating system. Improved control of these parameters was the focus of many developments that have occurred during the past few years, and is the focus of many current developments. These include incorporating empirical or real-time feedback looping, redesigning fundamental gun components and feedstock powders (e.g. chemical composition, size distribution, and shape), and rethinking power-supply design. There also have been major changes in gas-handling equipment. Mass-flow control and metering are replacing traditional analog gages, which enable digital output with feedback potential. Data logging is gaining acceptance; flawed areas within a coating are now attributable to an "event" in gas flow, for example. Similar control schemes have been adopted for the powder-feed operation, including a variety of devices that display instantaneous powder-feed rates. Powder feeders also have changed, with fluidized-bed feeders becoming common; these feeders permit smooth flow (less pulsing) of a wider range of powder types. In the area of power supply, controlled de-power supply systems incorporating heat exchangers have been designed specifically for use with plasma guns, and are becoming the standard in the industry. And while not yet commonly practiced throughout the industry, monitoring and logging current, voltage, cooling-system temperature at various locations (including the gun), gas parameters, and feed rate is a relatively straightforward task. A revolutionary development in plasma-spray technology that occurred in the 1980s is reduced-pressure atmosphere chamber spraying. Plasma spraying essentially in the absence of oxygen allows the coating/substrate system to be maintained at a high temperature during processing, resulting in interfacial diffusion, which produces a true metallurgical bond. Chamber plasma spraying is expected to be capable of producing coatings having unique properties in a wide range of applications. For example, it is possible to chamber spray refractory oxides to obtain fully dense, well-bonded coatings. It also is relatively easy to add a high-temperature metal alloy to the oxide to obtain a

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composite having good high-temperature wear resistance. Chamber spraying also can produce good coatings of reactive metals, such as titanium and zirconium. An extension of the technique involves spraying the interior of large pipes or tanks for handling chemicals, using the vessel itself as a reduced-pressure inert-gas chamber by excluding air during spraying. Enhanced coating characteristics (e.g., density and adhesion strength) and accompanying improved coating properties achieved in chamber spraying are related to increased particle velocity and the high temperature of the coating/substrate system attained during spraying. Another variation of chamber spraying is reverse-arc sputtering. The technique involves electrically connecting the target substrate to the spray-gun system, which establishes a transferred arc at the surface, thus effecting a highly efficient sputtercleaning process. This surface pretreatment combined with the high coating/substrate temperature results in excellent coating adhesion.

Versatility through process variety Combustion-flame spraying generally uses an oxyacetylene flame to melt and spray either powder or wire feedstock. Due to its lower flame temperature and particle velocity compared with plasma spraying, flame spraying produces a less dense coating having lower adhesion strength. However, flame spraying is simpler in principle and operation, and system and production costs are lower than for plasma spraying. An additional consideration is the possible use of less-skilled operators because the process is more forgiving. Commercially available wire combustion flame guns can be used to spray virtually any welding wire including composite wires. A variation of combustion-flame spraying is the spray-and-fuse method of surface hardening. This well-established technique enables flame-spray deposition of a hardfacing material, for example, with subsequent flame fusing. Although the process lacks some control, it is highly effective and is widely used. The hypervelocity oxyfuel (HVOF) gun represents a major development in thermalspray technology. Developed by several companies to obtain well-bonded, dense coatings, HVOF guns have in common a method to burn oxygen and fuel and carry the combustion products through a nozzle with subsequent free expansion. This arrangement results in hypersonic flame gas velocities, and by introducing the feedstock powder "up-wind", powder particles attain high heat and supersonic velocities, this permits particle flattening upon striking the substrate, thus forming a dense coating. Special particle-size distributions are required for HVOF spraying, creating challenges and significant opportunities for powder producers. HVOF sprayed metallic coatings often have properties superior to those of plasmasprayed coatings, and equal to or superior to coatings produced using the detonation gun. The aircraft industry is especially interested in the HVOF spraying process for producing wear-resistant coatings. Refinements in the process are expected in the future, which may extend its application into areas traditionally dominated by plasma spraying.

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Two-wire, electric-arc spraying represents an important method to achieve low-cost application of metallic coatings. Most welding wires can be electric-arc sprayed at high throughput (from 30 to 50 kg/hr). During the process, two consumable wires, through which an electric current is passed, form an electric arc at the point where they intersect. The arc melts the wires and the molten metal is atomized by a continuous flow of either high-velocity compressed air or nonoxidizing gases, such as carbon dioxide, nitrogen, or argon. Coatings formed using air atomization are relatively dense and have good adhesion. Those formed using inert-gas atomization (which can be carried out in a reducedpressure chamber) are very dense and well-bonded to the substrate. The Sonarc process combines two-wire, electric-arc and HVOF spraying; molten metal at the arc is atomized and rapidly propelled to the substrate by the HVOF flame. The introduction of hard reinforcement particles (e.g. alumina or silicon carbide) into the flame makes it is possible to form either a metal-matrix composite coating or a free-standing bulk shape. The high particle velocities attainable in the Sonarc process result in extremely dense composite materials.

New powders create new opportunities The enhanced quality and variety of feedstock powders is contributing significantly to the advancement of thermal-spray technology. New processes are being used to economically produce special metal-alloy and ceramic formulations (e.g. cemented chromium and tungsten carbides). For example, GTE Products Corp. has developed a new microatomization process in which metal is melted using a plasma torch and molten droplets are propelled at high velocities against a rapidly rotating substrate. The droplets are fragmented and rapidly solidified resulting in spherical powders tens of micrometers in diameter, which can be used as feedstock for plasma and HVOF spraying. Spherical powders are especially necessary in plasma and HVOF spraying to obtain even, nonpulsing powder injection into the flame.

5.21. Hardenability of Steels Abstract: The traditional route to high strength in steels is by quenching to form martensite that is subsequently reheated or tempered, at an intermediate temperature, increasing the toughness of the steel without too great a loss in strength. The ability of steel to form martensite on quenching is referred to as the hardenability. Therefore, for the optimum development of strength, steel must be first fully converted to martensite. To achieve this, the steel must be quenched at a rate sufficiently rapid to avoid the decomposition of austenite during cooling to such products as ferrite, pearlite and bainite.

The traditional route to high strength in steels is by quenching to form martensite which is subsequently reheated or tempered at an intermediate temperature, increasing the toughness of the steel without too great a loss in strength. Therefore, for the optimum development of strength, steel must be first fully converted to martensite. To achieve this, the steel must be quenched at a rate sufficiently rapid to avoid the decomposition of austenite during cooling to such products as ferrite, pearlite and bainite. The effectiveness of the quench will depend primarily on two factors:

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• •

the geometry of the specimen, and the composition of the steel.

A large diameter rod quenched in a particular medium will obviously cool more slowly than a small diameter rod given a similar treatment. Therefore, the small rod is more likely to become fully martensitic. It has already been shown that the addition of alloying elements to a steel usually move the TTT curve to longer times, thus making it easier to pass the nose of the curve during a quenching operation, i.e. the presence of alloying elements reduces the critical rate of cooling needed to make a steel specimen fully martensite. If this critical cooling rate is not achieved a steel rod will be martensitic in the outer regions which cool faster but, in the core, the slower cooling rate will give rise to bainite, ferrite and pearlite depending on the exact circumstances. The ability of steel to form martensite on quenching is referred to as the hardenability. This can be simply expressed for steel rods of standard size, as the distance below the surface at which there is 50% transformation to martensite after a standard quenching treatment, and is thus a measure of the depth of hardening.

Use of TTT and CCT Diagrams TTT diagrams- TTT diagrams provide a good starting point for an examination of hardenability, but as they are statements of the kinetics of transformation of austenite carried out isothermally, they can only be a rough guide. To take one example, the effect of increasing molybdenum, Figure 1. shows the TTT diagrams for a 0.4 %C 0.2% Mo steel and steel with 0.3 %C 2 % Mo, Figure 2. The 0.2% Mo steel begins to transform in about one second at 550°C, but on increasing the molybdenum to 2% the whole C-curve is raised and the reaction substantially slowed so that the nose is above 700°C, the reaction starting after 4 minutes. The latter steel will clearly have a greatly enhanced hardenability over that of the 0.2 Mo steel. CCT diagrams- The obvious limitations of using isothermal diagrams for situations involving a range of cooling rates through the transformation temperature range have led to efforts to develop more realistic diagrams, i.e. continuous cooling (CCT) diagrams. These diagrams record the progress of the transformation with falling temperature for a series of cooling rates. They are determined using cylindrical rods, which are subjected to different rates of cooling, and the onset of transformation is detected by dilatometry, magnetic permeability or some other physical technique. The products of the transformation, whether ferrite, pearlite or bainite, are partly determined from isothermal diagrams, and can be confirmed by metallographic examination. The results are then plotted on a temperature/cooling time diagram, which records, for example, the time to reach the beginning of the pearlite reaction over a range of cooling rates. This series of results will give rise to an austenite-pearlite boundary on the diagram and likewise lines showing the onset of the bainite transformation can be constructed. A schematic diagram is shown in Figure 3. in which the boundaries for ferrite, pearlite, bainite and martensite are shown for hypothetical steel. The diagram is best used by superimposing a transparent overlay sheet with the same scales and having

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lines representing various cooling rates drawn on it. The phases produced at a chosen cooling rate are those which the superimposed line intersects on the continuous cooling diagram. In Figure 3. two typical cooling curves are superimposed for the surface and the centre of an oil-quenched 95 mm diameter bar. In this example, it should be noted that the centre cooling curve intersects the bainite region and consequently some bainite would be expected at the core of the bar after quenching in oil.

Figure 1. TTT diagram of a molybdenum steel 0.4C 0.2Mo

Figure 2. TTT diagram of a molybdenum steel 0.3C 2.0Mo

Figure 3. Relation between cooling curves for the surface and core of an oil-quenched 95 mm diameter bar and the microstructure

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Hardenability Testing The rate at which austenite decomposes to form ferrite, pearlite and bainite is dependent on the composition of the steel, as well as on other factors such as the austenite grain size, and the degree of homogeneity in the distribution of the alloying elements. It is extremely difficult to predict hardenability entirely on basic principles, and reliance is placed on one of several practical tests, which allow the hardenability of any steel to be readily determined: • •

The Grossman test The Jominy end quench test

Effect of Grain Size and Chemical Composition on Hardenability The two most important variables which influence hardenability are grain size and composition. The hardenability increases with increasing austenite grain size, because the grain boundary area is decreasing. This means that the sites for the nucleation of ferrite and pearlite are being reduced in number, with the result that these transformations are slowed down, and the hardenability is therefore increased. Likewise, most metallic alloying elements slow down the ferrite and pearlite reactions, and so also increase hardenability. However, quantitative assessment of these effects is needed. There are a bewildering number of steels, the compositions of which are usually complex and defined in most cases by specifications, which give ranges of concentration of the important alloying elements, together with the upper limits of impurity elements such as sulfur and phosphorus. While alloying elements are used for various reasons, the most important is the achievement of higher strength in required shapes and sizes and often in very large sections which may be up to a meter or more in diameter in the case of large shafts and rotors. Hardenability is, therefore, of the greatest importance, and one must aim for the appropriate concentrations of alloying element needed to harden fully the section of steel under consideration. Equally, there is a little point in using too high a concentration of alloying element, i.e. more than that necessary for full hardening of the required sections. Alloying elements are usually much more expensive than iron, and in some cases are diminishing natural resources, so there is additional reason to use them effectively in heat treatment. Carbon has a marked influence on hardenability, but its use at higher levels is limited, because of the lack of toughness which results, the greater difficulties in fabrication and, most important, increased probability of distortion and cracking during heat treatment and welding. The most economical way of increasing the hardenability of plain carbon steel is to increase the manganese content, from 0.60 wt% to 1.40 wt%, giving a substantial improvement in hardenability. Chromium and molybdenum are also very effective, and amongst the cheaper alloying additions per unit of increased hardenabilily. Boron

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has a particularly large effect when it’s added to fully deoxidized low carbon steel, even in concentrations of the order of 0.001%, and would be more widely used if its distribution in steel could be more easily controlled. Hardenabilily data now exists for a wide range of steels in the form of maximum and minimum end-quench hardenability curves, usually referred to as hardenability bands. This data is, available for very many of the steels listed in specifications such as those of the American Society of Automotive Engineers (SAE), the American Iron and Steel Institute (AISI) and the British Standards.

5.22. The Formation of Martensite Abstract: Rapid quenching of austenite to room temperature often results in the formation of martensite, a very hard structure in which the carbon, formerly in solid solution in the austenite, remains in solution in the new phase. Unlike ferrite or pearlite, martensite forms by a sudden shear process in the austenite lattice which is not normally accompanied by atomic diffusion. The martensite reaction in steels is the best known of a large group of transformations in alloys in which the transformation occurs by shear without change in chemical composition. The generic name of martensitic transformation describes all such reactions.

Rapid quenching of austenite to room temperature often results in the formation of martensite, a very hard structure in which the carbon, formerly in solid solution in the austenite, remains in solution in the new phase. Unlike ferrite or pearlite, martensite forms by a sudden shear process in the austenite lattice which is not normally accompanied by atomic diffusion. Ideally, the martensite reaction is a diffusionless shear transformation, highly crystallographic in character, which leads to a characteristic lath or lenticular microstructure. The martensite reaction in steels is the best known of a large group of transformations in alloys in which the transformation occurs by shear without change in chemical composition. The generic name of martensitic transformation describes all such reactions. It should however be mentioned that there is a large number of transformations which possess the geometric and crystallographic features of martensitic transformations, but which also involve diffusion. Consequently, the broader term of shear transformation is perhaps best used to describe the whole range of possible transformations. The martensite reaction in steels normally occurs athermally, i.e. during cooling in a temperature range which can be precisely defined for a particular steel. The reaction begins at a martensitic start temperature Ms which can vary over a wide temperature range from as high as 500°C to well below room temperature, depending on the concentration of γ-stabilizing alloying elements in the steel. Once the Ms is reached, further transformation takes place during cooling until the reaction ceases at the Mf temperature. At this temperature all the austenite should have transformed to martensite but frequently, in practice, a small proportion of the austenite does not transform. Larger volume fractions of austenite are retained in some highly alloyed steels, where the Mf temperature is well below room temperature.

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To obtain the martensitic reaction it is usually necessary for the steel to be rapidly cooled, so that the metastable austenite reaches Ms. The rate of cooling must be sufficient to suppress the higher temperature diffusion-controlled ferrite and pearlite reactions, as well as other intermediate reactions such as the formation of bainite. The critical rate of cooling required is very sensitive to the alloying elements present in the steel and, in general, will be lower the higher total alloy concentration. Each grain of austenite transforms by the sudden formation of thin plates or laths of martensite of striking crystallographic character. The laths have a well-defined habit plane and they normally occur on several variants of this plane within each grain. The habit plane is not constant, but changes as the carbon content is increased. Martensite is a supersaturated solid solution of carbon in iron which has a bodycentred tetragonal structure, a distorted form of bcc iron. It is interesting to note that carbon in interstitial solid solution expands the fcc iron lattice uniformly, but with bcc iron the expansion is nonsymmetrical giving rise to tetragonal distortion. Analysis of the distortion produced by carbon atoms in the several types of site available in the fcc and bcc lattices, has shown that in the fcc structure the distortion is completely symmetrical, whereas in the bcc one, interstitial atoms in z positions will give rise to much greater expansion of iron-iron atom distances than in the x and y positions. Assuming that the fcc-bcc tetragonal transformation occurs in a diffusionless way, there will be no opportunity for carbon atoms to move, so those interstitial sites already occupied by carbon will be favored. Since only the z sites are common to both the fcc and bcc lattices, on transformation there are more carbon atoms at these sites causing the z-axis to expand, and the non-regular the martensite, as well as the shape deformation for a number of martensitic transformations including ferrous martensites. It is, however, necessary to have accurate data, so that the habit planes of individual martensite plates can be directly associated with a specific orientation relationship of the plate with the adjacent matrix. Martensitic planes in steel are frequently not parallel-sided; instead they are often perpendicular as a result of constraints in the matrix, which oppose the shape change resulting from the transformation. This is one of the reasons why it is difficult to identify precisely habit planes in ferrous martensite. However, it is not responsible for the irrational planes, but rather the scatter obtained in experiments. Another feature of higher carbon martensites is the burst phenomenon, in which one martensite plate nucleates a sequence of plates presumably as a result of stress concentrations set up when the first plate reaches an obstruction such as a grain boundary or another martensite plate. Perhaps the most striking advances in the structure of ferrous martensites occurred when thin foil electron microscopy was first used on this problem. The two modes of plastic deformation are needed for the in-homogeneous deformation part of the transformation, i.e. slip and twinning. All ferrous martensites show very high dislocation densities of the order of 1011 to 1012cm2, which are similar to those of very heavily cold-worked alloys. Thus it is usually impossible to analyze systematically the planes on which the dislocations occur or determine their Burgers vectors.

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The lower carbon (1.4wt%, the orientation relationship changes from Kurdjumov-Sachs to Nishiyama, and the habit plane changes to around {259}. The change is not detectable microscopically as the morphology is still lenticular plates which form individually and are heavily twinned. Detailed crystallographic analysis shows that this type of martensite obeys more closely the theoretical predictions than the {225} martensite. The plates are formed by the burst mechanism and often an audible click is obtained. The {259} martensite only forms at very high carbon levels in plain carbon steels, although the addition of metallic alloying elements causes it to occur at much lower carbon contents, and in the extreme case in a carbon-free alloy such as Fe-Ni when the nickel content exceeds about 29 wt%.

5.23. Induction Surface Hardening and Tempering Abstract: Induction hardening is primarily used for surface hardening. The heating process does not affect the core structure. It is possible to heat a material locally where it is functionally desired. Other sectors of the material remain untreated and it is easy to machine them. Main advantages of induction hardening are:

• • • •

Low distortion Low risk of scaling Localized hardening Good reproducibility of hardening process

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• • • • •

Easy integration in production line Fully automatic process easily attainable Easy to operate machines Less harmful to the environment compared to other hardening processes Use of unalloyed steels.

Applications for Induction Hardening Induction hardening is primarily used for surface hardening. The heating process does not affect the core structure. It is possible to heat a material locally where it is functionally desired. Other sectors of the material remain untreated and it is easy to machine them. Induction hardening is in most cases more economical compared to other heating processes. In some cases it is the only possible heat treatment process. Main advantages are: • • • • • • • • •

Low distortion Low risk of scaling (These two advantages may allow final machining before hardening) Localized hardening Good reproducibility of hardening process Easy integration in production line Fully automatic process easily attainable Easy to operate machines Less harmful to the environment compared to other hardening processes Use of unalloyed steels.

The surface hardness depends on the carbon content of the steel. Surface hardening increases the wear resistance and can be used to increase the strength of highly stressed components. It is advisable to design a new component suitable for induction hardening. The following applications are a small selection of induction hardened components and applications. The frequency range of power sources for induction hardening is as follows: • • •

High frequency above 100 kHz (HF) Intermediate frequency between 10kHz and 100 kHz (ZF) Medium frequency from 3 kHz to 10 kHz (MF)

Small components with low hardness depth are normally hardened with HF or ZF; bigger components with higher hardness depth are normally hardened with ZF or MF. An exact definition of the frequency application range in advance is not always possible. Machine types used for induction hardening are: • •

Vertical hardening machines Horizontal hardening machines

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• •

Indexing table machines Special machines

Component Group: CV-Joints The tendency for light components with increasing torque values leads to thin-walled joints with high requirements on the hardening process. The rolling paths in the bell area as well as the shaft (if there) are hardened. Depending on the requirements it is done with static rotation, progressive, progressive rotation or single shot hardening. The huge number of different joints requires flexible hardening machines, quick and reliable change of tooling for different components. The machine and hardening parameters have do be controlled continuously. Hardness and crack detecting machines can be integrated into the hardening installation. Outer and inner races are mostly single shot hardened. Thus process requires advanced power sources with high power output, short and precise heating time and a quick ramp up time to achieve a reproducible quality. Some of these components can be tempered by residual heating. Interconnecting shafts are hardened in single shot or progressive with rotation. The hardening process is horizontal or vertical clamping between centers. Induction tempering is also possible. The change of length during heating can be registered and corrective action can be taken.

Components Group: Steering Parts The shaft can be hardened progressively with rotation of the workpiece. The tooth part can be hardened with three different methods: • • •

Combined induction-conduction hardening Progressive induction hardening with rotation and ring-inductor Progressive induction hardening with form-inductor and pre-positioned tooth section

Pinion racks and power valves of power steering systems can be induction tempered after hardening. Indexing table machines with automatic loading/unloading systems reduce the cycle time. A suitable heating process and the correct frequency are essential for the achievable quality. Hardening and tempering without scale is done in a protective atmosphere. Polishing and cleaning of the component after hardening is not needed. A patented quenching curtain system at in- and outlet of a chamber minimizes the gas consumption. The oxygen concentration in the chamber is controlled to ensure the process reliability. The process is suitable for integration in production lines and it can be retrofitted in appropriate older equipment.

Components Group: Engine Parts

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The demand for lower emissions, lighter and smaller engines influence the design of camshafts drastically. Precise hardening of camshafts in one operating sequence requires precise CNC controlled hardening machines with special inductors and quench rings to position each cam very precisely. A low distance between cams may require the simultaneous hardening of several cams at once. Separate control of power and heating time is an advantage. It is nowadays standard that all essential process parameters are automatically controlled by CNC. Induction hardening of valves requires precise positioning between inductor and work piece. Frequency, work piece rotation and energy input into the work piece have to be controlled in small tolerances. A fast floor-to-floor time requires fully integrated indexing table machines often with hardening of two valves simultaneously. Hardness testing equipment can be integrated. Valve hardening with lower production rates can be done with a good reproducibility on universal hardening machines. The hammer area of rocker arms is hardened progressively or by oscillation. Exact and reproducible coupling between work piece and inductor are also essential for this work piece. Rocker arms in cast quality need a special heat treatment with preheating and special timings. Tempering with residual heat is possible-also bore hardening for bigger diameters.

Components Group: Drive Shafts Rear axle shafts are hardened progressively with rotation or in single shot. Horizontal and vertical hardening machines with one or more hardening stations are used for this process. Hardening of the radius increases the component strength. Depending on the design and length it is recommended to use a steady rest to reduce the runout. These shafts are single shot hardened. In comparison with a structural steel hardened in a furnace cheaper carbon steel can be used with higher strength after induction hardening.

Components Group: Axles and Rods The high production volume of motor shafts requires horizontal in-line machines and progressive hardening with rotation. The very high feed rate requires precise and fast control for the power source and machine to achieve exact and reproducible hardening zones at the beginning and end of the component. A tempering process can follow. Hardening requires a horizontal machine with continuous feed and work piece rotation. The low wall thickness of the hollow shaft calls for precise power and frequency to reduce bending and to avoid through hardening. The machine is ideal for integration in the production line. Process control in connection with a rejecting chute ensures that only good work pieces leave the machine.

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Piston rods are preferably hardened on a horizontal machine independent of the machining condition of the end face. In-line or separate machines with loading magazine can be realized -induction tempering after hardening is possible.

Components Group: Gears Induction hardening increases the flank durability and tooth ground strength like furnace hardening but grinding may be skipped or reduced which results in cost advantage.

5.24. Overview of Mechanical Working Processes: Part One Abstract: During the process of shape change which accompanies mechanical working the volume of the mass remains constant and an increase in length such as in rolling is accompanied by a decrease in thickness. As deformation is applied to a structure consisting of one kind of deformable grains, they will become elongated. At the same time mechanical properties become directional and the structure and properties are anisotropic.

It is proposed to deal with the effect of mechanical work on the structure and macroproperties of metals and to follow this with a classification of the processes used for mechanical working.

Effects of Mechanical Work on Metals During the process of shape change which accompanies mechanical working the volume of the mass remains constant and an increase in length such as in rolling is accompanied by a decrease in thickness. As deformation is applied to a structure consisting of one kind of deformable grains, they will become elongated. At the same time mechanical properties become directional and the structure and properties are anisotropic. The behavior of a duplex structure is very similar except that the two phases or types of grains, α and β, are likely to react differently to the deformation process. α may be soft and ductile, whilst β may be hard and brittle, will therefore tend to fracture and appear as orientated fragments or stringers in the longitudinal direction. A duplex structure will tend to become more anisotropic than a single-phase structure. At very high degrees of deformation the structure appears fibrous because the grains have been so elongated as to lose their individual characteristics. Deformation also affects mechanical properties, in that the hardness, ultimate tensile and yield stresses all increase to a maximum, whilst the ductility falls to a very low value. The toughness, as measured by the Izod or Charpy test, increases with working up to a maximum and then gradually decreases. It is found in practice that the hardness and strength of most metals increase by 2.5 to 3 times the annealed value as a result of cold working. All structural metals have approximately the same ductility as measured by percentage elongation. An annealed metal will have approximately 35% elongation;

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whilst a metal which has been cold worked 80% will have only approximately 2% elongation before failure in a tensile test. The best combination of properties is usually found in the longitudinal direction, and the worst in the short transverse direction.

The Effect of Heat on Cold-Worked Metals A metal sample which has been cold worked 80% will be hard and brittle, the grains will be elongated and there will be a considerable degree of anisotropy. If the sample is heated, a temperature will be reached at which new nuclei begin to form in the distorted grains. This occurs due to the fact that the thermal energy supplied allows the atoms to diffuse to sites and form stable nuclei. How much thermal energy is needed depends upon the amount of prior cold work carried out on the metal. Cold work increases the internal energy of the metal, and the greater the cold work the higher the residual internal energy. This means that less thermal energy is required to nucleate a heavily cold-worked metal than a lightly cold- worked one. It is important to understand the mechanism of nucleation and the factors which control the number of nuclei formed. It is recognized that nucleation will occur in those regions with the highest residual stresses, and these occur at multiple boundary intersections. The longer the time that the worked sample is held at a nucleating temperature, the greater the number of atoms that will diffuse to the nuclei and occupy positions of minimum energy. The volume around each nucleus will grow to visible size and after some time further growth will be prevented by the interference of one growth volume with another. These growth volumes become grains and the interstititial zones of distorted atomic pattern are the grain boundaries. The grains will be softer and much larger than the worked grains and the atomic orientation will be random as between grains, replacing the common forced orientation in the worked material. Each nucleus has grown to form one grain and this gives the recrystallised grain size. The greater the degree of cold works the smaller the recrystallised grain size. With no cold work there are no high stress centers so no recrystallisation on heating. With the critical amount of cold work there are a few and these grow excessively to give very large grains. If the metal is held at the recrystallisation temperature after it has completely recrystallised, diffusion of atoms still occurs and some grains grow at the expense of others. This is called grain growth. It is quite possible in an industrial process that quite an appreciable amount of grain growth occurs so that the final or annealed grain size is much coarser than the recrystallised grain size. Grain growth occurs by a diffusion process and all such processes are affected by time and temperature. It has been seen that diffusion is a linear function of time, but increasing temperature has a far more critical effect on diffusion since the rate is an exponential function of temperature. Increasing the temperature by 10°C doubles the diffusion rate, and if the sample is heated to a temperature substantially above the recrystallisation temperature the grain growth will result in a coarse structure.

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Final grain size after cold working and annealing is very important in industrial processes. If the grains are too coarse the metal will exhibit a rough surface finish on machining and an "orange peel" effect after pressing. Grain size also affects the toughness. The best structure for further working consists of small, uniform equiaxed grains. The most important factor in the industrial process is the final temperature in the furnace. This should be as low as possible, whilst ensuring complete recrystallisation in adequate time.

Hot Working of Metals Cold working followed by annealing can be compared to working at above the recrystallisation temperature. This is described as hot working and deformation of the grains is followed by instantaneous recrystallisation. The effects of deformation on structure and properties are therefore instantly removed. This is the idealized situation in hot working. In practice the effects of deformation are instantaneous but recrystallisation requires time and unless the hot deformation system is slow enough to allow complete recrystallisation, then evidence of working persists at the end of the process. This concept gives us the true definition of hot and cold working: • •

Hot working is working at such a temperature and strain rate that recrystallisation keeps pace with deformation. Cold working is working under conditions such that recrystallisation does not keep pace with deformation.

An interesting example of the application of the above principles to hot working is in the hot rolling of steel strip for deep drawing. The final material should consist of small equiaxed grains exhibiting the minimum amount of mechanical anisotropy. Steel consists of two phases, ferrite and cementite. The two phases, however, have totally different recrystallisation temperatures. The ferrite recrystallises at around 600°C whereas the cementite requires temperatures between 700°C and 900°C, depending upon the carbon content. Even if the ferrite recrystallises after hot working, the presence of the cementite in an oriented formation prevents the development of equiaxed grains and mechanical anisotropy persists giving "pan caked" grains. Two precautions are taken to avoid this, firstly the carbon content is limited to 0.1% maximum to reduce the amount of second phase, and secondly the form of the second phase is controlled so that it is in an innocuous form. Cementite may appear as a massive form on the ferrite grain boundaries, if either the cooling after hot rolling is very slow or the finishing temperature is very high. A lamellar form, i.e. pearlite, is obtained if the cooling rate is intermediate, resulting from a relatively high finishing temperature. Finally, it can occur as fine dispersed spheroids if the cooling rate is an optimum value achieved by a relatively low hotworking temperature.

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6. Cast Iron

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6.1. Cast irons Abstract: In this paper is presentation flake graphite iron. Finds use due to: its cheapness and ease of machining; low-melting temperature (1140-1200°C); ability to take good casting impressions; wear resistance; high damping capacity; a reasonable tensile strength of 108-340 MPa associated with a very high compressive strength, making it very suitable for applications requiring rigidity and resistance to wear. The different types vary from grey iron which is machinable to either mottled or white iron which is not easily machinable. The white irons of suitable composition can be annealed to give malleable cast iron. During the last thirty years much development work has taken place and it has been found worth while to add even expensive elements to the cheap metal because vastly improved properties result. The new irons formed by alloying or by special melting and casting methods are becoming competitors to steel.

Flake graphite iron finds use due to: 1. 2. 3. 4. 5. 6.

its cheapness and ease of machining; low-melting temperature (1140-1200°C); ability to take good casting impressions; wear resistance; high damping capacity; a reasonable tensile strength of 108-340 MPa associated with a very high compressive strength, making it very suitable for applications requiring rigidity and resistance to wear.

The different types vary from grey iron which is machinable to either mottled or white iron which is not easily machinable. The white irons of suitable composition can be annealed to give malleable cast iron. During the last thirty years much development work has taken place and it has been found worth while to add even expensive elements to the cheap metal because vastly improved properties result. The new irons formed by alloying or by special melting and casting methods are becoming competitors to steel. The various irons can be classified as shown in Fig. 1 based on the form of graphite and the type of matrix structure in which it is embedded. The metallurgical structure, composition and section of the casting largely govern the engineering properties. One of the differences between cast iron and steel is the presence of a large quantity of carbon, generally 2-4%, and frequently high silicon contents. While carbon in ordinary steel exists as cementite (Fe3C), in cast iron it occurs in two forms: • •

stable form-graphite; unstable form-cementite, analysed as combined carbon. CAST IRON Grey machinable iron

Flake Graphite

White, unmachinable iron no graphite

Spheroidal Graphite

Ferritic Pearlitic Austenitic

Pearlitic

Martensitic

Martensitic

Malleable iron temper carbon graphite Ferritic Blackheart

Thin Whiteheart

Pearlitic Whiteheart

Special Malleable

Figure 1. Classification of cast iron (Pearce)

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Graphite is grey, soft, and occupies a large bulk, hence counteracting shrinkage; while cementite is intensely hard, with a density of the same order as iron. On the relative amounts, shape and the distribution of these two forms of carbon largely depend the general properties of the iron. The factors mainly influencing the character of the carbon are: 1. The rate of cooling. 2. The chemical composition. 3. The presence of nuclei of graphite and other substances.

1. Rate of cooling. A high rate of cooling tends to prevent the formation of graphite, hence maintains the iron in a hard, unmachinable condition. If the casting consists of varying sections then the thin ones will cool at a much greater rate than the thick. Consequently, the slowly cooled sections will be grey and the rapidly cooled material will be chilled. These points are illustrated in Fig. 2, which shows the variation in hardness of a step casting.

Figure 2. The relation between the rate of cooling and hardness as indicated by sections of varying thickness 2. The effect of chemical composition. 1. Carbon lowers the melting-point of the metal and produces more graphite. Hence it favours, a soft, weak iron. 2. Silicon slightly strengthens the ferrite but raises the brittle transition temperature, Indirectly, however, it acts as a softener by increasing the tendency of the cementite to slip up into graphite and ferrite. Fig. 3 shows the relation between the carbon and silicon contents in producing the different irons for one rate of cooling. It will be noted that either a high carbon and low silicon or low carbon and high silicon content give grey iron; the fracture can, therefore, be misleading as to analysis, especially if the rate of cooling is not considered. The amounts of silicon, giving the maximum values for various properties, are also shown in Fig. 3. The percentage of silicon is varied according to the thickness of the casting.

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3. Sulphur and manganese. Sulphur can exist in iron, as either iron sulphide, FeS, or manganese sulphide, MnS. Sulphur as FeS tends to promote cementite producing a harder iron. When manganese is added, MnS is formed which rapidly coalesces and rises to the top of the melt. The first effect of the manganese is, therefore, to cause the formation of graphite due to its effect on the sulphur. The direct effect of manganese is to harden the iron, and this it will do when it exists in amounts greater than that required to combine with the sulphur-1 part sulphur to 1,72 part manganese. 4. Phosphorus has a little effect on the graphite-cementite ratio; but renders the metal very fluid indirectly through the production of a low-melting constituent, which is readily recognised in the micro-structure (Fig. 4). In the production of sound castings of heavy section, phosphorus should be reduced to about 0,3% in order to avoid shrinkage porosity. 5. Trace elements not normally considered in routine analyses can exert a profound influence upon the characteristics of cast iron. Examples are 0,1% of aluminium graphitises, antimony embrittles, lead, tellurium promotes carbide but reduces strength of iron; 0,003% of hydrogen can greatly affect soundness of castings and tends to coarsen graphite. Nitrogen behaves as a carbide stabiliser; oxygen has no specific effect.

Figure 4. Common grey iron showing Figure 3. Diagram indicating the ferrite (F), pearlite (P) and phosphide structures of iron resulting from variation eutectic (PH) (x250). Ferrite is associated of silicon and carbon contents with the graphite. Note banded structure in the phosphide eutectic The carbon equivalent value. From Fig. 5 it will be seen that the eutectic E is at 4,3% carbon and irons with a greater carbon content will (under suitable conditions) start freezing by throwing out kish graphite of large size. With carbon contents progressively less than 4,3% normal graphite is formed in diminishing quantities until a mottled or white iron range is reached. Naturally other elements, especially silicon and phosphorus, affect the composition of the eutectic point in a complex alloy and a carbon equivalent value is suggested as an index which converts the amount of these elements into carbon replacement values.

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Figure 5. Iron-cementite equilibrium diagram Carbon equivalent value (CE) = Total C% + 1/3 (Si% + P%) For a given cooling rate the carbon equivalent value, therefore, determines how close a given composition of iron is to the eutectic (CE 4,3) and therefore how much free graphite is likely to be present, and consequently the probable strength in a given section: the carbon equivalent value is also a useful guide to chilling tendency of a given section, although it must be borne in mind that pouring temperature, cooling rate and alloying elements have a marked influence.

6.2. Relation between CE structure and mechanical properties Abstract: A useful first attempt to relate composition and structure was shown in Fig. 3 of the article Cast Irons but it had limited use in the foundry. Figure 1 shows a more useful relationship between CE value, structure, tensile strength in 30 mm dia bars and section size. A cylindrical test bar of given dia cools more rapidly than a flat plate of equivalent thickness, hence the section is expressed as bar diameter or section thickness. Line H is the boundary of unmachinable irons while line P is the boundary between soft and pearlitic irons. Thus an iron of carbon equivalent 4,35 should not be made thicker than 20 mm as a bar or 10 mm as a plate to attain a pearlitic iron. To avoid an unmachinable chilled casting the bar should not be less than 8 mm dia or plate less than 4 mm thick.

A useful first attempt to relate composition and structure was shown in Fig. 3 of the article Cast Irons but it had limited use in the foundry. Figure 1 shows a more useful relationship between CE value, structure, tensile strength in 30 mm dia bars and section size. A cylindrical test bar of given dia cools more rapidly than a flat plate of equivalent thickness, hence the section is expressed as bar diameter or section thickness. Line H is the boundary of unmachinable irons while line P is the boundary between soft and pearlitic irons. Thus an iron of carbon equivalent 4,35 should not be made thicker than 20 mm as a bar or 10 mm as a plate to attain a pearlitic iron. To avoid an unmachinable chilled casting the bar should not be less than 8 mm dia or plate less than 4 mm thick.

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T S. MPa in 30 mm dia. bar

Figure 1. Diagram relating section size, CE value, tensile strength and structure (After BCIRA) A melting furnace usually produces iron of a constant CE value and silicon is the element normally used to control chill. Alloying elements are added to cast iron to confer special properties and also to control the chill.

Formation of graphite Flake. Neglecting the effect phosphorus, and the presence of primary austenite dendrites, the successive stages in the growth from the liquid of flake graphite is shown in Fig. 2a The eutectic begins to solidify at nuclei from each of which is formed a roughly spherical lump, referred to as a eutectic cell. In this cell there has been simultaneous growth of austenite and graphite, the latter being in continuous

Figure 2.

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contact with the liquid. The normal appearance of graphite in a micrograph suggests that the structure is made up of a number of separate flakes, but now it is considered that within each eutectic cell there is a continuous branched skeleton of graphite, like a cabbage. The skeleton is branched more frequently with a rapid radial growth of the cell such as occurs when increasing the rate of cooling of an iron which produces undercooling, and therefore finer graphite in the micrograph (Fig. 3).

Figure 3. Medium size graphite outlining dendrites (x60) The diameter of a eutectic cell, therefore, has a major effect on mechanical properties, e.g. the greater the number of cells per unit volume the higher the tensile strength, but soundness is affected adversely. Superheating or holding time of the molten iron reduce the number of nuclei, while inoculants such as ferro-silicon and also sulphur increase nuclei. Spheroidal. Fig. 2b show the growth of spherulitic graphite in a magnesium-treated iron. In this case the spherulitic graphite is quickly surrounded by a layer of austenite and growth of the spheroid occurs by diffusion of carbon from the liquid through the austenite envelope. If diffusion distances become large there will be a tendency for the remaining liquid to solidify as white iron eutectic, hence inoculation in this iron is highly desirable in order to increase the number of graphite centres. Temper carbon nodules. At the malleabilising temperature (800-950°C) the solid white iron consists of eutectic matrix of cementite, austenite and sulphide inclusions. Nucleation of graphite then occurs at austenite cementite interfaces and at sulphide inclusions. The cementite gradually dissolves in the austenite and the carbon diffuses to the graphite nuclei. The MnS tends to form a flake aggregate and the FeS a spherulitic nodule (Fig. 2c).

Micro-structure of cast iron In preparing the specimens care is required, otherwise, erroneous results might arise. The graphite is readily removed during polishing and in this case the cavities can be either burnished over or enlarged. The various types of micro-structure can be classified into groups without considering the presence of phosphorus.

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The graphite can vary in size and form as illustrated in Figs. 3-5. The coarse flaky graphite is found in common iron, while the fine curly type, frequently outlining the dendrites, is found in high-class iron, especially when superheated before casting. Spheroidal graphite is found in magnesium treated irons (Fig. 6). The nodular form is found in annealed irons in which the cementite has decomposed at 800-950°C. Thus we have:

Figure 4. Coarse graphite flakes. Matrix unetched (x 60)

Figure 5. Temper carbon in a malleable iron; ferrite crystals etched (x 100)

Figure 6. Enlarged view of graphite spheroid. Polarised light (x 600)

Figure 7. Hypo-eutectic white cast iron, cementite and pearlite(black) (x 100) BH =100

Figure 8. Hyper-eutectic white cast iron (x 100). White primary crystals of cementite in eutectic (cementite and pearlite)

Figure 9. Grey iron. High duty; pearlite and graphite (x 200)

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Pearlite + cementite (i.e. eutectic cementite in hypoeutectic irons (Fig. 7) and primary andeutectic cementite in hypereutectic irons (Fig. 8)

white, hard, unmachinable.

Cementite + graphite + pearlite

mottled, difficult to machine.

Graphite + pearlite (Fig. 9)

grey, machinable, high strength.

Graphite + pearlite + ferrite (Fig. 4 from the Cast Iron article) grey, soft, weaker, grey, very soft, easily machined.

Graphite + ferrite

The ferrite is of course much less pure than that in carbon steels.

Phosphide eutectic Most cast irons contain phosphorus in amounts varying from 0,03 to 1,5%, consequently another micro-constituent is frequently present in the structure, in addition to those phases mentioned above. It occurs in white irons as a laminated constituent (ternary eutectic), consisting of: Iron, 91,19%

Ferrite (with a little phosphorus).

Carbon, 1,92%

Cementite, Fe3C.

Phosphorus, 6,89 %

Iron phosphide, Fe3P.

The melting-point is in the region of 960°C, consequently it is the last constituent to solidify and forms islands in the interstices of the dendrites. Although this constituent is very brittle it does not unduly weaken the iron when in small amounts (up to 1%) due to the fact that continuous cells are not formed round the grains. The structure is illustrated in Fig. 4 from the Cast Iron article which shows the structure of the phosphide eutectic, together with graphite, ferrite and pearlite. Phosphorus will thus form this additional constituent in any of the "grouped" structures already discussed.

6.3. High-strenght iron Abstract: It has been shown that the structures of grey cast irons are similar to those of ordinary steels but with the addition of graphite flakes which break up the continuity of the iron. Thus with a totally pearlitic structure cast iron should approach in tensile strength and toughness the properties of a 0,95% carbon normalised steel; the limiting factor being the shape and distribution of the graphite and fineness of the pearlite.

It has been shown that the structures of grey cast irons are similar to those of ordinary steels but with the addition of graphite flakes which break up the continuity of the iron. Thus with a totally pearlitic structure cast iron should approach in tensile strength and toughness the properties of a 0.95%. carbon normalised steel; the limiting factor being the shape and distribution of the graphite and fineness of the pearlite (Fig. 9 from the article Relation between CE structure and mechanical

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properties). Such irons have tensile strengths of up to 370 MPa. Modification of the micro-structure and properties of cast iron can be brought about by: 1. The use of special melting and casting technique. 2. The addition of alloying elements. 3. Heat-treatment, particularly of white iron.

1. High-duty irons due to casting technique The gradual introduction of so-called semi-steel during 1914-18 marked the real commencement in improved properties. It is made by adding to the cupola steel scrap which slightly reduces the carbon content and in particular the amount of free graphite together with the production of a pearlitic matrix. Other methods consist of superheating the molten metals in a separate furnace, whereby the graphite is greatly refined. Alternatively, an iron which would normally cast white can be graphitised by inoculation with ferro silicon (75% Si), sometimes with addition strontium in the ladle to give strength of 370 MPa.

2. Addition of alloying elements The most common of the special elements added to cast iron are nickel, chromium, copper and molybdenum. Nickel tends to produce grey iron, in which respect it is less powerful than silicon. Consequently in castings of widely varying section the silicon can be reduced slightly and nickel added to prevent chilling in the thin sections, but still retaining a close structure in the thick ones. On the other hand, chromium, by forming carbides, acts in the opposite way to nickel, but at the same time it exerts a grain refining action. These elements, singly or together, are commonly found in motor cylinder irons. Molybdenum strengthens the matrix by promoting a fine pearlite, but it is used preferably with other elements such as nickel to produce acicular structures. A rough classification of the types of alloy iron is:

1. Pearlitic Irons 0,5-2% nickel (chromium up to 0,8% and molybdenum up to 0,6%). Used for many general castings. The addition of tin in amounts up to 0,1% promotes a fully pearlitic matrix. High carbon Ni-Cr-Mo cast iron is useful for resisting thermal shock in applications such as die-casting moulds and brake-drums. The nickel and chromium give the desired closeness of grain and molybdenum helps to strengthen the matrix. The considerable graphite reduces the tendency to "crazy crack". Chromium (0,6)molybdenum (0,6) irons are useful for engine liners, press sleeves, dies, etc., where wear resistance in relatively heavy sections is important. Cast iron with 1 % each of chromium and molybdenum is used for piston-ring pots which are heat-treated to give a high transverse breaking strength coupled with a high elasticity value.

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2. Acicular Irons. Carbon 2,9-3,2, nickel 1,5-2,0, molybdenum 0,3-0,6%. Copper can replace nickel up to 1-5%. This rigid, high-strength, shock-resisting material is used for diesel crankshafts, gears and machine columns. With the correct amounts of nickel and molybdenum correlated with the cooling rate of a particular casting the pearlitic change point can be suppressed and an acicular intermediate constituent (ferrite needles in austenitic matrix) can be produced with high mechanical properties. Acicular cast iron is very much tougher than any of the pearlitic cast irons of lower strength. The tensile strength of acicular cast iron with a carbon content of about 3,0% will vary from 380 to 540 MPa but these figures can be maintained in quite large sections. Phosphorus should not exceed about 0,15% in the presence of molybdenum, otherwise a compound is formed which impoverishes the matrix of molybdenum. Quite large variations in silicon content can be tolerated, but chromium in excess of 0,4% is harmful. The structure changes rapidly at 600-750°C and these irons should not be used at temperatures greater than 300°C.

3. Martensitic Irons. 5-7% nickel with other elements. Very hard irons used for resisting abrasion (Fig. 1), e.g. metal working rolls.

4. Austenitic Irons. Non-magnetic, with 11-33% nickel but below 20% it is necessary to add about 6% copper or 6% manganese to maintain fully austenitic structures e.g. Nomag irons contain 11% Ni with 6% Mn. These have a good resistance to corrosion and heat, e.g. Ni-Resist. The outstanding characteristics of the austenitic cast irons, as compared with ordinary cast iron, are: a) resistance to corrosion; b) marked resistance to heat; c) non-magnetic, with suitable compositions; d) a high electrical resistance coupled with a low temperature coefficient of resistance; e) a high coefficient of thermal expansion; f) no change points.

5. Spheroidal graphite cast iron. The production of spheroidal graphite as in Fig. 2 in the as-cast state is an outstanding development of a new iron, initially due to the use of cerium by Morrogh (BCIRA, 1946 BP 645862) and later, magnesium by the International Nickel Co. (1947 BP 630.070). The use of magnesium, to give 0,04-0,06% residual content proved to be the more adaptable and economic of the two processes. The production of spheroidal structure is prevented, however, by certain trace elements, e.g. 0,1 Ti, 0,009 Pb, 0,003 Bi, 0,004% Sb, but their effect can be eliminated by 0,005-0,01% cerium. For most raw materials the combined use of cerium and magnesium followed by ferro-silicon as an inoculent is used to produce spheroidal graphite iron.

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Remelting causes a reversion to flake graphite due to loss of magnesium. Magnesium treatment desulphurises the iron to below 0,02% before alloying with the iron, and for economic reasons the sulphur content should be as low as possible. The SG iron can be used with a pearlite matrix or ferrite after a short annealing or with an acicular or austenitic matrix when suitably alloyed. The stress strain curve is similar to that of steel, with measurable elongation. The ferrite grade of SG iron has a strength of 370 MPa with 17% El whereas a normalised pearlitic SG iron has a strength of 700 MPa with a minimum of 2% El. The strength can be increased to 925 MPa by special heat treatment or by the addition of alloying elements. Damping capacity is lower but shock, heat and growth resistance and weldability are higher than for flake graphite iron. SG iron can, therefore, compete successfully with malleable iron for thick sections, cast steel and alloy flake graphite cast iron. SG cast irons are not so section sensitive as normal iron, e.g. a variation of 25-150 mm section causes grey iron to change from 278 to 154 MPa whereas a SG iron would change from 664 to 587 MPa. A new iron contains fine vermicular graphite similar but finer than undercooled graphite. It has a worm-like form which enables high strengths to be obtained with 2-3% El. Very precise production control is necessary and this limits commercial production at the moment. The sulphur content must be below 0,002% and casting must be cooled rapidly.

Figure 1. Martensitic iron (Ni-hard). Cementite (white masses) in martensite austenite matrix (x 200) BH = 700

Figure 2. Spheroidal cast iron. Spheroidal graphite in pearlite matrix (x 200)

Stress relief of grey cast iron Stress is completely removed at 650°C, but grain growth commences at 550°C and is serious at 600°C. Current practice is to heat slowly to 475-500°C, hold at temperature for 1 hour per 25 mm section and cool in furnace to 300°C.

6.4. Classification of Cast Iron Abstract: The term cast iron, like the term steel, identifies à large family of ferrous alloys. Cast irons are multicomponent ferrous alloys. They contain major (iron, carbon, silicon), minor (0.01%) elements. Cast iron has higher carbon and silicon contents than steel. Because of the higher carbon content, the structure of cast iron, as opposed to that of steel, exhibits a rich carbon phase.

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Depending primarily în composition, cooling rate ànd melt treatment, cast iron can solidify according to the thermodynamically metastable Fe-Fe3C system or the stable Fe-Gr system.

The term cast iron, like the term steel, identifies a large family of ferrous alloys. Cast irons are multicomponent ferrous alloys. They contain major (iron, carbon, silicon), minor (0.01%) elements. Cast iron has higher carbon and silicon contents than steel. Because of the higher carbon content, the structure of cast iron, as opposed to that of steel, exhibits a rich carbon phase. Depending primarily on composition, cooling rate and melt treatment, cast iron can solidify according to the thermodynamically metastable Fe-Fe3C system or the stable Fe-Gr system. When the metastable path is followed, the rich carbon phase in the eutectic is the iron carbide; when the stable solidification path is followed, the rich carbon phase is graphite. Referring only to the binary Fe-Fe3C or Fe-Gr system, cast iron can be defined as an iron-carbon alloy with more than 2% C. Important notice is that silicon and other alloying elements may considerably change the maximum solubility of carbon in austenite (g). Therefore, in exceptional cases, alloys with less than 2% C can solidify with a eutectic structure and therefore still belong to the family of cast iron. The formation of stable or metastable eutectic is a function of many factors including the nucleation potential of the liquid, chemical composition, and cooling rate. The first two factors determine the graphitization potential of the iron. A high graphitization potential will result in irons with graphite as the rich carbon phase, while a low graphitization potential will result in irons with iron carbide. The two basic types of eutectics - the stable austenite-graphite or the metastable austenite-iron carbide (Fe3C) - have wide differences in their mechanical properties, such as strength, hardness, toughness, and ductility. Therefore, the basic scope of the metallurgical processing of cast iron is to manipulate the type, amount, and morphology of the eutectic in order to achieve the desired mechanical properties.

Classification Historically, the first classification of cast iron was based on its fracture. Two types of iron were initially recognised: • •

White iron: Exhibits a white, crystalline fracture surface because fracture occurs along the iron carbide plates; it is the result of metastable solidification (Fe3C eutectic) Gray iron: Exhibits a gray fracture surface because fracture occurs along the graphite plates (flakes); it is the result of stable solidification (Gr eutectic).

With the advent of metallography, and as the body of knowledge pertinent to cast iron increased, other classifications based on microstructural features became possible:

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• •

Graphite shape: Lamellar (flake) graphite (FG), spheroidal (nodular) graphite (SG), compacted (vermicular) graphite (CG), and temper graphite (TG); temper graphite results from ? solid-state reaction (malleabilization.) Matrix: Ferritic, pearlitic, austenitic, martensitic, bainitic (austempered).

This classification is seldom used by the floor foundryman. The most widely used terminology is the commercial one. A first division can be made in two categories: • •

Common cast irons: For general-purpose applications, they are unalloyed or low alloyed Special cast irons: For special applications, generally high alloyed.

The correspondence between commercial and microstructural classification, as well as the final processing stage in obtaining common cast irons, is given in Fig. 2. Special cast irons differ from the common cast irons mainly in the higher content of alloying elements (>3%), which promote microstructures having special properties for elevated-temperature applications, corrosion resistance, and wear resistance. A classification of the main types of special cast irons is shown in Fig. 1.

Fig. 1. Classification of special high - alloy cast iron

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Fig.2. Basic microstructures and processing for obtaining common commercial cast irons

6.5. Nodular Ductile Iron Abstract: Special cast irons differ from the common cast irons mainly in the higher content of alloying elements which promote microstructures having special properties for elevated-temperature applications, corrosion resistance, and wear resistance. This article explains how chemical composition, cooling rate, liquid treatment and heat treatment influencing on a structure and the expected mechanical properties of cast irons.

Historically, the first classification of cast iron was based on its fracture. Two types of iron were initially recognized:

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• •

White iron: Exhibits a white, crystalline fracture surface because fracture occurs along the iron carbide plates; it is the result of metastable solidification (Fe-C eutectic) Gray iron: Exhibits a gray fracture surface because fracture occurs along the graphite plates (flakes); it is the result of stable solidification (Gr eutectic)

Special cast irons differ from the common cast irons mainly in the higher content of alloying elements which promote microstructures having special properties for elevated-temperature applications, corrosion resistance, and wear resistance. The goal of the metallurgist is to design a process that will produce a structure that will yield the expected mechanical properties. This requires knowledge of the structure-properties correlation for the particular alloy under consideration as well as of the factors affecting the structure. When discussing the metallurgy of cast iron, the main factors of influence on the structure that one needs to address are: • • • •

Chemical composition Cooling rate Liquid treatment Heat treatment.

In addition, the following aspects of combined carbon in cast irons should also be considered: • • •

•

In the original cooling or through subsequent heat treatment, a matrix can be internally decarbonized or carburized by depositing graphite on existing sites or by dissolving carbon from them. Depending on the silicon content and the cooling rate, the pearlite in iron can vary in carbon content. This is a ternary system, and the carbon content of pearlite can be as low as 0.50% with 2.5% Si. The conventionally measured hardness of graphitic irons is influenced by the graphite, especially in gray iron. Martensite micro hardness may be as high as 66 HRC, but measures as low as 54 HRC conventionally in gray iron (58 HRC in ductile). The critical temperature of iron is influenced (raised) by silicon content, not by carbon content.

For common cast iron, the main elements of the chemical composition are carbon and silicon. High carbon content increases the amount of graphite or Fe3C. High carbon and silicon contents increase the graphitization potential of the iron as well as its castability. The manganese content varies as a function of the desired matrix. Typically, it can be as low as 0.1% for ferrule irons and as high as 1.2% for pearlitic irons, because manganese is a strong pearlite promoter. From the minor elements, phosphorus and sulfur are the most common and are always present in the composition. They can be as high as 0.15% for low-quality iron

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and are considerably less for high-quality iron, such as ductile iron or compacted graphite iron. The main effects of chemical composition to nodular (ductile) iron are similar to those described for gray iron, with quantitative differences in the extent of these effects and qualitative differences in the influence on graphite morphology. The carbon equivalent has only a mild influence on the properties and structure of ductile iron, because it affects graphite shape considerably less than in the case of gray iron. Nevertheless, to prevent excessive shrinkage, high chilling tendency, graphite flotation or a high impact transition temperature, optimum amounts of carbon and silicon must be selected. Minor elements can significantly alter the structure in terms of graphite morphology, chilling tendency, and matrix structure. Minor elements can promote the spheroidization of graphite or can have an adverse effect on graphite shape. The general influence of various elements on graphite shape. The elements in the first group - the spheroidizing elements - can change graphite shape from flake through compacted to spheroidal. The most widely used element for the production of spheroidal graphite is magnesium. The amount of residual magnesium required to produce spheroidal graphite is generally 0.03 to 0.05%. The precise level depends on the cooling rate. A higher cooling rate requires less magnesium. The amount of magnesium to be added in the iron is a function of the initial sulfur level. A residual magnesium level that is too low results in insufficient nodularity. This in turn results in a deterioration of the mechanical properties of the iron. If the magnesium content is too high, carbides are promoted. The presence of antispheroidizing minor elements may result in graphite shape deterioration, up to complete graphite degeneration. Therefore, upper limits are set on the amount of deleterious elements to be accepted in the composition of cast iron. These values can be influenced by the combination of various elements and by the presence of rare earths in the composition. Furthermore, some of these elements can be deliberately added during liquid processing in order to increase nodule count. Alloying elements have in principle the same influence on structure and properties as for gray iron. Because better graphite morphology allows more efficient use of the mechanical properties of the matrix, alloying is more common in ductile iron than in gray iron. Cooling Rate. When changing the cooling rate, effects similar to those discussed for gray iron also occur in ductile iron, but the section sensitivity of ductile iron is lower. This is because spheroidal graphite is less affected by cooling rate than flake graphite. The liquid treatment of ductile iron is more complex than that of gray iron. The two stages for the liquid treatment of ductile iron are: • •

Modification, which consists of magnesium or magnesium alloy treatment of the melt, with the purpose of changing graphite shape from flake to spheroidal. Inoculation (normally, postinoculation that is, after the magnesium treatment) to increase the nodule count. Increasing the nodule count is an

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important goal, because a higher nodule count is associated with less chilling tendency and a higher as-cast ferrite/pearlite ratio. Heat treatment is extensively applied on ductile iron because better advantage can be taken of the matrix structure than for gray iron. The heat treatments usually applied are as follows: • • • • •

Stress relieving Annealing to produce a feritic matrix Normalizing to produce a pearlitic matrix Hardening to produce tempering structures Austempering to produce a ferritic bainite.

The advantage of austempering is that it results in ductile irons with twice the tensile strength for the same toughness. Compacted graphite (CG) irons have a graphite shape intermediate between spheroidal and flake. Typically, compacted graphite looks like type IV graphite. The chemical composition effects are similar to those described for ductile iron. Carbon equivalent influences strength less obviously than for the case of gray iron, but than for ductile iron. The graphite shape is controlled, as in the case of ductile iron, through the content of minor elements. When the goal is to produce compacted graphite, it is easier from the stand point of controlling the structure to combine spheroidizing (magnesium, calcium, and/ or rare earths) and antispheroidizing (titanium and/or aluminum) elements. The cooling rate affects properties less for gray iron but more for ductile iron. In other words, CG iron is less section sensitive than gray iron. However, high cooling rates are to be avoided because of the high propensity of CG iron for chilling and high nodule count in thin sections. The usual microstructure of gray iron is a matrix of pearlite with graphite (flakes dispersed throughout). Foundry practice can be varied so that nucleation and growth of graphite flakes occur in a pattern that enhances the desired properties. The amount, size, and distribution of graphite are important. Cooling that is too rapid may produce so-called chilled iron, in which the excess carbon is found in the form of massive-carbides. Cooling at intermediate rates can produce mottled iron, in which carbon is present in the form of both primary cementite (iron carbide) and graphite. Very slow cooling of irons that contain large percentages silicon and carbon is likely to produce considerable ferrite and pearlite throughout the matrix, together with coarse graphite flakes. Flake graphite is one of seven types (shapes or forms) of graphite established in ASTM A 247. Flake graphite is subdivided into five types (patterns), which are designated by the letters A through E. Graphite size is established by comparison with an ASTM size chart, which shows the typical appearances of flakes of eight different sizes at 100x magnification.

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Type A flake graphite (random orientation) is preferred for most applications. In the intermediate flake sizes, type A flake graphite is superior to other types in certain wear applications such as the cylinders of internal combustion engines. Type B flake graphite (rosette pattern) is typical of fairly rapid cooling, such as is common with moderately thin sections (about 10 mm) and along the surfaces of thicker sections, and sometimes results from poor inoculation. The large flakes of type C flake graphite are typical of kish graphite that is formed in hypereutectic irons. These large flakes enhance resistance to thermal shock by increasing thermal conductivity and decreasing elastic modulus. On the other hand, large flakes are not conducive to good surface finishes on machined parts or to high strength or good impact resistance. The small, randomly oriented interdendritic flakes in type D flake graphite promote a fine machined finish by minimizing surface pitting, but it is difficult to obtain a pearlitic matrix with this type of graphite. Type D flake graphite may be formed near rapidly cooled surfaces or in thin sections. Frequently, such graphite is surrounded by a ferrite matrix, resulting in soft spots in the casting. Type E flake graphite is an interdendritic form, which has a preferred rather than a random orientation. Unlike type D graphite, type 6 graphite can be associated with a pearlitic matrix and thus can produce a casting whose wear properties are as good as those of a casting containing only type A graphite in a pearlitic matrix. Solidification of Gray Iron. In a hypereutectic gray iron, solidification begins with the precipitation of kish graphite in the melt. Kish grows as large, straight, undistorted flakes or as very thick, lumpy flakes that tend to rise to the surface of the melt because of their Sow relative density. When the temperature has been lowered sufficiently, the remaining liquid solidifies as a eutectic structure of austenite and graphite. Generally, eutectic graphite is finer than kish graphite.

6.6. Standard Terminology Related to Iron Castings Abstract: The most frequently used terms and definitions, related to iron castings, according to ASTM A 644 - 85 are listed bellow.

Austenitize - to convert the matrix of a ferrous alloy to austenite by heating above the transformation temperature. Batch - the component raw materials properly weighed, proportioned, and mixed for delivery to a processing unit. Also, the product output from a processing unit in which there is essentially no product output until all component materials are charged and processed. Carbide, primary - carbide precipitated in cast iron during solidification. Cast iron - a generic term for a series of alloys primarily of iron, carbon, and silicon in which the carbon I in excess of the amount which can be retained in solid solution in austenite at the eutectic temperature. Cementite - a very hard and brittle compound of iron and carbon corresponding to the empirical formula Fe3C, commonly known as iron carbide. Cementite, primary - cementite precipitated in cast iron during solidification. Also

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known as primary carbide. Chilled iron - a cast iron that would normally solidify as a gray cast iron which is purposely caused to solidify as white cast iron locally or entirely by accelerated cooling caused by contact with a metal surface, that is, a chill. Direct reduced iron - iron ores that have been reduced to essentially metallic iron by heat and reducing agents, but without melting, and processed into suitable shapes for use as a charge material in a melting operation. Dual metal - two metals of different composition that are fusion bonded at all interfacial surfaces by casting metal of one composition against metal of a second composition. Ductile iron - a cast iron that has been treated in the liquid state so as to cause substantially all of its graphitic carbon to occur as spheroids or nodules in the as-cast condition. Ferritize - to increase the quantity of ferrite in the matrix of a ferrous casting through an appropriate heat treatment. Ferritizing anneal - the process of producing a predominantly ferritic matrix in cast iron through an appropriate heat treatment. Graphite, compacted - a graphite shape that is intermediate between flake graphite and nodular graphite that typically appears in a polished section as thick flakes with blunt ends. Graphite, flake - an irregularly shaped particle of graphite, usually appearing in a polished section as curved plates, such as found in gray cast irons. Graphite, nodular - spheroidal shaped graphite typically found in ductile irons and compact clusters of graphite typically found in malleable irons. Graphite, primary - graphite precipitated in cast iron during solidification. Graphite rosette - arrangement of graphite flakes in which the flakes extend radially from centers of crystallization in gray cast iron. Graphite, spheroidal - spheroidal shaped graphite having a polycrystalline radial structure, usually found in ductile iron and to a controlled, limited extent in compacted graphite iron. Graphitize - to precipitate graphite in an iron-carbon alloy. Gray iron - cast iron that has a relatively large proportion of the graphitic carbon present in the form of flake graphite. The metal has a gray fracture. Heat - the total molten metal output from a single heating in a batch melting process or the total metal output from essentially a single heating in a continuous melting operation using basically constant charge and processing conditions and targeted at a fixed metal chemistry at the furnace spout. A heat can also be defined as a fixed time period for a continuous melting operation provided that it is shorter than the time period covered by the above definition. Inoculating alloy - an alloy added to molten iron for the principle purpose of nucleating a primary phase such as graphite. Inoculating alloys are frequently used to avoid the formation of primary carbide by enhancing the nucleation of graphite. Malleable, ferritic - a ferrous alloy that is cast as white is converted by an appropriate heat treatment to a microstructure of temper carbon embedded in a ferritic matrix essentially free of pearlite and carbide. Malleable iron - a cast iron of such composition that it solidifies as white iron, which upon proper heat treatment is converted to a metallic matrix with nodules of temper carbon. Malleable, pearlitic - a ferrous alloy that is cast as white iron but which is converted by an appropriate heat treatment to a microstructure of temper carbon embedded in a matrix containing a controlled quantity, form, and distribution of pearlite or tempered martensite. Malleableize - to convert white iron into malleable iron through an appropriate

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graphitizing heat treatment. Melt - the total molten metal produced in a single heat. Merchant pig iron - pig iron produced for commercial sale to foundries. Mottled iron - a cast iron containing a mixed structure of gray iron and white iron of variable proportions. The fracture has a mottled appearance. Nodular graphite - graphite in the form of nodules or spheroids in iron castings. Nodularity - the volumetric proportion of spheroidal or nodular graphite to total graphite in a ductile iron or a compacted graphite iron matrix. Nodularity, degree of - the volumetric proportion of spheroidal or nodular graphite to total graphite in a ductile iron matrix. Nodulizing alloy - an alloy added to molten iron for the primary purpose of causing the formation of spheroidal graphite during solidification. Pig iron - the high carbon iron product obtained by the reduction of iron ores, typically in a blast furnace or an electric furnace, and cast into uniform shapes having physical and chemical characteristics suitable for end as foundry melting stock. Sample - one or more portions of a liquid or solid material taken in an unbiased manner from a batch, heat, lot or process stream to be representative of the whole, for subsequent testing to determine the chemical, physical, mechanical, or other quality characteristics of the material, or combination thereof. Temper carbon - compact aggregates or nodules of graphite found in malleable iron as a result of heat treatment. Test bar - a bar-shaped coupon that is tested with or without subsequent preparation for the determination of physical or mechanical properties. Test coupon - specially designed casting, or portion thereof, that is used to provide a representative sample of the iron from which it was cast. Test lug - a sample produced as an appendage on a casting, that may be removed and tested to qualify the casting or the iron which it was produced. Test specimen - a test object, suitably prepared from a sample, for evaluation of the chemical, physical, mechanical, or metallurgical quality of the sample. Treated iron - molten cast iron to which all basic alloys and nodulizing alloys have been added but not necessarily all inoculating alloy additions. White iron - cast iron in which substantially all of the carbon is in solution and in the combined form. The metal has a white fracture.

6.7. High-Alloy White Irons Abstract: High-alloy white cast irons are an important group of materials whose production must be considered separately from that of ordinary types of cast irons. In these cast iron alloys, the alloy content is well above 4%, and consequently they cannot be produced by ladle additions to irons of otherwise standard compositions. They are usually produced in foundries specially equipped to produce highly alloyed irons.

High-alloy white cast irons are an important group of materials whose production must be considered separately from that of ordinary types of cast irons. In these cast iron alloys, the alloy content is well above 4%, and consequently they cannot be produced by ladle additions to irons of otherwise standard compositions. They are usually produced in foundries specially equipped to produce highly alloyed irons. The high-alloy white irons are primarily used for abrasion-resistant applications and are readily cast into the parts needed in machinery for crushing, grinding, and handling of abrasive materials. The chromium content of high-alloy white irons also

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enhances their corrosion-resistant properties. The large volume fraction of primary and/or eutectic carbides in their microstructures provides the high hardness needed for crushing and grinding other materials. The metallic matrix supporting the carbide phase in these irons can be adjusted by alloy content and heat treatment to develop the proper balance between the resistance to abrasion and the toughness needed to withstand repeated impact. While low-alloy white iron castings, which have alloy content below 4%, develop hardnesses in the range of 350 to 550 HB, the high-alloy irons range in hardness is from 450 to 800 HB. Specification ASTM A 532 covers the composition and hardness of the abrasionresistant white iron grades. Many castings are ordered according to these specifications. However, a large number of castings are produced with composition modifications for specific applications. It is most desirable that the designer, metallurgist, and foundry man work together to specify the composition, heat treatment, and foundry practice to develop the most suitable alloy and casting design for a specific application. The high-alloy white cast irons fall into two major groups: • •

Nickel-chromium white irons, which are low-chromium alloys containing 3 to 5% Ni and 1 to 4% Cr, with one alloy modification that contains 7 to 11% Cr, Chromium-molybdenum irons containing 11 to 23% Cr, up to 3% Mo and often additionally alloyed with nickel or copper.

A third group comprises the 25% or 28% Cr white irons, which may contain other alloying additions of molybdenum and/or nickel up to 1.5%. The nickel-chromium irons are also commonly identified as Ni-Hard types 1 to 4.

Nickel-Chromium White Irons The oldest group of high-alloy irons of industrial importance, the nickel-chromium white irons, or Ni-Hard irons, have been produced for more than 50 years and are very cost-effective materials for crushing and grinding. In these martensitic white irons, nickel is the primary alloying element because at levels of 3 to 5% it is effective in suppressing the transformation of the austenite matrix to pearlite, thus ensuring that a hard martensitic structure (usually containing significant amounts of retained austenite) will develop upon cooling in the mold. Chromium is included in these alloys, at levels from 1.4 to 4%, to ensure that the irons solidify carbidic, that is, to counteract the graphitizing effect of nickel. The optimum composition of a nickel-chromium white iron alloy depends on the properties required for the service conditions and the dimensions and weight of the casting. Abrasion resistance is generally function of the bulk hardness and the volume of carbide in the microstructure. When abrasion resistance is the principal requirement and resistance to impact loading is secondary, alloys having high carbon contents, ASTM A 532 class I type A (Ni-Hard 1), are recommended. When conditions of repeated impact are anticipated, the lower carbon alloys, class I type B

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( Ni-Hard 2) are recommended because they have less carbide and, therefore, greater toughness. A special grade, class J type C, has been developed for producing grinding balls and slugs. Here, the nickel-chromium alloy composition has been adapted for chill casting and specialized sand casting processes. The Class I type D (Ni-Hard 4) alloy is a modified nickel-chromium iron that contains higher levels of chromium, ranging from 7 to 11%, and increased levels of nickel, ranging from 5 to 7%. Carbon is varied according to the properties needed for the intended service. Carbon contents in the range of 3.2 to 3.6% are prescribed when maximum abrasion resistance is desired. When impact loading is expected, carbon content should be held in the range of 2.7 to 3.2%. Nickel content increases with section size or cooling time of the casting to inhibit pearlitic transformation. For castings of 38 to 50 mm thick, 3.4 to 4.2% Ni is sufficient to suppress pearlite formation upon mold cooling. Heavier sections may require nickel levels up to 5.5% to avoid the formation of pearlite. It is important to limit nickel content to the level needed for control of pearlite; excess nickel increases the amount of retained austenite and lowers hardness. Silicon is needed for two reasons. A minimum amount of silicon is necessary to improve fluidity of the melt and to produce a fluid slag, but of equal importance is its effect on as-cast hardness. Increased levels of silicon, in the range of 1 to 1.5%, have been found to increase the amount of martensite and the resulting hardness. Late additions of ferrosilicon (0.2% as 75% Si grade ferrosilicon) have been reported to increase toughness. It is important to note that higher silicon contents can promote pearlite and may increase the nickel requirement. Chromium is primarily added to offset the graphitizing effects of nickel and silicon in types A, B, and C alloys, ranges from 1.4 to 3.5%. Chromium content must be increased with increasing section size. In type D alloy, chromium levels range from 7 to 11% (typically 9%) for the purpose of producing eutectic carbides of the M7C3 chromium carbide type, which are harder and less deleterious to toughness. Manganese is typically held to a maximum of 0.8% even though 1.3% maximum is allowed according to ASTM A 532 specification. While it provides increased hardenability to avoid pearlite formation, it is a more potent austenite stabilizer than nickel, and promotes increased amounts of retained austenite and lower as-cast hardness. For this reason, higher manganese levels are undesirable. When considering the nickel content required to avoid pearlite in a given casting, the level of manganese present should be a factor. Copper increases both hardenability and the retention of austenite and therefore must be controlled for the same reason that manganese must be limited. Copper should be treated as a nickel substitute and, when properly included in the calculation of the amount of nickel required to inhibit pearlite, it reduces the nickel requirement. Molybdenum is a potent hardenability agent in these alloys and is used in heavysection castings to augment hardenability and inhibit pearlite.

High-Chromium White Irons

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The high-chromium white irons have excellent abrasion resistance and are used effectively in slurry pumps, brick molds, coal-grinding mills, shot-blasting equipment, and components for quarrying, hard-rock mining, and milling. In some applications they must also be able to withstand heavy impact loading. These alloyed white irons are recognized as providing the best combination of toughness and abrasion resistance attainable among the white cast irons. In the high-chromium irons, as with most abrasion-resistant materials, there is a trade-off between wear resistance and toughness. By varying composition and heat treatment, these properties can be adjusted to meet the needs of most abrasive applications. Specification ASTM A 532 covers the compositions and hardnesses of two general classes of the high-chromium irons. The chromium-molybdenum irons (Class II of ASTM A532) contain 11 to 23% Cr and up to 3.5% Mo and can be supplied either as-cast with an austenitic or austenitic-martensitic matrix, or heattreated with a martensitic matrix microstructure for maximum abrasion resistance and toughness. They are usually considered the hardest of all grades of white cast irons. Compared to the lower-alloy nickel-chromium white irons, the eutectic carbides are harder and can be heat-treated to achieve castings of higher hardness. Molybdenum, as well as nickel and copper when needed, is added to prevent pearlite and to ensure maximum hardness. The high-chromium irons (class III of ASTM A 532) represent the oldest grade of high-chromium irons, with the earliest patents dating back to 1917. These generalpurpose irons, also called 25% Cr and 28% Cr irons, contain 23 to 28% Cr with up to 1.5% Mo. To prevent pearlite and attain maximum hardness, molybdenum is added in all but the lightest-cast sections. Alloying with nickel and copper up to 1% is also practiced. Although the maximum attainable hardness is not as high as in the class II chromium-molybdenum white irons, these alloys are selected when resistance to corrosion is also desired.

6.8. Specifications for Ductile Iron Abstract: Standard specifications for engineering grades of ductile iron castings classify the grades according to the tensile strength of a test bar cut from a prescribed test casting. The International Standards Organization (ISO) specification ISO 1083:1976 and most national specifications also specify the ductility in terms of percentage of elongation and the 0.2% proof strength or offset yield strength. The impact values of those grades with the highest ductility are frequently specified in the ISO, UK, and German specifications, and a guide to microstructure is included in most specifications. Hardness is usually specified, but is only mandatory in SAE J434C.

Standard specifications for engineering grades of ductile iron castings classify the grades according to the tensile strength of a test bar cut from a prescribed test casting. The International Standards Organization (ISO) specification ISO 1083:1976 and most national specifications also specify the ductility in terms of percentage of elongation and the 0.2% proof strength or offset yield strength. The impact values of those grades with the highest ductility are frequently specified in the ISO, UK, and German specifications, and a guide to microstructure is included in most specifications. Hardness is usually specified, but is only mandatory in SAE J434C.

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The actual values of properties to be expected from good-quality ductile irons produced to meet any given specified grade will normally cover a range that more than satisfies the requirements of the specification. Specifications for the highest-strength grades usually mention the possibility of hardened-and-tempered structures, but for the most recently reported austempered ductile irons, which have the highest combinations of tensile strength and ductility, there are as yet only tentative unofficial specifications.

Factors That Affect Properties Graphite Structures. The amount and form of the graphite in ductile iron are determined during solidification and cannot be altered by subsequent heat treatment. All of the mechanical and physical properties of this class of materials are a result of the graphite being substantially or wholly in the spheroidal nodular shape, and any departure from this shape in a proportion of the graphite will cause some deviation from these properties. It is common to attempt to produce greater than 90% of the graphite in this form (>90% nodularity), although structures between 80 and 100% nodularity are sometimes acceptable. All properties relating to strength and ductility decrease as the proportion of nonnodular graphite increases, and those relating to failure, such as tensile strength and fatigue strength, are more affected by small amounts of such graphite than properties not involving failure, such as proof strength. The form of non-nodular graphite is important because thin flakes of graphite with sharp edges have a more adverse effect on strength properties than compacted forms of graphite with rounded ends. For this reason, visual estimates of percentage of nodularity are only a rough guide to properties. Graphite form also affects modulus of elasticity, which can be measured by resonant frequency and ultrasonic velocity measurements, and such measurements are therefore often a better guide to nodularity and its effects on other properties. A low percentage of nodularity also lowers impact energy in the ductile condition, reduces fatigue strength, increases damping capacity, increases thermal conductivity, and reduces electrical resistivity. Graphite Amount. As the amount of graphite increases, there is a relatively small decrease in strength and elongation, in modulus of elasticity, and in density. In general, these effects are small compared with the effects of other variables because the carbon equivalent content of spheroidal graphite iron is not a major variable and is generally maintained close to the eutectic value. Matrix Structure. The principal factor in determining the different grades of ductile iron in the specifications is the matrix structure. In the as-cast condition, the matrix will consist of varying proportions of pearlite and ferrite, and as the amount of pearlite increases, the strength and hardness of the iron also increase. Ductility and impact properties are principally determined by the proportions of ferrite and pearlite in the matrix. The matrix structure can be changed by heat treatments, and those most often carried out are annealing to produce a fully ferritic matrix and normalizing to produce a substantially pearlitic matrix. In general, annealing produces a more ductile matrix with a lower impact transition temperature than is obtained in as-cast ferritic irons.

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Normalizing produces a higher tensile strength with a higher amount of elongation than is obtained in fully pearlitic as-cast irons. Section Size. As section size decreases, the solidification and cooling rates in the mold increase. This results in a fine-grain structure that can be annealed more rapidly. In thinner sections, however, carbides may be present, which will increase hardness, decrease machinability, and lead to brittleness. To achieve soft ductile structures in thin sections, heavy inoculation, probably at a late stage, is desirable to promote graphite formation through a high nodule number. As the section size increases, the nodule number decreases, and micro segregation becomes more pronounced. This results in a large nodule size, a reduction in the proportion of as-cast ferrite, and increasing resistance to the formation of a fully ferritic structure upon annealing. In heavier sections, minor elements, especially carbide formers such as chromium, titanium, and vanadium, segregate to produce a segregation pattern that reduces ductility, toughness, and strength. The effect on proof strength is much less pronounced. It is important for heavy sections to be well inoculated and to be made from a composition low in trace elements. Composition. In addition to the effects of elements in stabilizing pearlite or retarding transformation (which facilitates heat treatment to change matrix structure and properties), certain aspects of composition have an important influence on some properties. Silicon hardens and strengthens ferrite and raises its impact transition temperature; therefore, silicon content should be kept as low as practical, even below 2%, to achieve maximum ductility and toughness. Nickel also strengthens ferrite, but has much less effect than silicon in reducing ductility. When producing as-cast grades of iron requiring fairly high ductility and strength such as ISO Grade 500-7, it is necessary to keep silicon low to obtain high ductility, but it may also be necessary to add some nickel to strengthen the iron sufficiently to obtain the required tensile strength. Almost all elements present in trace amounts combine to reduce ferrite formation, and high-purity charges must be used for irons to be produced in the ferritic as-cast condition. Similarly, all carbide-forming elements and manganese must be kept low to achieve maximum ductility and low hardness. Silicon is added to avoid carbides and to promote ferrite as-cast in thin sections. The electrical, magnetic, and thermal properties of ductile irons are influenced by the composition of the matrix. In general, as the amount of alloying elements increases, resistivity and the magnetic hardness of the material increase and thermal conductivity decreases.

Heat Treatment of Ductile Iron The first stage of most heat treatments designed to change the structure and properties of ductile iron consists of heating to, and holding at, a temperature between 850 and 950oC for about 1h plus 1h for each 25 mm of section thickness to homogenize the iron. When carbides are present in the structure, the temperature should be approximately 900 to 950oC, which decomposes the carbides prior to subsequent stages of heat treatment. The time may have to be extended to 6 or 8h if carbide-stabilizing elements are present. In castings of complex shape in which

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stresses could be produced by nonuniform heating, the initial heating to 600oC should be slow, preferably 50 to 100oC per hour. To prevent scaling and surface decarburization during this stage of treatment, it is recommended that a nonoxidizing furnace temperature be maintained using a sealed furnace; a controlled atmosphere may be necessary. Care must also be taken to support castings susceptible to distortion and to avoid packing so that castings are not distorted by the weight of other castings placed above them. The most important heat treatments and their purposes are: • • • • • •

Stress relieving - a low-temperature treatment, to reduce or relieve internal stresses remaining after casting Annealing - to improve ductility and toughness, to reduce hardness and to remove carbides Normalizing - to improve strength with some ductility Hardening and tempering - to increase hardness or to give improved strength and higher proof stress ratio Austempering - to yield bainitic structures of high strength, with some ductility and good wear resistance Surface hardening - by induction, flame, or laser to produce a local wearresistant hard surface

6.9. Malleable cast iron Abstract: Malleable cast iron is a heat-treated iron-carbon alloy, which solidifies in the as-cast condition with a graphite-free structure, i.e. the total carbon content is present in the cementite form (Fe3C). Two groups of malleable cast iron are specified (whiteheart and blackheart malleable cast iron), differentiated by chemical composition, temperature and time cycles of the annealing process, the annealing atmosphere and the properties and microstructure resulting therefrom.

Malleable cast iron is a heat-treated iron-carbon alloy, which solidifies in the as-cast condition with a graphite-free structure, i.e. the total carbon content is present in the cementite form (Fe3C). Two groups of malleable cast iron are specified, differentiated by chemical composition, temperature and time cycles of the annealing process, the annealing atmosphere and the properties and microstructure resulting therefrom.

Whiteheart malleable cast iron The microstructure of whiteheart malleable cast iron depends on section size. Small sections contain pearlite and temper carbon in ferritic substrate. In the large sections exists three different zones: • • •

surface zone which contains pure ferrite, intermediate zone which has pearlite, ferrite and temper carbon, core zone containing pearlite, temper carbon and ferritic inclusions.

The microstructure shall not contain flake graphite.

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Blackheart and pearlitic malleable cast iron The microstructure of blackheart malleable cast iron has a matrix essentially of ferrite. The microstructure of pearlitic malleable cast iron has a matrix, according to the grade specified, of pearlite or other transformation products of austenite. Graphite is present in the form of temper carbon nodules. The microstructure shall not contain flake graphite.

Malleable cast iron designation system The designation according to ISO 5922 (1981) of malleable cast iron consists of one letter designating the type of iron, two figures designating the tensile strength and two figures designating the minimum elongation. a. Letters designating the type of malleable cast iron can be: o W for whiteheart malleable cast iron, o B for blackheart malleable cast iron, o P for peariitic malleable cast iron. This letter is followed by a space. b. The first two figures designating the minimum tensile strength, in Newtons per square millimetre, of a 12 mm diameter test piece, divided by ten. For example if the minimum tensile strength were 350 N/mm², the designation would be 35. c. The next two figures designating the minimum elongation (L0 = 3d) as a percentage of a 12 mm diameter test piece. A nought (0) shall be the first figure when the value is less than 10%, for example if the minimum elongation is 4%, the designation is 04, and if the minimum elongation is 12%, the designation is 12. For example: The designation of a whiteheart malleable cast iron having a minimum tensile strength of 400 N/mm² and minimum elongation of 5% when measured on a 12 mm diameter test piece, would be W 40-05.

Chemical composition of malleable iron The chemical composition of malleable iron generally conforms to the ranges given in the Table 1. Small amounts of chromium (0.01 to 0.03%), boron (0.0020%), copper (≤ 1.0%), nickel (0.5 to 0.8%), and molybdenum (0.35 to 0.5%) are also sometimes present. Table 1. Chemical composition of malleable iron Element

Composition %

Carbon

2.16-2.90

Silicon

0.90-1.90

Manganese

0.15-1.25

Sulfur

0.02-0.20

Phosphorus

0.02-0.15

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Mechanical properties of malleable iron Malleable iron, like ductile iron, possesses considerable ductility and toughness because of its combination of nodular graphite and low-carbon metallic matrix. Because of the way in which graphite is formed in malleable iron, however, the nodules are not truly spherical as they are in ductile iron but are irregularly shaped aggregates. Malleable iron and ductile iron are used for some of the applications in which ductility and toughness are important. In many cases, the choice between malleable and ductile iron is based on economy or availability rather than on properties. In certain applications, however, malleable iron has a distinct advantage. It is preferred for thin-section castings: • • • •

for for for for

parts parts parts parts

that are to be pierced, coined, or cold formed, requiring maximum machinability, that must retain good impact resistance at low temperatures, and requiring wear resistance (martensitic malleable iron only).

Ductile iron has a clear advantage where low solidification shrinkage is needed to avoid hot tears or where the section is too thick to permit solidification as white iron (Solidification as white iron throughout a section is essential to the production of malleable iron). Malleable iron castings are produced in section thicknesses ranging from about 1.5 to 100 mm and in weights from less than 0.03 to 180 kg or more. The mechanical properties of test pieces of malleable cast iron shall be in accordance with the values listed below: Table 2. Mechanical properties of whiteheart malleable cast iron Designation

Diameter of test piece mm

Tensile strength N/mm²

0,2% proof stress N/mm²

W 35-04

9 - 15

340 - 360

-

5-3

230

W 38-12

9 - 15

320 - 380

170 - 210

15 - 8

200

W 40-05

9 - 15

360 - 420

200 - 230

8-4

220

W 45-07

9 - 15

400 - 480

230 - 280

10 - 4

220

Elongation Hardness (L0 = 3d) HB % min

Table 3. Mechanical properties of blackheart and pearlitic malleable cast iron Designation

Diameter of test piece mm

Tensile strength N/mm²

0,2% proof stress N/mm²

B 30-06

12 - 15

300

-

6

150 max

B 32-12

12 - 15

320

190

12

150 max

B 35-10

12 - 15

350

200

10

150 max

P 45-06

12 - 15

450

270

6

150-200

P 50-05

12 - 15

500

300

5

160-220

P 55-04

12 - 15

550

340

4

180-230

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Elongation Hardness (L0 = 3d) HB % min

P 60-03

12 - 15

600

390

3

200-250

P 65-02

12 - 15

650

430

2

210-260

P 70-02

12 - 15

700

530

2

240-290

P 80-01

12 - 15

800

600

1

270-310

Melting Practices Melting can be accomplished by batch cold melting or by duplexing. Cold melting is done in coreless or channel-type induction furnaces, electric arc furnaces, or cupola furnaces. In duplexing, the iron is melted in a cupola or electric arc furnace, and the molten metal is transferred to a coreless or channel-type induction furnace for holding and pouring. Charge materials (foundry returns, steel scrap, ferroalloys, and, except in cupola melting, carbon) are carefully selected, and the melting operation is well controlled to produce metal having the desired composition and properties. Minor corrections in composition and pouring temperature are made in the second stage of duplex melting, but most of the process control is done in the primary melting furnace. Molds are produced in green sand, silicate CO2 bonded sand, or resin bonded sand (shell molds). Equipment ranges from highly mechanized or automated machines to that required for floor or hand molding methods, depending on the size and number of castings to be produced. In general, the technology of molding and pouring malleable iron is similar to that used to produce gray iron. Heat treating is done in high-production controlled-atmosphere continuous furnaces or batch-type furnaces, again depending on production requirements. After solidification and cooling, the metal is in a white iron state, and gates, sprues, and feeders can be easily removed from the castings by impact. This operation, called spruing, is generally performed manually with a hammer because the diversity of castings produced in the foundry makes the mechanization or automation of spruing very difficult. After spruing, the castings proceed to heat treatment, while gates and risers are returned to the melting department for reprocessing.

6.10. Heat Treating of Gray Irons: Part One Abstract: Gray Irons are a group of cast irons that form flake graphite during solidification, in contrast to the spheroidal graphite morphology of ductile irons. The flake graphite in gray irons is dispersed in a matrix with a microstructure that is determined by composition and heat treatment. The usual microstructure of gray iron is a matrix of pearlite with the graphite flakes dispersed throughout. In terms of composition, gray irons usually contain 2.5 to 4% C, 1 to 3% Si, and additions of manganese, depending on the desired microstructure (as low as 0.1% Mn in ferritic gray irons and as high as 1.2% in pearlitic). Other alloying elements include nickel, copper, molybdenum, and chromium.

Gray Irons are a group of cast irons that form flake graphite during solidification, in contrast to the spheroidal graphite morphology of ductile irons. The flake graphite in gray irons is dispersed in a matrix with a microstructure that is determined by composition and heat treatment. The usual microstructure of gray iron is a matrix of pearlite with the graphite flakes dispersed throughout. In terms of composition, gray

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irons usually contain 2.5 to 4% C, 1 to 3% Si, and additions of manganese, depending on the desired microstructure (as low as 0.1% Mn in ferritic gray irons and as high as 1.2% in pearlitic). Other alloying elements include nickel, copper, molybdenum, and chromium. The heat treatment of gray irons can considerably alter the matrix microstructure with little or no effect on the size and shape of the graphite achieved during casting. The matrix microstructures resulting from heat treatment can vary from ferritepearlite to tempered martensite. However, even though gray iron can be hardened by quenching from elevated temperatures, heat treatment is not ordinarily used commercially to increase the overall strength of gray iron castings because the strength of the as-cast metal can be increased at less cost by reducing the silicon and total carbon contents or by adding alloying elements. The most common heat treatments of gray iron are annealing and stress relieving. Chemical composition is another important parameter influencing the heat treatment of gray cast irons. Silicon, for example, decreases carbon solubility, increases the diffusion rate of carbon in austenite, and usually accelerates the various reactions during heat treating. Silicon also raises the austenitizing temperature significantly and reduces the combined carbon content (cementite volume). Manganese, in contrast, lowers the austenitizing temperature and increases hardenability. It also increases carbon solubility, slows carbon diffusion in austenite, and increases the combined carbon content. In addition, manganese alloys and stabilizes pearlitic carbide and thus increases the pearlite content.

Annealing The heat treatment most frequently applied to gray iron, with the possible exception of stress relieving, is annealing. The annealing of gray iron consists of heating the iron to a temperature high enough to soften it and/or to minimize or eliminate massive eutectic carbides, thereby improving its machinability. This heat treatment reduces mechanical properties substantially. It reduces the grade level approximately to the next lower grade: for example, the properties of a class 40 gray iron will be diminished to those of a class 30 gray iron. The degree of reduction of properties depends on the annealing temperature, the time at temperature, and the alloy composition of the iron. Gray iron is commonly subjected to one of three annealing treatments, each of which involves heating to a different temperature range. These treatments are ferritizing annealing, medium (or full) annealing, and graphitizing annealing. Ferritizing Annealing. For an unalloyed or low-alloy cast iron of normal composition, when the only result desired is the conversion of pearlitic carbide to ferrite and graphite for improved machinability, it is generally unnecessary to heat the casting to a temperature above the transformation range. Up to approximately 595°C (1100°F), the effect of short times at temperature on the structure of gray iron is insignificant. For most gray irons, a ferritizing annealing temperature between 700 and 760°C (1300 and 1400°F) is recommended. Medium (full) annealing. It is usually performed at temperatures between 790 and 900°C (1450 and 1650°F). This treatment is used when a ferritizing anneal would be ineffective because of the high alloy content of a particular iron. It is

279

recommended, however, to test the efficacy of temperatures below 760°C (1400°F) before a higher annealing temperature is adopted as part of a standard procedure. Holding times comparable to those used in ferritizing annealing are usually employed. When the high temperatures of medium annealing are used, however, the casting must be cooled slowly through the transformation range, from about 790 to 675°C (1450 to 1250°F). Graphitizing Annealing. If the microstructure of gray iron contains massive carbide particles, higher annealing temperatures are necessary. Graphitizing annealing may simply serve to convert massive carbide to pearlite and graphite, although in some applications it may be desired to carry out a ferritizing annealing treatment to provide maximum machinability. The production of free carbide that must later be removed by annealing is, except with pipe and permanent mold castings, almost always an accident resulting from inadequate inoculation or the presence of excess carbide formers, which inhibit normal graphitization; thus, the annealing process is not considered part of the normal production cycle. To break down massive carbide with reasonable speed, temperatures of at least 870°C (1600°F) are required. With each additional 55°C (100°F) increment in holding temperature, the rate of carbide decomposition doubles. Consequently, it is general practice to employ holding temperatures of 900 to 955°C (1650 to 1750°F).

Normalizing Gray iron is normalized by being heated to a temperature above the transformation range, held at this temperature for a period of about 1 hour per inch of maximum section thickness, and cooled in still air to room temperature. Normalizing may be used to enhance mechanical properties, such as hardness and tensile strength, or to restore as-cast properties that have been modified by another heating process, such as graphitizing or the preheating and postheating associated with repair welding. The temperature range for normalizing gray iron is approximately 885 to 925°C (1625 to 1700°F). Austenitizing temperature has a marked effect on microstructure and on mechanical properties such as hardness and tensile strength. The tensile strength and hardness of a normalized gray iron casting depend on the following parameters: • • •

Combined carbon content Pearlite spacing (distance between cementite plates) Graphite morphology.

The graphite morphology does not change to any significant extent during normalization, and its effect on hardness and tensile strength is omitted in this discussion on normalizing. Combined carbon content is determined by the normalizing (austenitizing) temperature and the chemical composition of the casting. Higher normalizing

280

temperatures increase the carbon solubility in austenite (that is, the cementite volume in the resultant pearlite). A higher cementite volume, in turn, increases both the hardness and the tensile strength. The alloy composition of a gray iron casting also influences carbon solubility in austenite. Some elements increase carbon solubility, some decrease it, and others have no effect on it. The carbon content of the matrix is determined by the combined effects of the alloying elements. The other parameter affecting hardness and tensile strength in a normalized gray iron casting is the pearlite spacing. Pearlite spacing is determined by the cooling rate of the casting after austenitization and the alloy composition. Fast cooling results in small pearlite spacing, higher hardness, and higher tensile strength. Too high a cooling rate may cause partial or full martensitic transformation. The addition of alloying elements may change hardness and tensile strength significantly.

6.11. Heat Treating of Gray Irons: Part Two Abstract: The heat treatment of gray irons can considerably alter the matrix microstructure with little or no effect on the size and shape of the graphite achieved during casting. The matrix microstructures resulting from heat treatment can vary from ferrite-pearlite to tempered martensite. However, even though gray iron can be hardened by quenching from elevated temperatures, heat treatment is not ordinarily used commercially to increase the overall strength of gray iron castings because the strength of the as-cast metal can be increased at less cost by reducing the silicon and total carbon contents or by adding alloying elements.

Gray irons are a group of cast irons that form flake graphite during solidification, in contrast to the spheroidal graphite morphology of ductile irons. The flake graphite in gray irons is dispersed in a matrix with a microstructure that is determined by composition and heat treatment. The heat treatment of gray irons can considerably alter the matrix microstructure with little or no effect on the size and shape of the graphite achieved during casting. The matrix microstructures resulting from heat treatment can vary from ferritepearlite to tempered martensite. However, even though gray iron can be hardened by quenching from elevated temperatures, heat treatment is not ordinarily used commercially to increase the overall strength of gray iron castings because the strength of the as-cast metal can be increased at less cost by reducing the silicon and total carbon contents or by adding alloying elements.

Hardening and Tempering Gray irons are hardened and tempered to improve their mechanical properties, particularly strength and wear resistance. After being hardened and tempered, these irons usually exhibit wear resistance approximately five times greater than that of pearlitic gray irons. Furnace or salt bath hardening can be applied to a wider variety of gray irons than can either flame or induction hardening. In flame and induction hardening, a relatively large content of combined carbon is required because of the extremely short period available for the solution of carbon in austenite. In furnace or salt bath hardening, however, a casting can be held at a temperature above the

281

transformation range for as long as is necessary; even an iron initially containing no combined carbon can be hardened. Unalloyed gray iron of low combined carbon content must be austenitized for a longer time to saturate austenite with carbon. With increased time, more carbon is dissolved in austenite and hardness after quenching is increased. Because of its higher silicon content, an unalloyed gray iron with a combined carbon content of 0.60% exhibits higher hardenability than a carbon steel with the same carbon content. However, because of the effect of silicon in reducing the solubility of carbon in austenite, unalloyed irons with higher silicon contents necessarily require higher austenitizing temperatures to attain the same hardness. Manganese increases hardenability; approximately 1.50% Mn was found to be sufficient for through hardening a 38 mm section in oil or for through hardening a 64 mm section in water. Manganese, nickel, copper, and molybdenum are the recognized elements for increasing the hardenability of gray iron. Although chromium, by itself, does not influence the hardenability of gray iron, its contribution to carbide stabilization is important, particularly in flame hardening. Austenitizing. In hardening gray iron, the casting is heated to a temperature high enough to promote the formation of austenite, held at that temperature until the desired amount of carbon has been dissolved, and then quenched at a suitable rate. The temperature to which a casting must be heated is determined by the transformation range of the particular gray iron of which it is made. The transformation range can extend more than 55°C above the At (transformationstart) temperature. A formula for determining the approximate A, transformation temperature of unalloyed gray iron is: A (°C) = 730 + 28.0 (% Si) - 25.0 (% Mn) Chromium raises the transformation range of gray iron. In high-nickel, high-silicon irons, for example, each percent of chromium raises the transformation range by about 10 to 15°C. Nickel, on the other hand, lowers the critical range. In a gray iron containing from 4 to 5% Ni, the upper limit of the transformation range is about 710°C. Castings should be treated through the lower temperature range slowly, in order to avoid cracking. Above a range of 595 to 650°C, which is above the stress-relieving range, heating may be as rapid as desired. In fact, time may be saved by heating the casting slowly to about 650°C in one furnace and then transferring it to a second furnace and bringing it rapidly up to the austenitizing temperature. Quenching. Molten salt and oil are the quenching media used most frequently for gray iron. Water is not generally a satisfactory quenching medium for furnace-heated gray iron; it extracts heat so rapidly that distortion and cracking are likely in all parts except small ones of simple design. Recently developed water-soluble polymer quenches can provide the convenience of water quenching, along with lower cooling rates, which can minimize thermal shock.

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The least severe quenching medium is air. Unalloyed or low-alloy gray iron castings usually cannot be air quenched because the cooling rate is not sufficiently high to form martensite. However, for irons of high alloy content, forced-air quenching is frequently the most desirable cooling method. Tempering. After quenching, castings are usually tempered at temperatures well below the transformation range for about 1h per inch of thickest section. As the quenched iron is tempered, its hardness decreases, whereas it usually gains in strength and toughness.

Austempering In austempering, the microstructural end product of the gray iron matrix formed below the pearlite range but above the martensite range is an acicular or bainitic ferrite, plus varying amounts of austenite depending on the transformation temperature. The iron is quenched from a temperature above the transformation range in a hot quenching bath and is maintained in the bath at constant temperature until the austempering transformation is complete. In all hot quenching processes, the temperatures to which castings must be heated for austenitizing and the required holding times at temperature prior to quenching in the hot bath correspond to the temperatures and times used in conventional hardening, that is, temperatures between 840 and 900°C (1550 and 1650°F). The holding time depends on the size and chemical composition of the casting. Gray iron is usually quenched in salt, oil, or lead baths at 230 to 425°C for austempering. When high hardness and wear resistance are the ultimate aim of this treatment, the temperature of the quench bath is usually held between 230 and 290°C. The effect of iron composition on the holding time may be considerable. Alloy additions, such as nickel, chromium, and molybdenum, increase the time required for transformation.

Martempering Martempering is used to produce martensite without developing the high stresses that usually accompany its formation. It is similar to conventional hardening except that distortion is minimized. Nevertheless, the characteristic brittleness of the martensite remains in a gray iron casting after martempering, and martempered castings are almost always tempered. The casting is quenched from above the transformation range in a salt, oil, or lead bath: held in the bath at a temperature slightly above the range at which martensite forms (200 to 260°C or 400 to 500°F. for unalloyed irons) only until the casting has reached the bath temperature; and then cooled to room temperature. If a wholly martensitic structure is desired, the casting must be held in the hot quench bath only long enough to permit ii to reach the temperature of the bath. Thus, the size and shape of the casting dictate the duration of martempering.

Flame Hardening

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Flame hardening is the method of surface hardening most commonly to gray iron. After flame hardening, a gray iron casting consists of a hard, wear-resistant outer layer of martensite and a core of softer gray iron, which during treatment does not reach the At transformation temperature. Both unalloyed and alloyed gray irons can be successfully flame hardened. However, some compositions yield much better results than do others. One of the most important aspects of composition is the combined carbon content, which should be in the range of 0.50 to 0.70%, although irons with as little as 0.40% combined carbon can be flame hardened. In general, flame hardening is not recommended for irons that contain more than 0.80% combined carbon because such irons (mottled or white irons) may crack in surface hardening. Effects of Alloying Elements. In general, alloyed gray irons can be flame hardened with greater ease than can unalloyed irons, partly because alloyed gray irons have increased hardenability. Final hardness also may be increased by alloying additions. The maximum hardness obtainable by flame hardening an unalloyed gray iron containing approximately 3% total carbon, 1.7% Si, and 0.60 to 0.80% Mn ranges from 400 to 500 HB. This is because the Brinell hardness value for gray iron is an average of the hardness of the matrix and that of the relatively soft graphite flakes. Actually, the matrix hardness on which wear resistance depends approximates 600 HB. With the addition of 2.5% Ni and 0.5% Cr, an average surface hardness of 550 HB can be obtained. The same result has been achieved using 1.0 to 1.5% Ni and 0.25% Mo. Stress Relieving. Whenever practicable or economically feasible, flame-hardened castings should be stress relieved at 150 to 200°C.

Induction Hardening Gray iron castings can be surface hardened by the induction method when the number of castings to be processed is large enough to warrant the relatively high equipment cost and the need for special induction coils. Considerable variation in the hardness of the cast irons may be expected because of a variation in the combined carbon content. A minimum combined carbon content of 0.40 to 0.50% C is recommended for cast iron to be hardened by induction, with the short heating cycles that are characteristic of this process. Heating castings with lower combined carbon content to high hardening temperatures for relatively long periods of time may dissolve some free graphite, but such a procedure is likely to coarsen the grain.

6.12. Heat Treating of High Alloy Graphitic Irons Abstract: High-alloy cast irons are an important group of materials whose production should be considered separately from that of the ordinary types of cast irons. In these cast iron alloys, alloy content is well above 4% and, consequently, they cannot be produced by ladle additions to irons of otherwise standard compositions. The producing foundries usually have the equipment needed to handle the heat treatment and other thermal processing unique to the production of these alloys. The heat treatment practices for the following high-alloy graphitic irons are described:

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• • •

Austenitic gray and ductile irons High-silicon irons for heat resisting applications High-silicon irons for corrosion resisting applications

High-alloy cast irons are an important group of materials whose production should be considered separately from that of the ordinary types of cast irons. In these cast iron alloys, alloy content is well above 4% and, consequently, they cannot be produced by ladle additions to irons of otherwise standard compositions. The producing foundries usually have the equipment needed to handle the heat treatment and other thermal processing unique to the production of these alloys. The cast iron alloys discussed in this article are alloyed for increased abrasion resistance, for strength and oxidation resistance at elevated temperatures, and for improved corrosion resistance. They include the high-alloy graphitic irons and the high-alloy white irons. The heat treatment practices for the following high-alloy graphitic irons are described: • • •

Austenitic gray and ductile irons High-silicon irons for heat resisting applications High-silicon irons for corrosion resisting applications

The high-alloy graphitic cast irons have found special use primarily in applications requiring (1) corrosion resistance or (2) strength and oxidation resistance in hightemperature service. Those alloys used in applications requiring corrosion resistance comprise the nickel-alloyed (13 to 36% Ni) gray and ductile irons, and the highsilicon (14.5% Si) gray irons. The alloyed irons produced for high-temperature service comprise the austenitic, nickel-alloyed gray and nodular irons, the high-silicon (4 to 6% Si) gray and nodular irons and the aluminum-alloyed gray and nodular irons. Two groups of aluminumalloyed irons are recognized: the 1 to 7% Al irons and the 18 to 25% Al irons.

Austenitic Nickel-Alloyed Graphitic Irons These nickel-alloyed austenitic irons have found usefulness in applications requiring corrosion resistance, wear resistance, and high-temperature stability and strength. Additional properties of benefit are low thermal expansion coefficients, nonmagnetic properties, and cast iron materials having good toughness at low temperatures. The procedures and temperatures of the heat treatments for these ductile irons with nodular graphite are similar to those for gray (flake-graphite), corrosion-resistant austenitic cast irons. ASTM Specification A 436 defines eight grades of austenitic gray iron alloys, four of which are designed to be used in elevated-temperature applications and four types are used in applications requiring corrosion resistance.

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The ASTM Specification A 439 defines the group of austenitic ductile irons. There are nine alloys listed in the specification. The austenitic ductile iron alloys have similar compositions to the austenitic gray iron alloys but have been treated with magnesium to produce nodular graphite. The ductile iron alloys have high strength and ductility combined with the same desirable properties of the gray iron alloys. They provide resistance to frictional wear, corrosion resistance, strength and oxidation resistance at high temperatures, nonmagnetic characteristics and, in some alloys, low thermal expansivity at ambient temperatures. Heat Treatment of Austenitic Ductile Irons. Heat treatment of the nickel-alloyed austenitic irons serves to reduce residual stresses and to stabilize the microstructure for increased durability. Heat treatments are similar with the graphite in nodular form (ductile iron) or flake form (gray iron). Stress Relieving. For most applications, it is recommended that austenitic cast irons be stress relieved at 620 to 675°C (1150 to 1250°F), for 1 h per 25 mm (1 in.) of section, to remove residual stresses resulting from casting or machining, or both. Stress relieving should follow rough machining, particularly for castings that must conform to close dimensional tolerances, that have been extensively welded, or that are to be exposed to high stresses in service. Stress relieving does not affect tensile strength, hardness, or ductility. For large, relatively thin-section castings, moldcooling to below 315°C (600°F) is recommended rather than stress relief heat treatment. Spheroidize Annealing. Castings with hardness above 190 HB may be softened by heating to 980 to 1040°C (1800 to 1900°F) for 1/2 to 5 h except those alloys containing 4% or more chromium. Excessive carbides cause this high hardness and may occur in rapidly cooled castings and thin sections. Annealing dissolves or spheroidizes carbides. Although it lowers hardness, spheroidize annealing does not adversely affect strength. High-Temperature Stabilization. This treatment consists of holding at 760°C (1400°F) for 4 h minimum or at 870°C (1600°F) for 2 h minimum, furnace cooling to 540°C (1000°F), and then cooling in air. This treatment stabilizes the microstructure and minimizes growth and warpage in service. The treatment is designed to reduce carbon levels in the matrix and some growth and distortion often accompanies heat treatment. Thus, it is usually advisable to stabilize castings prior to final machining. Dimensional Stabilization. This treatment normally is limited to castings that require true dimensional stability, such as those used in precision machinery or scientific instruments. The treatment is not applicable to castings of type I alloys. Other alloys may be dimensionally stabilized by the following treatment: • • • • •

Heat to 870°C (1600°F), and hold for 2 h minimum plus 1 h per 25 mm (1 in.) of section Furnace cool, at a maximum rate of 50°C/h (100°F/h), to 540°C (1000°F) Hold at 540°C (1000°F) for 1 h per 25 mm (1 in.) of section, and then cool uniformly in air After rough machining, reheat to 455 to 480°C (850 to 900°F) and hold for 1 h per 25 mm (1 in.) of section, and cool uniformly in air Finish machine and reheat to 260 to 315°C (500 to 600°F), and cool uniformly in air.

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Solution Treating. Although this treatment is seldom used, quenching from high temperatures is capable of producing higher-than-normal strength levels and slightly higher hardnesses by dissolving some carbon in austenite at elevated temperatures and by preventing precipitation of the carbon by rapid cooling. This treatment consists of heating to 925 to 1010°C (1700 to 1850°F) and quenching in oil or water. Because no metallurgical phase change occurs, the possibility of cracking is lessened.

High-Silicon Irons for High-Temperature Service Graphitic irons alloyed with from 4 to 6% Si have provided good service, and low cost, in many elevated-temperature applications. These irons, whether gray or nodular, provide good oxidation resistance and stable ferritic matrix structures that will not go through a phase change at temperatures up to 815°C. The elevated silicon content of these otherwise normal cast iron alloys reduces the rate of oxidation at elevated temperatures, because it promotes the formation of a dense, adherent film at the surface, which consists of iron silicate rather than iron oxide. This layer is much more resistant to oxygen penetration and its effectiveness improves with increasing silicon content. High-Silicon Nodular Irons. The advent of ductile iron led to the development of high-silicon nodular irons, which currently represent the greatest tonnage of these types of irons being produced. Converting the eutectic flake graphite network to isolated graphite nodules further improved resistance to oxidation and growth. The higher strength and ductility of the nodular iron versions of these alloys qualifies them for more rigorous service. The high-silicon nodular iron alloys are designed to extend the upper end of the range of service temperatures viable for ferritic nodular irons. These irons are used to temperatures of 900°C. At 5 to 6% Si, oxidation resistance is improved and critical temperature is increased, but the iron can be very brittle at room temperature. For most applications, alloying with 0 to 1% Mo provides adequate strength at elevated temperatures and creep resistance. The high-silicon gray and nodular irons are predominantly, ferritic as-cast, but the presence of carbide stabilizing elements will result in a certain amount of pearlite and often intercellular carbides. These alloys are inherently more brittle than standard grades of iron and usually have higher levels of internal stress due to lower thermal conductivity and higher elevated-temperature strength. These factors should be taken into account where deciding on heat treatment requirements. For the high-silicon nodular irons, high-temperature heat treatment is advised in all cases to anneal any pearlite and stabilize the casting against growth in service. A normal graphitizing (full) anneal in the austenitic temperature range is recommended where undesirable amounts of carbide are present. For the 4 to 5% Si irons this will require heating to at least 900°C (1650°F) for several hours, followed by slow cooling to below 700°C (1300°F). At higher silicon contents (>5%), in which carbides readily break down, and in castings relatively carbide-free, subcritical annealing in the temperature range 720 to 790°C (1325 to 1450°F) for 4 h is effective in ferritizing the matrix. Compared to full annealing, the

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subcritically annealed material will have somewhat higher strength, but ductility and toughness will be reduced. High-Silicon Irons for Corrosion Resistance. Irons with high silicon content (14.5% Si) comprise a unique corrosion-resistant ferritic cast iron group. These alloys are widely used in the chemical industry for processing and for transporting highly corrosive liquids. The most common of the high-silicon iron alloys are covered in ASTM Specification A 518M. Because of the very brittle nature of high-silicon cast iron, castings are usually shaken out only after mold cooling to ambient temperature. However, some casting geometries demand hot shakeout so that the castings can be immediately stressrelieved and furnace cooled to prevent cracking. Castings are stress relieved by heating in the range of 870 to 900°C (1600 to 1650°F) followed by slow cooling to ambient temperatures to minimize the likelihood of cracking. Heat treatments have no significant effect on corrosion resistance.

6.13. Heat Treating of High-Alloy White Irons Abstract: The high-alloy white irons are primarily used for abrasion-resistant applications and are readily cast in the shapes needed in machinery used for crushing, grinding, and general handling of abrasive materials. The large volume of eutectic carbides in their microstructures provides the high hardness needed for crushing and grinding other materials. The metallic matrix supporting the carbide phase in these irons can be adjusted by alloy content and heat treatment to develop the proper balance between resistance to abrasion and the toughness needed to withstand repeated impact.

High-alloy cast irons are an important group of materials whose production should be considered separately from that of the ordinary types of cast irons. The producing foundries usually have the equipment needed to handle the heat treatment and other thermal processing unique to the production of these alloys. The high-alloy white irons are primarily used for abrasion-resistant applications and are readily cast in the shapes needed in machinery used for crushing, grinding, and general handling of abrasive materials. The large volume of eutectic carbides in their microstructures provides the high hardness needed for crushing and grinding other materials. The metallic matrix supporting the carbide phase in these irons can be adjusted by alloy content and heat treatment to develop the proper balance between resistance to abrasion and the toughness needed to withstand repeated impact. All high-alloy white irons contain chromium to prevent formation of graphite on solidification and to ensure the stability of the carbide phase. Most also contain nickel, molybdenum, copper, or combinations of these alloying elements to prevent the formation of pearlite in the microstructure. While low-alloyed pearlitic white iron castings develop hardness in the range 350 to 550 HB, the high-alloyed white irons range from 450 to 800 HB. ASTM Specification A 532 covers the composition and hardness of white iron grades used for abrasion-resistant applications. Many castings are ordered according to these specifications: however, a large number of castings are produced with modifications to composition for specific applications. It is most desirable that the

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designer, metallurgist, and foundry-man work together to specify the composition, heat treatment, and foundry practice to develop the most suitable alloy and casting design for a specific application. The high-alloy white cast irons fall into three major groups: • • •

The Ni-Cr white irons, which are low-chromium alloys containing 3 to 5% Ni and 1 to 4% Cr with one alloy modification which contains 7 to 11% Cr. The chromium-molybdenum irons containing 11 to 23% Cr, up to 3% Mo, and often additionally alloyed with nickel or copper. The 25% Cr or 28% Cr white irons, which may contain other alloying additions of molybdenum and/or nickel up to 1.5%

Nickel-Chromium White Irons One of the oldest groups of high-alloy irons of industrial importance, the Ni-Cr white irons, or Ni-Hard irons, have been produced for more than 50 years and are very cost-effective materials for crushing and grinding. In these martensitic white irons, nickel is the primary alloying element because at levels of 3 to 5% it is effective in suppressing the transformation of the austenite matrix to pearlite, and thus ensuring that a hard, martensitic structure will develop on cooling in the mold. Chromium is included in these alloys, at levels from 1.4 to 4% to ensure that the irons will solidify with carbides to counteract the graphitizing effect of nickel. The optimum composition of the Ni-Cr white iron alloy depends on the properties required for the service conditions and the dimensions and weight of the casting. Abrasion resistance is generally a function of the bulk hardness and the volume of carbide in Cr-Mo iron. Carbon is varied according to properties needed for the intended service. Carbon contents in the range of 3.2 to 3.6% are prescribed when maximum abrasion resistance is desired. Where impact loading is present, carbon content should be held in the range of 2.7 to 3.2%. Silicon is needed for two reasons. A minimum amount of silicon is necessary to improve fluidity and produce a fluid slag. But of equal importance is its effect on ascast hardness. Increased levels of silicon, in the range of 1 to 1.5%, have been found to increase the amount of martensite and the resulting hardness. Late additions of ferrosilicon have been reported to increase toughness. Note that higher silicon contents can promote pearlite and may increase the nickel requirement. Manganese is usually held to 0.8% max. While it provides increased hardenability to avoid pearlite formation, it is also a potent austenite stabilizer, more so than nickel, and will promote increased amounts of retained austenite and lower as-cast hardness. For this reason higher manganese levels are undesirable. In considering the nickel content required to avoid pearlite in a given casting, the level of manganese present should be a factor. Copper increases hardenability and the retention of austenite and, therefore, must be controlled for the same reason manganese is limited. Copper should be treated as a nickel substitute and, when properly included in the calculation of the amount of

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nickel required to inhibit pearlite in a given casting, it reduces the nickel requirement. Molybdenum is a potent hardenability agent in these alloys and is used in heavy section castings to augment hardenability and inhibit pearlite. Heat Treatment or Nickel-Chromium White Irons. Nickel-chromium white iron castings are given a stress-relief heat treatment because, properly made, they have a martensitic matrix structure, as-cast. Tempering is performed between 205 to 260°C (400 to 450°F) for at least 4 h. This tempers the martensite, relieves some of the transformation stresses, and increases the strength and impact toughness by 50 to 80%. Some additional martensite may form on cooling from the tempering temperature. This heat treatment does not reduce hardness or abrasion resistance. In the heat treatment of any white cast iron, care must be taken to avoid cracking by thermal shock; never place the castings in a hot furnace or otherwise subject them to rapid heating or cooling. The risk of cracking increases with the complexity of the casting shape and section thickness. An austenitizing heat treatment usually comprised heating at temperatures between 750 and 790°C (1380 and 1450°F) with a soak time of 8 h. Air or furnace cooling, not over 30°C/h, was conducted followed by a tempering/stress-relief heat treatment. Refrigeration heat treatment is the more commonly practiced remedy for low hardness today.

High-Chromium White Irons The oldest high-alloy white irons produced commercially were the high-chromium (28% Cr) white irons. The high-chromium white irons have excellent abrasion resistance and are used effectively in slurry pumps, brick molds, coal-grinding mills, rolling mill rolls, shot blasting equipment, and components for quarrying, hard-rock mining and milling. In some applications they must also be able to withstand heavy impact loading. These alloyed white irons are recognized as providing the best combination of toughness and abrasion resistance attainable among the white cast irons. Through variations in composition and heat treatment these properties can be adjusted to meet the needs of most abrasive applications. Special High-Chromium Iron Alloys for Corrosion Resistance. Alloys with improved resistance to corrosion, for applications such as pumps handling, are produced with high chromium contents (26 to 28% Cr) and low carbon contents (1.6 to 2.0% C). These high-chromium, low-carbon irons will provide the maximum chromium content in the matrix. Addition of 2% Mo is recommended for improving resistance to chloride-containing environments. Chromium causes the formation of an adherent, complex, chromium-rich oxide film providing resistance to scaling at temperatures up to 1040°C (1900°F). The high-chromium irons designated for use at elevated temperatures fall into one of three categories, depending upon the matrix structure: • •

The martensitic irons alloyed with 12 to 28% Cr The ferritic irons alloyed with 30 to 34% Cr

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•

The austenitic irons which in addition to containing 15 to 30% Cr also contain 10 to 15% Ni to stabilize the austenite phase

Carbon contents of these alloys range from 1 to 2%. Optimum performance is usually achieved with heat treated martensitic structures. As described in the previous section, alloying must be sufficient to ensure that a pearlite-free microstructure is obtained in heat treatment. Of necessity, the heat treatment requires an air quench from the austenitizing temperature. Faster cooling rates should not be used, because the casting can develop cracks due to high thermal and/or transformation stresses. Thus the alloy must have sufficient hardenability to allow air hardening. Over-alloying with manganese, nickel, and copper will promote retained austenite, which detracts from resistance to abrasion and spalling. Austenitization. There is an optimum austenitizing temperature to achieve maximum hardness, which varies for each composition. The austenitizing temperature determines the amount of carbon that remains in solution in the austenite matrix. Too high a temperature increases the stability of the austenite, and the higher retained austenite content reduces hardness. Low temperatures result in low-carbon martensite reducing both hardness and abrasion resistance. Class II irons containing 12 to 20% Cr are austenitized in the temperature range 950 to 1010°C (1750 to 1850°F). Class III irons containing 23 to 28% Cr are austenitized in the temperature range 1010 to 1090°C (1850 to 2000°F). Quenching. Air quenching (vigorous fan cooling) the castings from the austenitizing temperature to below the pearlite temperature range (that is, between 550 and 600°C, or 1020 and 1110°F) is highly recommended. The subsequent cooling rate should be substantially reduced to minimize stresses; still-air or even furnace cooling to ambient is common. Complex and heavy section castings are often placed back into the furnace, which is at 550 to 600°C, and allowed sufficient time to reach uniform temperature within the casting. After temperature is equalized, the castings are either furnace or still-air cooled to ambient temperature. Tempering. Castings can be put into service in the hardened (as cooled) condition without further tempering or subcritical heat treatments; however, tempering in the range of 200 to 230°C (400 to 450°F) for 2 to 4 h is recommended to restore some toughness in the martensitic matrix and to further relieve residual stresses. Sub critical Heat Treatment. Sub critical heat treatment (tempering) is sometimes performed, particularly in large heat-treated martensitic castings, to reduce retained austenite contents and increase resistance to spalling. The tempering parameters necessary to eliminate retained austenite are very sensitive to time and temperature and vary depending on the castings composition and prior thermal history. Typical tempering temperatures range from 480 to 540°C (900 to 1000°F) and times range from 8 to 12 h. Excess time or temperature results in softening and a drastic reduction in abrasion resistance. Annealing. Castings can be annealed to make them more machinable, either by sub-critical annealing or a full anneal. Subcritical annealing is accomplished by pearlitizing, via soaking in the narrow range between 690 and 705°C for from 4 to 12 h, which will produce hardness in the range 400 to 450 HB. Lower hardness can

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often be achieved with full annealing, whereby castings are heated in the range 955 to 1010°C followed by slow cooling to 760°C and holding at this temperature for 10 to 50 h depending on composition. Stress-Relieving. Very little information is available on the amount of stress relief that occurs with tempering. The predominant stresses present in heat-treated castings develop as a result of the volume change accompanying austenite to martensite transformation. Low-temperature tempering, in the range of 200 to 230°C, is particularly desirable because a substantial improvement (20%) in fracture toughness occurs when tempering the martens lie phase. Tempering at temperatures sufficient to significantly relieve stresses, that is, above 540°C, will substantially reduce abrasion resistance.

6.14. Heat Treating of Malleable Irons Abstract: Ferritic and pearlitic malleable irons are both produced by annealing white iron of controlled composition. Malleable irons have largely been replaced by ductile iron in many applications. This is due in part to the necessity of lengthy heat treatments for malleable iron and the difficulty in cooling thick sections rapidly enough to produce white iron. Malleable iron is still often preferred for thin section castings and parts that require maximum machinability and wear resistance.

Ferritic and pearlitic malleable irons are both produced by annealing white iron of controlled composition. Malleable irons have largely been replaced by ductile iron in many applications. This is due in part to the necessity of lengthy heat treatments for malleable iron and the difficulty in cooling thick sections rapidly enough to produce white iron. Malleable iron is still often preferred for thin section castings and parts that require maximum machinability and wear resistance. The annealing of malleable iron should be done in a furnace with a controlled atmosphere of dry nitrogen, hydrogen (1.5%), and carbon monoxide (1.5%). The dew point of this mixture should be between -40 and -70°C (-40 and -20°F). These conditions eliminate the possibility of decarburization and loss of learner carbon nodules below the casting surface. The annealing treatment involves three important steps: • •

•

The first causes nucleation of temper carbon. It is initiated during heating to a high holding temperature and occurs very early during the holding period. The second step consists of holding at 900 to 970°C (1650 to 1780°F); this step is called first-stage graphitization (FSG). During FSG, massive carbides are eliminated from the iron structure. Long holding periods at 955°C (1750°F) will reduce the solubility of nitrogen in iron (which should be kept at 80 to 120 ppm), thereby reducing the mechanical properties of the iron. This occurrence should be kept in mind for long, or "weekend" holding periods. When the carbides are eliminated, the iron is rapidly cooled to 740°C. The third step in the annealing treatment consists of slow cooling through the allotropic transformation range of the iron; this step is called second-stage graphitization (SSG). During SSG a completely ferrule matrix free of pearlite and carbides is obtained when the cooling rate is 2 to 28°C/h (3 to 50°F/h). This cooling rate, which depends on the silicon content of the iron and the temper carbon nodule count, may be increased to 85°C/min (150°F/min) by

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air quenching from 900°C (1650°F) to form a pearlitic matrix. Oil quenching from 900°C (1650°F) will produce a martensitic matrix. However, unless the temperature in the furnace is lowered to 845°C (1550°F) for at least 4 h (plus 1 h for each 25 mm or inch of section casting thickness), prior to uniform quenching in oil, the matrix microstructure will not be uniform in combined carbon. This nonuniformity reduces machinability. If the hardness is reduced by extended tempering, the resulting structure may not have a good response to selective hardening.

Hardening and Tempering of Malleable Iron A typical procedure for producing a hardened pearlitic malleable iron consists of: • • •

first air quenching castings after first-stage annealing, which results in retention of about 0.75% combined carbon in the matrix; second, reheating and holding for 1 hour at 885°C (l625°F) to reaustenitize the matrix and homogenize the combined carbon; and then quenching in heated and agitated oil, thereby developing a matrix consisting of martensite without bainite and having a hardness of 555 to 627 HB.

The appropriate austenitizing temperature for pearlitic malleable iron is 885°C (1625°F) and for ferritic malleable iron it is 900°C (1650°F). If direct oil quenching is used, caution must be exercised to prevent cracking due to high combined carbon. Air-quenched and pearlitic malleable iron has a matrix consisting of a ferrite ring around the temper carbon (which produces a lower yield strength) and partially broken lamely pearlite. The remaining lamellar pearlite reduces machinability to a limit of 240 HB. Increasing the austenitizing time and temperature increases the amount dissolved carbon, which is measured as combined carbon in the matrix after quenched to room temperature. Austenitizing temperatures in the range of 900 to 930°C (1650 to 1700°F) result in a more homogeneous austenite, which is desirable for more uniform martensite. Higher temperatures can result in a greater tendency toward distortion or cracking. Tempering of pearlite is time and temperature dependent. Tempering of martensite is primarily temperature dependent, while time being secondary. Hardened and tempered pearlitic malleable iron can also be produced from fully annealed ferritic malleable iron, the matrix of which is essentially carbon-free: graphite can be dissolved in austenite by holding at 900 to 930°C (1650 to 1700°F) for a time sufficiently long for the production of an austenite matrix of uniform carbon content. In general, the combined carbon content of the matrix produced by this procedure is slightly lower than they of a pearlitic malleable iron made by air quenching directly from 900°C (1650°F). Tempering treatments consist of cycles of no less than 2 h at temperature to ensure uniformity of product. Tempering times must also be adjusted for section thickness and quenched microstructures. Fine pearlite and bainite require longer tempering times than that for martensite. In general, final hardness is controlled with process controls approximately the same as those encountered in the heat treatment of

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medium-carbon and higher-carbon steels. This is particularly true when the specification requires final hardnesses in the range from 241 to 321 HB. The effects of tempering on the hardness of alloyed and unalloyed malleable irons illustrate the beneficial effects of alloying on as-quenched hardness and stability at elevated temperatures. During all tempering treatments, carbide has a tendency to decompose, with resulting deposition of graphite on existing temper carbon nodules. This tendency is least at the lower tempering temperatures or in suitably alloyed pearlitic malleable irons. Martempering and tempering develops mechanical properties similar to those resulting from conventional oil quenching and tempering: typical tensile strength 860 MPa (125 ksi), yield strength 760 MPa (110 ksi), and hardness 300 HB. Pearlitic malleable iron castings that are susceptible to cracking when quenched in warm oil (40 to 95°C, or 100 to 200°F) from the austenitizing temperature may be safely quenched in salt or oil at about 200°C (400°F). Elevator camshafts varying in length from 0.3 to 0.45 m (12 to 18 in.) and various sizes of wear-chain components are examples of martempered pearlitic malleable iron.

Bainitic Heat Treatment of Pearlitic Malleable Iron Both upper and lower bainite can be formed in pearlitic malleable iron with a marked increase in tensile strength and hardness but with a decrease in ductility. A pearlitic malleable iron (2.6C-1.4Si-0.5Mn-0.1S), annealed at 930°C (1700°F) for 16 h, air quenched and tempered at 680°C (1250°F) for 4 h, developed an ultimate tensile strength of 650 MPa (94.2 ksi), a yield strength of 460 MPa (66.5 ksi), and a 3.4% elongation at 217 HB. This same iron austenitized at 900°C (1650°F) in molten salt for 1 h, quenched in molten salt at 295°C (560°F) for 3 h, and air cooled gave an ultimate strength of 995 MPa (144.2 ksi), a yield strength of 920 MPa (133.4 ksi), and 388 HB.

Surface Hardening of Pearlitic Malleable Iron Fully pearlitic malleable iron may be surface hardened by either induction heating and quenching or flame heating and quenching. Laser and electron beam techniques also have been used for hardening selected areas on the surface of pearlitic and ferritic malleable iron castings that are free from decarburization. Generally, hardness in the range from 55 to 60 HRC is attainable, with the depth of penetration being controlled by the rate of heating and by the temperature developed at the surface of the part being hardened. In induction hardening, this is accomplished by the close regulation of power output, operating frequency, heating time, and alloy content of the iron. The maximum hardness obtainable in the matrix of a properly hardened part is 67 HRc; however, conventional hardness measurements show less than the true matrix hardness because of the temper carbon nodules that are averaged into the hardness.

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Generally, a casting with a matrix microhardness of 67 HRc will have average hardness of about 62 HRc, as measured with the standard Rockwell tester. Rocker arms and clutch hubs are examples of automotive production parts that are surface hardened by induction. Flame hardening requires close control for these applications in order to avoid distortion that would interfere with their operation. The two examples that follow describe the successful application of induction and flame hardening to other production parts.

6.15. Heat Treating of Nodular Irons: Part One Abstract: Nodular cast irons (or ductile, or spheroidal graphite iron) are primarily heat treated to create matrix microstructures and associated mechanical properties not readily obtained in the as-cast condition. Ascast matrix microstructures usually consist of ferrite or pearlite or combinations of both, depending on cast section size and/or alloy composition.

Nodular cast irons (or ductile, or spheroidal graphite iron) are primarily heat treated to create matrix microstructures and associated mechanical properties not readily obtained in the as-cast condition. As-cast matrix microstructures usually consist of ferrite or pearlite or combinations of both, depending on cast section size and/or alloy composition. The most important heat treatments and their purposes are: • • • • • •

Stress relieving, a low-temperature treatment, to reduce or relieve internal stresses remaining after casting Annealing, to improve ductility and toughness, to reduce hardness, and to remove carbides Normalizing, to improve strength with some ductility Hardening and tempering, to increase hardness or to improve strength and raise proof stress ratio Austempering, to yield a microstructure of high strength, with some ductility and good wear resistance Surface hardening, by induction, flame, or laser, to produce a locally selected wear-resistant hard surface

The normalizing, hardening, and austempering heat treatment, which involve austenitization, followed by controlled cooling or isothermal reaction, or a combination of the two, can produce a variety of microstructures and greatly extend the limits on the mechanical properties of ductile cast iron. These microstructures can be separated into two broad classes: • •

Those in which the major iron-bearing matrix phase is the thermodynamically stable body-centered cubic (ferrite) structure Those with a matrix phase that is a meta-stable face-centered cubic (austenite) structure. The former are usually generated by the annealing, normalizing, normalizing and tempering, or quenching and tempering processes. The latter are generated by austempering, an isothermal reaction process resulting in a product called austempered ductile iron (ADI).

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Other heat treatments in common industrial use include stress-relief annealing and selective surface heat treatment. Stress-relief annealing does not involve major micro-structural transformations, whereas selective surface treatment (such as flame and induction surface hardening) does involve microstructural transformations, but only in selectively controlled parts of the casting. The basic structural differences between the ferritic and austenitic classes are explained in the Fig 1 and 2. Figure 1 shows a continuous cooling transformation (CCT) diagram and cooling curves for furnace cooling, aircooling, and quenching. It can be seen from Fig 1 that slow furnace cooling results in a ferritic matrix (the desired product of annealing), whereas the cooling curve for air cooling, or normalizing, results in a pearlitic matrix, and quenching produces a matrix microstructure consisting mostly of martensite with some retained austenite. Tempering softens the normalized and quenched conditions, resulting in microstructures consisting of the matrix ferrite with small panicles of iron carbide (or secondary graphite).

Fig.1: CCT diagram showing annealing, normalizing and quenching; Ms stand martensite start, Mf for martensite finish. Figure 2 shows an isothermal transformation (IT) diagram for a ductile cast iron, together with a processing sequence depicting the production of ADI. In this process, austenitizing is followed by rapid quenching (usually in molten salt) to an intermediate temperature range for a time that allows the unique metastable carbon-rich (≈2% C) austenitic matrix (γH) to evolve simultaneously with nucleation and growth of a plate-like ferrite (α) or of ferrite plus carbide, depending on the austempering temperature and time at temperature. This austempering reaction progresses to a point at which the entire matrix has been transformed to the metastable product (stage I in Fig 2), and then

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that product is "frozen in" by cooling to room temperature before the true bainitic ferrite plus carbide phases can appear (stage II in Fig 2). In ductile cast irons the presence of 2 to 3 wt% Si prevents the rapid formation of iron carbide (Fe3C). Hence the carbon rejected during ferrite formation in the first stage of the reaction (stage I in Fig 2) enters the matrix austenite, enriching it and stabilizing it thermally to prevent martensite formation upon subsequent cooling. Thus the processing sequence in Fig 2 shows that the austempering reaction is terminated before stage II begins and illustrates the decrease in the martensite start (Ms) and martensite finish (Mf) temperatures as γH forms in stage I. Typical austempering times range from 1 to 4 h depending on alloy content and section size. If the part is austempered too long, undesirable bainite will form. Unlike steel, bainite in cast iron microstructures exhibits lower toughness and ductility.

Fig.2: IT diagram of a processing sequence for austempering.

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6.16. Heat Treating of Nodular Irons: Part Two Abstract: The most important heat treatments and their purposes are:

• • • • • •

Stress relieving, a low-temperature treatment, to reduce or relieve internal stresses remaining after casting Annealing, to improve ductility and toughness, to reduce hardness, and to remove carbides Normalizing, to improve strength with some ductility Hardening and tempering, to increase hardness or to improve strength and raise proof stress ratio Austempering, to yield a microstructure of high strength, with some ductility and good wear resistance Surface hardening, by induction, flame, or laser, to produce a locally selected wear-resistant hard surface

In this article annealing, normalizing, austempering, quenching and tempering of ductile cast iron are described.

Austenitizing Ductile Cast Iron The usual objective of austenitizing is to produce an austenitic matrix with as uniform carbon content as possible prior to thermal processing. For a typical hypereutectic ductile cast iron, an upper critical temperature must be exceeded so that the austenitizing temperature is in two-phase (austenite and graphite) field. This temperature varies with alloy content. The "equilibrium" austenite carbon content in equilibrium with graphite increases with an increase in austenitizing temperature. This ability to select (within limits) the matrix austenite carbon content makes austenitizing temperature control important in processes that depend on carbon in the matrix to drive a reaction. This is particularly true in structures to be austempered, in which the hardenability (or austemperability) depends to a significant degree on matrix carbon content. In general, alloy content, the original microstructure, and the section size determine the time required for austenitizing. The sections to follow on annealing, normalizing, quenching and tempering, and austempering discuss austenitizing when it is of concern.

Annealing Ductile Cast Iron When maximum ductility and good machinability are desired and high strength is not required, ductile iron castings are generally given a full ferritizing anneal. The microstructure is thus converted to ferrite, and the excess carbon is deposited on the existing nodules. This treatment produces ASTM grade 60-40-18. Amounts of manganese, phosphorus, and alloying elements such as chromium and molybdenum should be as low as possible if superior machinability is desired because these elements retard the annealing process. Recommended practice for annealing ductile iron castings is given below for different alloy contents and for castings with and without eutectic carbides:

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• •

•

Full anneal for unalloyed 2 to 3% Si iron with no eutectic carbide: Heat and hold at 870 to 900°C (1600 to 1650°F) for 1 h per inch of section. Furnace cool at 55°C/h (100°F/h) to 345°C (650°F). Air cool. Full anneal with carbides present: Heat and hold at 900 to 925°C (1650 to 1700°F) for 2 h minimum, longer for heavier sections. Furnace cool at 110°C/h (200°F/h) to 700°C (1300°F). Hold 2 h at 700°C (1300°F). Furnace cool at 55°C/h (100°F/h) to 345°C (650°F). Air cool. Subcritical anneal to convert pearlite to ferrite: Heat and hold at 705 to 720°C (1300 to 1330°F), 1 h per inch of section. Furnace cool at 55°C/h (100°F/h) to 345°C (650°F). Air cool. When alloys are present, controlled cooling times through the critical temperature range down to 400°C (750°F) must be reduced to below 55°C/h (100°F/h).

However, certain carbide-forming elements, mainly chromium, form primary carbides that are very difficult, if not impossible, to decompose. For example, the presence of 0.25% Cr results in primary intercellular carbides that cannot be broken down in 2 to 20 h heat treatments at 925°C (1700°F). The resulting matrix after pearlite breakdown is carbides in ferrite with only 5% elongation. Other examples of carbide stabilizers are molybdenum contents greater than 0.3%, and vanadium and tungsten contents exceeding 0.05%.

Hardenability of Ductile Cast Iron The hardenability of ductile cast iron is an important parameter for determining the response of a specific iron to normalizing, quenching and tempering, or austempering. Hardenability is normally measured by the Jominy test, in which a standard-sized bar (1 inch diameter by 4 inch in length) is austenitized and water quenched from one end. The variation in cooling rate results in micro-structural variations, giving hardness changes that are measured and recorded. The higher matrix carbon content resulting from the higher austenitizing temperature results in an increased hardenability (the Jominy curve is shifted to larger distances from the quenched end) and a greater maximum hardness. The purpose of adding alloy elements to ductile cast irons is to increase hardenability. Manganese and molybdenum are much more effective in increasing hardenablitty, per weight percent added, than nickel or copper. However, as is the case with steel, combinations of nickel and molybdenum, or copper and molybdenum, or copper, nickel, and manganese are more effective than the separate elements. Thus heavy-section castings that require through hardening or austempering usually contain combinations of these elements. Silicon, apart from its effect on matrix carbon content, does not have a large effect on hardenability.

Normalizing Ductile Cast Iron Normalizing (air cooling following austenitizing) can result in a considerable improvement in tensile strength and may be used in the production of ductile iron of ASTM type 100-70-03.

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The microstructure obtained by normalizing depends on the composition of the castings and the cooling rate. The composition of the casting dictates its hardenability that is, the relative position of the fields in the time-temperature CCT diagram. The cooling rate depends on the mass of the casting, but it also may be influenced by the temperature and movement of the surrounding air, during cooling. Normalizing generally produces a homogeneous structure of fine pearlite, if the iron is not too high in silicon content and has at least a moderate manganese content (0.3 to 0.5% or higher). Heavier castings that require normalizing usually contain alloying elements such as nickel, molybdenum, and additional manganese, for higher hardenability to ensure the development of a fully pearlitic structure after normalizing. Lighter castings made of alloyed iron may be martensitic or may contain an acicular structure after normalizing. The normalizing temperature is usually between 870 and 940°C (1600 and 1725°F). The standard time at temperature of 1 h per inch of section thickness or 1 h minimum is usually satisfactory. Longer times may be required for alloys containing elements that retard carbon diffusion in the austenite. For example, tin and antimony segregate to the nodules, effectively preventing the solution of carbon from the nodule sites. Normalizing is sometimes followed by tempering to attain the desired hardness and relieve residual stresses that develop upon air cooling when various parts of a casting, with different section sizes, cool at different rates. Tempering after normalizing is also used to obtain high toughness and impact resistance. The effect of tempering on hardness and tensile properties depends on the composition of the iron and the hardness level obtained in normalizing. Tempering usually consists of reheating to temperatures of 425 to 650°C (800 to 1200°F) and holding at the desired temperature for 1 h per inch of cross section. These temperatures are varied within the above range to meet specification limits.

Quenching and Tempering Ductile Cast Iron An austenitizing temperature of 845 to 925°C (1550 to 1700°F) is normally used for austenitizing commercial castings prior to quenching and tempering. Oil is preferred as a quenching medium to minimize stresses and quench cracking, but water or brine may be used for simple shapes. Complicated castings may have to be oil quenched at 80 to 100°C (180 to 210°F) to avoid cracks. The influence of the austenitizing temperature on the hardness of water-quenched cubes of ductile iron shows that the highest range of hardness (55 to 57 HRC) was obtained with austenitizing temperatures between 845 and 870°C (1550 and 1600°F). At temperatures above 870°C, the higher matrix carbon content resulted in a greater percentage of retained austenite and therefore a lower hardness. Castings should be tempered immediately after quenching to relieve quenching stresses. Tempered hardness depends on as-quenched hardness level, alloy content, and tempering temperature, as well as time. Tempering in the range from 425 to 600°C (800 to 1100°F) results in a decrease in hardness, the magnitude of which depends upon alloy content, initial hardness, and time. Vickers hardness of quenched ductile iron alloys change with tempering temperature and time.

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Tempering ductile iron is a two-stage process. The first involves the precipitation of carbides similar to the process in steels. The second stage (usually shown by the drop in hardness at longer times) involves nucleation and the growth of small, secondary graphite nodules at the expense of the carbides. The drop in hardness accompanying secondary graphitization produces a corresponding reduction in tensile and fatigue strength as well. Because alloy content affects the rate of secondary graphitization, each alloy will have a unique range of useful tempering temperatures.

Austempering Ductile Cast Iron When optimum strength and ductility are required, the heat treater has the opportunity to produce an austempered structure of austenite and ferrite. The austempered matrix is responsible for a significantly better tensile strength-toductility ratio than is possible with any other grade of ductile cast iron. The production of these desirable properties requires careful attention to section size and the time-temperature exposure during austenitizing and austempering. Section Size and Alloying. As section size increases, the rate of temperature change between the austenitizing temperature and austempering temperature decreases. Quenching and austempering techniques include the hot-oil quench (up to 240°C, or 460°F, only), nitrate/nitrite sail quenches, fluidized-bed method (for thin, small parts only), and, in tool-type applications, lead baths. In order to avoid high-temperature reaction products (such as pearlite in larger section sizes), salt bath quench severities can be increased with water additions or with alloying elements (such as copper, nickel, manganese, or molybdenum) that enhance pearlite hardenability. It is important to understand that these alloying elements tend to segregate during solidification so that a nonuniform distribution exists throughout the matrix. This has a potentially detrimental effect on the austempering reaction and therefore on mechanical properties. Ductility and impact toughness are the most severely affected. Manganese and molybdenum have the most powerful effect upon pearlite hardenability but will also segregate and freeze into intercellular regions of the casting to promote iron or alloy carbides. While nickel and copper do not affect hardenability nearly as much, they segregate to graphite nodule sites and do not form detrimental carbides. Combinations of these elements, which segregate in opposite fashions, are selected for their synergistic effect on hardenability. Austenitizing Temperature and Time. Usual schematic phase diagram shows that as austenitizing temperature increases, so does the matrix carbon content; the actual matrix carbon content depends in a complex way on the alloy elements present, their amount, and their location (segregation) within the matrix. The most important determinant of matrix carbon content in ductile irons is the silicon content; as silicon content increases for a given austenitizing temperature, the carbon content in the matrix decreases. Austenitizing temperatures between 845 and 925°C (1550 and 1700°F) are normal, and austenitizing times of approximately 2 h have been shown to be sufficient to recarburize the matrix fully. Austenitizing temperature, through its effect upon matrix carbon, has a significant effect on hardenability. The higher austenitizing temperature with its higher carbon content

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promotes increased hardenability, which causes a slower rate of isothermal austenite transformation. Austempering Temperature and Time. The austempering temperature is the primary determinant of the final microstructure and therefore the hardness and strength of the austempered product. As the austempering temperature increases, the strength and impact toughness vary. The attainment of maximum ductility at any given austempering temperature is a sensitive function of time. The initial increase in elongation occurs as stage I and elongation progresses to completion, at which point the fraction of austenite is a maximum. Further austempering merely serves to reduce ductility as the stage II reaction causes decomposition to the equilibrium bainite product. Typical austempering times vary from 1 to 4 h.

6.17. Casting Defects in Steels Abstract: Metal casters try to produce perfect castings. Few castings, however, are completely free of defects. Modern foundries have sophisticated inspection equipment can detect small differences in size and a wide variety of external and even internal defects. For example, slight shrinkage on the back of a decorative wall plaque is acceptable whereas similar shrinkage on a position cannot be tolerated. No matter what the intended use, however, the goal of modern foundries is zero defects in all castings.

Metal casters try to produce perfect castings. Few castings, however, are completely free of defects. Modern foundries have sophisticated inspection equipment can detect small differences in size and a wide variety of external and even internal defects. For example, slight shrinkage on the back of a decorative wall plaque is acceptable whereas similar shrinkage on a position cannot be tolerated. No matter what the intended use, however, the goal of modern foundries is zero defects in all castings. Scrap castings cause much concern. In industry, scrap results in smaller profits for the company and ultimately affects individual wages. Scrap meetings are held daily. Managers of all the major departments attend these meeting. They gather a castings that have been identified as scrap by in inspector. The defect(s) is circled with chalk. An effort is made to analyze the cause of the defect, and the manager whose department was responsible for it is directed to take corrective action to eliminate that specific defect in future castings. There are so many variables in the production of a metal casting that the cause is often a combination of several factors rather than a single one. All pertinent data related to the production of the casting (sand and core properties, pouring temperature) must be known in order to identify the defect correctly. After the defect is identified you should attempt to eliminate the defect by taking appropriate corrective action. The system used here for classifying defects is one based on a physical description of the defect under consideration. It is intended to permit an identification to be made either by direct observation of the defective casting or from a precise description of the defect, involving only the criteria of shape, appearance, location and dimensions.

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This unique system of classification, based upon the morphology of the defects, is more logical than one based upon causes since it requires no prior assumptions to be made. Seven basic categories of defects have been established, as listed below and for each basic category only one typical defect is being presented here.

1. Metallic Projections Joint flash or fins. Flat projection of irregular thickness, often with lacy edges, perpendicular to one of the faces of the casting. It occurs along the joint or parting line of the mold, at a core print, or wherever two elements of the mold intersect.

Possible Causes • •

Clearance between two elements of the mold or between mold and core; Poorly fit mold joint.

Remedies • • • •

Care in pattern making, molding and core making; Control of their dimensions; Care in core setting and mold assembly; Sealing of joints where possible.

2. Cavities Blowholes, pinholes. Smooth-walled cavities, essentially spherical, often not contacting the external casting surface (blowholes). The largest cavities are most often isolated; the smallest (pinholes) appear in groups of varying dimensions. In specific cases, the casting section can be strewn with blowholes of pinholes. The interior walls of blowholes and pinholes can be shiny, more or less oxidized or, in the case of cast iron, can be covered with a thin layer of graphite. The defect can appear in all regions of the casting.

Possible Causes Blowholes and pinholes are produced because of gas entrapped in the metal during the course of solidification: • •

• • • •

Excessive gas content in metal bath (charge materials, melting method, atmosphere, etc.); Dissolved gases are released during solidification; In the case of steel and cast irons: formation of carbon monoxide by the reaction of carbon and oxygen, presents as a gas or in oxide form. Blowholes from carbon monoxide may increase in size by diffusion of hydrogen or, less often, nitrogen; Excessive moisture in molds or cores; Core binders which liberate large amounts of gas; Excessive amounts of additives containing hydrocarbons; Blacking and washes which tend to liberate too much gas;

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• •

Insufficient evacuation of air and gas from the mold cavity; -insufficient mold and core permeability; Entrainment of air due to turbulence in the runner system.

Remedies • • • • • • •

Make adequate provision for evacuation of air and gas from the mold cavity; Increase permeability of mold and cores; Avoid improper gating systems; Assure adequate baking of dry sand molds; Control moisture levels in green sand molding; Reduce amounts of binders and additives used or change to other types; -use blackings and washes, which provide a reducing atmosphere; -keep the spree filled and reduce pouring height; Increase static pressure by enlarging runner height.

3. Discontinuities Hot cracking. A crack often scarcely visible because the casting in general has not separated into fragments. The fracture surfaces may be discolored because of oxidation. The design of the casting is such that the crack would not be expected to result from constraints during cooling.

Possible Causes Damage to the casting while hot due to rough handling or excessive temperature at shakeout.

Remedies • • •

Care in shakeout and in handling the casting while it is still hot; Sufficient cooling of the casting in the mold; For metallic molds; delay knockout, assure mold alignment, use ejector pins.

4. Defective Surface Flow marks. On the surfaces of otherwise sound castings, the defect appears as lines which trace the flow of the streams of liquid metal.

Possible Causes Oxide films which lodge at the surface, partially marking the paths of metal flow through the mold.

Remedies • • • •

Increase mold temperature; Lower the pouring temperature; Modify gate size and location (for permanent molding by gravity or low pressure); Tilt the mold during pouring;

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•

In die casting: vapor blast or sand blast mold surfaces which are perpendicular, or nearly perpendicular, to the mold parting line.

5. Incomplete Casting Poured short. The upper portion of the casting is missing. The edges adjacent to the missing section are slightly rounded, all other contours conform to the pattern. The spree, risers and lateral vents are filled only to the same height above the parting line, as is the casting (contrary to what is observed in the case of defect).

Possible Causes • •

Insufficient quantity of liquid metal in the ladle; Premature interruption of pouring due to workman’s error.

Remedies • • •

Have sufficient metal in the ladle to fill the mold; Check the gating system; Instruct pouring crew and supervise pouring practice.

6. Incorrect Dimensions or Shape Distorted casting. Inadequate thickness, extending over large areas of the cope or drag surfaces at the time the mold is rammed.

Possible Causes Rigidity of the pattern or pattern plate is not sufficient to withstand the ramming pressure applied to the sand. The result is an elastic deformation of the pattern and a corresponding, permanent deformation of the mold cavity. In diagnosing the condition, the compare the surfaces of the pattern with those of the mold itself.

Remedy Assure adequate rigidity of patterns and pattern plates, especially when squeeze pressures are being increased.

7. Inclusions or Structural Anomalies Metallic Inclusions. Metallic or intermetallic inclusions of various sizes which are distinctly different in structure and color from the base material, and most especially different in properties. These defects most often appear after machining.

Possible Causes • • •

Combinations formed as intermetallics between the melt and metallic impurities (foreign impurities); Charge materials or alloy additions which have not completely dissolved in the melt; Exposed core wires or rods;

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•

During solidification, insoluble intermetallic compounds form and segregate, concentrating in the residual liquid.

Remedies • • • • •

Assure that charge materials are clean; eliminate foreign metals; Use small pieces of alloying material and master alloys in making up the charge; Be sure that the bath is hot enough when making the additions; Do not make addition too near to the time of pouring; For nonferrous alloys, protect cast iron crucibles with a suitable wash coating.

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7. Mechanical Testing

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7.1. Resilience Abstract: The ability of a material to absorb energy when deformed elastically and to return it when unloaded is called resilience. This is usually measured by the modulus of resilience, which is the strain energy per unit volume required to stress the material from, zero stress to the yield stress. The toughness of a material is its ability to absorb energy in the plastic range. The ability to withstand occasional, stresses above the yield stress without fracturing is particularly desirable in parts such as freight-car couplings, gears, chains, and crane hooks. Toughness is a commonly used concept, which is difficult to pin down and define.

The ability of a material to absorb energy when deformed elastically and to return it when unloaded is called resilience. This is usually measured by the modulus of resilience, which is the strain energy per unit volume required to stress the material from, zero stress to the yield stress σ. The strain energy per unit volume for uniaxial tension is (1) From the above definition the modulus of resilience is (2) This equation indicates that the ideal material for resisting energy loads in applications where the material must not undergo permanent distortion, such as mechanical springs, is having a high yield stress and a low modulus of elasticity. Table 1 gives some values of modulus of resilience for different materials. Table 1. Modulus of resilience for various materials Material Medium-carbon steel High-carbon spring steel Duraluminium

E, psi

s0, psi

Modulus of resilience, Ur

6

45000

33,7

140000

320

18000

17,0

4000

5,3

300

300

2000

4,0

30⋅10

6

30⋅10

6

10,5⋅10 6

Cooper

16⋅10

Rubber

150

Acrylic polymer

6

0,5⋅10

Toughness The toughness of a material is its ability to absorb energy in the plastic range. The ability to withstand occasional, stresses above the yield stress without fracturing is particularly desirable in parts such as freight-car couplings, gears, chains, and crane hooks. Toughness is a commonly used concept, which is difficult to pin down and define. One way of looking at toughness is to consider that it is the total area under the stress-strain curve. This area is an indication of the amount of work per unit

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volume, which can be done, on the material without causing it to rupture. Figure 1.2 shows the stress-strain curves for high- and low-toughness materials. The highcarbon spring steel has a higher yield strength and tensile strength than the medium-carbon structural steel. However, the structural steel is more ductile and has a greater total elongation. The total area under the stresstrain curve is greater for the structural steel, and therefore it is a tougher material. This illustrates that toughness is a parameter that comprises both strength and ductility. The crosshatched regions in Fig. 1 indicate the modulus of resilience for each steel. Because of its higher yield strength, the spring steel has the greater resilience. Several mathematical approximations for the area under the stress-strain curve have been suggested. For ductile metals that have a stress-strain curve like that of the structural steel, the area under the curve can be approximated by either of the following equations: (3) (4) For brittle materials the stress-strain curve is sometimes assumed to be a parabola, and the area under the curve is given by

(5)

Figure 1. Comparison of stress-strain curves All these relations are only approximations to the area under the stress-strain curves. Further, the curves do not represent the true behavior in the plastic range, since they are all based on the original area of the specimen.

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7.2. True Stress - True Strain Curve Abstract: The engineering stress-strain curve does not give a true indication of the deformation characteristics of a metal because it is based entirely on the original dimensions of the specimen, and these dimensions change continuously during the test. Also, ductile metal which is pulled in tension becomes unstable and necks down during the course of the test. Because the cross-sectional area of the specimen is decreasing rapidly at this stage in the test, the load required continuing deformation falls off. The average stress based on original area like wise decreases, and this produces the fall-off in the stress-strain curve beyond the point of maximum load.

The engineering stress-strain curve does not give a true indication of the deformation characteristics of a metal because it is based entirely on the original dimensions of the specimen, and these dimensions change continuously during the test. Also, ductile metal which is pulled in tension becomes unstable and necks down during the course of the test. Because the cross-sectional area of the specimen is decreasing rapidly at this stage in the test, the load required continuing deformation falls off. The average stress based on original area likewise decreases, and this produces the fall-off in the stress-strain curve beyond the point of maximum load. Actually, the metal continues to strain-harden all the way up to fracture, so that the stress required to produce further deformation should also increase. If the true stress, based on the actual cross-sectional area of the specimen, is used, it is found that the stress-strain curve increases continuously up to fracture. If the strain measurement is also based on instantaneous measurements, the curve, which is obtained, is known as a true-stress-true-strain curve. This is also known as a flow curve since it represents the basic plastic-flow characteristics of the material. Any point on the flow curve can be considered the yield stress for a metal strained in tension by the amount shown on the curve. Thus, if the load is removed at this point and then reapplied, the material will behave elastically throughout the entire range of reloading. The true stress σ is expressed in terms of engineering stress s by (1) The derivation of Eq. (1) assumes both constancy of volume and a homogenous distribution of strain along the gage length of the tension specimen. Thus, Eq. (1) should only be used until the onset of necking. Beyond maximum load the true stress should be determined from actual measurements of load and cross-sectional area. (2) The true strain εmay be determined from the engineering or conventional strain e by (3)

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Figure 1. Comparison of engineering and true stress-strain curves This equation is applicable only to the onset of necking for the reasons discussed above. Beyond maximum load the true strain should be based on actual area or diameter measurements. (4) Figure 1 compares the true-stress-true-strain curve with its corresponding engineering stress-strain curve. Note that because of the relatively large plastic strains, the elastic region has been compressed into the y-axis. In agreement with Eqs. (1) and (3), the true-stress-true-strain curve is always to the left of the engineering curve until the maximum load is reached. However, beyond maximum load the high-localized strains in the necked region that are used in Eq. (4) far exceed the engineering strain calculated from Eq. (1). Frequently the flow curve is linear from maximum load to fracture, while in other cases its slope continuously decreases up to fracture. The formation of a necked region or mild notch introduces triaxial stresses, which make it difficult to determine accurately the longitudinal tensile stress on out to fracture. The following parameters usually are determined from the true-stress-true-strain curve.

True Stress at Maximum Load The true stress at maximum load corresponds to the true tensile strength. For most materials necking begins at maximum load at a value of strain where the true stress equals the slope of the flow curve. Let σu and εu denote the true stress and true strain at maximum load when the cross-sectional area of the specimen is Au. The ultimate tensile strength is given by

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(9) Eliminating Pmax yields (5)

True Fracture Stress The true fracture stress is the load at fracture divided by the cross-sectional area at fracture. This stress should be corrected for the, triaxial state of stress existing in the tensile specimen at fracture. Since the data required for this correction are often not available, true-fracture-stress values are frequently in error.

True Fracture Strain The true fracture strain εf is the true strain based on the original area A0 and the area after fracture Af (6) This parameter represents the maximum true strain that the material can withstand before fracture and is analogous to the total strain to fracture of the engineering stress-strain curve. Since Eq. (3) is not valid beyond the onset of necking, it is not possible to calculate εf from measured values of εf. However, for cylindrical tensile specimens the reduction of area q is related to the true fracture strain by the relationship (7)

True Uniform Strain The true uniform strain εu is the true strain based only on the strain up to maximum load. It may be calculated from either the specimen cross-sectional area Au or the gage length Lu at maximum load. Equation (3) may be used to convert conventional uniform strain to true uniform strain. The uniform strain is often useful in estimating the formability of metals from the results of a tension test. (8)

True Local Necking Strain The local necking strain εn is the strain required to deform the specimen from maximum load to fracture.

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The flow curve of many metals in the region of uniform plastic deformation can be expressed by the simple power curve relation (10) where n is the strain-hardening exponent and K is the strength coefficient. A log-log plot of true stress and true strain up to maximum load will result in a straight-line if Eq. (10) is satisfied by the data (Fig. 1). The linear slope of this line is n and K is the true stress at ε = 1.0 (corresponds to q = 0.63). The strain-hardening exponent may have values from n = 0 (perfectly plastic solid) to n = 1 (elastic solid) (see Fig. 2). For most metals n has values between 0.10 and 0.50 (see Table 1.). It is important to note that the rate of strain hardening dσ /dε, is not identical with the strain-hardening exponent. From the definition of n

or (11)

Figure 2. Log/log plot of true stress-strain curve

Figure 3. Various forms of power curve σ=K* ε

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n

Table 1. Values for n and K for metals at room temperature Metal

Condition

0,05% C steel

Annealed

SAE 4340 steel

Annealed

0,60% C steel

n

K, psi

0,26

77000

0,15

93000

o

0,10

228000

o

Quenched and tempered 1000 F

0,60% C steel

Quenched and tempered 1300 F

0,19

178000

Copper

Annealed

0,54

46400

70/30 brass

Annealed

0,49

130000

There is nothing basic about Eq. (10) and deviations from this relationship frequently are observed, often at low strains (10-3) or high strains (ε≈1,0). One common type of deviation is for a log-log plot of Eq. (10) to result in two straight lines with different slopes. Sometimes data which do not plot according to Eq. (10) will yield a straight line according to the relationship (12) Datsko has shown how ε0, can be considered to be the amount of strain hardening that the material received prior to the tension test. Another common variation on Eq. (10) is the Ludwig equation (13) where σ0 is the yield stress and K and n are the same constants as in Eq. (10). This equation may be more satisfying than Eq. (10) since the latter implies that at zero true strain the stress is zero. Morrison has shown that σ0 can be obtained from the intercept of the strain-hardening point of the stress-strain curve and the elastic modulus line by

The true-stress-true-strain curve of metals such as austenitic stainless steel, which deviate markedly from Eq. (10) at low strains, can be expressed by

where eK is approximately equal to the proportional limit and n1 is the slope of the deviation of stress from Eq. (10) plotted against ε. Still other expressions for the flow curve have been discussed in the literature. The true strain term in Eqs.(10) to (13) properly should be the plastic strain εp= εtotal- εE= εtotal- σ/E

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7.3. Steel properties at low and high temperatures Abstract: Aircraft and chemical processing equipment are now required to work at subzero temperatures and the behavior of metals at temperatures down to -150°C needs consideration, especially from the point of view of welded design where changes in section and undercutting at welds may occur.

Steel Properties at Low Temperatures Aircraft and chemical processing equipment are now required to work at subzero temperatures and the behavior of metals at temperatures down to -150°C needs consideration, especially from the point of view of welded design where changes in section and undercutting at welds may occur. An increase in tensile and yield strength at low temperature is characteristic of metals and alloys in general. Copper, nickel, aluminium and austenitic alloys retain much or all of their tensile ductility and resistance to shock at low temperatures in spite of the increase in strength. In the case of unnotched mild steel, the elongation and reduction of area is satisfactory down to -130°C and then falls off seriously. It is found almost exclusively in ferritic steels, however, that a sharp drop in Izo-d value occurs at temperatures around 0°C (see Figs. 1 and 2). The transition temperature at which brittle fracture occurs is lowered by: • • • • • • •

a decrease in carbon content, less than 0,15% is desirable a decrease in velocity of deformation a decrease in depth of `notch` an increase in radius of `notch`, e.g. 6 mm minimum an increase in nickel content, e.g. 9% a decrease in grain size; it is desirable, therefore, to use steel deoxidized with aluminium normalized to give fine pearlitic structure and to avoid the presence of bainite even if tempered subsequently an increase in manganese content; Mn/C ratio should be greater than 21, preferably 8.

Figure 1. (a) Yield and cohesive stress curves (b) Slow notch bend test (c) Effect of temperature on the Izod value of mild steel

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Figure 2. Effect of low temperatures on the mechanical properties of steel in plain and notched conditions Surface grinding with grit coarser than 180 and shot-blasting causes embrittlement at -100°C due to surface work-hardening, which, however, is corrected by annealing at 650-700°C for 1 h. This heat-treatment also provides a safeguard against the initiation of brittle fracture of welded structures by removing residual stresses. Where temperatures lower than -100°C or where notch-impact stresses are involved in equipment operating below zero, it is preferable to use an 18/8 austenitic or a non-ferrous metal. The 9% Ni steel provides an attractive combination of properties at a moderate price. Its excellent toughness is due to a fine-grained structure of tough nickel-ferrite devoid of embrittling carbide networks, which are taken into solution during tempering at 570°C to form stable austenite islands. This tempering is particularly important because of the low ferrite-austenite transformation temperatures. A 4% Mn Ni (rest iron) is suitable for castings for use down to -196°C. Care should be taken to select plates without surface defects and to ensure freedom from notches in design and fabrication. Fig. 3 shows tensile and impact strengths for various alloys.

Steel Properties at High Temperatures Creep is the slow plastic deformation of metals under a constant stress, which becomes important in: 1. The soft metals used at about room temperature, such as lead pipes and white metal bearings. 2. Steam and chemical plant operating at 450-550°C. 3. Gas turbines working at high temperatures.

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Figure 3. Tensile and impact strengths of various alloys at subzero temperatures Creep can take place and lead to fracture at static stresses much smaller than those which will break the specimen when loaded quickly in the temperature range 0,5-0,7 of the melting point Tm. The Variation with time of the extension of a metal under different stresses is shown in Fig. 4a. Three conditions can be recognized: • • •

The primary stage, when relatively rapid extension takes place but at a decreasing rate. This is of interest to a designer since it forms part of the total extension reached in a given time, and may affect clearances. The secondary period during which creep occurs at a more or less constant rate, sometimes referred to as the minimum creep rate. This is the important part of the curve for most applications. The tertiary creep stage when the rate of extension accelerates and finally leads to rupture. The use of alloys in this stage should be avoided; but the change from the secondary to the tertiary stage is not always easy to determine from creep curves for some materials.

The limited nature of the information available from the creep curve is clearer when a family of curves is considered covering a range of operating stresses.

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Figure 4. a) Family of creep curves at stresses increasing from A to C b) Stress-time curves at different creep strain and repture As the applied stress decreases the primary stage decreases and the secondary stage is extended and the extension during the tertiary stage tends to decrease. Modifying the temperature of the test has a somewhat similar effect on the shape of the curves. Design data are usually given as series of curves for constant creep strain (0,010,03%, etc.), relating stress and time for a given temperature. It is important to know whether the data used are for the secondary stage only or whether it also includes the primary stage (Fig. 4b).

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In designing plants that work at temperatures well above atmospheric temperatures, the designer must consider carefully what possible maximum strains he can allow and what the final life of the plant is likely to be. The permissible amounts of creep depend largely on the article and service conditions. Examples for steel are: Rate of Creep mm/min

Time, h

Maximum Permissible Strain, mm

Turbine rotor wheels, shrunk on shafts

10-11

100000

0,0025

Steam piping, welded joints, boiler tubes

10-9

100000

0,075

Superheated tubes

10-8

20000

0,5

In designing missiles data are needed at higher temperatures and stresses and shorter time (5-60 min) than are determined for creep tests. This data is often plotted as isochronous stress-strain curves.

Creep tests For long-time applications it is necessary to carry out lengthy tests to get the design data. It is dangerous to extrapolate from short time tests, which may not produce all the structural changes, e.g. spheroidation of carbide. For alloy development and production control short time tests are used.

Long time creep tests A uniaxial tensile stress is applied by the means of a lever system to a specimen (similar to that used in tensile testing) situated in a tubular furnace and the temperature is very accurately controlled. A very sensitive mirror extensometer (of Martens type) is used to measure creep rate of 1×10-8 strain/h. From a series of tests at a single temperature, a limiting creep stress is estimated for a certain arbitrary small rate of creep, and a factor of safety is used in design.

Short time tests The rupture test is used to determine time-to-rupture under specified conditions of temperature and stress with only approximate measurement of strain by dial gauge during the course of the experiments because total strain may be around 50%. It is a useful test for sorting out new alloys and has direct application to design where creep deformation can be tolerated but fracture must be prevented.

7.4. Charpy Impact Test for Metallic Materials Abstract: Charpy impact test method for metallic materials is specified by European EN 10045 standard. This specification defines terms, dimension and tolerances of test pieces, type of the notch (U or V), test force, verification of impact testing machines etc. The test consists of breaking by one blow from a swinging pendulum, under conditions defined by standard, a test piece notched in the middle and supported at each end. The energy absorbed is

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determined in joules. This absorbed energy is a measure of the impact strength of the material.

Charpy impact test method for metallic materials is specified by European EN 10045 standard. This specification defines terms, dimension and tolerances of test pieces, type of the notch (U or V), test force, verification of impact testing machines etc. For certain particular metallic materials and applications, Charpy impact test may be the subject of specific standards and particular requirements. The test consists of breaking by one blow from a swinging pendulum, under conditions defined by standard, a test piece notched in the middle and supported at each end. The energy absorbed is determined in joules. This absorbed energy is a measure of the impact strength of the material. The designations applicable to this standard are as indicated in the Table 1 and on the Figure1. Table 1. Characteristics of test piece and testing machine Reference (Figure 1)

Designation

Unit

1

Length of test piece

mm

2

Height of test piece

mm

3

Width of test piece

mm

4

Height below notch

mm

5

Angle of notch

6

Radius of curvature of base of notch

mm

7

Distance between anvils

mm

8

Radius of anvils

mm

9

Angle of taper of each anvil

Degree

10

Angle of taper of striker

Degree

11

Radius of curvature of striker

mm

12

Width of striker

mm

Energy absorbed by breakage KU or KV

Joule

-

Degree

320

Figure 1. Charpy impact test

Test pieces The standard test piece shall be 55 mm long and of square section with 10 mm sides. In the centre of the length, there shall be a notch. Two types of notch are specified:

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a. V notch of 45°, 2 mm deep with a 0,25 mm radius of curve at the base of notch. If standard test piece cannot be obtained from the material, a reduced section with a width of 7,5 mm or 5 mm shall be used, the notch being cut in one of the narrow faces. b. U notch or keyhole notch, 5 mm deep, with 1 mm radius of curve at the base of notch. The test pieces shall be machined all over, except in the case of precision cast foundry test pieces in which the two faces parallel to the plane of symmetry of the notch can be unmachined. The plane of symmetry of the notch shall be perpendicular to the longitudinal axis of the test piece. The tolerances of the specified dimensions of the test piece are given by standard as well. For the standard test piece, machining tolerance in length is 0.6 mm for both type of tests, and tolerances in height are 0.11 mm for U and 0.06 mm for V notch test piece. Tolerances for angle between plane of symmetry of the notch and longitudinal axis of test piece as well as for angle between adjacent longitudinal faces of test piece are ± 2° only. Comparison of results is only of significance when made between test pieces of the same form and dimensions. Machining shall be carried out in such a way that any alternation of the test piece, for example due to cold working or heating, is minimized. The notch shall be carefully prepared so that no grooves, parallel to the base of the notch, are visible to the naked eye. The test piece may be marked on any face not in contact with the supports or anvils and at a point at least 5 mm from the notch to avoid the effects of cold working due to marking.

Testing machine The testing machine shall be constructed and installed rigidly and shall be in accordance with European Standard 10 045 part 2. Standard test condition shall correspond to nominal machine energy of 300±10J at the use of a test piece of standard dimensions. The reported absorbed energy under these conditions shall be designated by the following symbols: • •

KU for a U notch test piece KV for a V notch test piece

Testing machines with different striking energies are permitted, in which case the symbol KU or KV shall be supplemented by an index denoting the energy of the testing machine. For example KV 150 denotes available energy of 150 J, and KU 100 denotes available energy of 100 J. KU 100 = 65 J means that: • • •

nominal energy is100 J standard U notch test piece is used energy absorbed during fracture is 65 J.

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For tests on a subsidiary V notch test piece, the KV symbol shall be supplemented by indices denoting first the available energy of the testing machine and second the width of the test piece, e.g.: • • •

KV 300 / 7,5: available energy 300 J, width of test piece 7.5 mm KV 150 / 5: available energy 150 J, width of test piece 5 mm KV 150 / 7,5 = 83 J denotes: o nominal energy 150 J o reduced section test piece of width 7,5 mm o energy absorbed during fracture: 83 J.

The test piece shall lie against the anvils in such a way that the plane of symmetry of the notch shall be no more than 0.5 mm from the plane of symmetry of the anvils. If the test temperature is not specified in the product standard, it shall be about 23°C. National standards corresponding to EN 10045-2 are DIN 51306 (1983), NFA 03-508 (1985), BS 131 Part 4 (1972) and international ISO 442 (1965).

7.5. Effect of Metallurgical Variables on Fatigue Abstract: The fatigue properties of metals are quite structure-sensitive. However, at the present time there are only a limited number of ways in which the fatigue properties can be improved by metallurgical means. By far the greatest improvements in fatigue performance result from design changes, which reduce stress concentration and from the intelligent use of beneficial compressive residual stress, rather than from a change in material. Nevertheless, there are certain metallurgical factors, which must be considered to ensure the best fatigue performance from a particular metal or alloy.

The fatigue properties of metals are quite structure-sensitive. However, at the present time there are only a limited number of ways in which the fatigue properties can be improved by metallurgical means. By far the greatest improvements in fatigue performance result from design changes, which reduce stress concentration and from the intelligent use of beneficial compressive residual stress, rather than from a change in material. Nevertheless, there are certain metallurgical factors, which must be considered to ensure the best fatigue performance from a particular metal or alloy. Fatigue tests designed to measure the effect of some metallurgical variable, such as special heat treatments, on fatigue performance are usually made with smooth, polished specimens under completely reversed stress conditions. It is usually assumed that any changes in fatigue properties due to metallurgical factors will also occur to about the same extent under more complex fatigue conditions, as with notched specimens under combined stresses. Fatigue properties are frequently correlated with tensile properties. In general, the fatigue limit of cast and wrought steels is approximately 50 percent of the ultimate tensile strength. The ratio of the fatigue limit (or the fatigue strength at 106 cycles) to the tensile strength is called the fatigue ratio. Several nonferrous metals such as nickel, copper, and magnesium have a fatigue ratio of about 0.35. While the use of correlations of this type is convenient, it should be clearly understood that these constant factors between fatigue limit and tensile

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strength are only approximations and hold only for the restricted condition of smooth, polished specimens which have been tested under zero mean stress at room temperature. For notched fatigue specimens the fatigue ratio for steel will be around 0,20 to 0,30. However, as yield strength is increased by the various strengthening mechanisms, the fatigue limit usually does not increase proportionately. Most high-strength materials are fatigue-limited. Several parallels can be drawn between the effect of certain metallurgical variables on fatigue properties and the effect of these same variables on tensile properties. The effect of solid-solution alloying additions on the fatigue properties of iron and aluminum parallels nearly exactly their effect on the tensile properties. Gensamer showed that the fatigue limit of a eutectoid steel increased with decreasing isothermal-reaction temperature in the same fashion as did the yield strength and the tensile strength. However, the greater structure sensitivity of fatigue properties, compared with tensile properties, is shown in tests comparing the fatigue limit of a plain carbon eutectoid steel heat-treated to coarse pearlite and to spheroidite of the same tensile strength. Even though the steel in the two structural conditions had the same tensile strength, the pearlitic structure resulted in a significantly lower fatigue limit due to the higher notch effects of the carbide lamellae in pearlite. There is good evidence that high fatigue resistance can be achieved by homogenizing slip deformation so that local concentrations of plastic deformation are avoided. This is in agreement with the observation that fatigue strength is directly proportional to the difficulty of dislocation cross slip. Materials with high stacking-fault energy permit dislocations to cross slip easily around obstacles, which promotes slip-band formation and large plastic zones at the tips of cracks. Both of these phenomena promote the initiation and propagation of fatigue cracks. In materials with low stacking-fault energy, cross slip is difficult and dislocations are constrained to move in a more planar fashion. This limits local concentrations of plastic deformation and suppresses fatigue damage. While the concept has been useful in understanding fatigue mechanisms, the ability to control fatigue strength by altering stacking-fault energy has practical limitations. A more promising approach to increasing fatigue strength appears to be the control of microstructure through thermomechanical processing to promote homogeneous slip with many small regions of plastic deformation as opposed to a smaller number of regions of extensive slip. The dependence of fatigue life on grain size varies also depending on the deformation mode. Grain size has its greatest effect on fatigue life in the low-stress, high-cycle regime in which stage 1 cracking predominates. In high stacking-faultenergy materials (such as aluminum and copper) cell structures develop readily and these control the stage 1 crack propagation. Thus, the dislocation cell structure masks the influence of grain size, and fatigue life at constant stress is insensitive to grain size. However, in a low slacking-fault-energy material (such as alpha brass) the absence of cell structure because of planar slip causes the grain boundaries to

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control the rate of cracking. In this case, fatigue life is proportional to grain diameter. In general, quenched and tempered microstructures result in the optimum fatigue properties in heat-treated low-alloy steels. However, at a hardness level above about Rc 40, a bainitic structure produced by austempering results in better fatigue properties than a quenched and tempered structure with the same hardness. Electron micrographs indicate that the poor performance of the quenched and tempered structure is the result of the stress-concentration effects of the thin carbide films that are formed during the tempering of martensite. For quenched and tempered steels the fatigue limit increases with decreasing tempering temperature up to a hardness of Rc 45 to Rc 55, depend on the steel. The fatigue properties at high hardness levels are extremely sensitive to surface preparation, residual stresses, and inclusions. The presence of only a trace of decarburization on the surface may drastically reduce the fatigue properties. Only a small amount of non-martensitic transformation products can cause an appreciable reduction in the fatigue limit. The influence of small amounts of retained austenite on the fatigue properties of quenched and tempered steels has not been well established. The results indicate that below a tensile strength of about 200,000 psi (~1400 MPa) the fatigue limits of quenched and tempered low-alloy steels of different chemical composition are about equivalent when the steels are tempered to the same tensile strength. This generalization holds for fatigue properties determined in the longitudinal direction of wrought products. However, tests have shown that the fatigue limit in the transverse direction of steel forcing may be only 60 to 70 percent of the longitudinal fatigue limit. It has been established that practically all the fatigue failures in transverse specimens start at nonmetalic inclusions. Nearly complete elimination of inclusions by vacuum melting produces a considerable increase in the transverse fatigue limit. The low transverse fatigue limit in steels containing inclusions is generally attributed to stress concentration at the inclusions, which can be quite high if an elongated inclusion stringer is oriented transverse to the principal tensile stress. However, the fact that nearly complete elimination of inclusions by vacuum melting still results in appreciable anisotropy of the fatigue limit indicates that other factors may be important. Further investigations of this subject have shown that appreciable changes in the transverse fatigue limit which cannot be correlated with changes in the type, number, or size of inclusions are produced by different deoxidation practices. Transverse fatigue properties appear to be one of the most structuresensitive engineering properties. The existence of a fatigue limit in certain materials, especially iron and, titanium alloys, has been shown to depend on the presence of interstitial elements. The S-N curve for a pure metal will be a monotonic function with N increasing as stress decreases. The introduction of a solute element raises the yield strength and since it is more difficult to initiate a slip band, the S-N curve is shifted upward and to the right. If the alloy has suitable interstitial content so it undergoes strain aging, there is an additional strengthening mechanism.

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Since strain aging will not be a strong function of applied stress, there will be some limiting stress at which a balance occurs between fatigue damage and localized strengthening due to strain aging. With enhanced strain aging, brought about by higher interstitial content or elevated temperature the fatigue limit is raised and the break in the curve occurs at a lower number of cycles. In quenched and tempered steels, which do not normally exhibit strain aging in the tension test, the existence of a pronounced fatigue limit presumably is the result of localized strain aging at the tip of the crack.

7.6. Hardness Testing Abstract: The hardness of a material is a poorly defined term which has many meanings depending upon the experience of the person involved. In general, hardness usually implies a resistance to deformation, and for metals the property is a measure of their resistance to permanent or plastic deformation. There are three general types of hardness measurements depending on the manner in which the test is conducted. These are:

• • •

scratch hardness indentation hardness, and rebound, or dynamic, hardness.

The hardness of a material is a poorly defined term which has many meanings depending upon the experience of the person involved. In general, hardness usually implies a resistance to deformation, and for metals the property is a measure of their resistance to permanent or plastic deformation. To a person concerned with the mechanics of materials testing, hardness is most likely to mean the resistance to indentation, and to the design engineer it often means an easily measured and specified quantity which indicates something about the strength and heat treatment of the metal. There are three general types of hardness measurements depending on the manner in which the test is conducted. These are: • • •

scratch hardness indentation hardness, and rebound, or dynamic, hardness.

Only indentation hardness is of major engineering interest for metals. Scratch hardness is of primary interest to mineralogists. With this measure of hardness, various minerals and other materials are rated on their ability to scratch one another. Scratch hardness is measured according to the Mohs’ scale. This consists of 10 standard minerals arranged in the order of their ability lo be scratched. The softest mineral in this scale is talc (scratch hardness 1), while diamond has a hardness of 10. The Mohs’ scale is not well suited for metals since the intervals are not widely spaced in the high-hardness range. Most hard metals fall in the Mohs’ hardness range of 4 to 8. In dynamic-hardness measurements the indenter is usually dropped onto the metal surface, and the hardness is expressed as the energy of impact. The Shore

326

seleroscope, which is the commonest example of a dynamic-hardness tester, measures the hardness in terms of the height of rebound of the indenter.

Brinell Hardness The first widely accepted and standardized indentation-hardness test was proposed by J. A. Brinell in 1900. The Brinell hardness test consists in indenting the metal surface with a 10-mm-diameter steel ball at a load of 3,000 kg mass ( 29400 N). For soft metals the load is reduced to 500 kg to avoid too deep an impression, and for very hard metals a tungsten carbide ball is used to minimize distortion of the indenter. The load is applied for a standard time, usually 30 s, and the diameter of the indentation is measured with a low-power microscope after removal of the load. The average of two readings of the diameter of the impression at right angles should be made. The Brinell hardness number (BHN) is expressed as the load P divided by the surface area of the indentation. This is expressed by the formula:

where

P - applied load, N D - diameter of ball mm d - diameter of indentation, mm t - depth of the impression, mm

It will be noticed that the units of the BHN are MPa. Unless precautions are taken to maintain P/D2 constant, which may be experimentally inconvenient, the BHN generally will vary with load. Over a range of loads the BHN reaches a maximum at some intermediate load. Therefore, it is not possible to cover with a single load the entire range of hardnesses encountered in commercial metals. The relatively large size of the Brinell impression may be an advantage in averaging out local heterogeneities. Moreover, the Brinell test is less influenced by surface scratches and roughness than other hardness tests. On the other hand, the large size of the Brinell impression may preclude the use of this test with small objects or in critically stressed parts where the indentation could be a potential site of failure.

Meyer Hardness Meyer suggested that a more rational definition of hardness than that proposed by Brinell would be one based on the projected area of the impression rather than the surface area. The mean pressure between the surface of the indenter and the indentation is equal to the load divided by the projected area of the indentation. Meyer proposed that this mean pressure should be taken as the measure of hardness. It is referred to as the Meyer hardness.

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Like the Brinell hardness, Meyer hardness has units of MPa. The Meyer hardness is less sensitive to the applied load than the Brinell hardness. For a cold-worked material the Meyer hardness is essentially constant and independent of load, while the Brinell hardness decreases as the load increases. For an annealed metal the Meyer hardness increases continuously with the load because of strain hardening produced by the indentation. The Brinell hardness, however, first increases with load and then decreases for still higher loads. The Meyer hardness is a more fundamental measure of indentation hardness; yet it is rarely used for practical hardness measurements. Meyer proposed an empirical relation between the load and the size of the indentation. This relationship is usually called Meyer’s law. P = kdn’ The parameter n’ is the slope of the straight line obtained when log P is plotted against log d, and k is the value of P at d = 1. Fully annealed metals have a value of n’ of about 2.5, while n’ is approximately 2 for fully strain-hardened metals. This parameter is roughly related to the strain-hardening coefficient in the exponential equation for the true-stress-true-strain curve. The exponent in Meyer’s law is approximately equal to the strain-hardening coefficient plus 2.

Vickers Hardness The Vickers hardness test uses a square-base diamond pyramid as the indenter. The included angle between opposite faces of the pyramid is 136°. This angle was chosen because it approximates the most desirable ratio of indentation diameter to ball diameter in the Brinell hardness test. Because of the shape of the indenter, this is frequently called the diamond-pyramid hardness test. The diamond-pyramid hardness number (DPH), or Vickers hardness number (VHN, or VPH), is defined as the load divided by the surface area of the indentation. In practice, this area is calculated from microscopic measurements of the lengths of the diagonals of the impression. The DPH may be determined from the following equation:

where P - applied load, kg L - average length of diagonals, mm θ - angle between opposite faces of diamond = 136° The Vickers hardness test has received fairly wide acceptance for research work because it provides a continuous scale of hardness, for a given load, from very soft metals with a DPH of 5 to extremely hard materials with a DPH of 1,500. The Vickers hardness test is described in ASTM Standard E92-72.

Rockwell Hardness Test

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The most widely used hardness test is the Rockwell hardness test. Its general acceptance is due to its speed, freedom from personal error, ability to distinguish small hardness differences in hardened steel, and the small size of the indentation, so that finished heat-treated parts can be tested without damage. This test utilizes the depth of indentation, under constant load, as a measure of hardness. A minor load of 10 kg is first applied to seat the specimen. This minimizes the amount of surface preparation needed and reduces the tendency for ridging or sinking in by the indenter. The major load is then applied, and the depth of indentation is automatically recorded on a dial gage in terms of arbitrary hardness numbers. The dial contains 100 divisions, each division representing a penetration of 0.00008 in (0.002 mm). The dial is reversed so that a high hardness, which corresponds to a small penetration, results in a high hardness number. This is in agreement with the other hardness numbers described previously, but unlike the Brinell and Vickers hardness designations, which have units of MPa, the Rockwell hardness numbers are purely arbitrary. Major loads of 60, 100, and 150 kg are used. Since the Rockwell hardness is dependent on the load and indenter, it is necessary to specify the combination which is used. This is done by prefixing the hardness, number with a letter indicating the particular combination of load and indenter for the hardness scale employed. A Rockwell hardness number without the letter prefix is meaningless. Hardened steel is tested on the C scale with the diamond indenter and a 150-kg major load. The useful range for this scale is from about RC 20 to RC 70. Softer materials are usually tested on the B scale with a 1/16-in-diameter steel ball and a 100-kg major load. The range of this scale is from RB 0 to RB 100. The A scale (diamond penetrator, 60-kg major load) provides the most extended Rockwell hardness scale, which is usable for materials from annealed brass to cemented carbides. Many other scales are available for special purposes. The Rockwell hardness test is a very useful and reproducible one provided that a number of simple precautions are observed. Most of the points filled below apply equally well to the other hardness tests: • • • • •

•

The indenter and anvil should be clean and well seated. The surface to be tested should be clean and dry, smooth, and free from oxide. A rough-ground surface is usually adequate for the Rockwell test. The surface should be flat and perpendicular to the indenter. Tests on cylindrical surfaces will give low readings, the error depending on the curvature, load, indenter, and hardness of the material. Theoretical and empirical corrections for this effect have been published. The thickness of the specimen should be such that a mark or bulge is not produced on the reverse side of the piece. It is recommended that the thickness be at least 10 times the depth of the indentation. The spacing between indentations should be three to five times the diameter of the indentation. The speed of application of the load should be standardized. This is done by adjusting the dashpot on the Rockwell tester. Variations in hardness can be

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appreciable in very soft materials unless the rate of load application is carefully controlled.

Microhardness Tests Many metallurgical problems require the determination of hardness over very small areas. The measurement of the hardness gradient at a carburized surface, the determination of the hardness of individual constituents of a microstructure, or the checking of the hardness of a delicate watch gear might be typical problems. The use of a scratch-hardness test for these purposes was mentioned earlier, but an indentation-hardness test has been found to be more useful. The development of the Knoop indenter by the National Bureau of Standards and the introduction of the Tukon tester for the controlled application of loads down to 25 g have made micro hardness testing a routine laboratory procedure. The Knoop indenter is a diamond ground to a pyramidal form that produces a diamond-shaped indentation with the long and short diagonals in the approximate ratio of 7:1 resulting in a state of plane strain in the deformed region. The Knoop hardness number (KHN) is the applied load divided by the unrecovered projected area of the indentation. The special shape of the Knoop indenter makes it possible to place indentations much closer together than with a square Vickers indentation, e.g., to measure a steep hardness gradient. The other advantage is that for a given long diagonal length the depth and area of the Knoop indentation are only about 15 percent of what they would be for a Vickers indentation with the same diagonal length. This is particularly useful when measuring the hardness of a thin layer (such as an electroplated layer), or when testing brittle materials where the tendency for fracture is proportional to the volume of stressed material.

Hardness at Elevated Temperatures Interest in measuring the hardness of metals at elevated temperatures has been accelerated by the great effort which has gone into developing alloys with improved high-temperature strength. Hot hardness gives a good indication of the potential usefulness of an alloy for high-temperature strength applications. In an extensive review of hardness data at different temperatures, Westbrook showed that the temperature dependence of hardness could be expressed by H = Ae-BT where

H = hardness, kg/mm2 T = test temperature, K A,B constants

Plots of log H versus temperature for pure metals generally yield two straight lines of different slope. The change in slope occurs at a temperature which is about one-half the melting point of the metal being tested. Similar behavior is found in plots of the logarithm of the tensile strength against temperature. Above mentioned figure shows

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this behavior for copper. It is likely that this change in slope is due to a change in the deformation mechanism at higher temperature. The constant A derived from the low-temperature branch of the curve can be considered to be the intrinsic hardness of the metal, that is, H at 0 K. Westbrook correlated values of A for different metals with the heat content of the liquid metal at the melting point and with the melting point. This correlation was sensitive to crystal structure. The constant B, derived from the slope of the curve, is the temperature coefficient of hardness. This constant was related in a rather complex way to the rate of change of heat content with increasing temperature. With these correlations it is possible to calculate fairly well the hardness of a pure metal as a function of temperature up to about one-half its melting point.

7.7. Magneto-Inductive Verification of Material Characteristics Abstract: The magneto-inductive method is ideally suitable for automatic 100% testing. Such characteristics as the surface hardness and case hardening depth can be verified without destruction using this method. This not only ensures greater reliability with regard to the quality of the individual products, but also guarantees that defects can be identified and remedied during the production process since the test is integrated into the production line. When testing hardened components, the test unit can be linked to a control computer performing SPC analyses and thus permitting automatic control of the hardening furnace.

Compliance with specified material characteristics is an important part of quality control when manufacturing high-quality industrial products. Even today, however, quality control is still only based on random samples in a large number of cases. The relatively large test effort required and the sometimes inevitable destruction of the test specimen frequently makes it impossible to carry out 100% tests using classical methods, such as determining the case hardening depth of surface-hardened components and the mechanical strength of sheet steel. The magneto-inductive method, on the other hand, is ideally suitable for automatic 100% testing. Such characteristics as the surface hardness and case hardening depth can be verified without destruction using this method. This not only ensures greater reliability with regard to the quality of the individual products, but also guarantees that defects can be identified and remedied during the production process since the test is integrated into the production line. When testing hardened components, the test unit can be linked to a control computer performing SPC analyses and thus permitting automatic control of the hardening furnace. Compared with classical materials testing, however, a new approach will be necessary when using the magneto-inductive method, for it is not an absolute-value method and must be calibrated with the aid of parameters which may have to be ascertained destructively. Physical principles of magneto-inductive testing. The magneto-inductive test method is based on an electromagnetic field built up by a field coil and modified by the presence of a test piece with conductivity and permeability μ. This makes the method suitable for all technological variables correlating with μ. The instrumentation comprises either a field coil and a sensor coil embracing the test specimen or several

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such coils of smaller diameter which, together with a highly permeable core, form what is known as a probe (Fig. 1).

Figure 1. Two basic arrangment for magneto-inductive testing In the first case, the specimen forms the core of this transformer-type arrangement; the interaction between field and specimen is most effective when the inner volume of the field coil has been filled completely. In the second case, the probe is placed on the surface of the specimen in order to carry out the test. The field strengths produced in the material in this way are smaller than those produced by the embracing coils. However, the test area is pinpointed more precisely with this method and these probes can also be used to test large parts, as well as parts with complex geometry. Permeability characteristics are extremely informative parameters for a variety of material states due, for instance, to hardening processes, particularly for materials with a relative permeability of well over 1. The properties of a hysteresis loop create additional effects in the sensor coil signal, its amplitude depending on the maximum magnetization while the existence of coercive field strength alters the signal phase relative to the exciting phase and the non-linearity of the hysteresis generates harmonics 3f, 5f, 7f, etc. The maximum information available can only be derived from the sensor signal if all these effects are included in the analysis. In addition, however, the form of the hysteresis depends on the test frequency f and on the field amplitude in a manner characteristic of the material state. This makes it clear that considerably more information about the material properties can be obtained by testing with several frequencies and amplitudes. The result is a multi-parameter test. Since different test frequencies also penetrate into the material to different depths, the test frequency can be used to achieve a certain degree of depth selection. Integral and superficial properties can be identified simultaneously by appropriately analyzing the results obtained with several test frequencies. Calibration based on representative parts with known properties is an essential prerequisite for magneto-inductive testing, since the precise shape and unambiguous relationship between the measuring signals and the required technological

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characteristic are dependent on the momentary test task and therefore not immediately known. The relationships are determined empirically using statistical methods (Fig. 2). Only the relationships between the physical variables μ,... (and the geometry within limits) and the magneto-inductive signal are predicted theoretically. The effect of primary influencing variables on μ and ... on the one hand and on technlogical characteristics on the other, however, cannot be predicted quantitatively. For this reason, the relationship between technological characteristics and the measuring signal can only be determined empirically in the test task concerned. The sequence of work for every test task is therefore as follows: setting the test parameters - calibration - testing.

Figure 2. Schematic illustration of the relationships prevailing in magneto-inductive testing. Properties of a computer-controlled test system. The Magnatest is a computercontrolled modular test unit with one or more test channels containing the actual test electronics and a computer unit as controller. The capabilities offered by this system include the following: • • • • •

100% on-line testing of material identity and hardness Case-hardening depth, strength and many other material properties Integration into CAQ systems Quantitative determination of test values for statistical process control Storage of test settings and results.

Two different evaluation methods have been implemented in the Magnatest. Group analysis is the classical method for sorting test specimens into two or more classes. The second evaluation method is based on a regression algorithm and is particularly suitable for test tasks involving quantifiable properties, i.e. properties which can be described numerically. The fact that regression analysis permits greater differentiation of the test results can be used to great advantage to control the production parameters. Such a control loop, in which SPC methods may be applied, can only fulfill its intended purpose - preventing production of reject parts - if it can also identify minute deviations from the product quality.

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This system can also be used for multi-parameter testing, i.e. combined testing with several test settings. The system helps the user to find the optimum test parameters. For single-parameter tests, up to 24 settings can initially be specified and the system then calculates the reliability, the selectivity of group analysis and regression correlation for each setting. The user can then select the required test setting. Up to 24 (individual) settings are also specified for multi-parameter tests. The combination of test parameters yielding the greatest reliability for the test task is then calculated from the 24 specified settings by a program running on a separate PC linked to the test system. Applications. The magneto-inductive method can basically be used to test all material properties associated with changes in conductivity or permeability, such as alloy fluctuations, strength, core hardness, surface hardness, case-hardening depth, heat treatment, residual austenite content, soft-spottiness, cementite accumulation and surface decarburization. However, each specific test task must be preceded by a test phase in which setting and optimization aids are used to determine whether the test property is reflected clearly enough in the magneto-inductive signal. Only a small number of parameters may vary in addition to the property under investigation if the test is to prove successful. Testing of the case-hardening depth is a typical application in which electromagnetic methods are effectively the only non-destructive methods available. The mean deviation for both methods equals 0.08 mm. Three sets of parameters were combined without optimizing the selection. The deviation was reduced to 0.054 mm, by using an optimized test setting. This shows that the reliability of this method can be improved "economically", i.e. without increasing the number of measured quantities and the associated disadvantages, such as longer test times. On-line testing of sheet steel is a promising application that is currently still in the trial stage. Four sensors are installed side by side so that the tensile strength of the sheet can be tested over the full width as it is transported. Maintaining a constant (small) distance between sheet and sensor is a factor of critical importance for tests on moving sheets. The mechanically ascertained strength values have been plotted against the results of magneto-inductive tests for a sample series. Two values were obtained for each specimen in the magneto-inductive tests, for they permit testing of both the upper and the lower sides. Any differences between the two sides of the sheet due to the rolling process can be identified in this way. Modern magneto-inductive test units can be used to determine technological material characteristics thanks to the implementation of complex algorithms for evaluation. The magneto-inductive method is predestined for use in automatic non-destructive 100% tests integrated into the production line. When using this method, it is important not to forget that it is not an absolute-value method and that a great deal depends on the care taken during calibration. If this point is noted, the accuracy and variability achieved with the new evaluation capabilities make magneto-inductive testing an ideal complement or, substitute for the random tests based on classical destructive or semi-destructive methods.

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7.8. Ultrasonic Testing of Safety Parts in Automobile Manufacturing Abstract: It is a matter of choice whether the ultrasonic testing of safety parts takes place on-line or off-line. Longitudinal and transverse flaws, internal and external flaws, such as shrink-holes, cracks, and inclusions etc., must be nondestructively tested according to the terms of delivery arranged between supplier and customer. In ultrasonic testing high-frequency sound waves produced by a piezo-crystal element are transferred into the test piece through a liquid coupling medium, usually water. If there are flaws in the test piece -- they are called discontinuities -- and if these are encountered by the sonic beam, the sound waves are reflected by these discontinuities.

The Principles of Ultrasonic Testing It is a matter of choice whether the ultrasonic testing of safety parts takes place online or off-line. Longitudinal and transverse flaws, internal and external flaws, such as shrink-holes, cracks, and inclusions etc., must be nondestructively tested according to the terms of delivery arranged between supplier and customer. In ultrasonic testing high-frequency sound waves produced by a piezo-crystal element are transferred into the test piece through a liquid coupling medium, usually water. If there are flaws in the test piece -- they are called discontinuities -- and if these are encountered by the sonic beam, the sound waves are reflected by these discontinuities. The reflected sound waves are picked up by the crystal element ("received") and transformed back into electric signals. The ultrasonic instrument processes the electric signals and represents the flaws as echos on the monitor. The time of flight, the amplitude and sometimes the echo shape are called on to evaluate the detected discontinuities.

Carrying Out the Test in the Factory For a statistical test or random sampling off-line it is usually enough to make a manual test with a portable ultrasonic testing device and one probe. The critical zones are sonically beamed and tested for discontinuities. This method is called direct contact mode and the probe is coupled directly to the surface of the material through a coupling medium. Automatic ultrasonic test machines designed for on-line tests in series usually employ the immersion technique. The safety part which is to be tested is automatically inserted into a water tank where several probes are mounted corresponding to the number of desired testing positions. For the automatic evaluation of the test results the electronics of such a machine include the so-called "monitor gates". These are placed in such a way that they cover the critical zones of the test piece. If echos enter these gates and if they exceed in amplitude a pre-adjustable threshold, a "bad" signal is produced. In this way pieces which are found to be flawed can be identified optically or accoustically, marked with color or appropriately sorted.

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In order to simplify communication with other units in the production line it is possible to use a computerized control system to coordinate the functions and the processes. Depending on the desired degree of automation and the geometrical forms of the test pieces, various solutions are possible for transporting the test pieces. All kinds of custom made equipment can be supplied from the simplest manual insertion to loading and unloading with the help of chutes, conveyor belts, transference with pick-up arms or intake and outtake with chain conveyors.

Test Examples for Various Safety Parts Testing ball bearings. Ball bearings are usually tested for longitudinal and transverse flaws by turning the piece 360°. As adjusting aids original pieces with artificially introduced flaws are used: e.g. a notch 0.3 mm deep and 0.5 mm wide. The immersion technique is used and the coupling medium is water. As a rule, test frequencies between 5 MHz and 10 MHz are employed. In the test cycle a function check is carried out by means of an additional through transmission cycle. This assures that no test piece leaves the testing machine untested. In practice, cycle times of about 3600 pieces per hour are reached with an automatic piece testing machine. Testing steering columns. Steering columns are among the safety parts which are especially subject to cracks because of the special processing by cold forming. The parts are tested for external and internal longitudinal cracks. In the transition zone to the pivot an additional quasitransverse flaw test is carried out. The testing is performed using the immersion technique. The coupling medium is thin-bodied oil. As a rule longitudinally focusing probes are used with a test frequency of 4 MHz. With one probe transverse flaws in the transition zone can be detected and with two more probes longitudinal flaws in the cylindrical zone. Testing of cardan shafts. In front-drive automobiles the cardan shafts are among the safety parts. Here it is necessary to detect internal and external longitudinal flaws in the cylindrical housing and tranverse flaws in the internal transition zones from the housing to the pivot. This automatic test is also carried out with the immersion technique using water as a coupling medium. The test frequencies are 4-5 MHz and focusing probes are used. All probes are fixed but they can be moved with an additional transverse drive in order to increase the testing density. The piece is turned in a water tank at least 360°. A function check of the transverse flaws takes place by means of the initial pulse and of the longitudinal flaws by means of a through-transmission cycle. This ultrasonic test is performed with automatic testing machines which are fully integrated into a production line and reach cycle times of about 8 seconds. Testing of pre-chamber housings. Especially in welding technology such welding techniques as electron-beam welding and laser welding have become established. But even today the technique must be controlled at every point, because not even the smallest welding defects can be permitted in safety parts. Like most automatic ultrasonic tests the testing of electron beam welding on the pre-chamber housings of diesel engines is carried out with the immersion technique. These parts are scanned spirally with special probes with a water delay line of about 20 mm corresponding to the admitted zone which is not welded.

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Testing pistons and piston pins. When testing bonding of piston ring supports one of the best known testing methods in automobile manufacture is employed. With the aid of special probes pistons are tested with the immersion technique using water or thin-bodied oil as a coupling medium. The special probes are adapted to the geometrical forms of the pistons and can usually be applied to particular geometrical zones. Testing valves. Valves of various dimensions should be tested for welding defects between the stellit welding and the valve head and/or the friction welding on the shaft of the valve. Bonding and welding defects should be detected. With the aid of probes of the types H5M with a water delay line these parts are tested using the immersion technique. The coupling medium is normally water. The water delay line is about 20 mm. The probes can be arranged in a variety of ways: the bonding tests as in the sketch or in the through-transmission mode; for the friction welding test a "tandem arrangement" is possible. Testing of drive shafts, external lamellar supports, sun gears and synchronizing rings. Because of the need to save weight and to use rational production methods and also with regard to the development of improved welding procedures gear parts today are often welded with laser or electron beams. Ultrasonics is used to test so called weld drifts applying the immersion technique; in the case of complex geometrical forms the sound beam is directed into the test piece with deflecting mirrors at 90°. The machine is calibrated on a master piece and in special cases on natural flaws. Usually, testing mechanics which include turntables are employed. For automatic testing a handling robot can insert the parts and sort them out. Testing of pins and wheel bearings. Forging defects and internal material flaws are not permissible in safety parts. The flaws that occur demand a probe arrangement appropriate to each case. This ultrasonic test is usually supplemented with an x-ray test (random sampling) and a magnetoscopic test. Depending on the location, size and position of the flaw, various probes with water delay lines are used. A wide range of automatic testing machines and even simple immersion technique testing machines with turntable or roller block can be employed. A selection of these testing machines depends on the parts to be produced and the desired cycle time. Testing of cup valve tappets. In high-powered motors and low-pollution engine units, cup valve tappets welded with laser beams are used, making the adjustment of the tappet clearance unnecessary. Welding defects and weld drifts should be detected. In the immersion technique these parts are usually tested with an immersion technique probe. A master part with an artificially introduced test flaw serves as a calibration aid in ultrasonic testing. In these cases a single-channel ultrasonic flaw detector is usually employed since one probe is normally enough. Testing oftrock axles. In heavy-transport trucks with their extreme requirements the testing of safety parts is just as necessary as in the case of cars. Truck axles are usually welded from profiles. Here, the problem is to detect welding defects in the complete weld area. The immersion technique is used with the probes placed at calculated distances and at defined incident angles of acoustic waves in order to include the complete area in accordance with the law of refraction. Calibration is performed on test flaws (drill holes and/or notches). A function check is carried out

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by means of a separate through-transmission cycle. For the evaluation of the test results, multi-channel test electronic instruments or packages are used.

8. Fracture Mechanics

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8.1. Fracture Abstract: Most alloys contain second phases which lose cohesion with the matrix or fracture and the voids so formed grow as dislocations flow into them. Coalescence of the voids forms a continuous fracture surface followed by failure of the remaining annulus of material usually on plane at 45° to the tension axis. The central fracture surface consists of numerous cup-like depressions generally called dimples. The shape of the dimples is strongly influenced by the direction of major stresses-circular in pure tension and parabolic under shear

Ductile fracture A pure and inclusion free metal can elongate under tension to give approx. 100% RA and a point fracture, Fig. 1. Most alloys contain second phases which lose cohesion with the matrix or fracture and the voids so formed grow as dislocations flow into them. Coalescence of the voids forms a continuous fracture surface followed by failure of the remaining annulus of material usually on plane at 45° to the tension axis. The central fracture surface consists of numerous cup-like depressions generally called dimples. The shape of the dimples is strongly influenced by the direction of major stresses-circular in pure tension and parabolic under shear. Dimple size depends largely on the number of inclusion sites. Fig. 2a shows typical dimples.

Figure 1. (a) Stages in ductile fracture from inclusions (b) Fracture toughness n thickness Some important features of ductile fracture can be summarised as follows:

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• • • •

Pure metals and solid solutions that are relatively free from second phase particles (including impurity particles) are usually more ductile than strong two-phase alloys. The local stress required for whole nucleation at particles depends on their resistance to cracking and the strength of their bond with the matrix. The local stress generated at the particles depends on the flow strength of the alloy, the applied strain and the shape and size of the particles. Growth of the holes, so that they coalesce to form a macroscopic fracture, depends on the applied stresses being tensile. Much higher ductilities are achieved in compressive straining.

In cleavage fracture the material fails along well defined crystallographic planes within the grain but the crack path is affected by grain boundaries and inclusions. Basically a cleavage fracture surface contains large smooth areas separated by cleavage steps and feathers, river markings and cleavage tongues which are the direct result of crack path disturbances-Fig. 2b. Intercrystalline fracture is characterized by separation of the grains to reveal a surface composed of grain boundary facets, Fig. 2c. This type of fracture is found in stress-corrosion, creep hot tearing and hydrogen embrittlement. Fatigue fractures are characterized by striations (Fig. 2d) representing the extent of crack propagation under each cycle of loading.

a) Dimples in a ductile fracture of mild steel (x5000)

b) Cleavage patterns in HS steel fracture (x12000)

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c) Intergranular fracture in low alloy d) Fatigue striations in Nimonic 80A steel (x1500) (x7000)(A.Strang) Figure 2. Fracture in Scanning electron micrographs

Compound stresses and brittle fracture The failure of some American all-welded ships during the Second World War has stimulated much work on the causes of brittle fracture in steel. In the tensile test plastic deformation involves shearing slip along crystal planes within the crystals, but in the presence of tension of equal magnitude in each principal direction, shearing stresses are absent, plastic deformation is prevented and a brittle fracture occurs as soon as the cohesive strength of the material is exceeded. Equal triaxial tension stresses do not arise frequently in practice, but it is common to find a triaxial tension superimposed on a unidirectional tension, and if the margin between cohesive strength and plastic yield strength is small, a brittle fracture may occur in a material ordinarily considered highly ductile. Compound stresses arise in a weld in very thick plate and in a tube under internal pressure and an axial tension. This is shown in Fig. 3 with cohesive stress-strain curves, B, N, and F. If the two curves intersect at Y, brittle fracture occurs preceded by plastic deformation, which decreases as the cohesive strength curve becomes lower with respect to the yield stress-strain curve. Orowan has shown that if the yield stress is denoted by Y, the strength for brittle fracture by B (both Y and B depend on the plastic strain), and the initial value of Y (for strain = 0) by Y0 we have the following relation: • • •

The material is brittle if B < Y0; The material is ductile but notch-brittle if Y0 < B < 3Y0 The material is not notch-brittle if 3Y < B.

The factor 3 takes into account the stress increase at a notch. Whether the material is notch-brittle or not depends on the very small margin between B and 3Y.

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Figure 3. (a) Yield and cohesive stress curves (b) Slow notch bend test (c) Effect of temperature on the Izod value of mild steel Carbon steel is an exceptional material, because 3Y and B are so close, and this is the reason why the results of Izod tests seem to be so erratic, and why notch brittleness is so sensitive to slight variations of composition, previous treatment and temperature. Brittle fracture is characterised by the very small amount of work absorbed and by a crystalline appearance of the surfaces of fracture, often with a chevron pattern pointing to the origin of fracture, due to the formation of discontinuous cleavage cracks which join up (Fig. 4). It can occur at a low stress of 75-120 MPa with great suddenness; the velocity of crack propagation is probably not far from that of sound in the material in this type of fracture plastic deformation is very small, and the crack need not open up considerably in order to propagate, as is necessary with a ductile failure.

Figure 4. Steel brittle fracture surface with chevron markings. Micrograph shows discontinuous cracks ahead of main crack The work required to propagate a crack is given by Griffith`s formula:

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(1) where: σ = tensile stress required to propagate a crack of length c γ = surface energy of fracture faces E = Young`s modulus Orowan modified the Griffith theory to include a plastic strain energy factor, p, since some plastic flow is always found near the fracture surface: (2) When the temperature is above the brittle-ductile transition temperature, p is large and the stress, σ, required to make the crack grow will also be large. Below the transition temperature the metal is brittle and p will be smaller. The stress necessary to cause crack growth, therefore, will be reduced. The reason for the increasing speed of crack propagation, once a crack has started, is clear from both Griffith`s and Orowan`s equations: as the crack grows in length, the stress required for propagation continually decreases.

8.2. Fracture Mechanics Abstract: Fracture mechanics approaches require that an initial crack size be known or assumed. For components with imperfections or defects (such as welding porosities, inclusions and casting defects, etc.) an initial crack size may be known. Alternatively, for an estimate of the total fatigue life of a defect-free material, fracture mechanics approaches can be used to determine propagation. Strain-life approaches may then be used to determine initiation life, with the total life being the sum of these two estimates.

The fatigue life of a component is made up of initiation and propagation stages. This is illustrated schematically in Fig. 1

Figure 1. Initiation and propagation portions of fatigue life

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The size of the crack at the transition from initiation to propagation is usually unknown and often depends on the point of view of the analyst and the size of the component being analyzed. For example, for a researcher equipped with microscopic equipment it may be on the order of a crystal imperfection, dislocation, or a 0,1 mmcrack, while to the inspector in the field it may be the smallest crack that is readily detectable with nondestructive inspection equipment. Nevertheless, the distinction between the initiation life and propagation life is important. At low strain amplitudes up to 90% of the life may be taken up with initiation, while at high amplitudes the majority of the fatigue life may be spent propagating a crack. Fracture mechanics approaches are used to estimate the propagation life. Fracture mechanics approaches require that an initial crack size be known or assumed. For components with imperfections or defects (such as welding porosities, inclusions and casting defects, etc.) an initial crack size may be known. Alternatively, for an estimate of the total fatigue life of a defect-free material, fracture mechanics approaches can be used to determine propagation. Strain-life approaches may then be used to determine initiation life, with the total life being the sum of these two estimates.

Linear Elastic Fracture Mechanics Background Linear elastic fracture mechanics (LEFM) principles are used to relate the stress magnitude and distribution near the crack tip to: ƒ ƒ ƒ

Remote stresses applied to the cracked component The crack size and shape The material properties of the cracked component

Historical Overview In the 1920s, Griffith formulated the concept that a crack in a component will propagate if the total energy of the system is lowered with crack propagation. That is, if the change in elastic strain energy due to crack extension is larger than the energy required to create new crack surfaces, crack propagation will occur. Griffith`s theory was developed for brittle materials. In the 1940s, Irwin extended the theory for ductile materials. He postulated that the energy due to plastic deformation must be added to the surface energy associated with the creation of new crack surfaces. He recognized that for ductile materials, the surface energy term is often negligible compared to the energy associated with plastic deformation. Further, he defined a quantity, G, the strain energy release rate or "crack driving force," which is the total energy absorbed during cracking per unit increase in crack length and per unit thickness. In the mid-1950s, Irwin made another significant contribution. He showed that the local stresses near the crack tip are of the general form

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(1)

where r and θ are cylindrical coordinates of a point with respect to the crack tip (see Fig. 2) and K is the stress intensity factor. He further showed that the energy approach (the "G" approach above) is equivalent to the stress intensity approach and that crack propagation occurs when a critical strain energy release rate, G, (or in terms of a critical stress intensity, Kc) is achieved.

Figure 2. Location of local stresses near a crack tip in cylindrical coordinates

LEFM Assumptions Linear elastic fracture mechanics (LEFM) is based on the application of the theory of elasticity to bodies containing cracks or defects. The assumptions used in elasticity are also inherent in the theory of LEFM: small displacements and general linearity between stresses and strains. The general form of the LEFM equations is given in Eq. 1. As seen, a singularity exists such that as r, the distance from the crack tip, tends toward zero, the stresses go to infinity. Since materials plastically deform as the yield stress is exceeded, a plastic zone will form near the crack tip. The basis of LEFM remains valid, though, if this region of plasticity remains small in relation to the overall dimensions of the crack and cracked body.

Loading Modes There are generally three modes of loading, which involve different crack surface displacements (see Fig. 3). The three modes are:

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Mode 1: opening or tensile mode (the crack faces are pulled apart) Mode 2: sliding or in-plane shear (the crack surfaces slide over each other) Mode 3:

tearing or anti-plane shear (the crack surfaces move parallel to the leading edge of the crack and relative to each other)

The following discussion deals with Mode 1 since this is the predominant loading mode in most engineering applications. Similar treatments can readily be extended to Modes 2 and 3.

Figure 3. Three loading modes

Stress Intensity Factor The stress intensity factor, K, which was introduced in Eq. 1, defines the magnitude of the local stresses around the crack tip. This factor depends on loading, crack size, crack shape, and geometric boundaries, with the general form given by

(2)

where: σ = remote stress applied to component (not to be confused with the local stresses, σij, in Eq. 1) a = crack length f (a/w) = correction factor that depends on specimen and crack geometry Figure 4 gives the stress intensity relationships for a few of the more common loading conditions. Stress intensity factors for a single loading mode can be added algebraically. Consequently, stress intensity factors for complex loading conditions of the same mode can be determined from the superposition of simpler results, such as those readily obtainable from handbooks.

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Figure 4. Stress intensity factor for (a) Center-cracked plate loaded in tension, (b) Edge-cracked plate loaded in tension, (c) Double-edge-cracked plate loaded in tension (d) Cracked beam in pure bending

8.3. Fracture Toughness Abstract: Materials develop plastic strains as the yield stress is exceeded in the region near the crack tip. The amount of plastic deformation is restricted by the surrounding material, which remains elastic. The size of this plastic zone is dependent on the stress conditions of the body.

Plastic Zone Size Materials develop plastic strains as the yield stress is exceeded in the region near the crack tip (see Fig. 1). The amount of plastic deformation is restricted by the surrounding material, which remains elastic. The size of this plastic zone is dependent on the stress conditions of the body.

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Figure 1. Yielding near crack tip. Plane stress and plane strain conditions. In a thin body, the stress through the thickness (σz) cannot vary appreciably due to the thin section. Because there can be no stresses normal to a free surface, σz = 0 throughout the section and a biaxial state of stress results. This is termed a plane stress condition (see Fig. 2).

Figure 2. Plane stress and plane strain conditions In a thick body, the material is constrained in the z direction due to the thickness of the cross section and εz = 0, resulting in a plane strain condition. Due to Poisson`s effect, a stress, σz, is developed in the z direction. Maximum constraint conditions exist in the plane strain condition, and consequently the plastic zone size is smaller than that developed under plane stress conditions. Monotonic plastic zone size. The plastic zone sizes under monotonic loading have been estimated to be

plane stress

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(1)

plane strain

where r is defined as shown in Fig. 3.

Figure 3. Monotonic plastic zone size Cyclic plastic zone size. The reversed or cyclic plastic zone size is four times smaller than the comparable monotonic value. As the nominal tensile load is reduced, the plastic region near the crack tip is put into compression by the surrounding elastic body. As shown in Fig. 4, the change in stress at the crack tip due to the reversed loading is twice the value of the yield stress. Equation 2 become

plane stress (2) plane strain

The cyclic plastic zone size is smaller than the monotonic and more characteristic of a plane strain state even in thin plates. Thus LEFM concepts can often be used in the analysis of fatigue crack growth problems even in materials that exhibit considerable amounts of ductility. The basic assumption that the plastic zone size is small in relationship to the crack and the cracked body usually remains valid.

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Figure 4. Reversed plastic zone size

Fracture Toughness As the stress intensity factor reaches a critical value (Kc), unstable fracture occurs. This critical value of the stress intensity factor is known as the fracture toughness of the material. The fracture toughness can be considered the limiting value of stress intensity just as the yield stress might be considered the limiting value of applied stress. The fracture toughness varies with specimen thickness until limiting conditions (maximum constraint) are reached. Recall that maximum constraint conditions occur in the plane strain state. The plane strain fracture toughness, KIc is dependent on specimen geometry and metallurgical factors. ASTM Designation E-399, Standard Method of Test for Plane Strain Fracture Toughness of Metallic Materials, sets forth accepted procedures for determining this value. It is often difficult to perform a valid test for KIc. For example, a valid test using a thin plate of high toughness material often cannot be performed. Rather the value, Kc at the given conditions is obtained. The fracture toughness depends on both temperature and the specimen thickness. The following example shows the importance of the fracture toughness in designing against unstable fracture.

8.4. Macroscopic Aspects of Fracture Abstract: Fracture can be viewed on many levels, depending on the size of the fractured region that is of interest. At the macroscopic level fracture occurs over dimensions that are of the order of the size of flaws or notches (1 mm or greater). From the principles of fracture mechanics it is possible to determine macroscopic fracture criteria in terms of the nominal fracture strength, the flaw length, and the critical amount of plastic work required to initiate unstable fracture -- the fracture toughness.

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Fracture is an inhomogeneous process of deformation that makes regions of material to separate and load-carrying capacity to decrease to zero. It can be viewed on many levels, depending on the size of the fractured region that is of interest. At the atomistic level fracture occurs over regions whose dimensions are of the order of the atomic spacing (10-7 mm); at the microscopic level fracture occurs over regions whose dimensions are of the order of the grain size (about 10-3 mm); and at the macroscopic level fracture occurs over dimensions that are of the order of the size of flaws or notches (1 mm or greater). At each level there are one or more criteria that describe the conditions under which fracture can occur. For example, at the atomistic level fracture occurs when bonds between atoms are broken across a fracture plane and new crack surface is created. This can occur by breaking bonds perpendicular to the fracture plane, a process called cleavage, or by shearing bonds across the fracture plane, a process called shear. At this level the fracture criteria are simple; fracture occurs when the local stresses build up either to the theoretical cohesive strength σc ≈ E/10 or to the theoretical shear strength τc ≈ G/10, where E and G are the respective elastic and shear module. The high stresses required to break atomic bonds are concentrated at the edges of inhomogeneities that are called micro cracks or macro cracks (flaws, notches, cracks). At the microscopic and macroscopic levels fracture results from the passage of a crack through a region of material. The type of fracture that occurs is characterized by the type of crack responsible for the fracture. Few structural materials are completely elastic; localized plastic strain usually precedes fracture, even when the gross fracture strength is less than the gross yield strength. Fracture in these instances is initiated when a critical amount of local plastic strain or plastic work occurs at the tin of a flaw. From the principles of fracture mechanics it is possible to determine macroscopic fracture criteria in terms of the nominal fracture strength, the flaw length, and the critical amount of plastic work required to initiate unstable fracture -- the fracture toughness.

Types of Fracture That Occur Under Uniaxial Tensile Loading Cleavage fractures occur when a cleavage crack spreads through a solid under a tensile component of the externally applied stress. The material fractures because the concentrated tensile stresses at the crack tip are able to break atomic bonds. In many crystalline materials certain crystallographic planes of atoms are most easily separated by this process and these are called cleavage planes. Under uniaxial tensile loading the crack tends to propagate perpendicularly to the tensile axis. When viewed in profile, cleavage fractures appear "flat" or "square", and these terms are used to describe them. Most structural materials are polycrystalline. The orientation of the cleavage plane(s) in each grain of the aggregate is usually not perpendicular to the applied stress so that, on a microscopic scale, the fractures are not completely flat over distances larger than the grain size.

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In very brittle materials cleavage fractures can propagate continuously from one grain to the next. However, in materials such as mild steel the macroscopic cleavage fracture is actually discontinuous on a microscopic level; most of the grains fracture by cleavage but some of them fail in shear, causing the cleaved grains to link together by tearing. Shear fracture, which occurs by the shearing of atomic bonds, is actually a process of extremely localized (inhomogeneous) plastic deformation. In crystalline solids, plastic deformation tends to be confined to crystallographic planes of atoms which have a low resistance to shear. Shear fracture in pure single crystals occurs when the two halves of the crystal slip apart on the crystallographic glide planes that have the largest amount of shear stress resolved across them. When the shear occurs on only one set of parallel planes, a slant fracture is formed. In polycrystalline materials the advancing shear crack tends to follow the path of maximum resolved shear stress. This path is determined by both the applied stress system and the presence of internal stress concentrators such as voids, which are formed at the interface between impurity particles (e.g., nonmetallic inclusions) and the matrix material. Crack growth takes place by the formation of voids and their subsequent coalescence by localized plastic strains. Shear fracture in thick plates and round tensile bars of structural materials begins in the center of the structure (necked region) and spreads outwards. The macroscopic fracture path is perpendicular to the tensile axis. On a microscopic scale the fracture is quite jagged, since the crack advances by shear failure (void coalescence) on alternating planes inclined at 30-45° to the tensile axis. This form of fracture is commonly labeled normal rupture (since the fracture path is normal to the tensile axis) or fibrous fracture (since the jagged fracture surface has a fibrous or silky appearance). Normal rupture forms the central (flat) region of the familiar cup-cone pattern. The structure finally fails by shear rupture (shear lip formation) on planes inclined at 45° to the tensile axis. This form of fracture is less jagged, appears smoother, and occurs more rapidly than the normal rupture which precedes it. Similarly noncleavage fracture in thin sheets of engineering materials occurs exclusively by shear rupture and the fracture profile appears similar to the slant fracture. Under certain conditions the boundary between adjacent grains in the polycrystalline aggregate is weaker than the fracture planes in the grains themselves. Fracture then occurs intergranularly, by one of the processes mentioned above, rather than through the grains (transgranular fracture). Thus there are six possible modes of fracture: transgranular cleavage, transgranular shear rupture, transgranular normal rupture, and intergranular cleavage, intergranular shear rupture, intergranular normal rupture. Fracture takes place by that mode which requires the least amount of local strain at the tip of the advancing crack. Both the environmental fracture exclusively by one particular mode over a large variety of conditions and the state of applied (nominal) stress mid strain determine the type of fracture which occurs, and only a few materials and structures operating conditions (i.e., service temperature, corrosive environment, and so on).

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Furthermore, under any given condition more than one mode of fracture can cause failure of a structural member and the fracture is described as "mixed". This implies that the relative ease of one type of crack propagation can change, with respect to another type, as the overall fracture process lakes place. For example, normal rupture, cleavage, and shear rupture are all observed on the fracture surfaces of notched mild steel specimens broken in impact at room temperature. In order to analyze the fracture process under various types of stress systems, it is necessary to establish a coordinate system with respect to both the fracture plane, the direction of crack propagation, and the applied stress system. One of the difficulties encountered by engineers and scientists who are interested in a particular aspect of the fracture problem is the large mass of notation and coordinate systems used by other workers who have investigated similar problems. Three distinct modes of separation at the crack tip can occur: • • •

Mode I -- The tensile component of stress is applied in the y direction, normal to the faces of the crack, either under plane-strain (thick plate, t large) or plane-stress (thin plate, t small) conditions. Mode II -- The shear component of stress is applied normal to the leading edge of the crack either under plane-strain or plane-stress conditions. Mode III -- The shear component of stress is applied parallel to the leading edge of the crack (antiplane strain).

Summary 1. Fracture is an inhomogeneous form of deformation which can be viewed on different levels. On an atomistic level it occurs by the breaking of atomic bonds, either perpendicular to a plane (cleavage) or across a plane (shear). On a microscopic level cleavage occurs by the formation and propagation of microcracks and the separation of grains along cleavage planes. Shear fracture (rupture) usually occurs by the formation of voids within grains and the separation of material between the voids by intense shear. On a macroscopic level cleavage occurs when a cleavage crack spreads essentially perpendicular to the axis of maximum tensile stress. Shear fracture occurs when a fibrous crack advances essentially perpendicular to the axis of maximum tensile stress (normal rupture) or along a plane of maximum shear stress (shear rupture). Fracture is said to be transgranular when microcrack propagation and void coalescence occur through the grains and intergranular when they occur along grain boundaries. More than one mode of crack propagation can contribute to the fracture of a structure. In general, cleavage fracture is favored by low temperatures. 2. Fracture occurs in a perfectly elastic solid when the stress level at the tip of a preinduced flaw reaches the theoretical cohesive stress E/10 and a sufficient amount of work γs is done to break atomic bonds and create free surface.

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3. When the yield stress of a material is less than E/10 (i.e., in a ductile material or in the vicinity of a stopped crack in a partially brittle one), plastic flow occurs near the crack tip and the stress level in the plastic zone is less than E/10. Consequently the crack cannot advance directly as an elastic Griffith crack.

8.5. Fracture of Steel: Part One Abstract: There are thousands of steels available today, each one characterized by a particular trade name or alloy composition. Although a quantitative value of fracture toughness parameters (e.g., NDT temperature and KIC) for each grade would greatly facilitate the selection of a material for a particular application, these parameters are available for only a very few of the steels. There are primarily two reasons for this. First, because a wide range of microstructures can be obtained in a steel of given alloy composition, simply by variations in thermomechanical treatment. Secondly, because the concentration of fabrication defects (i.e., blow holes, inclusions, and so on) is extremely sensitive to mill practice and can vary between heats of steel of the same composition or even in different parts of the same billet.

There are literally thousands of steels available today, each one characterized by a particular trade name or alloy composition. Although a quantitative value of fracture toughness parameters (e.g., NDT temperature and KIC) for each grade would greatly facilitate the selection of a material for a particular application, these parameters are available for only a very few of the steels. There are primarily two reasons for this. First, because a wide range of microstructures can be obtained in a steel of given alloy composition, simply by variations in thermomechanical treatment. Secondly, because the concentration of fabrication defects (i.e., blow holes, inclusions, and so on) is extremely sensitive to mill practice and can vary between heats of steel of the same composition or even in different parts of the same billet. Since it is microstructure and defect concentration that primarily determine toughness, rather than composition per se, a large variation in toughness can be produced in a given steel simply by varying the thermomechanical treatment and fabrication practice. A detailed understanding of the fracture of steel therefore requires an understanding of both the physical metallurgical aspects of the material (e.g., what microstructure will result from a given heat treatment) as well as an understanding of how this particular microstructure affects the toughness of a structure of given geometry.

The Fracture of Ferritic-Pearlitic Steels Ferritic-pearlitic steels account for most of the steel tonnage produced today. They are iron-carbon alloys that generally contain 0.05-0.20% carbon and a few per cent of other alloying elements that are added to increase yield strength and toughness. In these steels the microstructure consists of BCC iron (ferrite), containing about 0.01% carbon and soluble alloying elements, and Fe3C (cementite). In very low carbon steels the cementite particles (carbides) lie in the ferrite grain boundaries and grains, but when the carbon content is greater than about 0.02%, most of the Fe3C forms a lamellar structure with some of the ferrite. This lamellar structure is called

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pearlite and it tends to exist as "grains" or nodules, dispersed in the ferrite matrix. In low carbon (0.10-0.20%) steel (i.e., mild steel) the pearlite accounts for between 10-25% of the microstructure. Although the pearlite grains are very hard, they are so widely dispersed that the ferrite matrix can deform around them with little difficulty. It should be noted, however, that the ferrite grain size generally decreases with increasing pearlite content because the formation of pearlite nodules during the transformation interferes with ferrite grain growth. Consequently the pearlite can indirectly raise σy by raising d-1/2. From the point of view of fracture analysis, two ranges of carbon content are of most interest in the low carbon steels: (1) steels containing less than 0.03% carbon where the presence of pearlite nodules has little effect on toughness, and (2) steels containing higher carbon contents where the pearlite does have a direct effect on toughness and the shape of the Charpy curve. The effect of processing variables. It has been pointed out that the impact properties of water-quenched steels are superior to those of annealed or normalized steels because the fast cooling rate prevents the formation of grain boundary cementite and causes a refinement of ferrite grain size. Many commercial grades of steel are sold in the "hot-rolled" condition and the rolling treatments have a considerable effect on impact properties. Rolling to a lower finishing temperature (controlled rolling) lowers the impact-transition temperature. This results from the increased cooling rate and corresponding reduced ferrite grain size. Since thick plates cool more slowly than thin ones, thick plates will have a larger ferrite grain size and hence are more brittle than thin ones after the same thermomechanical treatment. Therefore, post rolling normalizing treatments are frequently given in order to improve the properties of rolled plate. Hot rolling also produces an anisotropic or directional toughness owing to combinations of texturing, pearlite banding, and the alignment of inclusions and grain boundaries in the rolling direction. Texturing is not considered to be important in most low carbon steels. Pearlite bands (due to phosphorous segregation during casting) and elongated inclusions are dispersed on too coarse a scale to have an appreciable effect on notch toughness at the low temperature end of the Charpy transition temperature range. The effect of ferrite-soluble alloying elements. Most alloying elements that are added to low carbon steel produce some solid solution hardening at ambient temperature and thereby raise the lattice friction stress σi. It is important to appreciate that equation cannot be used to predict the lower yield stress unless the resultant grain size is known. This, of course, depends on factors such as normalizing temperature and cooling rate. The importance of this type of approach is that it allows prediction of the extent that individual alloying elements will decrease toughness by increasing σi, since NDT increases by about 2°C per ksi increase in σi.

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Regression analyses for NDT temperatures or other Charpy transition temperatures have not been reported at this time and it is only possible to discuss the effects of the individual alloying additions on a qualitative basis. Manganese. Most commercial steels contain about 0.5% manganese to serve as a deoxidizer and to tie up sulfur as manganese sulfide, thereby preventing the occurrence of hot-cracking. In low carbon steels this effect is outweighed by the ability of manganese: • • •

to decrease the tendency for the formation of films of grain boundary cementite in air-cooled or furnace-cooled specimens containing 0.05% carbon, thereby lowering the value of γm; to cause a slight reduction in ferrite grain size; to produce a much finer pearlite structure.

The first two of these effects account for the lowering of the NDT temperature with increasing Mn additions. The third effect as well as the first cause the Charpy curves to become sharper. In steels containing higher carbon contents the effect of manganese on the 50% transition temperature is less pronounced, probably because the amount of pearlite rather than the distribution of grain boundary cementite is the most important factor in determining this transition temperature when the pearlite content is high. It should also be noted that if the carbon content is relatively high (greater than 0.15%) a high manganese content may have a detrimental effect on the impact properties of normalized steels because the high hardenability of the steel causes the austenite to transform to the brittle upper bainite structure rather than ferrite or pearlite. Nickel. Nickel, like manganese, is able to improve the toughness of iron carbon alloys. The magnitude of the effort is dependent on carbon content and heat treatment. In very low (about 0.02%) carbon steels, nickel additions up to 2% are able to prevent the formation of grain boundary cementite in hot-rolled and normalized alloys and cause a substantial decrease in the initiation-transition temperature TS(N), and a sharpening of the Charpy curves. Further additions of nickel produce substantially smaller improvements in impact properties. In alloys containing carbon contents lower than this, such that carbides are not present after normalizing, nickel has a smaller effect on the transition temperature. The principal beneficial effect of nickel additions to commercial steels containing about 0.1% carbon results from the substantial grain-size refinement and reduction of free nitrogen content after normalizing. The reasons for this behavior are not clear at present; it may be related to the fact that nickel is an austenite stabilizer and consequently lowers the temperature at which the austenite decomposition will take place. Phosphorous. In pure iron-phosphorus alloys, intergranular embrittlement can occur from the segregation of phosphorous at ferrite grain boundaries, which lowers the value of γm. Also, phosphorus additions produce a significant increase in σi and a coarsening of ferrite grain size since phosphorus is a ferrite stabilizer. These effects combine to make phosphorus an extremely effective embrittling agent, even when fracture occurs transgranularly.

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Silicon. Silicon is added to some commercial steels to deoxidize or "kill" them, and in this respect the silicon produces beneficial effects on impact properties. When manganese and aluminum are present, a large fraction of the silicon is dissolved in the ferrite and this raises σi by solid solution hardening. This effect, coupled with the fact that silicon additions raise ky, causes the 50% transition temperature to increase by about 44°C per wt per cent silicon in iron-carbon alloys of constant grain size. In addition, silicon, like phosphorus, is a ferrite stabilizer and hence promotes ferrite grain growth. The net effect of silicon additions in normalized alloys is to raise the average energy-transition temperature by about 60°C per wt per cent silicon added. Aluminum. The effect of alloying or killing a steel with aluminum is twofold. First, the aluminum combines with some of the nitrogen in solution to form AlN. The removal of this free nitrogen leads to a decrease in transition temperature because σi is decreased and γm/ky is increased, as described above. Second, the AlN particles that form interfere with ferrite grain growth and consequently refine the ferrite grain size. These combined effects cause the transition temperature to decrease about 40°C per 0.1% aluminum added. However, additions of aluminum greater than that required to tie up the nitrogen have little effect. Oxygen. Oxygen additions promote intergranular fracture in iron alloys. These fractures are thought to result from the segregation of oxygen to ferrite grain boundaries. In alloys that contain a high oxygen content (greater than 0.01 %), fracture occurs along the continuous path provided by the embrittled grain boundary. In alloys of lower oxygen content, cracks are nucleated at the grain boundary and then propagate transgranularly. The problem of oxygen embrittlement can be solved by the addition of deoxidizing elements such as carbon, manganese, silicon, aluminum, and zirconium, which react with the oxygen to form oxide particles, thereby removing the oxygen from the boundary region. These oxide particles are beneficial in their own right because they retard the growth of the ferrite grains, thereby increasing d-1/2.

8.6. Fracture of Steel: Part Two Abstract: In steels containing large-volume fractions of pearlite, deformation in the pearlite can initiate microcleavage crack formation at low temperatures and/or high strain rates. Since the fracture path is primarily along the cleavage plane in the ferrite plates (although there is some intercolony fracture), this indicates that there is some preferred orientation between ferrite plates in adjacent colonies within a prior austenite grain.

The effect of carbon additions between 0.3 and 0.8% In hypoeuteetoid steels containing between 0.3 and 0.8% carbon, proeutectoid ferrite is the continuous phase and forms primarily at austenite grain boundaries. The pearlite forms inside the austenite grains and makes up between 35-100% of the microstructure. More than one colony (set of parallel ferrite and cementite plates) forms within each austenite grain so that the pearlite is polycrystalline. Since the strength of the pearlite is greater than that of the proeutectoid ferrite, the pearlite constrains the flow of the ferrite. The yield strength and strain-hardening rate of these steels increase with increasing pearlite (carbon) content because the

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constraint effect increases with increasing amounts of the hard aggregate and because pearlite refines the size of the proeutectoid grains. In steels containing large-volume fractions of pearlite, deformation in the pearlite can initiate microcleavage crack formation at low temperatures and/or high strain rates. Since the fracture path is primarily along the cleavage plane in the ferrite plates (although there is some intercolony fracture), this indicates that there is some preferred orientation between ferrite plates in adjacent colonies within a prior austenite grain. The Fracture of Bainitic Steels. The addition of 0.05% molybdenum and boron to low carbon (0.1%) steels is able to suppress the austenite-ferrite transformation, which normally occurs between 700° and 850°C, without affecting the kinetics of the austenite-bainite transformation which then takes place between 675° and 450°C. Bainite formed between 675° and about 525°C is called "upper bainite" and bainite formed between 525° and 450°C is called "lower bainite". Both structures consist of acicular ferrite and dispersed carbides. The tensile strength of these un-tempered bainites increases from 85,000 to 170,000 psi (585 - 1170 MPa) as the transformation temperature drops from 675° to 450°C. Since the transformation temperature is determined by the amount of alloying elements (e.g., Mn and Cr) that are present, these elements exert an indirect effect on the yield find tensile strengths. The high strengths obtained in these steels is the result of two effects: • •

the progressive refinement of the bainitic ferrite plate size as the transformation temperature is lowered, and the fine carbide dispersion, which occurs within the grains of the lower bainite. Fracture characteristics of these steels is strongly dependent on the tensile strength and hence on the transformation temperature.

Two effects should be noted. First, at a given tensile strength level the impact properties of tempered lower bainite are far superior to that of untempered upper bainite. The reason for this behavior is that in upper bainite, as in pearlite, the cleavage facets traverse several bainite grains and the "effective grain size" for fracture is the prior austenite grain size rather than the ferrite grain size. In lower bainite the cleavage planes in the acicular ferrite are not aligned so that the effective grain size for quasicleavage fracture is the ferrite needle size. Since this is one to two orders of magnitude smaller than the prior austenite grain size, the transition temperature of the lower bainite is much below that of upper bainite, at the same strength level. A second feature that is important is the distribution of the carbides. In upper bainite these lie along grain boundaries and may promote brittleness by lowering γm as described previously in connection with furnace-cooled ferritic steels. In tempered lower bainite the carbides are more uniformly distributed in the ferrite and raise γm by interfering with cleavage cracks and promoting tearing as in the case of spherodized pearlites.

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A second effect that should be noted is the variation of transition temperature with tensile strength in the untempered alloys. In the upper bainite a decrease in transformation temperature produces a refinement of ferrite needle size and this raises Rp0.2. Tensile strength levels of 120,000 psi (830 MPa) or greater are obtained in lower bainite and the transition temperature decreases with increasing tensile strength. Because the fracture stress of the upper bainite is dependent on austenite grain size, and since the carbide particles are already large, tempering has little effect on tensile and impact properties.

The Fracture of Martensitic Steels The addition of carbon and other alloying elements to steel retards the transformation of austenite to either ferrite and pearlite, or bainite, and if the cooling rate after austenitizing is sufficiently rapid, the austenite will transform to martensite by a shear process that requires no measurable diffusion of atoms. The features that are pertinent to the fracture of martensite are as follows: • •

•

Because the transformation occurs at very low temperatures (200°C or lower) the size of the tetragonal ferrite or martensite needle is very small, at least in two of its three dimensions. Because the transformation occurs by shear, the carbon atoms do not have time to diffuse out of their lattice position in the austenite and hence the ferrite is supersaturated with carbon; this causes the martensite to have an elongated (bodycentered tetragonal) crystal structure and leads to lattice expansion. The martensitic transformation occurs over a range of temperatures because the formation of the first martensite plates increases the difficulty of transforming the remaining austenite. Thus transformation structures can be mixtures of martensite and retained austenite.

To produce stable steel that can be satisfactorily used in engineering applications, it is necessary to temper it. Three stages of tempering occur in high (greater than 0.3%) carbon martensites, tempered for about one hour in various ranges as follows: 1. At temperatures up to about 100°C some of the supersaturated carbon precipitates out of the martensite to form very fine particles of epsilon (hexagonal) carbide, which are dispersed in a martensite that consequently has a decreased carbon content. 2. Between 100° and 300°C any retained austenite is able to transform to bainite and epsilon carbide. 3. In the third stage of tempering, which begins about 200°C, depending on carbon content and alloy composition, the epsilon carbides dissolve and the low carbon martensite loses both its tetragonality and its carbon. As the temperature of tempering increases up to the eutectoid temperatures (723°C), the carbide precipitates coarsen and Rp0.2 decreases. Tempering just below the eutectoid temperature causes the cementite particles to assume a relatively large (1-10 μp) spheroidal shape, similar to that obtained by annealing a pearlitic structure for long periods of time.

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Fracture of medium strength steels (620 MPa < Rp0.2 < 1240 MPa) In addition to the removal of residual stress there are two effects associated with tempering that increase notch toughness. The first is the transformation of retained austenite. The austenite should be transformed at low temperatures (around 300°C) to the tough, acicular lower bainite. If it is transformed by tempering at a higher temperature, say 600°C, the brittle pearlite structure will form. Consequently steels that are to be tempered at 550°-600°C are first tempered at around 300°C to avoid this problem. This procedure is called "double tempering". Secondly, there is the decrease in yield strength and the increase in dispersed carbide content (γm increases), both of which cause the impact-transition tempering range to be lowered as the tempering temperature is increased. Tensile ductility and Cv (max) increase, at the same strength level, as the microstructure is refined. Temper embrittlement is reversible. If the tempering temperature is raised above the critical range, the transition temperature is lowered, but it can be raised back again if the material is reheat treated in the critical range. The presence of trace elements appears to be responsible for the embrittlement. The most important of these are antimony, phosphorus, tin, and arsenic, with manganese and silicon having a small effect. Molybdenum reduces temper brittleness when other alloying elements are present. Nickel and chromium appear to have little effect.

Fracture of high strength steels (Rp0.2 > 1240 MPa) High strength steels are produced by basically one of three processes; quenching and tempering, deforming the austenite before quenching and tempering (ausforming), or annealing and aging to produce precipitation hardening (e.g., maraging). In addition, further increases in strength can be achieved by straining and retempering or by straining during tempering. The high strength level of these steels makes them extremely brittle, especially when particular environments such as water vapor or hydrogen are present.

The Fracture of Stainless Steels Stainless steels are basically iron-chromium and iron-chromium-nickel alloys to which small amounts of other elements have been added to improve mechanical properties and corrosion resistance. Their resistance to corrosion arises from the formation of an impervious layer of chromium oxide on the metal surface, which, in turn, prevents any further oxidation of the metal. Consequently these steels are corrosion-resistant in oxidizing atmospheres, which strengthen this layer, but are susceptible to corrosion in a reducing environment, which breaks down the layer. The corrosion resistance (in oxidizing environment) increases with increasing chromium content and also with increasing nickel content. The latter element increases the overall passivity of the iron. Carbon is also added to improve mechanical properties (yield and tensile strength) and to stabilize the austenitic stainless steel. Generally speaking, the stainless steels can be classified by their microstructures:

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• •

•

Martensitic. These are iron-chromium alloys that can be austenitized and subsequently heat-treated to form martensite. They normally contain about 12% chromium and 0.15% carbon (type 410). Ferritic. These alloys contain about 14-18% chromium and 0.12% carbon (type 430) and are completely ferritic since chromium is a ferrite stabilizer; the austenite phase is completely suppressed in alloys containing more than 13% Cr. Austenitic. Nickel is a strong austenite stabilizer and alloys containing 8% nickel and 18% chromium (type 300) are austenitic at room temperature and below, as well as at high temperatures. These steels, like the ferritic grade, cannot be hardened by martensitic transformation.

The fracture characteristics of ferritic and martensitic stainless steels are similar to those of other ferritic or martensitic steels at the same strength level, grain size, and so on. Austenitic stainless steels have a FCC structure and consequently do not fracture by cleavage, even at cryogenic temperatures. After heavy cold rolling (80%), 310 type steels have an extremely high yield strength combined with a notch sensitivity ratio of 1.0 at temperatures as low as -253°C and consequently are used in missile systems for storage tanks for liquid hydrogen. Similarly 301 type stainless can be used down to -183°C (e.g., for liquid oxygen storage tanks), but below this temperature the austenite is unstable and deforms to brittle, untempered martensite if any plastic deformation occurs at the low temperatures. Most austenitic stainless steels are used in corrosive environments. When they are heated in the temperature range 500-900°C (e.g., during welding), chromium carbide precipitates at austenite grain boundaries, resulting in a depletion of chromium from the region near to the boundaries. This depleted layer is very susceptible to corrosive attack (particularly in hot chloride environments), and localized corrosion, in the presence of applied stress, leads to inter-granular brittle fracture. To alleviate this problem, small quantities of elements which are stronger carbide formers than chromium, such as titanium or niobium are commonly added. These elements combine with the carbon to form alloy carbides, which prevents chromium depletion and subsequent susceptibility to stress corrosion cracking. This process is called "stabilizing". Austenitic stainless steels are used extensively in high-temperature applications (e.g., pressure vessels) where both corrosion resistance and creep resistance are required. Some of these steels are susceptible to cracking in the heat-affected zone near welds during postwelding heat treatments and/or elevated temperature service. The cracking is the result of precipitation of niobium or titanium carbides in the grains and grain boundaries when the weld is reheated.

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8.7. Fracture Toughness of High-Strength Steels at Low Temperatures Abstract: According to available information on the fracture toughness of high-strength alloys at low temperatures, the effect of low temperatures on toughness is generally dependent on the alloy base. Alloy steels normally exhibit decreasing fracture toughness as the testing temperature is decreased through transition temperature range, when the structure contains ferrite or tempered martensite. The transition temperature is influenced by the alloy content, grain size and heat treatment.

Current and developing applications for materials at low temperatures include structures, vehicles, and pipeline equipment for arctic environments, storage and transport equipment for liquefied fuel gasses, oxygen and nitrogen, and superconducting machinery, devices and electrical transmission systems. Most of these applications relate to the production and distribution of energy and have attained greater prominence because of the current energy shortage. According to available information on the fracture toughness of high-strength alloys at low temperatures, the effect of low temperatures on toughness is generally dependent on the alloy base. For many aluminum alloys, the fracture toughness tends to increase or remain generally constant as the testing temperature is decreased. Titanium alloys tend to have lower toughness as the testing temperature is decreased, but the effect is influenced by the alloy content and heat treatment. Certain titanium alloys retain good toughness at very low temperatures. Alloy steels normally exhibit decreasing fracture toughness as the testing temperature is decreased through transition temperature range, when the structure contains ferrite or tempered martensite. The transition temperature is influenced by the alloy content, grain size and heat treatment. Carbon and low alloy steels represent body-center-cubic (bcc) atomic lattices and exhibit toughness transition temperature ranges either above, at, or below room temperature depending on a number of factors. At temperatures above the transition temperature, the alloy has substantially better toughness than at lower temperatures. Furthermore, the lower strength steels generally are strain-rate sensitive, while the higher strength steels are not strain-rate sensitive The curves for parent metal and welds in ASTM A517F steel plate indicate that the weld metal and heat-affected zones have lower transition temperatures than the parent metal. However, the weld metal in the specimens of A542 steel had higher transition temperatures than the parent metal. The fracture toughness of A533 Grade B Class 1 steel has been studied extensively for nuclear reactor pressure vessels. This study has shown that for testing temperatures above -100oF (approx. 70oC), the toughness increases substantially as the testing temperature is increased. Thus the required thickness of the specimens must be increased in order to increase the constraint that is necessary at the crack tip to simulate plane-strain conditions at the initiation of fracture. Results of fracture toughness tests on three ASTM forging steels may have similar general trends in the toughness data, but the compositions, grain sizes, and other factors have marked effects on the transition temperatures. Test results for HY-130

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steel indicate that this steel in the temperature range down to -320oF (approx. 195oC) is not strain-rate sensitive. These steels are not intended for use at temperatures in or below the transition temperature range, and there is no accepted method for indicating the specific transition temperature from a transition temperature curve. Furthermore, there is no accepted method relating the transition temperature to a safe minimum service temperature for structural components. However, if Kic data are obtained for given alloy at low temperatures, the critical crack sizes may be estimated in the lowtemperature range at the maximum service stress of the structure. The effects of variations in composition on a series of Ni-Cr-Mo-V steels has been studied in order to show the effects of the alloying elements on the low-temperature fracture toughness. Bars of these steels were quenched and tempered to about 170 ksi yield strength (approx. 1170 MPa) and tested as precracked bend specimens. The effects of carbon and nickel content were the most significant. An increase of carbon content and nickel content from 0.28 to 0.41 raised the transition temperature based on KQ data. Increasing the nickel content from 1.26 to 6.23 percent decreased the KQ transition temperature. This represents one of the major attributes of nickel additions to the alloy steels. The specimens of D6ac steel were austenitized at about 1650oF, furnace cooled to 975oF, and quenched in oil or molten salt according to several different procedures, to simulate quenching of the welded forging that comprise the F111 wing carythrough structure. The high-toughness specimens were quenched in oil, while the medium-toughness specimens were quenched in salt. Regardless of the quench, the yield strength of the specimens was approximately 217 ksi (1495 MPa) after tempering twice at 1000 to 1025oF. The fracture roughness tests were very sensitive indicators of the effect of the variation in quenching rate on the toughness. The specimens that had the highest toughness at room temperature also had the highest toughness at -65oF (-54oC). Available fracture toughness data at low temperature for other alloy steels: AISI 4340, 300M, HP9-4-20, HP9-4-25, and 18 Ni (200) maraging steel usually have the trend of decreasing toughness as the testing temperature is decreased. The one exception is HP9-4-25 in the temperature range +75 to -75oF (+24 to -59oC). At lower temperature, the expected trend would be for the toughness to drop as indicated for HP9-4-20 in the range from -100 to -320oF (-73 to -195oC). The data obtained by Steigerwald for AISI 4340 steel and by Wessel for the HP9-4-20 alloy steel were obtained before ASTM Method E 399 was available and are designated as KQ values. The 18Ni (200) grade maraging steel also exhibits considerable reduction in toughness as the testing temperature is reduced from -100 to -320oF (-73 to 195oC), but at -320oF, this heat of the 200 grade retained a toughness of about 80 ksi in.1/2. From limited information on the toughness of the 200 grade, it appears that there is considerable range in results of KIc tests at room temperature. This level of toughness at -320oF probably can be achieved only if the alloy has a toughness of about 160 ksi in.1/2 or over at 75oF (+24oC). The effect of low temperatures on the static and dynamic fracture toughness of bend specimens of 18Ni (200) maraging steel is a straight line relationship between the

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Kic values and the testing temperature in the range from 75 to -320oF. At -320oF, the KIc value was about 40 ksi in.1/2, and the alloy is not strain-rate sensitive in the low-temperature range. Tests results on part-through surface-crack specimens of 200 grade maraging steel has shown that these heats had high toughness at 75oF and also retained relatively good toughness at -320oF. With optimum welding conditions, the weld metal also retains good strength and toughness at -320oF.

8.8. Temper Embrittlement Abstract: Temper embrittlement is inherent in many steels and can be characterized by reduced impact toughness. The state of temper embrittlement has practically no effect on other mechanical properties at room temperature. Many alloy steels have two temperature intervals of temper embrittlement. For instance, irreversible temper brittleness may appear within the interval of 250-400°C and reversible temper brittleness, within 450-650°C.

Temper embrittlement is inherent in many steels and can be characterized by reduced impact toughness. The state of temper embrittlement has practically no effect on other mechanical properties at room temperature. Figure 1 shows schematically the effect of temperature on impact toughness of alloy steel which is strongly liable to temper embrittlement. Many alloy steels have two temperature intervals of temper embrittlement. For instance, irreversible temper brittleness may appear within the interval of 250-400°C and reversible temper brittleness, within 450-650°C.

The impact toughness of quenched steel after tempering at 250-400°C is lower than that obtained on tempering at temperatures below 250°C. If brittle steel tempered at 250-400°C is heated above 400°C and transferred into a tough state, a second

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tempering at 250-400°C cannot return it to the brittle state. The rate of cooling from the tempering temperature within 250-400°C has no effect on impact toughness. Steel in the state of irreversible temper embrittlement has a bright intercrystalline fracture at boundaries of former austenitic grains. This type of brittleness is inherent to some extent to all steels, including carbon grades. For that reason mediumtemperature tempering is, as a rule not employed in practice, though it can ensure a high yield limit. Irreversible temper embrittlement is thought to be due to the formation of carbides on decomposition of martensite, in particular, precipitation of carbides in the form of films at grain boundaries. At higher temperatures of tempering, this film disappears and cannot be restored on repeated heating at 250-400°C. Silicon in low-alloy steels can prevent irreversible temper embrittlement by retarding the decomposition of martensite. The embrittlement on high-temperature tempering may manifest itself in two different ways: • •

as a result of heating at 450-600°C (irrespective of the rate of subsequent cooling) and effect of temperature, and as a result of tempering at temperatures above 600°C with subsequent slow cooling within the range of 600-450°C.

A high-rate cooling from a tempering temperature above 600°C, for instance, watercooling, can prevent the appearance of temper embrittlement. On the other hand, a quick cooling on tempering at 450-600°C cannot prevent temper embrittlement. Thus, entering the dangerous temperature interval from either "below" (on heating and holding at that temperature) or from "above" (on slow cooling) can produce the same result. The most important feature of embrittlement on high-temperature tempering is that the process is reversible. If a steel embrittled through tempering at a temperature above 600°C with subsequent slow cooling or through tempering at 450-600°C (with any rate of cooling) is again heated above 600°C and cooled quickly, its impact toughness will restore to the initial value. If the steel then again enters the dangerous interval of tempering temperatures, it is again embrittled. A new heating at a temperature above 600°C, followed with quick cooling, can eliminate the embrittling effect, and so on. This is why the phenomenon discussed is called reversible embrittlement. Carbon steels with less than 0.5% Mn are not prone to reversible temper embrittlement. The phenomenon can only appear in alloy steels. Alloying elements may have different effects on steel after tempering at the steel proneness to temper embrittlement. Unfortunately, the most widely used alloying elements, such as chromium, nickel, and manganese, promote temper embrittlement. When taken separately, they produce a weaker effect than in the case of combined alloying. The highest embrittling effect is observed in Cr-Ni and Cr-Mn steels. Small additions of molybdenum (0.2-0.3%) can diminish temper embrittlement, while greater additions enhance the effect.

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A fundamental fact is that alloy steels of very high purity are utterly unsusceptible to temper embrittlement which is caused by the presence of various impurities, in the first place of phosphorus, tin, antimony and arsenic, in commercial steels. The rate and degree of development of temper embrittlement depend on the temperature and time of holding steel within the dangerous temperature interval (450-600°C). With a certain temperature of tempering within this interval, the initial stages of embrittlement appear appreciably sooner than at a higher or a lower temperature. Many scientists adhered for a long time to the "solution precipitation" hypothesis, according to which the loss in impact toughness was caused by precipitation of some phases, such as phosphides, at grain boundaries. These phases were thought to pass into the á-solution on heating up to approximately 650°C and to precipitate from the solution and embrittle the steel on slow cooling; quick cooling should prevent the precipitation of embrittling phases. As has been found by electron-microscopic analysis, however, there are no special precipitates at grain boundaries in embrittled steel, so that the "solution precipitation" hypothesis turned to be inconsistent. Another hypothesis explained temper embrittlement by an increased concentration of impurities in boundary layers of the solid solution. This was proved by an increased etchability of grain boundaries in embrittled steel by picric acid. The hypothesis on the leading role of impurity segregates has been fully confirmed in the recent years by a brilliant series of research work using Auger spectroscopy, a method enabling determination of concentrations of elements in monatomic surface layers. Using this method makes it possible to detect segregations of phosphorus and other impurity elements at the fracture surface in embrittled steel and measure their concentrations (as also the concentrations of alloying elements) at the fracture surface. It has also been shown that the development of temper embrittlement is directly linked with the rise of impurity concentration near the prior austenite boundaries. Owing to equilibrium segregation, the concentration of harmful impurities at the surface of a fracture may exceed tens or hundreds times their average concentration in the steel. The concentration of impurities in commercial purity steels is usually a few thousandths or hundredths of a percent, but amounts to a few percent at the surface of fracture. As the temperature increases, the diffusion process of grain boundary segregation is accelerated, with the absolute value of equilibrium segregation being simultaneously decreased owing to thermal motion. At temperatures above 600-650°C, the segregation of impurities either disappears fully (Sb) or drops to a very low level (P). On subsequent cooling of the steel in water, the segregates have no time to restore. The role of alloying elements in the development of temper embrittlement is not less than that of impurities. The segregation of harmful impurities in iron-carbon alloys is so small that causes no temper embrittlement. In the presence of alloying elements (Ni, Cr or Mn), the segregation of impurities increases appreciably. In this process, the alloying elements themselves, which cause no equilibrium segregation in highpurity steels, segregate at grain boundaries in the presence of harmful impurities.

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Therefore, we can assume that an alloying element and impurity interact with each other in the á-solution and thus mutually promote their segregation. It can be also assumed that if atoms of an impurity and alloying element attract one another stronger than atoms of that impurity and iron, the segregation of the impurity and alloying element will be mutually enhanced. Namely in this way behave P and Ni, P and Cr, Sb and Ni, Sb and Mn and other "impurity - alloying element" pairs. A second alloying element can additionally enhance segregation of an impurity. For instance, nickel and chromium, when present together in steel, can cause a greater segregation of antimony than might be expected from simple summation of their separate effects. An increased concentration of harmful impurities in boundary layers of the solid solution, which may be caused by the effect of alloying additions, weakens the intergranular bondage and is one of the main causes why alloy steels containing Ni, Cr or Mn are highly susceptible to temper embrittlement. The main measures to prevent temper embrittlement are as follows: • • • •

of the content of harmful impurities in steel; accelerated cooling from the temperature of high-temperature tempering (above 600°C); alloying of steel with small additions of molybdenum (0.2-0.3%); and subjecting the metal to high-temperature thermo-mechanical treatment.

8.9. The Embrittlement and Fracture of Steels: Part One Abstract: Most groups of alloys can exhibit failure by cracking in circumstances where the apparent applied stress is well below that at which failure would normally be expected. Steels are no exception to this, and probably exhibit a wider variety of failure mechanisms than any other category of material. While ultimate failure under excessive stress must occur and can be reasonably predicted by appropriate mechanical tests, premature failure is always dangerous, involving a considerable element of unpredictability.

Most groups of alloys can exhibit failure by cracking in circumstances where the apparent applied stress is well below that at which failure would normally be expected. Steels are no exception to this, and probably exhibit a wider variety of failure mechanisms than any other category of material. While ultimate failure under excessive stress must occur and can be reasonably predicted by appropriate mechanical tests, premature failure is always dangerous, involving a considerable element of unpredictability. However, a detailed knowledge of structure and of the distribution of impurities in steels is gradually leading to a much better understanding of the origins and mechanisms of the various types of cracks encountered. Furthermore, the now wellestablished science of fracture mechanics allows the quantitative assessment of growth of cracks in various stress situations, to an extent that it is now frequently possible to predict the stress level to which steel structures can be confidently subjected without the risk of sudden failure.

Cleavage fracture in iron and steel Cleavage fracture is familiar in many minerals and inorganic crystalline solids as a crack propagation frequently associated with very little plastic deformation and

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occurring in a crystallographic fashion along planes of low indices, i.e. high atomic density. This behavior would appear to be an intrinsic characteristic of iron but it has been shown that iron, highly purified by zone refining and containing minimal concentrations of carbon, oxygen and nitrogen, is very ductile even at extremely low temperatures. For example, at 4.2 K reductions in area in tensile tests of up to 90 % have been observed with iron specimens of the highest available purity. As the carbon and nitrogen content of the iron is increased, the transition from ductile to brittle cleavage behavior takes place at increasing temperatures, until in some steels this can occur at ambient and above-ambient temperatures. Clearly, the significant variables in such a transition are of great basic and practical importance.

Factors influencing the onset of cleavage fracture The propagation of a cleavage crack in iron and steel requires much less energy than that associated with the growth of a ductile crack. There are several factors, some interrelated, which play an important part in the initiation of cleavage fracture: • • • • •

The temperature dependence of the yield stress The development of a sharp yield point Nucleation of cracks at twins Nucleation of cracks at carbide particles Grain size.

All body-centered cubic metals, including iron, show a marked temperature dependence of the yield stress, even when the interstitial impurity content is very low, i.e. the stress necessary to move dislocations, the Peierls-Nabarro stress, is strongly temperature dependent. This means that as the temperature is lowered the first dislocations to move will do so more rapidly as the velocity is proportional to the stress, and so the chances of forming a crack nucleus, e.g. by dislocation coalescence, will increase. The interstitial atoms, carbon and nitrogen, will cause the steel to exhibit a sharp yield point either by the catastrophic breakaway of dislocations from their interstitial atom atmospheres (Cottrell-Bilby theory), or by the rapid movement of freshly generated dislocations (Oilman-Johnson theory). In either case, the conditions are suitable for the localized rapid movement of dislocations as a result of high stresses, which provides a favorable situation for the nucleation of cracks by dislocation coalescence. The nucleation and the propagation of a cleavage crack must be distinguished clearly. Nucleation occurs when a critical value of the effective shear stress is reached, corresponding to a critical grouping, ideally a pile-up, of dislocations which can create a crack nucleus, e.g. by fracturing a carbide particle. In contrast, propagation of a crack depends on the magnitude of the local tensile stress, which must reach a critical level. Simple models of slip-nucleated fracture assume either interaction of dislocations or cracks formed in grain boundary carbides. However, recently it has been realized that both these structural features must be taken into account in deriving an expression for the critical fracture stress.

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This critical stress does not appear to be temperature dependent. At low temperatures the yield stress is higher, so the crack propagates when the plastic zone ahead of the crack is small, whereas at higher temperatures, the yield stress being smaller, a larger plastic zone is required to achieve the critical local tensile stress. This tensile stress has been determined for a wide variety of mild steels. The scatter probably arises from differences in test temperature and carbide dimensions. This is conclusive evidence for the role of finer grain sizes in increasing the resistance to crack propagation. Regarding grain boundary carbide size, effective crack nuclei will occur in particles above a certain critical size so that, if the size distribution of carbide particles in particular steel is known, it should be possible to predict its critical fracture stress. Therefore, in mild steels in which the structure is essentially ferrite grains containing carbide particles, the particle size distribution of carbides is the most important factor. In contrast, in bainitic and martensitic steels the austenite grains transform to lath structures where the lath width is usually between 0.2 and 2 ìm. The laths occur in bundles or packets with low angle boundaries between the laths. Larger misorientations occur across packet boundaries. In such structures, the packet width is the main micro structural feature controlling cleavage crack propagation.

Practical aspects of brittle fracture At the onset of fracture, elastic energy stored in the stressed steel is only partly used for creation of the new surfaces and the associated plastic deformation and the remainder provides kinetic energy to the crack. The phenomenon of brittle fracture became particularly prevalent with the introduction of welding as the major steel fabrication technique. Previously, brittle cracks often stopped at the joints of riveted plates but the steel structures resulting from welding provided continuous paths for crack propagation. Added to this, incorrect welding procedures can give rise to high stress concentrations and also to the formation of weld-zone cracks which may initiate brittle fracture. While brittle failures of steels have been experienced since the latter half of the nineteenth century when steel began to be used widely for structural work, the most serious failures have occurred later, as the demand for integral large steel structures has greatly increased, e.g. in ships, pipelines, bridges and pressure vessels. Spectacular failures took place in many of the all-welded Liberty ships produced during the Second World War, when nearly 1500 incidents involving serious brittle failure were recognized and nineteen ships broke completely in two without warning. Despite our increasing understanding of the phenomenon and the great improvements in steel making and in welding since then, serious brittle failures still occur. Brittle fractures of thick-walled steel pressure vessels are reminder that human error and lack of scientific control can be disastrous. Bearing in mind the temperature dependence of the failure behavior, and the widening use of steels at low temperatures, e.g. in Arctic pipelines, for storage of liquid gases etc., it is increasingly necessary to have steels with very low transition temperatures and high fracture toughness.

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While there are many variables to consider in achieving this end, including the detailed steel-making practice, the composition including trace elements and the fabrication processes involved, the most important is probably grain size refinement. The development of high strength low alloy steels (HSLA) or micro-alloyed steels, in the manufacture of which controlled rolling plays a vital part, has led to the production of structural steels with grain sizes combined with good strength levels and low transition temperatures. In these steels, to which small concentrations ( E/150) have such low notch toughness that brittle fracture can occur at nominal stresses in the elastic range at all temperatures and strain rates when flaws ace present. High-strength steel, aluminum and titanium alloys fall into this category. At low temperature fracture occurs by brittle cleavage, while at higher temperatures fracture occurs by low-energy rupture. It is under these conditions that fracture mechanics analysis is useful and appropriate. The notch toughness of low- and medium-strength bcc metals, as well as Be, Zn, and ceramic materials is strongly dependent on temperature. At low temperature the fracture occurs by cleavage while at high temperature the fracture occurs by ductile rupture. Thus, there is a transition from notch brittle to notch tough behavior with increasing temperature. In metals this transition occurs at 0.1 to 0.2 of the absolute melting temperature Tm, while in ceramics the transition occurs at about 0.5 to 0.7 Tm.

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A well-defined criterion is to base the transition temperature on the temperature at which the fracture becomes 100 percent cleavage. This point is known as nil ductility temperature (NDT). The NDT is the temperature at which fracture initiates with essentially no prior plastic deformation. Below the NDT the probability of ductile fracture is negligible.

Metallurgical Factors Affecting Transition Temperature Changes in transition temperature of over 55°C (100°F) can be produced by changes in the chemical composition or microstructure of mild steel. The largest changes in transition temperature result from changes in the amount of carbon and manganese. This transition temperature is lowered about 5.5°C (10°F) for each increase of 0.1 percent manganese. Increasing the carbon content also has a pronounced effect on the maximum energy and the shape of the energy transition-tempera lure curves. The Mn/C ratio should be at least 3/1 for satisfactory notch toughness. A maximum decrease of about 55°C (100°F) in transition temperature appears possible by going to higher Mn/C ratios. Phosphorus also has a strong effect in raising the transition temperature. The role of nitrogen is difficult to assess because of its interaction with other elements. It is, however, generally considered to be detrimental to notch toughness. Nickel is generally accepted to be beneficial to notch toughness in amounts up to 2 percent and seems to be particularly effective in lowering the ductility transition temperature. Silicon, in amounts over 0.25 percent, appears to raise the transition temperature. Molybdenum raises the transition almost as rapidly as carbon, while chromium has little effect. Notch toughness is particularly influenced by oxygen. For high-purity iron it was found that oxygen contents above 0.003 percent produced intergranular fracture and corresponding low energy absorption. Grain size has a strong effect on transition temperature. An increase of one ASTM number in the ferrite grain size (actually a decrease in grain diameter), results in a decrease in transition temperature of 16°C (30°F) for mild steel. Decreasing the grain diameter from ASTM grain size 5 to ASTM grain size 10 can change the 10 ft/lb Charpy V-notch transition temperature from about 39°C to -33°C (70°F to -60°F). The energy absorbed in the impact test of an alloy steel at a given test temperature generally increases with increasing tempering temperature. However, there is a minimum in the curve in the general region of 200 to 320°C (400 to 600°F). This has been called 260°C (500°F) embritilement, but because the temperature at which it occurs depends on both the composition of the steel and the tempering time, a more appropriate name is tempered-martensite embrittlement.

Drop-Weight Test and Other Large-Scale Tests Probably the chief deficiency of the Charpy impact test is that the small specimen is not always a realistic model of the actual situation. Not only does the small specimen lead to considerable scatter, but a specimen with a thickness of 10 mm (0.394 in)

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cannot provide the same constraint as would be found in a structure with a much greater thickness. The most logical approach to this problem is the development of tests that are capable of handling specimens at least several inches thick. The development of such tests and their rational method of analysis has been chiefly the work of Pellini and his coworkers at the Naval Research Laboratory. The basic need for large specimens resulted from the inability to produce fracture in small laboratory. The first development was the explosion-crack-starter test which featured a short, brittle weld bead deposited on the surface of a 14x14x1 in steel plate. The plate was placed over a circular die and dynamically loaded with an explosive charge. The brittle weld bead introduces a small natural crack in the test plate similar to a welddefect crack. Tests are carried out over a range of temperature and the appearance of the fracture determines the various transition temperatures. Below the NDT the fracture is a flat fracture running completely to the edges of the test plate. Above the nil ductility temperature a plastic bulge forms in the center of the plate, but the fracture is still a flat elastic fracture out to the plate edge. At still higher temperature the fracture does not propagate outside of the bulged region. The temperature at which elastic fracture no longer propagates to the edge of the plate is called the fracture transition elastic (FTE). The FTE marks the highest temperature of fracture propagation by purely elastic stresses. At yet higher temperature the extensive plasticity results in a helmet-type bulge. The temperature above which this fully ductile tearing occurs is the fracture transition plastic (FTP).

8.13. Brittle Fracture and Impact Testing: Part Two Abstract: This article describes how NDT, FTE, FTP are used in engineering design through the fracture analysis diagram (FAD). Temperature dependence of yield strength, tensile strength, and fracture strength is explained, as well as influence of various initial flaw sizes and the dynamic tear test (DT), as a highly versatile test both for low-strength ductile materials and high-strength low-toughness materials.

The first part of this article has introduced a number of terms dealing with brittle fracture, such as NDT, FTE, FTP, etc. The tests for determining these transition temperatures have been described. Before seeing how they are used in engineering design through the fracture analysis diagram, we redefine these transition points through reference to basic properties of the tension lest. The subambient temperature dependence of yield strength σo (Rp0.2) and ultimate tensile strength σu in a bcc metal are shown in Fig.1. For an unnotched specimen without flaws the material is ductile until a very low temperature, point A, where σo = σu. Point A represents the NDT temperature for a flaw-free material. The curve BCD represents the fracture strength of a specimen containing a small flaw (a < 0.1). The temperature corresponding to point C is the highest temperature at which the fracture strength σf ≈ σo. Point C represents the NDT for a specimen with a small crack or flaw.

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Fig.1. Temperature dependence of yield strength (σo), tensile strength (σu), and fracture strength for a steel containing flaws of different sizes The presence of a small flaw raises the NDT of a steel by about 200°F (110°C). Increasing the flaw size decreases the fracture stress curve, as in curve EF, until with increasing flaw size a limiting curve of fracture stress HJKL is reached. Below the NDT the limiting safe stress is 5,000 to 8,000 psi ( 35 to 55 MPa). Above the NDT the stress required for the unstable propagation of a long flaw (JKL) rises sharply with increasing temperature. This is the crack-arrest temperature curve (CAT). The CAT defines the highest temperature at which unstable crack propagation can occur at any stress level. Fracture will not occur for any point to the right of the CAT curve. The temperature above which elastic stresses cannot propagate a crack is the fracture transition elastic (FTE). This is defined by the temperature when the CAT curve crosses the yield-strength curve (point K). The fracture transition plastic (FTP) is the temperature where the CAT curve crosses the tensile-strength curve (point L). Above this temperature the material behaves as if it is flaw-free, for any crack, no matter how large, cannot propagate as an unstable fracture. Data obtained from the DWT and other large-scale fracture tests have been, assembled by Pellini and coworkers into a useful design procedure called the fracture analysis diagram (FAD). The NDT as determined by the DWT provides a key data point to start construction of the fracture analysis diagram. For mild steel below the NDT the CAT curve is flat. A stress level in excess of 5,000 to 8,000 psi (35 to 55 MPa) causes brittle fracture regardless of the size of the initial flaw. Extensive correlation between the NDT and Robertson CAT tests for a variety of structural steels have shown that the CAT curve bears a fixed relationship to the NDT temperature. Thus, the NDT+30°F provides a conservative estimate of the CAT curve at stress of σo/2. NDT+60°F provides an estimate of the CAT at σ = σo, that is, the FTE and NDT+120°F provides an estimate of the FTP. Therefore, for structural steels, once the NDT has been determined, the entire scope of the CAT curve can be established well enough for engineering design.

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The curve that has been traced out on Fig. 2 represents the worse possible case for large flaws in excess of 24 in.

Fig.2 Fracture-analysis diagram showing influence of various initial flaw sizes. One can envision a spectrum of curves translated upward and to the left for smaller, less severe flaws. Correlation with service failures and other tests has allowed the approximate determination of curves for a variety of initial flaw sizes. Thus, the FAD provides a generalized relationship of flaw size, stress, and temperature for lowcarbon structural steels of the type used in ship construction. The fracture analysis diagram can be used several ways in design. One simple approach would be to use the FAD to select a steel which has an FTE that is lower than the lowest expected service temperature. With this criterion the worst expected flaw would not propagate so long as the stress remained elastic. However, this procedure may prove to be too expensive and overconservative. A slightly less conservative design against brittle fracture, but still a practical approach, would be to design on the basis of an allowable stress level not exceeding σo/2. From Fig. 2 we see that any crack will not propagate under this stress so long as the temperature is not below NDT+30°F. If for example, the service temperature is not expected to be below 10°F, we would select a steel whose NDT is 10° - 30°, that is 20°F. The dynamic tear test (DT) can be used to construct the FAD. Below the NDT the fracture is brittle and has a flat, featureless surface devoid of tiny shear lips. At temperatures above the NDT there is a sharp rise in energy for fracture and the fracture surfaces begin to develop shear lips. The shear lips become progressively more prominent as the temperature is increased to the FTE. Above the FTE the fracture is ductile, void coalescence-type fracture. The fracture surface is a fibrous slant fracture. The upper shelf of energy represents the FTP. The lower half of the DT energy curve traces the temperature course of the CAT curve from NDT to FTE.

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The DT test is a highly versatile test because it is equally useful with low-strength ductile materials which show a high upper energy shelf and with high-strength lowtoughness materials which have a low value of upper shelf energy. The large size of the DT specimen provides a high degree of triaxial constraint and results in a minimum of scatter. Extensive correlations are being developed between DT results and fracture toughness and Cv test data.

8.14. Fatigue of Metals: Part One Abstract: It has been recognized since 1830 that a metal subjected to a repetitive or fluctuating stress will fail at a stress much lower than that required to cause fracture on a single application of load. Failures occurring under conditions of dynamic loading are called fatigue failures, presumably because it is generally observed that these failures occur only after a considerable period of service. Fatigue has become progressively more prevalent as technology has developed a greater amount of equipment, such as automobiles, aircraft, compressors, pumps, turbines, etc., subject to repeated loading and vibration. Today it is often stated that fatigue accounts for al least 90 percent of all service failures due to mechanical causes.

It has been recognized since 1830 that a metal subjected to a repetitive or fluctuating stress will fail at a stress much lower than that required to cause fracture on a single application of load. Failures occurring under conditions of dynamic loading are called fatigue failures, presumably because it is generally observed that these failures occur only after a considerable period of service. Fatigue has become progressively more prevalent as technology has developed a greater amount of equipment, such as automobiles, aircraft, compressors, pumps, turbines, etc., subject to repeated loading and vibration. Today it is often stated that fatigue accounts for al least 90 percent of all service failures due to mechanical causes. A fatigue failure is particularly insidious because it occurs without any obvious warning. Fatigue results in a brittle-appearing fracture, with no gross deformation at the fracture. On a macroscopic scale the fracture surface is usually normal to the direction of the principal tensile stress. A fatigue failure can usually be recognized from the appearance of the fracture surface, which shows a smooth region, due to the rubbing action as the crack propagated through the section, and a rough region, where the member has failed in a ductile manner when the cross section was no longer able to carry the load. Frequently the progress of the fracture is indicated by a series of rings, or "beach marks", progressing inward from the point of initiation of the failure. Three basic factors are necessary to cause fatigue failure. These are: • • •

maximum tensile stress of sufficiently high value, large enough variation or fluctuation in the applied stress, and sufficiently large number of cycles of the applied stress.

In addition, there are a host of other variables, such as stress concentration, corrosion, temperature, overload, metallurgical structure, residual stresses, and combined stresses, which tend to alter the conditions for fatigue. Since we have not yet gained a complete understanding of what causes fatigue in metals, it will be necessary to discuss each of these factors from an essentially empirical standpoint. Because of the mass of data of this type, it will be possible to describe only the highlights of the relationship between these factors and fatigue.

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Stress Cycles At the outset it will be advantageous to define briefly the general types of fluctuating stresses which can cause fatigue. Figure 1 serves to illustrate typical fatigue stress cycles. Figure 1a illustrates a completely reversed cycle of stress of sinusoidal form. For this type of stress cycle the maximum and minimum stresses are equal. Tensile stress is considered positive, and compressive stress is negative. Figure 1b illustrates a repeated stress cycle in which the maximum stress σmax (Rmax) and minimum stress σmin (Rmin) are not equal. In this illustration they are both tension, but a repeated stress cycle could just as well contain maximum and minimum stresses of opposite signs or both in compression. Figure 1c illustrates a complicated stress cycle which might be encountered in a part such as an aircraft wing which is subjected to periodic unpredictable overloads due to gusts.

Figure 1. Typical fatigue stress cycles. (a) Reversed stress; (b) repeated stress; (c) irregular or random stress cycle. A fluctuating stress cycle can be considered to be made up of two components, a mean, or steady, stress σm (Rm), and an alternating, or variable, stress σa. We must also consider the range of stress σr. As can be seen from Fig. 1b, the range of stress is the algebratic difference between the maximum and minimum stress in a cycle.

The S-N Curve

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The basic method of presenting engineering fatigue data is by means of the S-N curve, a plot of stress S against the number of cycles to failure N. A log scale is almost always used for N. The value of stress that is plotted can be σa, σmax, or σmin. The stress values are usually nominal stresses, i.e., there is no adjustment for stress concentration. The S-N relationship is determined for a specified value of σm, R (R=σmin/σmax), or A (A=σa/σm). Most determinations of the fatigue properties of materials have been made in completed reversed bending, where the mean stress is zero. It will be noted that this S-N curve is concerned chiefly with fatigue failure at high numbers of cycles (N > 105 cycles). Under these conditions the stress, on a gross scale, is elastic, but as we shall see shortly the metal deforms plastically in a highly localized way. At higher stresses the fatigue life is progressively decreased, but the gross plastic deformation makes interpretation difficult in terms of stress. For the low-cycle fatigue region (N < 104 or 105 cycles) tests are conducted with controlled cycles of elastic plus plastic strain instead of controlled load or stress cycles. The usual procedure for determining an S-N curve is to test the first specimen at a high stress where failure is expected in a fairly short number of cycles, e.g., at about two-thirds the static tensile strength of the material. The test stress is decreased for each succeeding specimen until one or two specimens do not fail in the specified numbers of cycles, which is usually at least 107 cycles. The highest stress at which a runout (non-failure) is obtained is taken as the fatigue limit. For materials without a fatigue limit the test is usually terminated for practical considerations at a low stress where the life is about 108 or 5x108 cycles. The S-N curve is usually determined with about 8 to 12 specimens.

Statistical Nature of Fatigue A considerable amount of interest has been shown in the statistical analysis of fatigue data and in reasons for the variability in fatigue-test results. Since fatigue life and fatigue limit are statistical quantities, it must be realized that considerable deviation from an average curve determined with only a few specimens is to be expected. It is necessary to think in terms of the probability of a specimen attaining a certain life at a given stress or the probability of failure at a given stress in the vicinity of the fatigue limit. To do this requires the testing of considerably more specimens than in the past so that the statistical parameters for estimating these probabilities can be determined. The basic method for expressing fatigue data should then be a three-dimensional surface representing the relationship between stress, number of cycles to failure, and probability of failure. In determining the fatigue limit of a material, it should be recognized that each specimen has its own fatigue limit, a stress above which it will fail but below which it will not fail, and that this critical stress varies from specimen to specimen for very obscure reasons. It is known that inclusions in steel have an important effect on the fatigue limit and its variability, but even vacuum-melted steel shows appreciable scatter in fatigue limit.

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The statistical problem of accurately determining the fatigue limit is complicated by the fact that we cannot measure the individual value of the fatigue limit for any given specimen. We can only test a specimen at a particular stress, and if the specimen fails, then the stress was somewhere above the fatigue limit of the specimen. The two statistical methods which are used for making a statistical estimate of the fatigue limit are called probit analysis and the staircase method. The procedures for applying these methods of analysis to the determination of the fatigue limit have been well established.

Effect of Mean Stress on Fatigue Much of the fatigue data in the literature have been determined for conditions of completely reversed cycles of stress, σm = 0. However, conditions are frequently met in engineering practice where the stress situation consists of an alternating stress and a superimposed mean, or steady, stress. There are several possible methods of determining an S-N diagram for a situation where the mean stress is not equal to zero.

Cyclic Stress-Strain Curve Cyclic strain controlled fatigue, as opposed to our previous discussion of cyclic stress controlled fatigue, occurs when the strain amplitude is held constant during cycling. Strain controlled cyclic loading is found in thermal cycling, where a component expands and contracts in response to fluctuations in the operating temperature. In a more general view, the localized plastic strains at a notch subjected to either cyclic stress or strain conditions result in strain controlled conditions near the root of the notch due to the constraint effect of the larger surrounding mass of essentially elastically deformed material. Since plastic deformation is not completely reversible, modifications to the structure occur during cyclic straining and these can result in changes in the stress-strain response. Depending on the initial state a metal may undergo cyclic hardening, cyclic softening, or remain cyclically stable. It is not uncommon for all three behaviors to occur in a given material depending on the initial state of the material and the test conditions. Generally the hysteresis loop stabilizes after about 100 cycles and the material arrives at an equilibrium condition for the imposed strain amplitude. The cyclically stabilized stress-strain curve may be quite different from the stress-strain curve obtained on monotonic static loading. The cyclic stress-strain curve is usually determined by connecting the tips of stable hysteresis loops from constant-strainamplitude fatigue tests of specimens cycled at different strain amplitudes. Under conditions where saturation of the hysteresis loop is not obtained, the maximum stress amplitude for hardening or the minimum stress amplitude for softening is used. Sometimes the stress is taken at 50 percent of the life to failure. Several shortcut procedures have been developed.

Low-Cycle Fatigue Although historically fatigue studies have been concerned with conditions of service in which failure occurred at more than 104 cycles of stress, there is growing recognition of engineering failures which occur at relatively high stress and low

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numbers of cycles to failure. This type of fatigue failure must be considered in the design of nuclear pressure vessels, steam turbines, and most other types of power machinery. Low-cycle fatigue conditions frequently are created where the repeated stresses are of thermal origin. Since thermal stresses arise from the thermal expansion of the material, it is easy to see that in this case fatigue results from cyclic strain rather than from cyclic stress.

8.15. Analyzing Failures of Metal Components: Part One Abstract: Many elements of fracture have been used to describe and categorize the types of fractures encountered in the laboratory and in service. These elements include loading conditions, rate of crack growth, and macroscopic and microscopic appearance of fracture surfaces. Failure analysis often find itself useful to classify fractures on a macroscopic scale as ductile fractures, brittle fractures, fatigue fractures and fractures resulting from the combined effects of stress and environment.

Many elements of fracture have been used to describe and categorize the types of fractures encountered in the laboratory and in service. These elements include loading conditions, rate of crack growth, and macroscopic and microscopic appearance of fracture surfaces. Failure analysis often find itself useful to classify fractures on a macroscopic scale as ductile fractures, brittle fractures, fatigue fractures and fractures resulting from the combined effects of stress and environment. The last group includes stress-corrosion cracking and liquid-metal embrittlement, interstitial embrittlement, corrosion fatigue and stress rupture.

Ductile Fractures Ductile fractures are characterized by tearing of metal accompanied by appreciable gross plastic deformation and expenditure of considerable energy. Ductile tensile fractures in most materials have a gray, fibrous appearance and are classified on a macroscopic scale as either flat-face (square) or shear-face (slant) fractures.

Brittle Fractures Brittle fractures are characterized by rapid crack propagation with less expenditure of energy than with ductile fractures and without appreciable gross plastic deformation. Brittle tensile fractures have a bright granular appearance, are of the flat-face type, and are produced under plain-strain conditions with little or no necking.

Fatigue Fractures Fatigue fractures result from cyclic loading, and appear brittle on a macroscopic scale. They are characterized by incremental propagation of cracks until the cross section has been reduced to where it can no longer support the maximum applied load and fast fracture ensues. In the fatigue fracture the fracture surface consists of three distinct zones: a fairly smooth, multiple origin fatigue zone containing "ratchet marks", a low cycle, rougher fatigue zone and a single cycle final fracture zone.

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When designing modern equipment to operate in severe environments, a designer is confronted with many complex problems in selecting and evaluating materials, processing, expected loadings and design stresses. Components in turbines, reactors, missiles, submarines and cryogenic equipment may be subjected to such conditions as extremely high or low temperature, corrosive liquids, high vacuum, progressive deterioration due to radiation damage and surface wear. Materials selection must often be confined to a small group of metals for outstanding resistance in one characteristic, such as inertness to the environment in chemical processing equipment. However, many other factors must be considered such as strength, toughness, fabricability and wear resistance, before selection and design can be finalized. Detailed analysis of failures encountered in developing a prototype (or in a service component) is vital before appropriate changes can be made to assure a reliable product. In general, service failures may arise from many causes. For mechanical equipment, these causes might be broken down roughly into three categories, as follows: Design inadequacies. Sharp corners or abnormal stress-raisers, inadequate fasteners, wrong material or heat treatment, unforeseen conditions of service, and lack of accurate stress analysis are included. Processing and fabrication. About half of these may be due to metallurgical factors such as quench cracks, improper heat treatment, forging or casting defects, nonmetallic inclusions; the other half are due to misalignments, weld flaws, improper machining or assembly, grinding cracks, cold straightening, and the like. Environmental and service deterioration. These include overloads, chemical attack, wear, corrosion, diffusion, and improper maintenance. A "failure" usually occurs as: 1. fracture 2. excessive deformation 3. deterioration The failure mechanism is usually a material failure that is controlled by the entire environment and history. Failures of Category I (design considerations) result from mistakes or incompetence of the designer. Regarding failures due to flaws developed by processing or fabrication (Category II), few, if any, standard tests cover all of the possible inherent defects that may be induced by such operations as casting, forging, welding, machining, grinding, heat treating, plating, chemical diffusion, or careless assembly operations. Category III failures, caused by deterioration, can not be predicted by standard tests that evaluate materials. In some instances unforeseen vibrations or overload conditions may develop to cause failure. In others, service induced damage may develop fatigue failure. Many service conditions involve extremely rapid rates of heating, or include radiation damage, ablation, corrosion or various types of wear.

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Deterioration during service in an aggressive environment needs to be given special consideration. There are many types of surface disintegration, chemical activity or metal transfer that affect stability of the component. These are influenced by the time, temperature and dosage of the critical factors in the environment.

8.16. Analyzing Failures of Metal Components: Part Two Abstract: When studying a failure, great care must be used to avoid destroying important evidence. Detailed studies often require careful documentation of the service history (time, temperature, loadings and environment), supplemented by chemical analysis and electron micrographs. Further study of the sequence of events leading up to the failure, plus knowledge of the location, markings and condition of all adjacent parts after the incident, are necessary to confirm the analysis beyond reasonable doubt. Of course, there is always the possibility of an unforeseen loading, unreported collision, or unanticipated vibration that may develop to cause premature failure. These factors may usually be diagnosed by careful examination and gathering of facts.

Analysis of failure When studying a failure, great care must be used to avoid destroying important evidence. Detailed studies often require careful documentation of the service history (time, temperature, loadings and environment), supplemented by chemical analysis and electron micrographs. Further study of the sequence of events leading up to the failure, plus knowledge of the location, markings and condition of all adjacent parts after the incident, are necessary to confirm the analysis beyond reasonable doubt. Of course, there is always the possibility of an unforeseen loading, unreported collision, or unanticipated vibration that may develop to cause premature failure. These factors may usually be diagnosed by careful examination and gathering of facts. As supplementary information, we have appended a list of classifications as to cause of failure, and general guidelines to be following in determining the major cause or causes. Initial Observations. Make a detailed study by visual inspection of the actual component that failed, preferably at the failure site as soon after the failure as possible. Profuse color photographs are essential to record every detail for later review. Background Data. Gather all available data concerned with specifications and drawings, component design, fabrication, repairs, maintenance and service use. Concentration on obtaining facts pertinent to all possible failure modes is essential. Particular attention to environmental details, including normal service loads (as well as accidental overloads and cyclic loading) and resulting stress, temperature variations and gradients are desirable. Supplementary Laboratory Studies. Make tests to verify that the material in the component actually possesses the specified composition, dimensions, processing and properties. Supplementary studies may be needed (for example, composition of

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corrosion products, simulated service or environmental tests, stress analysis, determination of microstructure and development of cracks, dynamic strain measurements, elevated or low temperature tests and surface replicas) to define the factors which contributed to the failure. Electron probe X-ray analysis can be useful in examining inclusions, microsegregation, or the composition of oxide or surface contaminants. Frequently, examination of a fracture face with a low-power binocular microscope can reveal the type and cause of failure. However, the observer must be familiar with a wide variety of fracture textures, and be able to approach the examination without a preconceived idea of the cause for the fracture. Further confirmation can often be obtained by studies of cracks and structures at higher magnification, using small samples from the regions of failure. Synthesis of Failure. List not only all positive facts and evidence, but also all the negative responses to the questions that may be asked about the failure. Sometimes it is important to know that specific things did not happen or certain evidence did not appear to determine what could have happened. From a tabulation of these data, the actual failure should be synthesized to include all available items of the evidence. The cause will usually be classified in one of the categories outlined in the following list, and corrective action or applied research guidance can be recommended. Appropriate solutions may involve redesign, change of alloy and/or processing, quality control, protection against environment, changes in maintenance schedules, or restrictions on service loads or service life.

Classification of Failure Causes

Failures Due to Faulty Processing 1. Flaws due to faulty composition (inclusions, embrittling impurities, wrong material). 2. Defects originating in ingot making and casting (segregation, unsoundness, porosity, pipes, nonmetallic inclusions). 3. Defects due to working (laps, seams, shatter cracks, hot-short splits, delamination, and excess local plastic deformation). 4. Irregularities and mistakes due to machining, grinding, or stamping (gouges, burns, tearing, fins, cracks, embrittlement). 5. Defects due to welding (porosity, undercuts, cracks, residual stress and lack of penetration, under bead cracking, heat affected zone). 6. Abnormalities due to heat treating (overheating, burning, quench cracking, grain growth, excessive retained austenite, decarburization, precipitation). 7. Flaws due to case hardening (intergranular carbides, soft core and wrong heat cycles). 8. Defects due to surface treatments (cleaning, plating, coating, chemical diffusion, hydrogen embrittlement). 9. Careless assembly (mismatch of mating parts, entrained dirt or abrasive, residual stress, gouges or injury to parts and the like). 10. Parting line failures in forging due to poor transverse properties.

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Failures due to Faulty Design Considerations or Misapplication of Material 1. Ductile failure (excess deformation, elastic or plastic; tearing or shear fracture). 2. Brittle fracture (from flaw or stress raiser of critical size). 3. Fatigue failure (load cycling, strain cycling, thermal cycling, corrosion fatigue, rolling contact fatigue, fretting fatigue). 4. High-temperature failure (creep, oxidation, local melting, warping). 5. Static delayed fractures (hydrogen embrittlement, caustic embrittlement, environmentally stimulated slow growth of flaws). 6. Excessively severe stress raisers inherent in the design. 7. Inadequate stress analysis, or impossibility of a rational stress calculation in a complex part. 8. Mistake in designing on basis of static tensile properties, instead of the significant material properties that measure the resistance of the material to each possible failure mode.

Failure Due to Deterioration During Service- Conditions 1. Overload or unforeseen loading conditions. 2. Wear (erosion, galling, seizing, gouging, cavitations). 3. Corrosion (including chemical attack, stress corrosion, corrosion fatigue), dezincification, graphitization of cast iron, contamination by atmosphere. 4. Inadequate or misdirected maintenance or improper repair (welding, grinding, punching holes, cold straightening, and so forth). 5. Disintegration due to chemical attack or attack by liquid metals or platings at elevated temperatures. 6. Radiation damage (sometimes must decontaminate for examination which may destroy vital evidence of cause of failure), varies with time, temperature, environment and dosage. 7. Accidental conditions (abnormal operating temperatures, severe vibration, sonic vibrations, impact or unforeseen collisions, ablation, thermal shock and so forth).

8.17. Fatigue crack growth Abstract: The aircraft industry has leaded the effort to understand and predict fatigue crack growth. They have developed the safe-life or fail-safe design approach. In this method, a component is designed in a way that if a crack forms, it will not grow to a critical size between specified inspection intervals. Thus, by knowing the material growth rate characteristics and with regular inspections, a cracked component may be kept in service for an extended useful life.

The major share of the fatigue life of the component may be taken up in the propagation of crack. By applying fracture mechanics principles it is possible to predict the number of cycles spent in growing a crack to some specified length or to final failure. The aircraft industry has leaded the effort to understand and predict fatigue crack growth. They have developed the safe-life or fail-safe design approach. In this

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method, a component is designed in a way that if a crack forms, it will not grow to a critical size between specified inspection intervals. Thus, by knowing the material growth rate characteristics and with regular inspections, a cracked component may be kept in service for an extended useful life. This concept is shown schematically in Fig. 1.

Figure 1. Extended service life of a cracked component

Fatigue Crack Growth Curves Typical constant amplitude crack propagation data are shown in Fig. 2. The crack length, a, is plotted versus the corresponding number of cycles, N, at which the crack was measured.

Figure 2. Constant amplitude crack growth data As shown, most of the life of the component is spent while the crack length is relatively small. In addition, the crack growth rate increases with increased applied stress. The crack growth rate, da/dN, is obtained by taking the derivative of the above crack length, a, versus cycles, N, curve. Two generally accepted numerical approaches for obtaining this derivative are the spline fitting method and the incremental polynomial method. These methods are explained in detail in many numerical methods textbooks. Values of log da/dN can then be plotted versus log ΔK, for a given crack length, using the equation

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(1) where Δσ is the remote stress applied to the component as shown in Fig. 3.

Figure 3. Remote stress range A plot of log da/dN versus log ΔΚ, a sigmoidal curve, is shown in Fig. 4. This curve may be divided into three regions. At low stress intensities, Region I, cracking behavior is associated with threshold, ΔKth, effects. In the mid-region, Region II, the curve is essentially linear. Many structures operate in this region. Finally, in the Region III, at high ΔK values, crack growth rates are extremely high and little fatigue life is involved.

Figure 4. Three regions of crack growth rate curve

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Region II Most of the current applications of LEFM concepts to describe crack growth behavior are associated with Region II. In this region the slope of the log da/dN versus log ΔK curve is approximately linear and lies roughly between 10-6 and 10-3 in/cycle. Many curve fits to this region have been suggested. The Paris equation, which was proposed in the early 1960s, is the most widely accepted. In this equation (2) where C and m are material constants and ΔK is the stress intensity range Kmax Kmin. Values of the exponent, m, are usually between 3 and 4. These range from 2,3 to 6,7 with a sample average of m = 3,5. In addition, tests may be performed. ASTM E647 sets guidelines for these tests. The crack growth life, in terms of cycles to failure, may be calculated using Eq. (2). The relation may be generally described by

Thus, cycles to failure, Nf, may be calculated as

(3)

where ai is the initial crack length and af is the final (critical) crack length. Using the Paris formulation,

(4)

Because ΔK is a function of the crack length and a correction factor that is dependent on crack length [see Eq. (1)], the integration above must often be solved numerically. As a first approximation, the correction factor can be calculated at the initial crack length and Eq. (4) can be evaluated in closed form. As an example of closed form integration, fatigue life calculations for a small edgecrack in a large plate are performed below. In this case the correction factor, f(g) does not vary with crack length. The stress intensity factor range is (5)

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Substituting into the Paris equation yields (6) Separating variables and integrating (for m2) gives

(7)

Before this equation may be solved, the final crack size, af, must be evaluated. This may be done using as follows:

(8)

For more complicated formulations of ΔK, where the correction factor varies with the crack length, a, iterative procedures may be required to solve for af in Eq. (8). It is important to note that the fatigue-life estimation is strongly dependent on ai, and generally not sensitive to af (when ai«af). Large changes in af result in small changes of Nf as shown schematically in Fig. 5.

Figure 5. Effect of final crack size on life

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8.18. Structural Features of Fatigue Abstract: Studies of the basic structural changes that occur when a metal is subjected to cyclic stress have found it convenient to divide the fatigue process into the following stages:

• • • •

Crack initiation includes the early development of fatigue damage which can be removed by a suitable thermal anneal. Slip-band crack growth involves the deepening of the initial crack on planes of high shear stress. This frequently is called stage I crack growth. Crack growth on planes of high tensile stress involves growth of welldefined crack in direction normal to maximum tensile stress. Usually called stage II crack growth. Ultimate ductile failure occurs when the crack reaches sufficient length so that the remaining cross section cannot support the applied load.

Studies of the basic structural changes that occur when a metal is subjected to cyclic stress have found it convenient to divide the fatigue process into the following stages: • • • •

Crack initiation includes the early development of fatigue damage which can be removed by a suitable thermal anneal. Slip-band crack growth involves the deepening of the initial crack on planes of high shear stress. This frequently is called stage I crack growth. Crack growth on planes of high tensile stress involves growth of welldefined crack in direction normal to maximum tensile stress. Usually called stage II crack growth. Ultimate ductile failure occurs when the crack reaches sufficient length so that the remaining cross section cannot support the applied load.

The relative proportion of the total cycles to failure that are involved with each stage depends on the test conditions and the material. However, it is well established that a fatigue crack can be formed before 10 percent of the total life of the specimen has elapsed. There is, of course, considerable ambiguity in deciding when a deepened slip band should be called a crack. In general, larger proportions of the total cycles to failure are involved with the propagation of stage II cracks in low-cycle fatigue than in long-life fatigue, while stage I crack growth comprises the largest segment for low-stress, high-cycle fatigue. If the tensile stress is high, as in the fatigue of sharply notched specimens, stage I crack growth may not be observed at all. An overpowering structural consideration in fatigue is the fact that fatigue cracks usually are initiated at a free surface. In those rare instances where fatigue cracks initiate in the interior there is always an interface involved, such as the interface of a carburized surface layer and the base metal. Fatigue has certain things in common with plastic flow and fracture under static or unidirectional deformation. The work of Gough has shown that a metal deforms under cyclic strain by slip on the same atomic planes and in the same

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crystallographic directions as in unidirectional strain. Whereas with unidirectional deformation slip is usually widespread throughout all the grains, in fatigue some grains will show slip lines while other grains will give no evidence of slip. Slip lines are generally formed during the first few thousand cycles of stress. Successive cycles produce additional slip bands, but the number of slip bands is not directly proportional to the number of cycles of stress. In many metals the increase in visible slip soon reaches a saturation value, which is observed as distorted regions of heavy slip. Cracks are usually found to occur in the regions of heavy deformation parallel to what was originally a slip band. Slip bands have been observed at stresses below the fatigue limit of ferrous materials. Therefore, the occurrence of slip during fatigue does not in itself mean that a crack will form. A study of crack formation in fatigue can be facilitated by interrupting the fatigue test to remove the deformed surface by electro polishing. There will generally be several slip bands which are more persistent than the rest and which will remain visible when the other slip lines have been polished away. Such slip bands have been observed after only 5 percent of the total life of the specimen. These persistent slip bands are embryonic fatigue cracks, since they open into wide cracks on the application of small tensile strains. Once formed, fatigue cracks tend to propagate initially along slip planes, although they later take a direction normal to the maximum applied tensile stress. Fatigue-crack propagation is ordinarily transgranular. An important structural feature, which appears to be unique to fatigue deformation, is the formation on the surface of ridges and grooves called slip-band extrusions and slip-band intrusions. Extremely careful metallography on tapered sections through the surface of the specimen has shown that fatigue cracks initiate at intrusions and extrusions. W. A. Wood, who made many basic contributions to the understanding of the mechanism of fatigue, suggested a mechanism for producing slip-band extrusions and intrusions. He interpreted microscopic observations of slip produced by fatigue as indicating that the slip bands are the result of a systematic buildup of fine slip movements, corresponding to movements of the order of 10-7cm rather than steps of 10-5 to 10-4 cm, which are observed for static slip hands. Such a mechanism is believed to allow for the accommodation of the large total strain (summation of the micro strain in each cycle) without causing appreciable strain hardening. The notch would be a stress raiser with a notch root of atomic dimensions. Such a situation might well be the start of a fatigue crack. This mechanism for the initiation of a fatigue crack is in agreement with the facts that fatigue cracks start at surfaces and that cracks have been found lo initiate at slipband intrusions and extrusions. Extensive structural studies of dislocation arrangements in persistent slip bands have brought much basic understanding to the fatigue fracture process. The stage I crack propagates initially along the persistent slip bands. In a polycrystalline metal the crack may extend for only a few grain diameters before the crack propagation changes to stage II. The rate of crack propagation in stage I is generally very low, on the order of angstroms per cycle, compared with crack

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propagation rates of microns per cycle for stage II. The fracture surface of stage I fractures is practically featureless. Failure to observe striations on a fatigue surface may be due to a very small spacing that cannot be resolved with the observational method used, insufficient ductility at the crack tip to produce a ripple by plastic deformation that is large enough to be observed or obliteration of the striations by some sort of damage lo the surface. Since stage II cracking does not occur for the entire fatigue life, it does not follow that counting striations will give the complete history of cycles to failure. Stage II crack propagation occurs by a plastic blunting process that is illustrated in Fig. 1.

Fig.1 Plastic blunting process for growth of stage II fatigue crack At the start of the loading cycle the crack tip is sharp (Fig. 1 a). As the tensile load is applied the small double notch at the crack tip concentrates the slip along planes at 45° to the plane of the crack (Fig. 1b). As the crack widens to its maximum extension (Fig. 1c) it grows longer by plastic shearing and at the same time its tip becomes blunter. When the load is changed lo compression the slip direction in the end zones is reversed. The crack faces are crushed together and the new crack surface created in tension is forced into the plane of the crack (Fig. 1e).

From the History of Iron and Steel Making: Part One Abstract: The Voelklingen Ironworks was founded in 1883 and developed into one of the leading iron and steel works in Germany and Europe. The singular compactness of the ironworks blast-furnace unit with six blast furnaces and the inclined elevators for transporting iron ore and coke, unparalleled in the world, form a skyline that has made a lasting impression on the German Saar valley for more than 100 years. After its closure in 1986, it was declared as a historical monument, and in 1994 was awarded the status of a World Cultural Heritage by UNESCO. Today, Voelklingen Ironworks is the only plant worldwide surviving in its original form from the heyday of iron and steel industry.

The Voelklingen Ironworks was founded in 1883 and developed into one of the leading iron and steel works in Germany and Europe. The singular compactness of the ironworks blast-furnace unit with six blast furnaces and the inclined elevators for transporting iron ore and coke, unparalleled in the world, form a skyline that has made a lasting impression on the German Saar valley for more than 100 years.

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After its closure in 1986, it was declared as a historical monument, and in 1994 was awarded the status of a World Cultural Heritage by UNESCO. Today, Voelklingen Ironworks is the only plant worldwide surviving in its original form from the heyday of iron and steel industry.

Figure 1. Voelklingen Ironworks today: a historical monument, awarded the status of a World Cultural Heritage by UNESCO.

Early history of industrial iron and steel making plant development 1881. Businessman Carl Roehling buys the non-operating facility in Voelklingen. New infrastructure improvements such as construction of the Saarbruecken-Trier railway, the construction of Saar Coal Channels and the purchase of patent rights for the Thomas steel process ensures rapid rise of the company. The Englishman Sidney Gilchrist Thomas`s invention permits the production of steel from high phosphorus minette iron-ore from Lothringen (Loraine). Minette means "minor ore", which has an iron content of only around 30%. 1883. The first blast furnace in the Voelklingen Ironworks goes into operation. It was the largest in the Saar iron and steel region with its 200 m³ capacity. The second blast furnace was started in 1885, and then the blast furnace group was completed in a quick succession, from 1888 (furnace 3) to 1907 (furnace 7).

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Figure 2. Voelklingen Ironworks plant outlook in the beginning of 20th century. 1897. The first coke oven battery was erected directly adjacent to the blast furnace in Voelklingen. 1900. Two years earlier experiments with the use of blast furnace gas for the driving of power engines were successfully completed: the first blast furnace gas blast engine with a power of 600 HP went into operation. Blast furnace gas itself was used for propelling internal combustion engines. 1907. The induction furnace developed by Hermann Roechling and Wilhelm Rodenhauser for the production of high quality stainless steels goes into operation. 1911. The inclined ore lift is built in Voelklingen. This was a globally unique system for supplying all blast furnaces with raw materials by using a single rail system, developed after seven years of construction work. More than 300 hoppers were in use day and night for the decades, and there was more than 6 km of tracks. However, a serious disadvantage of this system became evident much later: the blast furnaces could not be extended above 27 meters in height. Therefore, production at Voelklingen Ironworks could not keep pace with the development of the iron and steel industry in the last quarter of 20th century. 1915. The construction of Siemens-Martin steel works expands production capacity with the production of top quality stainless steels.

The blast furnace group Blast furnace gas mains lend contour to the silhouette of the ironworks. They channeled off the gas from the six ironworks blast furnaces. The blast furnaces themselves are hidden behind scaffolding, pipes and chimneys.

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Figure 3. Blast furnace gas and wind pipes. Some 130 tons of pig iron were tapped off in a 2 hour rhythm, or about 1100 tons per blast furnace daily. The blast wind heaters are arranged in groups of three in front of the blast furnaces. They heated the combustion air blast used for smelting iron to 1100°C. In the lower area of the furnace where glowing coke and hot air met temperatures rose up to 2000°C.

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Figure 4. Tapping off iron. The charging platform was at a height of 30 meters, and the blast furnace men filled furnaces from there, and controlled the haulage system. Giant motors drove the haulage cable pulleys, the thick steel pulled heavily loaded hoppers up to the platform. The system was controlled by a switching station located on the charging platform.

Figure 5. A section of the charging platform.

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As it is well known today, the blast furnace operation needed a raw material with a high carbon content for use in the blast furnaces, in order to draw off oxygen from the iron ore and to generate heat needed for the smelting process. At that time, it became already evident that coal is unsuitable for this purpose, due to its too low caloric value and too high sulphur content. Therefore, coke plant was added to the ironworks, and the first coking ovens were lit in Voelklingen in 1897. The sheet steel silo that towers above the coking plant dates back to this period, and it is one of the oldest surviving structures of the Voelklingen Ironworks. Another peculiarity in the vicinity of blast furnaces was the Craftman`s lane (Handwerkergasse), the ironworks building department. Bricklayers, locksmiths, carpenters and other skilled workers had their workshops there. The art of improvisation was in demand: the group of skilled workers had to be constantly ready to carry out repairs, build and make extensions.

Figure 6. A tough place for work: a foundryman in 20`s.

Recycling and environment As technology pace went forward, production of iron and steel manufacturing byproducts was pushed ahead. In 1912, Voelklingen Ironworks` product range already included Thomas slag as a fertilizer, ammoniac, benzene and various tar products. The processing of waste materials from the coking plant in particular -- in the so called coal by-product operations -- proves to be important source of income for the company. In 1927, Roechling company, the owner of Voelklingen Ironworks, built a cement works for the further processing and economic use of blast furnace slag. The waste product slag was also increasingly used in road and house building. Approximately in the same time, sintering technology offered new opportunity to recycle waste products from the smelting process, i.e. ore dust, blast furnace dust.

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The largest sintering plant in the world was built in Voelklingen in 1928. Materials with a grain size that is too fine for use in the blast furnaces were heated to 1200°C to form a sinter cake in the sinter plant and then broken into the proper grain sizes. In spite of these recycling efforts, emissions were still high, thus hampering environment. The blast furnace blew 32 tons of dust daily into the atmosphere. No one could hang laundry in the garden when wind was coming from the south-west. Every evening the people of Voelklingen had to wipe a thick layer of brown dust from their window sills.

From the History of Iron and Steel Making: Part Two Abstract: In the Voelklingen Ironworks, which was founded in 1883 and now is awarded the status of a World Cultural Heritage by UNESCO, history of creating the wind element can be seen directly. The heart of the entire plant is over 6000 m² (65000 sq feet) large blowerhouse, in which gigantic machines produced the blast necessary for iron making.

Fire, water and wind are the essential elements of the ironworks. The wind fans the fire and gives it the power to melt the iron; the cooling effects of the water holds the destructive power of fire under control. In the Voelklingen Ironworks, which was founded in 1883 and now is awarded the status of a World Cultural Heritage by UNESCO, history of creating the wind element can be seen directly. The heart of the entire plant is over 6000 m² (65000 sq feet) large blowerhouse, in which gigantic machines produced the blast necessary for iron making.

Figure 1. The Voelklingen Ironworks, now a World Cultural Heritage protected by UNESCO.

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The blasting engines The striking idea for the new technology was late in coming, but it was to revolutionize iron production: in 1878 German engineers Otto and Langen constructed the first gas motor. The mechanical engineering works Deutz builds the first the first blast furnace gas engine in 1894. It was an ingenious innovation: at last it was now possible to use the gas, which was produced by the blast furnace in gigantic quantities in the process of smelting iron, to drive engines. The iron and steel industry was now able to recycle waste material. Voelklingen Ironworks puts the new technology to use at once. The first large scale gas engine was ordered from M.A.N. in Nuremberg in 1899, as a generator unit for electricity, and went into operation in 1901. A total of 30 gas engines were in operation in Voelklingen Ironworks, and they were not only used not only as blast engines and power generators, but also to drive pumps and rolling mills. In July 1903 Ironworks ordered the oldest surviving machine in the hall from the Augsburg-Nuernberg mechanical engineering works "at a price of 300 000 marks, transported by rail to Voelklingen station, fully installed and assembled and including one week of test operation". The gas engine was a twin blaster: two units were linked by a flywheel. This was latter put to a good use: when an irreparable damage occurred on a part of the twins in 1968, this part served as a "spare parts warehouse". Another three gas engines were ordered from Thyssen AG in 1906, and the blast capacity of the Ironworks was dramatically increased with the commissioning of the engines in 1908. At that time, there was still no crane available for installing the colossus, and Hermann Roehling, the owner of the Ironworks, included the instruction that "the heaviest parts are to be transported using our available tools" in the contract. Those machines were producing blast until 80`s, when the Ironworks were shut down.

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Figure 2. Blasting engine: a front look. The rest of the engines were acquired until early 40`s. They could produce either electricity or air blast as required. When the blast furnaces needed less air blast, the machines were used to generate electricity for the Ironworks own power grid. To do this, the gas engines had to be driven at a higher speed, and the operators then received a wage bonus.

Figure 3. A view on the blasting hall.

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These ten blast engines produced up to 110000 cubic meters of blast air per hour for each blast furnace in the Steelworks. The cold airflow was forced into the blast air heaters along six pipelines. Hot stones in the air heaters heated the air up to 1100°C. The hot air entered the smelting zone of the furnace through 16 blast openings. This raised the temperature of the glowing mass up to 2000°C. The oxygen blown in combined with the carbon from the coke and left the furnace flue as blast furnace gas. The blast furnace gas was then cleaned of dust and ash and taken back to the blasting hall, where it was used to drive the blasting engines. The circulation loop was closed.

Working in the blasting hall The blasting hall had a 12-men shift. Each blast engine has its own operator. The flywheels rotated and dispersed an uninterrupted fine oil spray into the hall which was inhaled by the operators while they worked. In addition to this came the noise of the blast engines, monotonous rhythm of the engines and hum of the flywheels as they rotated.

Figure 4. A blast engine flywheel. The men worked extra shifts when malfunctions occurred: heavy machine parts and outsized tools had to be moved. There were no fixed break times and operators generally ate the food they brought with them towards the middle of the shift. They also kept their eyes on the machine while they ate: there was a simple table and chair next to the each blast engine. Only from the mid 1970`s, a break room offered protection from noise and oil.

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The operators in the blasting hall worked for a long time worked a three-shift system. This was a constant round of early, day and night shifts, 56 hours a week without a single day off. When the pattern changed a "long shift" had to be done, working through from midday on Sunday to Monday morning. In 1960 a four-shift system was introduced. After working seven days a worker had a couple days free-time for himself and his family. Starting up and shutting down, oiling and monitoring the blast was the daily routine of the machine operators. But each shift has its own special tasks. The early shift was responsible for cleaning the machines. All parts were cleaned with a mixture of oil and petroleum. The cleaning of the cellar and the maintenance of spare parts was the task of day shift. The night shift ended by scrubbing the hall floor with potassium soap.

Figure 5. Blast engine hall: working place of the operators.

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