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METALS

[Adopsi dari: Zbigniew D Jastrzebski, “The Nature And Properties of Engineering Materials”, John Wiley & Sons, ISBN 0-471-63693-2, 1987, CHAPTER 8.]

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FIGURE 8-2 Macrostructure of solidification in ingot.

Segregation. During solidification of the alloy, variations in concentration occur in certain regions of the melt, resulting in segregation. Two main types of segregation can be distinguished: macrosegregation and microsegregation. In macrosegregation, composition changes in the melt occur over large distances. This may be caused by gravity segregation, normal segregation, and inverse segregation. Gravity segregation is caused by differences in densities of the components of the melt. In normal segregation, the solute is rejected at an advancing solid—liquid interface because of different diffusion rates in solid and in liquid. Inverse segregation is caused by outward move ment of the impurity-enriched interdendritic liquid, thereby increasing the average impurity Content at the head of the ingot.

Microsegregation refers to composition variation over very small distances, usually one or several millimeters. Foreign particles insoluble in both solid and liquid, such as particles of slag and furnace refractories remaining from extractive metallurgical processes, may be trapped during solidification, resulting in segregation at the grain boundaries. As the result of the composition changes in the liquid during freezing, constitutional supercooling may develop during crystal growth. This can occur both in stirred and in unstirred melts, causing microsegregation effects that run in the growth direction on a scale of about 100 µm. To reduce or eliminate constitutional supercooling and the associated microsegregation effects proper control of temperature gradient, cooling rates, and degree of mixing must be observed.

A steep temperature gradient. a slow freezing rate, and a high degree of mixing should be maintained if constitutional supercooling is to be eliminated.

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Shrinkage. Most commercially important metals and alloys undergo substantial contraction in volume when cooled from the liquid to the solid state. The resultant shrinkage is equal to the difference in volume between the liquid at the pouring temperature and the solid metal at room temperature. The total change in volume is the combined effect of the contraction of the liquid when cooled from its casting temperature to its freezing point, the contraction during freezing from liquid to solid, and the contraction of the solid metal when cooled from its freezing point to room temperature (Fig. 8-3).

Shrinkage may cause the formation of cavities in a casting, adversely affecting the properties of the metal. Such cavities, particularly in the form of piping, may cause considerable damage to the metal since, if air penetrates into the cavity, it will oxidize the metal (Fig. 8-4).

The presence of oxide layers inside the ingot will prevent the welding out of cavities during subsequent rolling or forging operations of the ingot. To minimize piping, molds with tapered walls of the big end-up type are used in conjunction with a hot-top or insulated riser as shown in Fig. 8-1. Solidification of the metal at the top is delayed, permitting the liquid metal to penetrate into any cavities and pores that may have already formed in frozen parts of the ingot.

Gas Porosity. Many metals dissolve considerable amounts of gases in their liquid state. The solubility of these gases is much less in the solid than in the liquid metal and decreases as temperature decreases, the exact opposite of the solubility of gases in aqueous solutions. Consequently, the gases will escape from the melt during solidification, forming blowholes and pinholes in the metal. This results in undesirable porosity that may considerably impair the properties of the solid metal or alloy. Pinholes are very small blowholes formed mostly on the surface of the casting; inside the casting blowholes are much larger. Blowholes are not as serious in ingots as in castings because, for the most part, they will be eliminated from the ingot on subsequent rolling and forging. Since the presence of blowholes is concealed from direct observation, castings are less reliable in important applications than wrought metals unless thoroughly examined for porosity.

The number of blowholes in the solid metal depends on the melting and casting technique as well as on the affinity of the molten metal for the dissolved gases.

To remove or reduce the gases from the molten metal, vacuum degassing Is used in steel-making operations. This can be accomplished by stream degassing, during which the molten steel streams through the nozzle into a large vacuum chamber, where it loses gas. The amount of dissolved gas can also be greatly reduced by chemical action such as the removal of oxygen as carbon monoxide by carbon deoxidation or by using deoxidizers such as aluminum, manganese, or silicon, which react with the oxygen in steel forming

FIGURE 8-3 Contraction on solidifying and cooling metals and alloys.

Sp

eci

fic v

olu

me

→

So

lidu

s

Liqu

idu

s M

eltin

g p

oin

t

Pure metal

Alloy with freezing range

Temperature →

FIGURE 8-4 Ingot metal showing piping. Big end down without hot top.

Pipe

Centerline shrinkage

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oxides. This latter method has the disadvantage of leaving in the metal small particles of oxidized products such as alumina or silica.

Impurities and Inclusions. Impurities in the form of oxides, sulfides, and silicates may be present in metals either in solid solution or as separate particles called inclusions. Such impurities generally impair the properties of the metal or alloy but, in some cases, they may actually prove advantageous. The presence of sulfur in steel, lead in brass or steel, and tellurium in copper has beneficial effects on the machining properties of these metals.

8-2 RAPID SOLIDIFICATION PROCESSES

Rapid solidification processes (RSP) involve high cooling rates ranging from 102 to 1010 K s−1 (degrees Kelvin per second). Such processes have been used to produce very fine-grained structural metals, metallic glasses, and very fine reactive powders for a wide range of applications. Such high cooling rates can only be achieved by spreading a thin layer of liquid in a perfect contact with a highly conductive substrate such as metal or sapphire (Fig. 8-5). The rate of quenching is determined by the liquid—substrate heat transfer and the thermal conductivity of the liquid layer. For cooling rates of 102, 106, and 1010 K s−1 the dimensions of a section cannot exceed 10mm, 0.1 mm, and 10 µm, respectively. In conventional solidification processes, cooling rates are about 10−1 to 10−2 K s−1. The lowest cooling rates of about 10−5 to 10−6 K s−1 are present in very large sand castings.

Amorphous or glassy solids can also be produced by methods in which the liquid state is bypassed completely. These may involve the deposition of solid from the vapor phase by thermal evaporation, decomposition of gaseous compounds by radiofrequency discharges, sputtering, or deposition from salt solution by electrolysis. These techniques provide a very high and effective quenching rate and give amorphous solids and metallic glasses that cannot be formed by liquid quenching. The consequence of high cooling rates is a supercooling of the melt by hundreds of degrees before any solid phase can form. This may cause a number of unusual structural and morphological effects such as refinement of dendrites and grain size, extension of solid solubility limits, reduction of segregation, formation of metastable crystalline phases and finally complete noncrystallization (metallic glasses).

FIGURE 8-5 The flow patterns of impinging melt during solidification.

(Annual Review of Materials Science, Vol 10, p. 380, Fig. 6b)

The extension of solid solubility limits of alloying elements results in highly supersaturated solid solutions which upon subsequent annealing decompose, producing a large concentration of uniform and finely dispersed particles. This subsequently increases the yield stress and ultimate tensile strength (UTS) of the alloy because of the increased dislocation−dispersoid interaction. In addition to very fine grain structures one can expect an increase in vacancy concentrations and dislocation density, incomplete ordering, and thermal stresses. The grain size of course depends on the quenching rate which may vary from 106 to 109 K s−1, giving grains with diameters from 100 to 0.1 µm, respectively. Pure metals tend to produce coarser grains of about 1—3 µm in diameter.

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Metallic Glasses. Metallic glasses are produced at relatively high cooling rates, 105 K s−1 and higher, which can be attained by one of the RSR processes such as split quenching, ≈109 K s−1 melt; casting, 106 K s−1; and water quenching, ≈ 103 K s−1 (Fig. 8-5). Metallic glasses of commercial interest are metal—metalloid alloys combining various amounts of iron, nickel, and cobalt with about 15% to 30% of such elements as boron, silicon, carbon, and phosphorus. The glass-forming ability (GFA) of metals and alloys is enhanced in the eutectic region of the phase equilibrium diagram at which the liquidus temperature T1 reaches a minimum and the interval between the liquid temperature T, and the glass transition temperature Tg is the shortest. The ratio Tg/T1, called the reduced transition temperature Trg, can be used as a measure of the glass-forming ability of an alloy. The higher the Trg the easier it is to form metallic glass. The ability to form the metallic glass is also dependent on quenching conditions and on the alloy composition. (See Fig. 8-6.)

For pure metals Trg is only 0.25 requiring cooling rates of about 1012 s−1, making it difficult to achieve glass by the present available quenching technique. However, for transition binary metal alloys Trg is about 0.5, making it easier to produce glasses. A still lower Trg, about 0.66, is shown for a ternary alloy. Generally the stability of metallic glasses increases upon alloying as has been found in such systems as Fe80P17C3, (Fe, Co, Ni, Mo, Cr)75P16B6Al 3 Fe80B17Si3, and (Zr, Ti)60Be59.

Metallic glasses are very hard and strong. Crystalline alloys seldom attain a strength greater than 1% of their theoretical strength because of the presence of dislocations, grain boundaries, and other defects which lowers the strength. The disordered atomic structure of metallic glasses prevents the existence of such defects resulting in strong and hard materials. The chemical homogeneity of the single phase nature of metallic glasses as compared to their crystalline products gives them a much higher corrosion resistance. For example, the glassy alloys containing chromium and phosphorus show much better corrosion resistance than crystalline 18Cr—8Ni stainless steel. Similarly rapidly quenched aluminum alloys, to which Cr, Mn, and Mg have been added, exhibit superior corrosion resistance to seawater, exceeding greatly the conventional resistance of aluminum (Mn and Mg) alloys. Metallic glasses can find a number of important technological applications such as high-strength fibers in composites, catalysts, protective coatings and others. Because of their high electrical resistivity and the absence of grain boundaries, iron-based metallic glasses are excellent magnetic materials for transformer cores, various electronic devices, magnetic tape heads of high permeability and saturation magnetism, with high wear resistance and small magnetostriction.

FIGURE 8-6 Time—temperature transformation (TTT) curves of metals with various reduced glass temperatures Trg = Tg/Tm. (Annual Review of

Materials Science, Vol.10, p. 372, Fig. 3.)

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FORMING PROCESSES The ingots or billets produced during solidification are fabricated by plastic deformation in order to obtain the desired shape and configuration of the final products. Castings are used without further deformation, and they retain the shape and properties resulting from the casting procedure.

Although practically any metal or alloy can be cast, there are limited materials of adequate castability. The greatest difficulty in production of a serviceable casting arises from shrinkage of metals during cooling from the liquid to the solid state at room temperature. The best casting alloys are those that have low shrinkage and short freezing ranges, so that they solidify by forming a solid shell at the mold walls, the thickness of which increases steadily with time. For sound castings favorable temperature gradients must be established in the freezing metal. Freezing must start at the farthest point within the feeding range of a particular riser and continue directionally toward the riser without the freezing of any intervening section before feeding is complete. This requires a proper design for the mold and adequate control of the fluidity of the molten metal, as well as uniform temperature gradients. To meet these requirements, various casting methods are being used according to the type of metal or alloy to be cast and the properties desired in the casting. The methods used are sand casting, die casting, permanent mold casting, centrifugal casting, investment casting, and continuous casting.

The continuous casting process has been used recently for steel up to 6.5x103 mm2 in section. The liquid steel is poured into a vertical water-cooled mold from which it is gradually withdrawn as a solid below the die. A variable-speed roll draws the metal at a rate adjustable up to about 1 m/s. The casting speed is controlled so that a desirable grain structure, usually columnar, is produced. The rate of crystal growth depends on the casting speed and on the angle between casting and the growth direction. The Continuous casting processes are also being developed for other metals and alloys as. for example, horizontal casting of copper alloys.

8-3 MECHANICAL WORKING

The products manufactured by plastic or mechanical deformation are referred to as wrought metals and alloys. Subjecting metals to mechanical working during shaping operations improves the properties of the wrought products over that of the cast product. Plastic deformation, when carried out below the recrystallization temperature, results in work hardening, which improves the mechanical properties of metals and also contributes to the development of a microstructure tougher or more desirable than that formed during casting operations. This mechanical working of metals can be accomplished by various processes such as forging, rolling, extrusion, stamping, drawing, and pressing.

In all these processes the flow of metal is caused by application of the external force or pressure necessary to push or pull a piece of metal through a die. The pressure required to produce such plastic flow is determined largely by the yield stress of the material which, in turn, controls the load capacity of the machinery required to accomplish desirable changes in shape. The pressure used to overcome the inherent material resistance during homogeneous plastic deformation can then be given as

εσ yP ==== (8.1)

where ε = the strain resulting during the deformation

yσ = the yield stress of the material

It is convenient to use the value of the natural strain

2

1lnA

A====ε (8-2)

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where A1 and A2 are the initial and final cross sections of the workpiece, respectively. Thus

2

1lnA

AP yσ==== (8-3)

The pressure so calculated is generally much lower than that required under actual working operations. This is because in addition to the resistance during homogeneous plastic flow, there are losses caused by friction between the workpiece and the die and internal losses caused by inhomogeneous plastic deformation and other factors, such as work hardening. For example, the value of P as calculated by Equation 8-3 is only 30% to 55% of the actual working pressure required for extrusion, 50% to 70% of the pressure required for wire drawing, and 80% to 90% of the pressure required for strip rolling. The total work W required for the actual forming operation can be estimated as the sum of three terms:

W = Wp + Wf + W1 (8.4)

where Wp = the work used to accomplish homogeneous plastic deformation Wf = the work used to overcome frictional resistance W1 = the work used for internal losses within a metal

Forging. Forging is a major method of shaping very large castings by forces that may be either steadily increased in a press or applied by hammering. Since mainly compressive forces are in operation during forging, tensile necking and fracture are avoided. Several forging methods are available, but all of them may be grouped into two main categories: pressing and hammering. Pressing is the relatively slow deformation of metal to the required shape between mating dies. Pressing is preferred for homogenizing large ingots because the deformation zone extends throughout the cross section. In hammering the more rapid deformation is restricted mainly to the surface region. Drop-forging, hammering, and swaging can be performed with lighter equipment, although a drop-hammer itself may weigh several tons. In swaging a small cylindrical rod is forged by passage through a rotary machine in which it is hammered by converging dies. This is particularly useful for shaping more brittle metals without fracture. Large forgings require very large forces and they are usually worked at elevated temperatures (hot working), even though the equipment required is very heavy. Forging is usually applied to produce objects of irregular shape.

Rolling. Rolling is widely used for both hot working and cold working of many metals and alloys. In rolling, the metal is continuously passed between rotating rollers (Fig. 8-7). The rolling process is much more economical than forging, since it is quicker, consumes less power, and produces long items of uniform cross section.

Deformations are restricted to a small volume at any given time, so the loads can be relatively low. The ingot is drawn continuously through the rolls by the friction force between the surfaces of the rolls and the metal. The metal is subjected to a longitudinal tension force that combined with the normal force produces the shear stress τ, which causes deformation.

pf ××××====τ (8-5)

where p = the normal stress f = the coefficient of friction

FIGURE 8-7 Rolling.

d1 d2

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In cold rolling of wide strips it is possible to attain speeds up to 25 m/s for thin strips. When thin hard strips are rolled, it is not possible to reduce the thickness below a certain limit because greater deformation of the rolls than of the strips can occur. To obtain the thinnest strip possible, the metal should be annealed, the coefficient of friction should be as low as possible, and small-diameter rolls should be used.

Extrusion. Extrusion involves pressing a billet or slug of metal by movement of a ram through an orifice or nozzle (Fig. 8-8a). During extrusion, both compressive and shear forces are developed; however, there are no tensile forces present. It is therefore possible to deform the metal very heavily without fracture. Under the influence of the applied force, the metal is continuously deformed as it is pushed through a die and is shaped into a long bar of desired cross section. In the case of tubes, the extrusion is carried out through a round die. and a mandrel, integral with the ram, is pushed through the previously pierced billet. Two main extrusion processes are direct extrusion and inverted extrusion. In direct extrusion, the pressure ram is advanced toward the die assembly; in indirect extrusion, the die is moved down the container bore. This latter movement is characterized by a simpler flow and lower extrusion pressures; however, the operation must employ a hollow ram along which the extruded product can pass after being formed in the die (Fig. 8-8b). Consequently, in indirect extrusion, frictional losses are much smaller, since there is no relative movement of billet and container and the deformation is confined only to a zone near the die. The efficiency of the extrusion is measured by the ratio of the work necessary for deformation to the total work expended. This latter is the sum of the work necessary for deformation, for overcoming frictional resistance at the billet—container and the billet—die interfaces, and the work expended in straining not required for shape change. Initially, extrusion was mainly used for lower-melting, nonferrous metals but, with improvements in lubricants and development of powerful presses, the extrusion of higher-melting metals such as steel becomes more and more important. Most carbon and stainless steels can now be extruded at temperatures of about 1200°C (2190°F) using glass as a lubricant. Glass is cheap, is relatively easy to use, is stable under most extrusion conditions, does not react with iron, and helps to dissolve any oxide layer on the metal surface. Although extrusion was first limited to hot working operations, recent development of powerful presses and high-pressure lubricants makes cold extrusion more and more important in metal-working operations. Modern extrusion processes can now be run automatically, making complex sections, particularly in metals such as brass and aluminum. The mechanical properties of cold extruded products are very good.

FIGURE 8-8 (a) Direct extrusion of tubes. (b) Inverse extrusion

Container

Container

Billet

Tube

(a)

(b)

Container

Billet

Billet

Ram

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Other Forming Processes. Wire drawing, deep drawing, and swaging are forming processes used for a variety of applications. Wire drawing involves pulling the stock through a tungsten carbide die having a tapered bore. Large quantities of rods, tubes, and wires are produced by this technique. Usually the reduction of the rod area is limited to about 30% so that, for a fine wire, the process consists of a long series of successive wire drawing operations through subsequent dies of suitable diameters. Deep drawing is a forming process used to shape metals into bowls, cups, panels, and similar shapes. In deep drawing the metal is deformed by flowing rapidly inward and then turning the corner to become the wall of the cup. The amount of drawing in a single operation should be such that the ratio of height to diameter of the article formed does not exceed one. If this ratio exceeds one the article must be made by redrawing in one or more steps. Recently, high-pressure and very high speed processes of formation have been introduced. Under very high pressure the ductility of metal considerably increases, permitting the combination of otherwise separate operations. Explosive forming involves the application of very high pressures ranging from 700 MPa (101.5 ksi) to over 7 GPa (1015 ksi) over a very short period of time (microseconds). In most industrial processes this is accomplished by producing an intense pressure pulse in a liquid medium caused by detonation of an explosive or explosion of a wire by a sudden passage of very high currents. The shock wave generated in liquid forces a metal blank against a shaping die, causing the required deformation. The pressure wave produces uniaxial compression involving shear and volume changes. The strain rates are about 105 s−1 and sometimes as high as 108 s−1.

8-4 SUPERPLASTICITY

Certain alloys when deformed in tension at elevated temperatures exhibit extensive plastic deformation showing a strain which may exceed 1000% and more without necking. This phenomenon, called superplasticity, presents a new way of fabricating certain metals and alloys by methods so far used only for glass and plastics. However, the extensive use of superplasticity in forming processes has been restricted by the fact that the strain rates in the region of superplastic deformation are much lower than that used in conventional hot working processes. Considerable efforts are made to determine all the conditions leading to extended superplasticity over higher ranges of strain rates. Superplasticity is a high-temperature phenomenon occurring at about half of the absolute melting point (0.5 Tm) of the material that has characteristic microstructure consisting of a uniform ultrafine equiaxed grains smaller than 10 µm and stable during the superplastic deformation.

The flow stress causing superplastic strain is highly sensitive to the strain rate, ε& , and can be described by equation

mK εσ & ==== (8-6) where σ = the flow stress

ε& = the strain rate K = the material constant that depends on temperature and microstructure m = the strain sensitivity factor

On plotting log stress versus log strain rate (Fig. 8-9a) a characteristic sigmoidal curve is obtained. The slope of the curve at any particular strain rate is given by

εδσδ&log

log====m (8.7)

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FIGURE 8-9 Superplastic behavior of an alloy. (a) A plot of log stress versus log strain rate, ε& . (b) Changes in the strain sensitivity index m with strain rate

ε& . (After D. Lee, Acta Metallurgica, 17, 1057. 1969.)

The value of m changes with the strain rate (see Fig. 8-9b) and in regions I and III, m < 0.3. In superplastic region II, where large elongation without necking occurs, m > 0.3, reaches its maximum of about 0.6 up to 1.0, and then drops to its initial values less than 0.3 for regions I and III. The strain rate sensitivity factor m is a measure of the necking resistance of the material during superplastic deformation.

(a) (b)

FIGURE 8-10 Microstructures of zinc—aluminum alloy (78% Zn—22% Al) at 800 X. (a) Alloy in a superplastic state showing very fine duplex microstructure produced from a single-phase alloy on quenching from above 275°C in cold water. (b) Alloy

in conventional metal state (not superplastic) showing coarse microstructure resulting on slow cooling from above 275°C. (The Western Electric Engineer. Vol.

XV. No. 1. January 1971).

The maximum ductility depends on the imposed strain rate, the temperature, and the initial grain size. The highest ductility as measured by elongation is usually found at an intermediate strain rate, and it decreases in both higher and lower strain rate regions. Increasing the deformation temperature reduces the flow stress and causes higher strain rates between regions I and Il and between regions II and Ill. A proper heat treatment of a superplastically formed component changes its microstructure from very fine equiaxed grains to coarser ones and the alloy ceases to be superplastic (Fig. 8-10).

Strain rate, s−1

Str

ain

sen

sitiv

ity in

dex

m

Str

ess,

MP

a II I

(a)

(b)

III

| 10−1

104 −

103 −

102 −

10 −

1 −

0.9 −

0.6 −

0.3 −

| 10−7

| 10−6

| 10−5

| 10−2

| 10−4

| 10−3

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Table 8.1. Some selected superplatic alloys

Material Maximum m

Maximum elongation,

%

Test Temperature

oC

Al-17% Cu 0.7 600 400

Al-10.7% Zn-0.93% Mg-0.4% Zr

0.9 1550 550

Ti-6% Al-4% V 0.85 1000 750 - 1000

Ni-26.2% Fe-34.9% Cr-0.58% Ti 0.5 1000 795 - 855

Fe-0.1% C-26% Cr-6,5% Ni 0.62 1000 700 - 1020

Sn-Pb eutectic 0.59 1080 20

Mechanism of Superplastic Deformation. The microstructure of superplastic alloys shows that there are no essentially large changes in the shape and size of grains and in the microstructure during superplastic deformation. The same ultrafine equiaxed grain structure that was present prior to deformation has been retained. The ability to maintain an equiaxed grain structure depends on the interface mobility and the extent of grain boundary sliding. The most important role in superplastic deformation is the effect of boundary sliding and grain boundary migration. During superplastic deformation considerable changes in the texture of an alloy occur that are the result of the cumulative effects of slip, extensive grain boundary sliding, dislocation motion, diffusional processes, and recrystallization. Enhanced dislocation climb caused by a supersaturation of vacancies of interstitials produced by the α−β phase changes contribute to superplastic deformation. Since interfacial sliding is a major cause of the superplastic deformation it may give rise to cavitation.

Cavitation is caused by the localization of flow along the grain and inter-phase boundaries by the process of sliding. Such sliding can be blocked at various irregularities but most effectively by hard particles that act as stress concentration elevators. This may lead to the formation of cracks which can produce small intergranular voids during superplastic stretch forming. Cavitation may limit the superplastic ductility of the material as well as reduce the final mechanical properties of components formed by superplastic forming operations. Cavitation arises under the combined action of stress and grain boundary slidings. Thermodynamic stability of a cavity occurs where the radius of void is given by

σγ2====r (8-8)

where σ = the applied stress γ = the surface energy

The value of r required to stabilize the size of the voids is about 10 nm (100 Å).

A number of alloys demonstrating extensive superplastic formability such as Cu—Zn−Ni, α/β brasses, aluminum bronze. Al—Zn—Mg−Zr, Al—Cu—Zr, Mg—Zn—0.5Zr, and others show a tendency to cavitation during superplastic stretching of the material. Superplastic ductilities and local fracture strains can be increased under hydrostatic pressure. By applying hydrostatic pressure during superplastic stretching, hígh-temperature creep cavitation is reduced and any adverse cavitation effects on subsequent room temperature properties of the superplastically formed component can

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be prevented. A critical ratio of hydrostatic pressure P to flow stresses cσ , cP σ/ , can be determined for superplastic forming conditions above which no cavitation occurs. For example, in the 7475 aluminum alloy, with superplastic forming parameters T = 516°C (961°F) and strain rate ε& = 2 X 10−4 s−1 , cavitation may be prevented when the critical ratio is cP σ/ > 0.6. Superplastic forming was accomplished by applying gas pressure on one side of the aluminum sheet causing it to form down into the die at a nearly constant rate.

8-5 JOINING METALS

Materials may be joined together by mechanical means, welding, brazing, or soldering.

Mechanical joining methods consist of riveting, bolting, or screwing; they are limited to materials, mainly metals and wood, that can be drilled or punched without danger of cracking. Mechanical methods of joining have many disadvantages, such as the need for drilling holes that lower the inherent strength of the material and an increased tendency to crevice corrosion. For these reasons metals are joined whenever possible by welding, brazing, or soldering, which will now be considered briefly.

Welding is the intimate union of two parts of a material by heating them until a molten or plastic state is reached, with or without the application of mechanical pressure. Welding is universally applied for joining parts of the same metal; it is sometimes used for joining dissimilar metals. All metal welding processes can be divided into pressure welding and fusion welding methods. Nearly all of them require heating and only a few commercially useful processes known as cold welding involve pressure only. Heat for welding may be generated by electrical, chemical, and mechanical means and, accordingly, numerous welding processes are classified this way. The fundamental welding problem is simply the removal of the adsorbed gas and oxide layers from the surfaces to be joined in order to get an intimate contact between the surfaces of the mating parts. New welding processes are being developed constantly for specific industrial applications to obtain improved weld joint properties, better appearance, increased speed of welding, and lowered cost.

Pressure Welding. Pressure welding, frequently referred to as solid phase welding, is essentially a process of joining metal specimens that have been usually heated, within the joining area, to a highly plastic state, so as to form a complete union after hammering or pressure is applied. The combined action of heating and pressure causes extensive plastic deformation of the metal surfaces. This leads to the breaking or cracking of the oxide films present on the metal surfaces, thereby bringing the metal parts into intimate contact. The diffusion of atoms from one metal surface to another is thus facilitated; furthermore, the bond between them may be as strong as the original metal itself. It is important in all cases that the two surfaces to be joined should be clean and free from any external impurities.

The most widely used pressure welding processes are forge welding and resistance welding. A simple type of forge welding can be illustrated by the familiar blacksmith method of joining metals by hammering them at a red heat, FIGURE 8-11 Resistance spot welding.

Pressure Pressure

Electrode

Workpiece

Electrode Spot weld

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but now steady pressure exerted by special machinery is generally used. Resistance welding involves heating the metal surfaces to be joined by passing an electric current across them. This is accompanied by mechanical pressure, which results in the intimate union of the two parts (Fig. 8-11).

If the pressure welding is carried out at low temperatures it is known as cold welding. This requires very extensive deformation and very high pressures to break up the oxide films. However, it does not involve much transfer of matter across the interface. The bond results from the interaction of the atomic forces of the metals. New developments in solid phase welding are friction welding, high-energy welding (explosive welding), ultrasonic welding, and diffusion bonding.

Fusion Welding. Fusion welding is accomplished by the fusion of metal to effect the desired union between two pieces. To produce a strong bond, the liquid metal must first wet the surface of the solid metal. Good wetting is favored if the liquid metal and the solid metal show some mutual solubility or form the inter-metallic compounds. The fused metal generally is obtained as a deposit from an external source, called the welding rod or filler rod although, in some cases, the surfaces are merely fused together alone to form the joint. Fusion welding can be accomplished by a variety of methods selected according to the metallurgical and chemical characteristics of the metal to be joined, the required properties of the weld, and the specific conditions of the workpiece. These methods consist of gas welding and electrical welding. In gas welding an oxyacetylene flame is occasionally employed. The oxyacetylene flame may have a neutral, oxidizing, or reducing character, depending on the proportions of acetylene and air, or oxygen, used. This, in turn, depends on the type of metal to be welded but, in general, a neutral flame is most widely used. An extra metal is added from the welding rod, but no fluxes are used.

Electric-arc welding methods include carbon-arc welding, atomic hydrogen welding, inert-arc welding, and multiarc welding. An electric arc is struck between the carbon or metal electrode and the work material to be welded. The intense heat of the arc causes the fusion of the base metal, as well as of the metal filler rod, in case an additional amount of the metal is needed for the weld. Practically all welding with the metallic arc is done with a shielded arc in which the metallic electrode is covered with a thick layer of fluxing ingredients. The shielding effect is due to the fluxing ingredients, which melt and form a slag covering over the weld and thus are able to protect the molten metal from the air (Fig. 8-12).

For reactive metals inert-arc, multiarc, and atomic hydrogen welding are used. In inert-arc welding (heliarc) either a tungsten or a carbon electrode is used. The inert gas, helium or argon, is fed through a nozzle which surrounds the electrode in the head of the torch and completely protects the arc and also the pool of molten metal from contact with the atmosphere. In atomic hydrogen welding a stream of hydrogen is passed through the arc struck between two tungsten electrodes. The intense heat dissociates molecular hydrogen into atomic hydrogen, which recombines on the metal surface to be welded with the evolution of the same amount of heat as that absorbed during dissociation. This method is used extensively for aluminum welding to prevent its oxidation.

FIGURE 8-12 Arc welding with shielded electrode. Many electrode coatings contain cellulose which, on

decomposition, yields gases, forming a shield.

Coated electrode

AC or DC power

Gaseous shield

Arc stream Slag

Molten pool

Extruded coating

Base metal Weld deposit

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Plasma-arc welding is a further development of the inert-arc welding in which a gas is ionized when passing through the electric arc and impinges on the surface to be welded. Very high temperatures are developed, resulting in the speedy welding of two parts.

A recent development in fusion welding is the employment of radiation energy, as in laser and electron beam welding. The essential feature of these processes is that the radiant energy is narrowly focused on the workpiece, and heat is generated only where the focused beam is intercepted. This reduces the possibility of distortion and results in improved weld properties. Using a laser radiation beam, it is then possible to drill very fine holes in diamonds or in other very hard materials. Highly reflective surfaces are not easily welded. In electron beam welding the energy is supplied by the impact of a focused beam of electrons. Operating in a vacuum is an additional advantage with welding metals that are sensitive to contamination from the atmosphere.

Brazing and Soldering. Brazing and soldering are metal-joining processes which use another metal or alloy that has a melting point substantially lower than that of the metal to be joined. In brazing, copper or silver alloys are widely used; in soldering, low-melting alloys of tin and lead are used. Brazing is done at temperatures that are above 535°C but are always lower than the melting point of the metal to be joined. Soldering differs from brazing in that the introduced molten metal has a softening point much lower than that of the brazing alloy. Solders are usually lead—tin alloys of varying composition, but compositions of 50% lead—tin or 60% to 40% lead—tin are most widely employed.

The essential requirements for obtaining good joints in brazing or soldering are proper joint design, preparation of the joint surfaces, fluxing, assembling, heating, and final cleaning. Since there is no melting of the parent metal, the molten brazing alloy either flows between heated surfaces of the joint members or is melted in place between such surfaces. Cleaned surfaces are essential to bring about adhesion of the brazed weld to the parent metal. This adhesion is obtained by using fluxes that dissolve the oxide film on the metal surface and at the same time prevent the oxidation of the metal during preheating. Fluxes should have a melting point below that of the brazing alloy for easy wetting and cleaning of the surfaces before the molten brazing alloy is applied. Fluxes are either acids or compounds that decompose into acidic substances that corrode the metal to be joined. Brazing fluxes are usually mixtures of borax with boric acid or borax with zinc chloride.

8-6 POWDER METALLURGY

Powder metallurgy is the process of producing metallic powders and converting them, sometimes in conjunction with nonmetallic powders, into ingots or finished articles by compacting and sintering. Powder metallurgy is now used extensively whenever porous parts are needed, whenever the parts have intricate shapes that can be easily processed and sintered, whenever the alloy or mixture of metals cannot be made in any other way, or whenever the metals have very high melting points and are therefore difficult to cast and machine. Nearly all commercially used metals can be shaped by the powder technique, which is ideally suited to the preparation of components with a controlled amount of porosity distributed in a specified manner. An organic binder is frequently added during the mixing stage but volatilizes at a low temperature, facilitating the formation of interconnected pores.

Manufacture of Metal Powders. Since metals are not brittle enough to be powdered easily by crushing and grinding, metallic powders must be produced by special methods. These methods involve the reduction of their corresponding oxides and salts, the thermal dissociation of metal compounds, electrolysis, atomization, and the crushing of certain hard metal hydrides.

Lecture Note

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The atomization method is now the major method of manufacture because it can make powders from alloys as well as from pure metals. The method depends on forcing a molten metal through an orifice and then breaking up this stream with a jet of water or gas. The metal particles formed by the jet solidify very rapidly, so that even in air only very superficial oxidation takes place, provided the metal is not too reactive.

Substantial amounts of fine powders quenched at cooling rate of 104 K s−1 or higher can be produced by plasma spraying and centrifugal atomization in helium and also from finely chopped melt spun ribbons. Such powders are characterized by high reactivity when hot pressed and sintered with an other powder forming consolidated products of much higher yield and ultimate tensile strength than that of alloys obtained by more conventional methods. For example, the development of aluminum alloys such as Al—Zn—Mg—Cu through consolidation of atomized pre-alloyed powder produced the material with yield strength as high as 850 MPa (124 ksi) attained after heat treatment. This is considerably higher than the yield strengths of conventional high-strength alloys such as 7001 and 7005 aluminum alloys wrought from ingot. Mechanical alloying has been used to produce complex oxide dispersion strengthened (OD-s) alloys that are difficult or impossible to manufacture by conventional melting and casting technique. By mechanical mixing, ultrafine powders can be blended with fine refractory oxides to form an OD-s alloy. On further attrition in a specific solvent, for example, isopropyl alcohol, further diminution of the particles occurs and the reactivity of the mixture is enhanced.

Compacting and Sinlerlng. The shaping of articles from metallic powders is accomplished by compacting and sintering, the principles of which were discussed in Chapter 6, Section 6-7. The finer the particles, the greater will be the strength of the compact. Usually polydisperse powders of the proper size range give stronger compact and higher sintering rates than monodisperse powders of the same particle size. Powder metals are often compacted in mechanical or hydraulic presses. where a vertical stroke exerts a force to consolidate the loose powder into a shaped green compact. The pressure between two rotating rolls can also be utilized for powder compaction (Fig. 8-13).

Hot pressing has been widely used for refractory-type metals and ceramics with high melting points and high hardness and strength. Using the pressure in conjunction with high temperature leads to high densities of the sintered product. Hot pressing is usually carried out in inert graphite or ceramic molds to avoid contamination of the powder mass.

Structures and properties of the sintered materials are determined by residual porosity and permeability. Grain size is second in importance to porosity with respect to uniqueness of structure. The strength of the sintered powder compact is given by

kPp e−−−−==== 0ασ (8-9)

where pσ = the strength of the porous body

0α = the strength of the homogeneous material k = the coefficient dependent on the shape of the pores P = the porosity

The variety of applications of powder metallurgy ranges from the most exotic innovations (e.g., in rockets) to parts in our washing machines and automobiles (Fig. 8-13). Ferrous materials, copper, nickel, aluminum, tita nium. and other metallic powders are widely used with a variety of alloying elements to impart desired characteristics, structures, and shapes. Refractory metals, such as tungsten, molybdenum, tantalum, and niobium, and reactive metals, such as beryllium, are extensively shaped by powder metallurgy technique. For example, tungsten products up to nearly theoretical density can be obtained by powder metallurgy technique. Another example is that of beryllium, which becomes brittle during casting and hot-working operations. All wrought forms are

Lecture Note

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produced from the beryllium powder, which is compacted and sintered at about 1100°C in an atmosphere of argon or nitrogen.

(a) (b)

FIGURE 8-13 Compacting and sintering of Mond (carbonyl) nickel powder, size less than 7 µm. (a) Pressed at 40.000 psi (276 MPa), 68.8% density, 413 x. (b) Pressed at 40.000 psi (276 MPa) and sintered 1 h in hydrogen at 1250°,. 87.5%

density, 413 x. Dark spots are the remaining pores.

FERROUS METALS Ferrous metals cover a wide variety of industrially important products that can be grouped into wrought irons, cast irons, carbon steels, and alloy steels. 8-7 IRON—CARBON EQUILIBRIUM DIAGRAM

Many properties of cast irons and carbon steels, as well as their microstructure, can be explained by reference to the iron—carbon phase equilibrium diagram (Fig. 8-14). Strictly speaking, the diagram refers to the iron—iron carbide system, but the existing phase relationship can be expressed in terms of the percentage carbon content. The diagram is divided into a number of phase fields, each occupied by either a single phase or a mixture of two phases. Curve ABCD is the liquidus line above which there is only one liquid phase consisting of iron and dissolved carbon. Curve AEPGCH is the solidus line below which the various iron-carbon compositions are completely solid. The regions between these two lines represent mixtures of solid and liquid.

Both α iron, having a body-centered cubic structure, and γ iron, with face-centered cubic structure, are capable of dissolving certain amounts of carbon, which vary according to temperature. The solid solution of carbon in α iron is called ferrite, whereas the solid solution of carbon in γ iron is called austenite. The terms α and γ iron are also used to designate ferrite and austenite, respectively. Austenite is capable of dissolving far more carbon than ferrite, and it may contain up to 2.11% carbon (point G), whereas ferrite can only dissolve up to 0.025% carbon (point L). The solubility of carbon changes with the temperature in both austenite and ferrite in the manner indicated by curves GK for austenite and LI for ferrite.

The third solid phase existing in the diagram is cementite, which is a compound of composition Fe3C, corresponding to 6.69% by weight carbon. Cementite is hard and brittle, whereas ferrite is relatively soft and malleable.

The transformation of one solid phase into another, as from austenite to ferrite or to cementite, occurs at certain critical temperatures indicated on the diagram by lines LKM, IK, and GK. These lines are usually referred to as A1, called the lower critical temperature line, and A3 and Acm , called the upper critical temperature lines. These lines represent the transition temperatures for both upward and downward transformations

Lecture Note

17/36

occurring under equilibrium conditions. In practice, however, equilibrium is seldom obtained and the critical temperature lines will be shifted to higher positions on heating and to lower positions on cooling. The degree of shift from the equilibrium conditions will depend on the rate of heating or cooling.

FIGURE 8-14 Carbon—iron phase diagram.

(Adapted from Metals Handbook, 1973 edition, Vol. 8. )

The diagram also indicates the existing ranges of composition for iron, steel, and cast iron. Compositions up to 0.008% carbon are regarded as commercially pure iron, those from 0.008 to 2% carbon represent steel, and those above 2% carbon represent cast iron. Steels are further subdivided into hypoeutectoid steel up to 0.77% carbon and hypereutectoid steel from 0.77% to 2% carbon (see also Fig. 8-15).

Austenite is not stable below lines A3 (IK) and Acm (GK). For compositions containing less than 0.77% carbon, austenite on cooling begins to transform into ferrite, and the remaining austenite increases its carbon content along the line A3 until point K (0.77% carbon) is reached. For compositions between 0.77% and 2.1% carbon, cementite will separate out and the austenite composition will vary along the line Acm (GK) until again point K is reached. At point K, called the eutectoid, austenite will transform into pearlite, which is an intimate mixture of ferrite and cementite. This is called the eutectoid transformation, characterized by the decomposition of a solid solution into two different solid phases according to the reaction

1600

| 6.0

1500 −

1400 −

1300 −

1200 −

1100 −

1000 −

900 −

800 −

700 −

600 −

500 −

400 −

300 −

200 −

100 −

0 − | 0.5

| 1.0

| 1.5

| 2.0

| 2.5

| 3.0

| 3.5

| 4.0

| 4.5

| 5.0

| 5.5

| 6.5 Fe

Weight percent carbon

A B E

1495 1538

F 1394

(γ-Fe) Austenite

G 1154 1148

F

Solubility of graphite in liquid Fe

Cementite (Fe3C) Acm A3

D

2.08

2.11

H

α + Fe3C

738 727 K L

912 I

α-Fe Ferrite

A1

230

(δ-Fe) P

C

4.30 6.69

0.77

Co 770

Lecture Note

18/36

solid phase A → solid phase B + solid phase C (8-10) (austenite) (ferrite) (cementite)

There are three phases, austenite, ferrite, and cementite, existing in equilibrium at the eutectoid point. It follows from the phase rule that the number of degrees of freedom is zero, F = C — P + I = 2 —3 + 1 = 0, and the eutectoid point, like that of the eutectic, is nonvariant. The eutectoid reaction is one of the most important transformations occurring in the iron—carbon system; it is also found in many other alloy systems.

FIGURE 8-15 Microstructures of carbon steels on slow cooling.

At a temperature of 1148°C, the eutectic transformation (point C) takes place. The eutectic liquid with 4.30% carbon begins to freeze to a eutectic mixture of austenite and cementite called ledeburite. On further cooling, the eutectic austenite transforms gradually to cementite, and its composition changes along line GK until again the eutectoid point is reached. Here all the remaining austenite will finally transform to pearlite. Another type of transformation is the peritectic, which occurs at a temperature of 1495°C (point P). The peritectic reaction involves the transformation of a liquid phase and a solid phase into another solid phase of different composition according to the reaction

liquid (0.5% C) + δ iron (0.08% C) → γ iron (0.18% C) (8-11)

The peritectic reaction affects only the solidification of steels with less than 0.55% carbon, and it is of little practical importance.

Consider a steel containing 0.4% carbon represented on the partial carbon—iron diagram (Fig. 8-15) by line xx. When cooling from the liquid region along line xx, freezing begins at the liquidus line. Complete solidification of an alloy occurs at the solidus line AG. As cooling proceeds, no changes occur until line A3 is reached, where the precipitation of ferrite from the solid phase of austemte begins. Further cooling results in an increase in ferrite, whereas the amount of austenite decreases and its composition varies along line IK. The composition of ferrite varies along line IL reaching a maximum carbon content of 0.025% at point L. At any point in region ILK, the percentages of austenite and ferrite can be calculated from the lever rule. For example, at 750°C (point b), austenite contains 0.5% carbon and ferrite contains 0.02% carbon (see the horizontal dashed line b’bb”). Hence the amount of austenite will be

%79%100%48.0

%38.0

%02.0%5.0

%02.0%4.0 ====××××====−−−−−−−−

and that of ferrite will be 100% — 79% = 21%.

(a)

(b)

(c)

(d)

(e)

(f)

Pearlite Pearlite

Cementite

Frrite-α and γ-austenite

Frrite-α

I A3

a X

d

b’ b b” c

γ-Austenite and cementite

Pearlite and cementite

Frrite-α and pearlite

γ-Austenite

L K

c’ c c” f

0.02

X

0.008 0.4 0.5 0.77 1

Lecture Note

19/36

At the eutectoid temperature of 727°C, the remaining austenite undergoes transformation to pearlite, which consists of 88.8% ferrite and 11.2% cementite. The composition of the pearlite can be found by considering the section between points L and M divided by point K as shown in Fig. 8-14.

Further cooling below the eutectoid temperature causes practically no further changes in the microstructure of the steel, since the solubility of carbon in ferrite (line LN) changes relatively very little. At any point below 727°C the microstructure of the steel will consist of a pearlite matrix in which primary crystals of ferrite are embedded. The amounts of the microconstituents at point c (the horizontal dashed line c’cc’’) can be calculated from lever rule,

%27.49%100%019.0%77.0

%40.0%77.0 ferrite ofamount ====××××

−−−−−−−−====

%73.50%100%019.0%77.0

%019.0%40.0 pearlite ofamount ====××××

−−−−−−−−====

where 0.019% is the carbon content of ferrite at point c’.

It can also be seei that a steel of the eutectoid composition (0.77% C) will consist only of pearlite. whereas a steel with a carbon content larger than 0.77% will consist of pearlite and cementite. The microstructures of 0.4% and 1.0% carbon steels at various points of the carbon—iron phase diagram are given in Fig. 8-15 by sketches a—f.

8-8 TIME—TEMPERATURE TRANSFORMATIONS

The iron—carbon equilibrium diagram shows only the phases and the resulting microstructure corresponding to equilibrium conditions. In practice, however, equilibrium is seldom obtained, since the cooling rates are usually much higher than those necessary to maintain equilibrium. As the cooling rate increases, the experimentally observed transformation temperatures are lowered, and metastable (nonequilibrium) phases may be formed. For example, at very high rates of cooling in the steel range a metastable phase called martensite can develop which, of course, has no place in the equilibrium diagram. The presence of certain alloying elements also affects considerably the transformation range of the various phases and may result in a completely new diagram.

The eutectoid reactions are also greatly affected by the rate of cooling. The course of the eutectoid transformation may be represented for each temperature by a series of isothermal curves. These are obtained by plotting the percentage of austenite transformed to pearlite versus time for any specific temperature (Fig. 8-16).

FIGURE 8.16 Isothermal reaction curves (0.77% carbon steel).

(J. Wulff et al. Metallurgy for Engineers, John Wiley & Sons, Inc., New York,1952.)

A series of such isothermal reaction curves can be used to plot a time—temperature transformation (TTT) diagram, as shown in Fig. 8-17.

Au

sten

ite t

ran

sfo

rmat

ion

%

Time, sec

100 95 −

50 −

5 − | 10

| 100 1000

500oC

595oC 705oC

3 2 1

1’

1” 2”

2’ 3’

3”

Lecture Note

20/36

FIGURE 8-17 Temperature-time transformation diagram or Bain (S) curves. (J. Wulff

et.aI. Metallurgy for Engineers, John Wiley & Sons, Inc., New York, 1952.)

The resulting curve is known as the Baine or S curve. It is seen from the S curve that the greatest speed of transformation corresponds to the nose of the knee of the S curve. The rate of transformation decreases below the temperature of the nose, owing to the slowing down of atomic diffusion. whereas above this temperature it decreases because of the increased stability of austenite.

Above the critical temperature A1, austenite (Fig. 8-18a) will be stable and there will be no transformation to pearlite. In addition to the variations in the rate of transformation with temperature, there are variations in the structure of the transformation products. Transformations at temperatures between approximately 727 and 535°C result in the characteristic lamellae microstructure of pearlite (Fig. 8-18b). At a temperature just below A1, nucleation of cementite from austenite will be very slow, but diffusion and growth of nuclei will proceed at maximum speed, so that there will be few large lamellae and the pearlite wIl be coarse. However, as the transformation temperature is lowered, the perlite becomes more and more fine. Transformation at temperatures between 535 and 260°C results in the formation of bainite consisting, like pearlite, of a ferrite matrix in which particles of cementite are embedded. The individual particles are much finer than in pearlite, and bainite cannot be resolved by ordinary microscopic examination. Bainite has an acicular microstracture, and It is harder than pearlite and tougher than martensite.

(a) (b) (c)

FIGURE 8-18 (a) Austenitic microstructure In steel (18% Cr-8% Ni), x250. (b) Pearlite microstructure in 0.77% C steel, X 1300. (c) Martensitic microstructure

in carbon steel x1500.

Martensite (Fig. 8-18c) is regarded as a supersaturated solid solution of carbon in ferrite, but it has a body-centered tetragonal structure with carbon atoms located interstitially. Martensite is formed by the diffusionless transformation of austenite on rapid cooling to a temperature about 260°C, designated as Ms. Since the martensite transformation is accompanied by an increase in volume, the temperature must be lowered in order to

Te

mpe

ratu

re o C

Pearlite 705 − 595 − 535 −

Ms

Mf

Bainite

A1 (727oC)

| 10

| 100

| 1000

| 10000

Time, sec

3 3’ 3” 2 2’ 2”

1 1’ 1”

260 −

Lecture Note

21/36

cause the transformation to proceed to completion. Finally, at a temperature designated as Mf, austenite will be wholly transformed to martensite. The starting temperature of the martensite transformation, Mf, and the finishing temperature, Ms, are represented on the TTT diagram as horizontal lines (Figs. 8-17 and 8-19). The positions of these lines vary somewhat with the carbon content of the steel. The martensite transformation differs from the other transformations in that it is not time dependent and occurs almost instantaneously; the proportion of austenite transformed depends only on the temperature to which it is cooled.

Martensite, as formed in steels, is distinguished by very high hardness, but it also has a tendency to brittleness. The hardness depends on the carbon content; steels containing 0.15% carbon or less do not harden appreciably on quenching. With higher carbon contents above 0.2% the martensite is so hard that the steel is brittle unless properly tempered.

8-9 CONTINUOUS COOLING TRANSFORMATIONS

During heat treatment operations, the steel is continuously cooled from above the critical temperature range (A3) to room temperature. Thus the transformation of austenite does not occur isothermally, as assumed in the TTT diagram, but over a certain period during which the temperature drops from, say, T1 to T2. Thus the average temperature of the transformation (T1 + T3)/2 will be lower during continuous cooling than during isothermal cooling and, consequently, the transformation of austenite will be somewhat delayed. These will cause the S curve to be shifted toward lower temperatures and longer transformation times during continuous cooling as compared to those of the TTT curve (Fig. 8-19).

FIGURE 8-19 Continuous cooling transformation curves for the eutectoid steel (0.77% C). Dashed curves are TTT curves, whereas solid curves shifted to the right and down are continuous cooling curves. Cooling curve B results in the microstructure composed of pearlite and martensite. Curve D corresponds to cooling rates during annealing, whereas curve C corresponds to cooling rates during normalizing. Cooling rates (curve A) higher than the critical cooling rate result in martensitic structure, or in bainite (austempering) structure, or martempering, according to the temperature of quenching.

Co

Lecture Note

22/36

The minimum cooling rate required to produce martensite in a given steel is determined by the position of the nose of the S curve. The cooling rate required to avoid the nose of the S curve is called the critical cooling rate; any cooling rate (curve A) greater than that will result in complete transformation of an austenite to martensite. If the cooling rate is slower than the critical cooling rate, the cooling curve (B) will intersect the pearlite zone. This results in the transformation of some austenite to pearlite, and only the remaining austenite will transform to martensite when the M, temperatures are reached. The microstructure of such a steel will then be composed of fine pearlite and martensite. Cooling curves C and D show that at such cooling rates austenite will transform completely to pearlite before any martensite can be formed.

In transformation of hypoeutectoid (C <0.77%) and hypereutectoid (C > 0.77%) steels, austenite transformation always begins with separation of the proeutectoid phase, either ferrite or cementite, the amount depending on the carbon content of the steel. This results in nucleation of the proeutectoid ferrite in hypoeutectojd steels or the procutectoici cementite in hypereutectoid steels at the grain boundaries of the austenite, as illustrated in Fig. 8-15. On further cooling the amount of ferrite or cementite increases until the eutectoid temperature is reached at which the austenite will have the eutectoid composition of 0.77% C. Any further cooling will result in the transformation of austenite to pearlite, so the microstructure will be the proeutectoid ferrite and pearlite. On quenching from the austenite range, the proeutectoid ferrite may not be able to form completely and austenite of higher C content will change directly to martensite (Fig. 8-20).

FIGURE 8-20 Continuous cooling transformation curves in proeutectoid

steel (C less than 0.77%).

The diagram also shows that bainite cannot be obtained in any appreciable quantity by continuous cooling to room temperature because it is sheltered by an overhanging pearlite nose. Thus austenite will be completely transformed to pearlite before the temperature of the bainite formation is reached. The bainite microstructure can be obtained only by cooling the steel from the austenite range to a temperature between 535 and 260°C at a rate greater than the critical cooling rate. This is accomplished in practice by quenching the steel in a bath of molten salt or a low-melting alloy and holding it at this temperature for some definite time. Then the steel is removed from the quenchant and cooled to room temperature.

Austenite A3

A1 A + F

A + F + C

Ferrite +

Cementite

MS

50% Martensite

| 10

| 102

| 103

| 104

| 105

| 106

Time, sec

800 −

600 −

400 −

200 −

Tem

per

atu

re o C

Lecture Note

23/36

8-10 HEAT TREATMENT OF CARBON STEELS

Carbon steels are those ferrous alloys in which carbon is the main element used to control their mechanical properties. Other elements, such as aluminum, boron, chromium, cobalt, nickel, titanium, tungsten, vanadium, and zirconium, are also commonly found in carbon steels, but their content is not specified and their presence is due to manufacturing processes. In addition, carbon steels contain up to 1.65% manganese, 0.60% silicon, and 0.4% copper.

Carbon steels are divided into three classes, according to their carbon content. Low-carbon steels contain from 0.08 to 0.35% carbon, medium-carbon steels contain from 0.35 to 0.50% carbon, and high-carbon steels have a carbon content greater than 0.5%.

Annealing and Normalizing. Steel, whether wrought or cast, is in a more or less coarsely granular condition. Heat treatment is necessary, particularly for cast steels, to break up the coarse structure and bring about refinement of grain size as well as uniform distribution of the constituents present. This is accomplished by annealing and normalizing. Full annealing consists of heating the steel to a temperature above its critical range, holding it there for a sufficient period of time, and slowly cooling it (usually in the furnace) to obtain an equilibrium structure. Heating above the upper critical temperature causes pearlite and any excess of cementite or ferrite to transform to austenite. This is known as austenitizing. Hypoeutectoíd steels are heated above the upper critical temperature (A3 line), whereas hypereutectoid steels are heated only above the lower critical temperature (A1 line), as seen in Fig. 8-21. Annealed hypereutectoid steels contain cementite and can never really be soft, but they are in the best possible condition for machining. Other heat treatment processes in which the steel is heated below the lower critical temperature are called stress relief and spheroidizing, in contrast to full annealing, which always consists of heating above the critical temperature.

FIGURE 8-21 Heat treatment temperatures for carbon steels. (Metals

Handbook, 1948 edition, American Society for Metals, Metals Park, Ohio.)

Spheroidizing is the formation of spheroids, which are embedded in a matrix of ferrite, from cementite. Particles are large enough to be readily visible under the light microscope. Martensite, bainite, and even pearlite can be changed to this structure, which results in considerable softening of the steel. Normalizing consists of heating the steel above its upper critical temperature, and then cooling it in still air. It differs from full annealing in that the rate of cooling is more rapid and there is no extended soaking period. Normalizing produces a uniform structure and refines the grain size of the steel, which may have been unduly coarsened at the forging temperature. If a steel is too soft

Te

mpe

ratu

re

o C

Fine grain temperature range

Normalizing Hardening Annealing Fine grain

temperature range

A1

Cementite +

Pearlite

A3

Ferrite +

Austenite

Spheroidizing Ferrite +

Pearlite Tempering

950 −

900 −

850 −

800 −

750 −

700 −

650 −

0 0.77% Carbon, % 2.0

Acm

Lecture Note

24/36

and ductile for machining, a normalizing process may achieve the required strength and ductility. The process is faster than full annealing and is often more suitable for producing a satisfactory steel.

Quenching. The hardening of steel requires the formation of martensite. This is accomplished by heating to a temperature high enough for steel to become austenitic, then cooling fast enough, usually by quenching in water or oil, to secure complete transformation to martensite. The composition of the steel to be hardened, the quenching technique used, and the design for heat treatment are all very important factors affecting the properties of the final product. Sharp corners, reentrant angles, and drastic changes in thickness should be avoided whenever possible in order to keep temperature gradients at a minimum throughout the specimen. In the fully quenched state steel containing more than 0.2% carbon has such low ductility as to render the material useless for engineering applications, and it must be softened to some desired degree by tempering before use.

Tempering. Tempering is a controlled heat treatment consisting of reheating martensite, which is extremely brittle when produced during quenching of carbon steel. To improve its properties and to impart greater toughness, the quenched steel is subject to tempering by heating to various temperatures below the critical temperature A1. Three stages of tempering are distinguished. In the first stage the quenched steel is heated to a temperature of 80 to 160°C (176 to 320°F), during which martensite loses some ε-carbide. This latter separates into very small particles and becomes less tetragonal. During the second stage. 230 to 280°C (446 to 536°F), any retained austenite transforms to bainite, and large dimensional changes occur. Finally, in the third stage, 260 to 360°C (500 to 680°F), ε-carbide changes to cementite platelets, producing a structure of ferrite and cementite. This is accompanied by a marked softening of the steel. The ε-carbide particles are extremely small in size, but they can be identified with the aid of the electron microscope. The ε-carbide is a closely packed hexagonal structure of composition of about Fe5C2. The tempered steel becomes much more ductile, but at the same time, there is an unavoidable loss of hardness.

In the process called martempering the steel is quenched to a temperature just above the martensite formation Ms and held there long enough to obtain the uniform temperature, but not long enough for bainite to form (see Fig.8-19). Then the steel is cooled in air. The resulting microstructure is martensitic, but the steel shows an improved ductility, and no tempering is necessary because martensite has been formed without the production of high internal stresses.

Hardenablllty. Hardenability is an index of the depth to which a given steel will harden when quenched in a particular manner, or it is the ease with which martensite can be produced in the steel. Hardenability is determined by a highly standardized test such as the Jominy end-quench test, but it is best assessed by the diameter of the largest steel cylinder that can be hardened all the way to its center by water quenching.

During quenching, a piece of steel will cool more rapidly at the surface than in the interior. This may cause the interior of the piece to be hardened to a lesser degree than the surface since, in a less rapidly cooled portion, some pearlite may be formed before martensite. For a piece of steel to be hardenable throughout, the critical cooling rate should be such that martensite forms before any pearlite has a chance to appear. This can be achieved by making the cross section very thin or by retarding the transformation of austenite to pearlite. This latter is equivalent to displacement of the S curve to the right on the TTT diagram so that the gate at the nose is considerably widened.

High ratios of surface to mass tend to produce greater depths of hardening because the rate of cooling depends on the speed with which heat leaves the specimen. Other factors are the surface conditions and the austenite grain size. Any scale formed on the quenched

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piece impairs hardenability, because it reduces the rate of heat removal from the interior. Large grain sizes tend to displace the S curve to the right, thus increasing hardenability, whereas fine grain sizes act in the opposite direction to reduce hardenability. The most effective method of producing deep hardenability in steel is to add certain alloying elements such as manganese or nickel (see Section 8-12). These displace the S curve to the right, thereby decreasing the rate of transformation of austenite and increasing the hardenability of the quenched specimen, even for large cross sections and under less drastic quenching conditions. This increased hardenability makes it possible for oil to be used as a quenching medium instead of water, thus reducing the tendency to warpage or cracking. Medium-carbon steels are capable of being through-hardened by the formation of martensite if the section size is small, perhaps up to 10 mm. High-carbon steels are always used in the fully hardened state, and they are used in many applications in which high hardness is essential.

8-11 CAST IRONS

In the carbon—iron equilibrium diagram (Fig. 8-14) cementite is considered a stable phase. However, cementite is far from being a stable compound and tends to decompose in the iron—carbon system to graphite and ferrite according to the reaction

Fe3C → 3Fe + C (8-12)

The tendency to form graphite flakes is governed mainly by the carbon content, the silicon content, and the cooling rate during solidification. Slow cooling results in the separation of carbon as graphite, producing a so-called gray cast iron that is not malleable. The carbon content of gray cast iron varies from 2.5% to 4% and the silicon content varies from 1% to 3%. Silicon increases the rate of graphite formation (Fig. 8-22a).

(a) (b)

(c) (d)

FIGURE 6.22 Optical photomicrographs of various cast irons. (a) Gray cast iron; the dark graphite flakes are embedded in an a ferrite matrix, X 244. (b) Malleable iron:

dark graphite rosettes (temper carbon) in an a ferrite matrix, X 150. (c) White cast iron; the white cementite regions are surrounded by dark pearlite, X 400. (d) Nodular

(ductile) cast iron: the dark graphite nodules are surrounded by a ferrite matrix, X 150.

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Gray cast iron is well suited to the production of intricate castings because of its high fluidity at casting temperatures and its low shrinkage. Other out standing advantages of gray cast iron are its excellent damping capacity, high resistance to wear and seizure. low notch sensitivity, and machinability. Because of its low cost and ready availability, gray cast iron deserves primary consideration whenever a cast metal is being selected.

If the total carbon content is held between 2.0% and 2.5% and the silicon content below 1%, and the casting is cooled rapidly, the carbon will remain in combined form as cementite. This latter accounts for the very high hardness and wear resistance of white cast iron. The name, white cast iron, stems from the white appearance of a fractured surface in contrast to the gray-colored fracture of gray cast iron. Because of its extreme hardness, white cast iron is almost unmachinable and has limited use. It is used only in those applications in which a very hard and wear-resisting surface is required (Fig. 8-22c).

When white cast iron is heated at 930°C (1706°F) for almost 50 h and then slowly cooled to room temperature, malleable cast iron results (Fig. 8-22b). This is because cementite decomposes to ferrite and temper carbon in the form of rosettes scattered in a ferritic matrix. Such a microstructure accounts for the appreciable ductility and good machinability of malleable cast iron, which compares favorably with cast low-carbon steels. When magnesium or cerium are added to the liquid cast iron, graphite precipitates in the form of spheroidal particles. The product is called ductile or nodular cast iron, showing considerable ductility while retaining the strength. machinability, and castability of gray cast iron (Fig. 8-22d).

8-12 EFFECT OF ALLOYING ELEMENTS

Alloying elements greatly affect the equilibrium diagram and alter the rate at which transformation reactions occur. The iron—carbon diagram may be profoundly altered; therefore its application should be limited to plain-carbon and low-alloy steels. For high-alloy steels it is necessary to consider the iron—chromium, iron—chromium—nickel, and iron—chromium--carbon systems.

Alloying elements can be divided into two classes: (1) those like nickel, manganese, copper, cobalt, carbon, and nitrogen that stabilize austenite. And (2) those like chromium, molybdenum, silicon, tungsten, vanadium, tin, columbium, and titanium that stabilize ferrite. These latter are called ferrite stabilizers or carbide-forming elements. since they easily combine with carbon to form carbides, whereas the first are called austenite formers. Silicon, although a ferrite former, is not a carbide-forming element but, on the contrary, it is a very effective graphitizing agent. Confusion may arise from the fact that these two groups of elements increase the stability of austenite with respect to the temperature (Ms) of spontaneous martènsite transformation. This is due to the ability of elements such as chromium, nickel, manganese, and molybdenum to form solid solutions with both austenite and ferrite. The elements, however, differ in their solubility in austenite or ferrite. Nickel, which stabilizes austenite, dissolves in austenite in all proportions but, inferrite, it dissolves only up to 25 to 30%. Chromium, a ferrite former, has only a limited solubility in austenite, up to 12.8%, but dissolves in all proportions in ferrite. In the presence of carbon the solubility of chromium in austenite increases and, for 0.5% carbon, the solubility of chromium, in austenite increases to 20%. Silicon is soluble in austenite only up to 2.1%; in ferrite it dissolves up to 18%.

The effect of nickel is to lower the critical transformation temperature, whereas the effect of chromium is to raise it. Both these elements shift the S curve ro the right, thus delaying the transformation of austenite to pearlite. This permits the formation of

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martensitic microstructures with slower rates of cooling, thereby increasing the hardenability of the steel (Fig. 8-23).

This principle is used in improving the hardenability of low-alloy engineering steels (AISI steels), which contain at least 0.3% C. Because of the presence of alloying elements such as nickel, chromium, manganese, or molybdenum, a complete transformation of austenite to martensite can be achieved in heavy cross sections, even in oil quenching. This makes it possible to achieve maximum strength in large forgings and to reduce the adverse effects of thermal stresses arising on very rapid cooling. Low-alloy, high-strength structural steels containing only up to 0.25% are not hardened to martensite. The purpose of the alloying elements is to improve the general overall properties such as corrosion resistance and durability.

FIGURE 8-23 Effect of alloying elements on the Bain curve. (E. C. Bain, Alloying

Elements in Steel, American Society for Metals, Metals Park, Ohio, 1939.)

With an increasing amount of alloying elements it is possible to control the basic microstructure of steels and, accordingly, to modify the properties. For example, on increasing the manganese content up to 12% while maintaining the carbon content of 1.1%, austenite becomes metastable at room temperature. Thus the austenitic microstructure results; such steel is referred to as austenitic-manganese steel. It exhibits exceptional hardening at the surface when subjected to impact or heavy cold work because austenite. as a result of its instability, transforms spontaneously to wear-resistant martensite, whereas the inner core of the steel remains austenitic. The combination of an extremely hard surface with a relatively soft but tough core makes the steel particularly suitable for jaw crushers, grinders, rails, and the teeth of power shovels, which require high abrasive resistance and high hardness as well as high impact resistance. Because of this extensive work hardening, austenitic manganese steel is very difficult to machine and is generally used in the form of castings.

Another example is a very important group of high-alloy steels known as stainless steels and heat-resistant steels.

8-13 STAINLESS AND HEAT-RESISTANT STEELS

The main alloying element in stainless steel is chromium, which must be present in sufficient quantity to make the steel corrosion resistant. Investigation has shown that for complete corrosion resistance the minimum chromium content should be over 11%.

Time, log scale

Tem

per

atu

re, o C

A1 Chromium steel

A1 Carbon steel 727oC

A1 Nickel steel

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Other alloying elements are also frequently added to enhance certain specific properties of the steel.

Stainless steels are usually divided into ferritic, martensitic, and austenitic types, according to their characteristic microstructure. The ferritic and martensitic stainless steels are straight-chromium steels designated as the type 400 series. The austenitic stainless steels are chromium—nickel alloys designated as the type 300 series.

Martensitic stainless steels contain from 11.5% to 18% chromium and from 0.15% to 1.20% carbon. The contents of chromium and carbon, also some times of nickel (only present in types 414 and 431 in an amount from 1.25% to 2.25%), are proportioned so as to permit hardening by heat treatment. Ferritic stainless steels are characterized by a chromium-to-carbon ratio that does not permit hardening by heat treatment. Chromium varies from 11.5% to 27% and carbon from 0.08% to 0.2% maximum. Both martensitic and ferritic stainless steels are magnetic under all conditions and offer good resistance to oxidation and corrosion. In ferritic stainless steels the main microconstituent is ferrite, present at both high and low temperatures. Consequently, there is no phase change on cooling or heating the alloy through the critical temperature range, and ferritic steels are not hardenable by heat treatment. In the martensitic stainless steels the same phase changes take place as in carbon steels and low-alloy engineering steels during their hardening. Austenite is a stable phase at higher temperature and transforms on rapid cooling to martensite, making the steel hardenable by suitable heat treatment.

The contents of carbon and chromium are the principal variables in determining the stability of austenite and ferrite at various temperatures. This is illustrated on the equilibrium phase diagram of the iron—chromium system with 0.1% carbon, as shown in Fig. 8-24.

FIGURE 8-24 The iron—carbon—chromium system with 0.1% C. (MetallurgicaJ

Aspects of Corrosion- and Heat-Remtance of Republic Enduro Stainless Steels, 1952.)

Chromium and iron are mutually soluble in the liquid state and, on cooling, form a continuous series of solid solution. The stability of austenite increases appreciably with the carbon content, whereas the stability of ferrite increases with the chromium content. Hence the γ loop representing the austenite region and the two-phase boundary region containing austenite and ferrite becomes larger as the carbon content is increased. At zero carbon content the limit for the existence of austenite is 12% chromium and somewhat larger (13% chromium) for the two-phase boundary region. As the carbon

Tem

per

atu

re o C

| | | | | | 0 5 10 15 20 25 30 35

Chromium %

1500 −

1300 −

1100 −

900 −

700 −

500 −

300 −

Liquid

Liquid + Ferrite

Ferrite

Ferrite + Austenite

Liquid + Austenite + Ferrite

Austenite

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content increases up to 0.35%, the austenite becomes stable up to about 15% chromium content, whereas the two-phase boundary region extends to even more than 20% chromium. It follows that to obtain ferritic stainless steel, the ratio of chromium to carbon must be sufficiently high to suppress the formation of austenite. Consequently, ferritic stainless steels are characterized by high chromium content with a reasonably low carbon content not exceeding 0.07% to 0.15%. On the other hand, martensitic stainless steels are characterized by a lower chromium-to- carbon ratio, so that austenite will be formed at high temperatures, allowing for hardening to martensite on cooling. Figure 8-24 shows that a steel with 18% chromium and 0.1% carbon remains essentially ferritic at all temperature ranges, so that hardening by quenching is not possible. A steel with 15% chromium is in a two-phase region above 1000°C (1830°F) and, when cooled from this temperature, will produce a microstructure made up of martensite and whatever ferrite exists. With still lower chromium content, the steel will consist only of the austenite phase above 1000°C (1832°F) and will produce a completely martensitic microstructure when quenched.

Both ferritic and martensitic stainless steels are magnetic under all conditions. Generally, the corrosion resistance and oxidation resistance of ferritic stainless steels are higher than those of martensitic steels, because of the higher chromium and lower carbon content. The martensitic stainless steels show their maximum corrosion resistance in the fully hardened condition, because then all chromium carbides are dispersed uniformly in the matrix, but ferritic steels exhibit their maximum corrosion resistance in the annealed state.

Chromium—Nickel Austenitic Stainless Steels. Austenitic stainless steels are essentially chromium—nickel—iron alloys, generally of a composition varying from 16% to 26% chromium and 6% to 22% nickel, with a maximum of up to 0.25% carbon. Other alloying elements such as molybdenum, vanadium, and titanium are also added to develop or intensify certain specific properties. The most widely used grade of austenitic stainless steel is type 18-8, containing from 17% to 19% chromium, 8% to 10% nickel, and up to 0.20% carbon. When 2% to 4% molybdenum is added to the basic composition, types 316 and 317 are obtained, which show an improved corrosion resistance against pitting and a higher temperature strength.

FIGURE 8-25 Iron—carbon—chromium phase diagram (18% Cr, 0.1% C). The phases are shown as they occur after heat treatment. (Metallurgical Aspects of Corrosion- and Heat-Resistance of Republic Enduro Stainless Steels, 1952.)

As nickel is progressively added to a ferritic stainless steel containing 18% chromium, the stability of the austenite phase increases until the steel becomes austenitic at all

1500 −

1300 −

1100 −

900 −

700 −

500 −

300 −

Liquid

Ferrite Liquid + Ferrite

Ferrite +

Austenite Austenite

Austenite +

Ferrite +

Carbide

Ferrite + Carbide

Austenite +

Carbide

| 0

| 2

| 4

| 6

| 8

| 10

| 12

| 14

| 16

Tem

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o C

Nickel %

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temperatures. For a steel with 0.1% carbon, a completely on austenitic structure is obtained with 7% nickel (Fig. 8-25). The structure, however, becomes complicated because of the presence of chromium carbides. These have very small solubility below a temperature of 1000°C on (1830°F) and precipitate preferentially along the grain boundaries. The microstructure of a steel with less than 7% nickel will consist of austenite, ferrite, and chromium carbides, whereas that of a steel with a minimum of 7% nickel will consist of austenite and precipitated carbides. Since chromium carbides, such as Cr27C3, contain appreciable quantities of chromium, the areas immediately adjacent to the grain boundaries will be impoverished in chromium content and the corrosion resistance of the steel will be considerably decreased. When the 18-8 steel is heated again to temperatures above 1000°C (1830°F), the carbides will pass into solution, resulting in a single-phase austenitic region. On quenching in water the precipitation of carbides can be suppressed. and the steel will have a completely austenitic structure. The quenching of austenitic steel from temperatures above 1000°C (1830°F) to room temperature is called stabilization, and steel in such a stabilized condition is the softest and most malleable material, while also showing maximum resistance to intergranular corrosion.

The stability of austenite increases as the nickel content is increased and decreases with the increase in chromium content or with the addition of other strongly carbide-forming elements such as molybdenum, titanium, vanadium, and columbium. Austenitic stainless steel is nonmagnetic in a stabilized condition, but the effect of cold work causes the steel to become slightly magnetic through the formation of some ferrite. As a result of the instability of austenite, various fabricated pieces of equipment may be in need of annealing after fabrication, particularly if localized sections of the steel are heated, for example, during a welding operation. The phenomenon known as welding decay is the result of intensive carbide precipitation in the regions adjacent to the weld, where the temperature of the steel ranges from 425 to 870°C (797 to 1598°F). Such steel can be restored to its normal condition only by stabilizing treatment. However, this is impractical in many circumstances, or even impossible when the steel has to be used in this temperature range. Therefore, it is essential to prevent the precipitation of chromium carbides in such cases; otherwise the material would find only limited application. This is accomplished by adding elements such as titanium or columbium. These, being much stronger carbide formers than chromium, combine with all the carbon in the steel, thus preventing the precipitation of chromium carbides. Carbide precipitation itself is not eliminated, but the undesirable effects of chromium carbide precipitation are minimized. In order to obtain complete stabilization titanium additions should be at least five to six times the carbon content of the steel and columbium additions eight to ten times the carbon content.

Another method of obtaining a stable material would be to eliminate completely the carbon from the steel. This is impractical and too expensive, but lowering the carbon content below 0.03% is possible and will decrease the danger of chromium carbide precipitation. Extra-low-carbon austenitic stain less steels (such as 316 E. L.) containing a maximum of 0.03% carbon have been developed, which give satisfactory performance without special heat treatment of the fabricated and welded pieces.

Austenitic stainless steels harden appreciably during cold working, but they cannot be hardened by heat treatment. Recently austenitic stainless steels (types 17-7 PH, 17-7 PH, AM 350, and others) that can be hardened by precipitation hardening have been developed. This is accomplished by adding such alloying elements as copper, titanium, beryllium, and aluminum, whose solubility in austenite decreases considerably as the temperature decreases causing their precipitation on suitable heat treatment. Recently, low-nickel, high-manganese austenitic stainless steels have been developed, designated as type 201 steel.

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The most recent development in this category is the “high-nickel maraging steels” in which solid-state precipitation is superimposed on martensitic reactions. The term “maraging” is derived because the martensitic transformation obtained on cooling this steel from about 830°C (1525°F) in air is followed by aging at 500°C (932°F) during which precipitation of such inter-metallic compounds as Ni3Mo occurs. This results in the yield strengths as high as 1.38 to 1.72 GPa (250 ksi), at the same time maintaining good fracture toughness. The capability for precipitation hardening is provided by adding alloying elements such as Ti, Mo, or Co, which form intermetallic compounds with nickel (Ni3Ti and Ni3Mo). The nickel content of such steels may range from 18% to 25%, while the carbon content is kept below 0.03%. The resultant low-carbon martensite is quite strong by itself, showing the yield strength of about 700 MPa (100 ksi), which is raised up to 1.38 (200 ksi) to 1.72 GPa (250 ksi) on precipitation hardening.

NONFERROUS METALS AND ALLOYS Methods similar to those for ferrous metals are used in strengthening non-ferrous metals and alloys. In the case of metals that do not show allotropic or transformation the available methods, in addition to cold work, reducing the grain’s size, and solid-solution hardening, are precipitation hardening and dispersion hardening processes.

8-14 PRECIPITATION HARDENING

As discussed previously (Chapter 3), the essential requirement for precipitation to occur in solid solution is the decreasing solubility of a solute with decreasing temperature. This results in a supersaturated solid solution that, being unstable, tends to decompose according to the relation

(8.13)

Figure 8-26a shows a phase diagram illustrating the type of alloy and conditions for precipitation hardening.

FIGURE 8-26 (a) Al—Cu partial equilibrium diagram. Only the part rich in aluminum is

shown. (Adapted from Metals Handbook, 1948 edition. By permission of American Society for Metals.) (b) Precipitation hardening process showing schematically resulting

microstructures. (1) Solid solution treatment: heat to a temperature of about 540°C (1004°F) (2) Quenching: cool rapidly in water at 20°C (78°F). (3) Age or precipitation hardening: reheat to a temperature of about 200°C (392°F) and hold 2 h. Fine precipitate results. (4)

Overaging results if temperature is too high and/or the time is too long. Fine particles coalesce to form coarse noncoherent precipitate causing a considerable decrease in strength.

Supersaturated α solid solution

Saturated α solid solution

β precipitation

Tem

per

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Solid solution

Time Cu, wt%

Liquid

α + liquid

α + θ (CuAl2) 5.65%

1

3

2

α 548oC

4.5%

Tem

per

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700 −

600 −

500 −

400 −

300 −

200 −

100 −

0 − | 1

| 2

| 3

| 4

| 5

| 6

| 7

| 8

| 9

| 10

| 11

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In general, the second-phase β precipitate is formed by a sequence of thermal treatments. First, an alloy of suitable composition, as indicated by line xx in Figure 8-26, is subject to a solution heat treatment. This consists of heating the alloy to a temperature above the solvus line but below the solidus line in region α (point 1), and holding at this temperature for a sufficiently long time so that complete homogenization of the solid solution takes place. Then the alloy is quenched to about room temperature (point 2 on the diagram) so that precipitation of a solute in excess of equilibrium concentration does not occur and a supersaturated solution results. On quenching, the high concentration of vacancies corresponding to the equilibrium concentration at the solution temperature (point 1) is retained. These vacancies slowly precipitate or condense out, producing clusters and dislocation loops that form the sites at which precipitate particles can nucleate. Furthermore, the quenched-in vacancies enhance the diffusion rate of the solute, thereby promoting the formation of precipitate.

Quenching does not cause significant changes in mechanical properties, although some lattice strains may be present because of thermal stresses. The quenched alloy is relatively soft and may be worked to the desired shape without precipitation occurring. The third step, called age hardening or precipitation hardening, involves reheating the alloy to an elevated temperature somewhere below the solvus line (around point 3) and holding at this temperature, called aging temperature, for a certain period of time to develop the necessary amount and kind of precipitate that will impart a desired strength (see Fig. 8-26b).

The increase in the yield strength of the precipitation-hardenable alloy depends primarily on the amount of precipitate and its characteristics, such as the particles’ size, shape, and distribution. As the temperature of precipitation is lowered, the precipitate particles become smaller and more numerous and the strengthening effect is greater. With increasing temperature particle size increases and the precipitate continues to coalesce as aging progresses.

Precipitation hardening is a very important method of strengthening many solid solution alloys such as a variety of aluminum and magnesium alloys, copper—beryllium alloys, high-nickel-base alloys, and some stainless steels. The most effective way of strengthening such alloys is to use a solute element with a sharp change of solubility with temperature and in sufficient atomic concentration to produce large volumes of fine precipitate uniformly distributed within the grains. The aging temperature depends on the composition, structure, and type of alloy. At early stages of precipitation, the electrical resistivity of alloys increases but, at later stages of precipitation, the resistivity may decrease. The use of an age-hardenable alloy is restricted to temperatures in service, during which time overaging may not become excessive. It is apparent that the precipitation-hardenable alloy should never be used at temperatures above its solvus line since, at such temperatures, the precipitate will dissolve and complete homogenization of the solid solution will occur, completely eliminating the effect of hardening. To avoid these problems, a dispersion hardening process has been developed that involves dispersing some fine insoluble phase, usually refractory oxides, throughout the base metal matrix. This method is considered in detail in Chapter 14. The effect of precipitation hardening of an aluminum—4.5% copper alloy is shown in the following table:

Treatment Yield Strength MPA

Tensile Strength MPa

Elongation %

1 - 2 105 240 30

Age hardened, 3 331 415 20

Overaged, 4 70 120 20

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8-15 COPPER AND COPPER ALLOYS

Copper and copper alloys are widely used in many industrial applications because of their high thermal and electrical conductivity, their good corrosion resistance under a wide range of operating conditions, their ease of fabrication, and their ready availability. Commercially pure copper is highly malleable and ductile and can be hardened by cold work or by the addition of alloying elements. Zinc, aluminum, tin, and silicon are among the elements most frequently used, although other elements may be used for specific purposes. Zinc is by far the most important alloying element and yields a series of important copper—zinc alloys known as brasses.

Brasses contain from 5% to 45% zinc with or without smaller additions of one or more of the other elements. The system coper-zinc is illustrated by the phase equilibrium diagram on Fig. 8-27. With copper zinc forms a series of solid solutions known as a brasses containing varying amounts of zinc from 5% to 37%. The α brasses have the same face-centered cubic crystal structure as copper, which accounts for their ductility and the ease with which they may be cold-worked. The α brasses can also be hot-worked, but the presence of impurities such as lead and bismuth causes hot shortness. When the zinc content exceeds 37.5%, a new β phase appears that is hard and brittle. Alloys of compositions from 37.5% to 44% zinc are composed of two phases, α and β, whereas those with zinc content of from 44% to 46% consist of only a β phase. The single-phase β alloys are not used in engineering applications because of their brittleness, but the two-phase alloys have good hot-working properties over a wide range of temperatures and are extensively used as hot-working materials.

FIGURE 8-27 Copper—zinc equilibrium diagram. (Metals Handbook, 1948

edition, American Society for Metals, Metals Park, Ohio.)

The name bronze is given to a variety of copper alloys with or without tin. Many commercial bronzes contain from 2% to 13% tin and are harder and stronger than brasses and generally show better corrosion resistance. The addition of 0.2% to 0.4% phosphorus improves the strength, particularly under cold-worked conditions. Other types of commercial bronzes contain such alloying elements as aluminum, silicon, and beryllium. Beryllium bronze, known also as beryllium copper, is one of the few strongly precipitation hardenable copper alloys with about 2% beryllium. A yield strength of 900 MPa (130 ksi) and a tensile strength of 1.21 GPa (175 ksi) can be obtained.

Copper—Nickel Alloys. Copper—nickel alloys, also called cupronickels, are copper-base alloys containing from 5% to 30% nickel. Because of the mutual solubility of copper and nickel in the solid state, cupronickels form a series of solid solution alloys that can be easily cold-worked. Their mechanical properties as well as their corrosion and heat resistance increase progressively with increasing nickel content. Wrought alloys

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can be considerably strengthened by cold work, whereas cast alloys may contain alloying elements such as iron, silicon, aluminum, and tin, which make the alloys age hardenable. The wrought alloys are usually available in three compositional ranges, 90% Cu—10% Ni, 80% Cu-20% Ni, and 70% Cu—30% Ni, and they find applications as corrosion-resistant alloys for use in salt water and brackish water.

8-16 NICKEL AND NICKEL ALLOYS

Nickel and high-nickel alloys have been used widely because of their superior corrosion resistance and high-temperature strength. Commercially pure nickel is almost as hard as a low-carbon steel, but it cannot be hardened by heat treatment. The addition of small amounts of alloying elements can make the metal heat treatable by the age-hardening method. Nickel work-hardens rapidly when cold-worked.

Nickel—copper alloys contain nickel in an amount greater than 50% and, like copper—nickel alloys, are of the solid solution type, with FCC structure. The most important alloy in this group is Monel, which has a composition of 63% to 70% nickel and up to 30% copper. Another very important group of nickel-base alloys called superalloys comprises various nickel—chromium—iron alloys such as Inconels, René-s, Uldimets, and others (see Table 2A) and may be produced in cast, wrought, and powder forms. They are solid solution alloys having an austenitic matrix FCC containing solid solution strengtheners such as V, Cr, Mo, W, Co, and others. Most of these superalloys can be strengthened by the γ’ precipitate coherent with the austenitic matrix. The precipitate is based on the ordered Ni3Al structure in which Ni atoms are located at face centers and Al atoms or atoms of other elements such as Ti, Nb, and Ta at cube corners. Alloys containing 2—6% Nb are strengthened by γ” precipitate of Ni3Nb having a body-centered tetragonal structure. The γ” precipitate is much finer than the γ’, and it may convert to an ordered Ni3Nb phase having an orthorhombic crystal structure. The presence of precipitate and the resultant coherency strain make it difficult for dislocations to penetrate the precipitate particles, thereby increasing the critical resolved shear strength. The γ’ phase is quite stable at elevated temperatures as high as 900°C (1650°F), maintaining a virtually constant yield strength. The shape and the structural stability of the γ’ precipitate depend on the misfit parameter δ as defined by Equation 4-5. When the misfit parameter is small or the particle size is small, the γ’ particles tend to assume spherical shape. With a large misfit and large precipitate continued aging results in a precipitate of cuboidal shape. Nb and Ti tend to increase the misfit whereas Fe and Mo tend to decrease it. For misfit nearly zero all γ’ precipitate is very small and spherical.

(a) (b)

FIGURE 8-28 Optical micrographs of(a) stellite 6B (1.1% C, 30% Cr, 4.5% W, 1.5% Mo, 3% max Ni, 3% max Fe, balance Co). Round large spheres are M7C3; small spheres at grain boundaries are M23C6. FCC matrix. Hot-rolled sheet, X 413. (b)

Hastelloy G (44.5% Ni. 0.03% C, 1.5% Mn. 22.3% Cr. 6.5% Mo. 20% Fe, 0.1% cb=Ta. FCC matrix. Precipitate is M6C. Hot-rolled sheet, X 150.

Lecture Note

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Most nickel-base superalloys contain carbides of types MC, M6C, and M23C6 dispersed in the matrix or located in the grain boundaries (Fig. 8-28). Carbides of the MC type such as TaC, NbC, TiC, and VC have FCC structure and have a tendency to decompose with increasing temperature. Carbides M6C and M23C6 have a complex cubic structure and have a tendency to form along grain boundaries; with high Cr content they are stable in the range 870—980°C (1600—1800°F). The grain boundaries are rich in elements such as boron, zirconium, and magnesium. Boride (M3B2) particles form a hard, refractory precipitate that delays the onset of grain boundary tearing during creep.

A number of alloys known as Hastelloys and Chlorimets, particularly those containing chromium and molybdenum, are also solid solution types showing superior corrosion resistance to many corrosive environments as well as high-temperature strength and oxidation resistance. Hastelloys B and B-2 are nickel—molybdenum—iron alloys, Hastelloy C-4 and Illium alloys are nickel—chromium—molybdenum−iron alloys, and Hastelloy D is a nickel—silicon alloy.

8-17 ALUMINUM, MAGNESIUM, AND THEIR ALLOYS

Aluminum and aluminum alloys are characterized not only by lightness, but also by a high strength-to-weight ratio, good corrosion resistance, high thermal and electrical conductivity, nontoxicity, and ease of fabrication. They are used in both the wrought and cast condition in many applications, such as in the aircraft construction, in chemical and process industries, in building industries, and in domestic uses.

The strength of pure aluminum is increased by cold work and alloying, but this also reduces its corrosion resistance. The advantage of increased corrosion and resistance is counteracted by the greater metal thickness required and the increased cost of the superpure material, which is quite weak. One of the main limitations of aluminum is its low melting point, 660°C (1220°F), and its inability to be used at temperatures above 300°C (570°F) (because of creep) in spite of its superior oxidation resistance. Special alloys with better high temperature properties have been developed. On the other hand, aluminum, with a face-centered cubic lattice, maintains its ductility at a very low temperature. Aluminum can be cold worked easily by a variety of methods.

High-purity aluminum (99.95% Al) is very soft and ductile, but it is considerably strengthened by the addition of even slight amounts of alloying elements such as silicon, manganese, iron, and copper. Commercially pure aluminum and some of its solid-solution alloys can be hardened only by cold work, but many age-hardenable alloys have been developed that contain alloying elements such as copper, zinc, or magnesium silicide. The most widely used are Al—Cu alloys (2000 series) and Al—Zn--Mg alloys (7000 series). They are strengthened by solution heat treatment to place all soluble elements in a supersaturated matrix followed by precipitation hardening (aging) at room or at mildly elevated temperatures. Alloys of the 1000, 3000, and 4000 series are strengthened only by cold work. Alloys of the 5000 (Al—Mg) and 6000 (Al—Mg—Si) series are lower in strength than those of the 2000 and 7000 series. The alloys of the 6000 series are hardened by solution heat treatment and subsequent precipitation aging with particles bearing Mg—Si.

MagnesIum and Its Alloys. Magnesium for engineering applications is usually alloyed with aluminum (6% to 12%), zinc (0% to 3%), and manganese (0.2%), and with small amounts of silicon, cadmium, tin, lithium, zirconium, and cerium. All these alloys are age hardenable. but the strength achieved by age hardening is not as great as that obtained in aluminum alloys.

8-18 TITANIUM AND TITANIUM ALLOYS

Titanium exists in two crystalline states: the low-temperature α phase, which has a HCP structure and c/a ratio of 1.58 stable below 883°C (1621°F), and the high-temperature β

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phase, which is stable above 828°C (1621°F) and has BCC structure. The phase composition and the transformation temperature can be modified by addition of alloying elements. Aluminum is an α stabilizer whereas vanadium is the β stabilizer.

Titanium finds numerous industrial applications because of its corrosion resistance and high strength/density ratio. For higher strength requirements there are a number of titanium alloys; the most widely used is the Ti—6 wt% AI—4 wt% V alloy which exists in two phases, α and β. The phase is stable above 1000°C (1832°F); below 1000°C, α and β phases coexist making it a two-phase alloy.

The successful use of the Ti—6AI—4V alloy, which can be conveniently produced by a powder metallurgy technique, depends on the microstructure which can be controlled and adequately modified by proper mechanical and thermal treatments to obtain the optimum strength, ductility, and toughness. On rapid quenching from 1000°C (1832°F) the Ti—6A1—4V alloy gives a HCP martensitic microstructure which contains vanadium in supersaturation. On subsequent aging at a temperature at which appreciable diffusion can take place, α’ is formed by the precipitation of β on the martensite plate boundaries and dislocation. Under moderate cooling rates the precipitation of α from the β field forms the α—β region. The TÍ—6Al—4V alloy is not used under martensitic conditions because it cannot be effectively hardened in a thickness greater than 25 mm (1 in.); it is then hot worked in the α—β region to produce a uniformly distributed α phase in a finely divided form. This is followed by annealing at 700°C (1292°F). The annealed Ti—6A1—4V alloy owes its strength of 895 MPa (130 ksi) and toughness K1C = 44—66 Mpa.m1/2 (40—60 ksi.in.1/2) to the presence of substitutional and interstitial elements forming solid solutions in the α and β phases. Aluminum is the most effective solid-solution strengthener.

Other high-strength α and α—β alloys are Ti—6Al—6V—2Sn used in air frames and the super-α Ti−6Al−2Sn—4Zr−6Mo alloy used in jet engines. Martensite forms more readily in these alloys than in the Ti—6A1—4V alloy. Commercially pure titanium and Ti—6Al—4V alloy are also used in many applications as implant and fixation devices.