Volume 01 - Properties and Selection Irons, Steels, and High-Performance Alloys Part 2 pptx

Fig Effect of carbon equivalent on the tensile strength of flake, compacted, and spheroidal graphite irons cast into 30 mm (1.2 in.) diam bars Source: Ref 10 Although increasing the silicon content decreases the pearlite to ferrite ratio in the as-cast state, both the strength and hardness of as-cast and annealed CG irons improve This is because of the hardening of ferrite by silicon For the same reasons, elongation in the annealed condition decreases, but increases for the as-cast state (Ref 10) Although increasing the phosphorus content slightly improves strength, a maximum of 0.04% P is desirable to avoid lower ductility and impact strength The pearlite/ferrite ratio, and thus the strength and hardness of CG irons, can be increased by the use of a number of alloying elements such as copper, nickel, molybdenum, tin, manganese, arsenic, vanadium, and aluminum (Ref 6, 14) The effect of copper and molybdenum on the tensile properties of CG irons is shown in Fig After annealing to a fully ferritic structure, it is possible to increase the yield point of CG iron by 24% when using 1.5% Ni (Table 3) This is because of the strengthening of the solid solution by nickel (Ref 6) The reader is cautioned, however, that additions of copper, nickel, and molybdenum may increase nodularity (Ref 6) Table Effect of heat treatment and alloying with nickel on the tensile properties of CG iron measured on a 25 mm section size Heat treatment Iron Matrix(a) 60% F Yield strength MPa As-cast Tensile strength ksi MPa 47.1 263 38.1 Hardness, HB Nickel, % 2.8 153 ksi 325 Elongation, % Annealed(b) 100% F 294 42.6 231 33.5 5.5 121 Normalized(c) 90% P 423 61.3 307 44.5 2.5 207 As-cast 427 61.9 328 47.6 2.5 196 1.53 Annealed(b) 100% F 333 48.3 287 41.6 6.0 137 1.53 Normalized(c) 90% P 503 73 375 54.4 2.0 235 1.53 (a) F, ferrite; P, pearlite (b) Annealed, h at 900 °C (1650 °F), cooled in furnace to 690 °C (1275 °F), held 12 h, cooled in air (c) Normalized, h at 900 °C (1650 °F), cooled in air Fig Effect of (a) copper and (b) molybdenum on the tensile properties of CG iron Source: Ref 13, 14 In order to compare the quality of different types of irons, several quality indexes can be used, such as the product of tensile strength and elongation (TS × El) or the ratio of tensile strength to Brinell hardness (TS/HB) Higher values of these indexes will characterize a better iron Using the data given in Ref 10, some typical values were calculated for these indexes for unalloyed CG irons Figure 10 compares the TS × El product and the TS/HB ratio for unalloyed and aluminum-alloyed CG irons It can be seen that when 2% Si is replaced by 2% Al, a much better quality CG iron is produced (Ref 11) Fig 10 Quality indexes for Fe-C-Si and Fe-C-Al irons TS, tensile strength; El, elongation Source: Ref 11 Effect of Structure One of the most important variables influencing the tensile properties of CG irons is nodularity As nodularity increases, higher strength and elongation are to be expected, as shown in Table and Fig 11, although nodularity must be maintained at levels under 20% for the iron to qualify as CG iron However, spheroidal graphite contents of up to 30% and even more must be expected in th in sections of castings with considerable variation in wall thickness Table Properties of CG iron as a function of nodularity Nodularity, % Tensile strength Elongation, % Thermal conductivity, W/(m · K) Shrinkage, % MPa ksi 10-20 320-380 46-55 2-5 50-52 1.8-2.2 20-30 380-450 55-65 2-6 48-50 2.0-2.6 40-50 450-500 65-73 3-6 38-42 3.2-4.6 Fig 11 Correlation between nodularity and elongation for ferritic CG iron Source: Ref 16 As previously discussed, the pearlite/ferrite ratio can be increased by causing alloying elements Another way of increasing or decreasing this ratio is by using heat treatment The influence of heat treatment on mechanical properties of CG irons with 20% nodularity is given in Table Effect of Section Size Like all other irons, CG irons are rather sensitive to the influence of cooling rate, that is, to section size, because it affects both the pearlite/ferrite ratio and graphite morphology As mentioned before, a higher cooling rate promotes more pearlite and increased nodularity A typical example of the influence of section size on the microstructure of CG iron is provided in Fig 12 Although CG iron is less section sensitive than FG iron (as shown in Fig 13 for tensile strength), the influence of cooling rate may be quite significant (see the article "Compacted Graphite Irons" in Casting, Volume 15 of ASM Handbook, formerly 9th Edition Metals Handbook) When the section size decreases until it is below 10 mm (0.4 in.), the tendency to increased nodularity and for higher chilling must be considered This is particularly true for overtreated irons While it is possible to eliminate the carbides that result from chilling by heat treatment, it is impossible to change the graphite shape, which remains spheroidal, with the associated consequences Other factors influencing the cooling of castings, such as shakeout temperature, can also influence properties Fig 12 Influence of section size on the microstructure of CG iron produced by inmold Samples from series E3.3 (a) 3.2 mm (0.125 in.) (b) 6.4 mm (0.250 in.) (c) 12.7 mm (0.500 in.) (d) 25.4 mm (1.000 in.).(e) 50.8 mm (2.000 in.) 100× Fig 13 Influence of section size on the tensile strength of CG iron Source: Ref 15 Compressive Properties The stress-strain diagram for compression and tensile tests of CG iron is shown in Fig 14 It can be seen that an elastic behavior occurs up to a compression stress of 200 MPa (30 ksi) Some compressive properties of the 179 HB as-cast ferritic CG iron in Table are compared with those of SG iron in Table It can be seen that the 0.1% proof stress in compression for CG iron is 76 MPa (11 ksi) higher than the 0.1% proof stress in tension, while for SG iron the difference is only 23 MPa (3.3 ksi) Compressive strengths up to 1400 MPa (203 ksi) have been reported for ferritic annealed CG irons (Ref 8) Table Comparison of tensile and compressive properties of CG and SG irons Property CG iron SG iron Tensile strength, MPa (ksi) 380 (55) 370 (54) 420 (61) 0.1% proof stress, MPa (ksi) 246 (35.7) 224 (32.5) 261 (37.8) 0.2% proof stress, MPa (ksi) 242 (35.1) 236 (34.2) 273 (39.6) 0.1% proof stress, MPa (ksi) 322 (46.7) 247 (35.8) 284 (41.2) 0.2% proof stress, MPa (ksi) 350 (50) 250 (36.3) 287 (41.6) Compressive stress Source: Ref Fig 14 Stress-strain curves in compression and tension for CG iron with 4.35 carbon equivalent Source: Ref Shear Properties For a pearlitic CG iron, the shear strength on 20 mm (0.8 in.) diam specimens was measured at 365 MPa (53 ksi), with a shear-to-tensile strength ratio of 0.97 (Ref 17) Ratios of 0.90 for SG iron and of 1.1 to 1.2 for FG iron have been reported Materials exhibiting some ductility have ratios lower than 1.0 (Ref 8) Modulus of Elasticity As is evident from Fig and 14, CG irons exhibit a clear zone of proportionality, both in tension and in compression Typical values for both static and dynamic (resonance frequency method) measurements are given in Table Dynamic tests give slightly higher numbers In general, the moduli of elasticity for CG iron are similar to those of high-strength FG irons and can even be higher as nodularity increases The elasticity modulus measured by the tangent method depends on the level of stress, as shown in Fig 15 A comparison of the stress dependency of the elasticity modulus for different types of cast irons is shown in Fig 16 Poisson's ratios of 0.27 to 0.28 have been reported for CG irons (Ref 8) Fig 15 Stress dependency by E-modulus for two heat-treated CG irons Source: Ref Fig 16 Influence of stress level on the elasticity modulus (a) in tension and (b) in compression for pearlitic FG, CG, and SG irons Source: Ref Impact Properties While SG iron exhibits substantially greater toughness at low pearlite contents, pearlitic CG irons have impact strengths equivalent to those of SG irons (Fig 17) Charpy impact energy measurements at 21 °C (70 °F) and -41 °C (-42 °F) showed that CG irons produced from an SG-base iron absorbed greater energy than those made from gray iron-base iron (Ref 10) This is attributed to the solute hardening effects of tramp elements in the gray iron Fig 17 Effect of pearlite content on the 21 °C (70 °F) Charpy V-notch impact strength of as-cast CG irons compared to that of SG iron Source: Ref 10 The results from dynamic tear tests were similar, although greater temperature dependence was observed A comparison of the dynamic tear energies of CG cast irons is presented in Fig 18 It is noted that significant differences in the values obtained occur in the ferritic condition, but that equivalent values are obtained when the matrix structure is primarily pearlitic Fig 18 Dynamic tear energy versus temperature for (a) CG and (b) SG irons Source: Ref 10 Studies on crack initiation and growth under impact loading conditions showed that, in general, the initiation of matrix cracking was preceded by graphite fracture at the graphite-matrix interface, or through the graphite, or both The most dominant form of graphite fracture appeared to be that occurring along the boundaries between graphite crystallites (Ref 18) Matrix cracks were usually initiated in the ferrite by transgranular cleavage (graphite was nearly always surrounded by ferrite), although in some instances intergranular ferrite fracture appeared to be the initiating mechanism Matrix crack propagation generally occurred by a brittle cleavage mechanism, transgranular in ferrite, and interlamellar in pearlite In general, the impact resistance of CG irons increases with carbon equivalent and decreases with phosphorous or increasing pearlite As may be seen in Table 6, cerium-treated CG irons seem to exhibit a higher impact energy than magnesiumtitanium-treated irons It is thought that this may be attributed to TiC and TiCN inclusions present in the matrix of magnesium-titanium-treated CG irons (Ref 6) Fig Isothermal transformation diagram for 1080 steel containing 0.79 wt% C and 0.76 wt% Mn Specimens were austenitized at 900 °C (1650 °F) and had an austenitic grain size of ASTM No The Ms, M50, and M90 temperatures are estimated Source: Ref The kinetics of pearlite formation, based on the nucleation and growth of spherical pearlite colonies, are described by (Ref 6): f(t) = - exp [ - πNG3t4/3] (Eq 2) where f(t) is the volume fraction of pearlite formed at any time t at a given temperature, N is the nucleation rate of the colonies, and G is the growth rate of the colonies Equation describes well the slow initial and final transformation rates and the more rapid intermediate transformation rates observed for isothermal pearlite formation With decreasing temperature below A1, N and G increase, and the transformation of austenite to pearlite accelerates, as shown in Fig During pearlite formation, in addition to carbon atom diffusion, irons atoms must also transfer across the interface between the austenite and pearlite This short-range iron atom transfer is necessary to accomplish the crystal structure changes among the austenite, ferrite, and cementite At a critical low temperature, this atom-by-atom short-range diffusion is no longer possible, and the iron atoms accomplish the crystal structure change by shearing or cooperative displacement (Ref 7) This change in transformation mechanism results in a new type of microstructure, referred to as bainite The ferrite crystals assume elongated morphologies, and the cementite is no longer continuous and lamellar (Ref 8) The bainite that forms at temperatures just below those at which pearlite forms (Fig 4) is termed upper bainite In medium- and high-carbon steels, it typically consists of groups of ferrite laths with coarse cementite particles between the laths The bainite that forms at lower temperatures is termed lower bainite and consists of large needlelike plates that contain high densities of very fine carbide particles Figures and show examples of upper and lower bainite, respectively In these micrographs the bainite morphology is dominated by the ferrite phase, and the carbides are too fine to be resolved by the light microscope In some low- and medium-carbon steels (generally, those alloyed with manganese, molybdenum, and silicon), bainitic microstructures with ferrite and austenite (or martensite formed from the austenite) will form instead of the classic ferrite-carbide bainitic structures (Ref 9) Fig Light micrograph showing patches of upper bainite (dark) formed in 4150 steel partially transformed at 460 °C (860 °F) Courtesy of F.A Jacobs Fig Light micrograph showing lower bainite (dark plates or needles) formed in 4150 steel Courtesy of F.A Jacobs References cited in this section G Krauss, Physical Metallurgy and Heat Treatment of Steel, in Metals Handbook Desk Edition, H.E Boyer and T.L Gall, Ed., American Society for Metals, 1985, p 28-2 to 28-10 M Hillert, The Formation of Pearlite, in Decomposition of Austenite by Diffusional Processes, V.F Zackay and H.I Aaronson, Ed., Interscience, 1962, p 197-247 W.A Johnson and R.F Mehl, Reaction Kinetics in Processes of Nucleation and Growth, Trans AIME, Vol 135, 1939, p 416-458 J.W Christian and D.V Edmonds, The Bainite Transformation, in Phase Transformations and Ferrous Alloys, A.R Marder and J.I Goldstein, Ed., The Metallurgical Society, 1984, p 293-325 R.F Hehemann, Ferrous and Nonferrous Bainite Structures, in Metals Handbook, 8th ed., Vol 8, American Society for Metals, 1973, p 194-196 B.L Bramfitt and J.G Speer, A Perspective on the Morphology of Bainite, Metall Trans A, to be published in 1990 Proeutectoid Ferrite and Cementite Steels with a lower carbon content than the eutectoid composition (hypoeutectoid) steels) and a higher carbon content than the eutectoid composition (hypereutectoid steels) form ferrite and cementite, respectively, prior to pearlite The structures formed upon cooling between Ar3 and Ar1, and Arcm and Ar1 are referred to as proeutectoid ferrite and proeutectoid cementite, respectively Figure shows a low-carbon steel ferrite-pearlite microstructure that formed during air cooling from the austenite phase field (Fig 1) The sequence of transformation to this microstructure is described below When the temperature of the specimen reaches the Ar3 temperature, proeutectoid ferrite begins to form The ferrite crystals, or grains, as they are referred to by metallurgists, nucleate on austenite grain boundaries and grow, by rearrangement of iron atoms, from the fcc austenite structure into the bcc ferrite structure at the austenite-ferrite interface Carbon atoms, because of their low solubility in the ferrite, are rejected into the untransformed austenite When the steel reaches the Ar1 temperature, most of the microstructure has transformed to proeutectoid ferrite, and the carbon content of the remaining austenite has been enriched to about 0.77 wt%, which is exactly the composition required for the pearlite reaction Thus, the balance of the austenite transforms to pearlite, as described in the preceding section Fig Light micrograph showing microstructure of proeutectoid ferrite (white) and pearlite (dark) in a 0.17C1.2Mn-0.19Si steel In Fig 7, the ferrite appears white, and the boundaries between ferrite grains of different orientation appear as dark lines The pearlite, in contrast to Fig 3, appears uniformly black because the interlamellar spacing of the pearlite in this example is too fine to be resolved by the light microscope Figure shows an example of proeutectoid cementite in a hypereutectoid steel The cementite has formed as a thin network along the grain boundaries of the austenite, and the balance of the microstructure is martensite, which formed when the specimen was quenched from a temperature between Acm and A1 The cementite and its interfaces are preferred sites for fracture initiation and propagation, and as a result, proeutectoid cementite networks make hypereutectoid steels extremely brittle Intercritical annealing treatments that break up and spheroidize the cementite are therefore used to increase toughness Fig Light micrograph showing cementite network on prior-austenite grain boundaries in an Fe-1.12C-1.5Cr alloy Courtesy of T Ando Processing: Ferrite-Pearlite Microstructures Although some steel products are directly cast to shape, most are wrought or subjected to significant amounts of hot and/or cold work during their manufacture Figure depicts some of the primary processing steps that result in ferritepearlite microstructures In some cases, processing produces finished products, such as plate and hot-rolled strip In other cases, as discussed in the following section, further processing is applied; for example, steel bars are subsequently forged and heat treated, and hot-rolled strip is cold rolled and annealed Fig Temperature-time schedules for primary processing of steels cast by various technologies The thermomechanical processing temperature ranges are shown in Fig relative to the critical temperatures identified in the Fe-C phase diagram (Fig 1) All the hot rolling is done with steels in the austenitic condition, and because of the large section sizes and equipment design, the austenite transforms to microstructures of ferrite and pearlite Depending on the casting technique, some steel products undergo several cycles of austenite to ferrite-pearlite transformation, as shown in Fig Each step in a cycle requires the nucleation and growth of new phases and offers the possibility of controlling grain size and the distribution of various microstructural components Traditionally, hot working has been used to reduce large ingots to products with reduced cross sections and special shapes, but there has been little attempt to control microstructure other than the use of slightly reduced finish hot-rolling temperatures This approach has been substantially modified in recent years by the implementation of two major changes in process design: the casting of smaller and smaller sections, and the use of hot rolling to control microstructure and properties as well as to reduce section size The progression of casting technology from ingot to continuous (or strand) to thin slab to direct strip casting is shown in Fig Continuous casting eliminates the soaking and breakdown hot rolling of large ingots Thin slab casting eliminates the roughing hot work applied to thick slabs that are produced by either ingot or continuous casting Direct strip casting, which is still under development, would eliminate all hot work (Ref 10) A considerable savings of time and energy and improved surface quality result from the new casting techniques The control of microstructure and properties during hot-rolling involves a thermomechanical processing technique known as controlled rolling Controlled rolling is used to enhance the toughness and strength of microalloyed low-carbon plate and strip steels by grain refinement (see Ref 11, 12, 13 and the article "High-Strength Structural and High-Strength LowAlloy Steels" in this Volume) The fine grain sizes produce significantly increased strength, from the 210 MPa (30 ksi) yield strengths that are typical of conventionally hot-rolled low-carbon steels to yield strengths between 345 and 550 MPa (50 and 80 ksi) The key to the use of controlled rolling is the formation of fine austenite grains that transform upon cooling to very fine ferrite grains Deformation of austenite induces strains, which, at high temperatures, are rapidly eliminated by recrystallization, followed by grain growth of the austenite However, at low deformation temperatures, grain growth is considerably retarded If the temperature is low enough, even recrystallization is suppressed, especially in steels to which small amounts of alloying elements such as niobium have been added The niobium, which is soluble at high temperatures, precipitates out as fine niobium carbonitrides at low austenitizing temperatures These fine precipitate particles stabilize the deformation substructure of the deformed austenite and prevent recrystallization Upon cooling, ferrite grains nucleate on the closely spaced grain boundaries of the unrecrystallized austenite and form very fine grain microstructures A number of thermomechanical processing schedules have been developed to produce low-carbon steels of high strength and toughness Figure 10 shows schematically a number of treatments developed by Japanese steelmakers (Ref 12) Similar controlled rolling schedules are used worldwide to produce fine-grain low-carbon steels A processing parameter introduced in Fig 10 is the austenite recrystallization temperature (TR, which is primarily dependent on the amount of deformation and alloying Fig 10 Temperature-time schedules for thermomechanical processing of steels (a) Normal processing (b) Controlled rolling of carbon-manganese steel (c) Controlled rolling of niobium-containing steel, finishing above Ac3 (d) Controlled rolling of niobium-containing steel, finishing below Ac3 Source: Ref 11 Controlled rolling is generally performed on microalloyed high-strength low-alloy (HSLA) steels, which have small amounts (generally less than 0.10%) of carbide- and nitride-forming elements such as niobium, titanium, and/or vanadium The various microalloying elements have different temperature-dependent solubility products (Ref 11, 13), and hot-rolling parameters must be adjusted to fit specific alloy compositions (see the section "Controlled Rolling" in the article "High-Strength Structural and High-Strength Low-Alloy Steels" in this Volume) References cited in this section 10 A.W Cramb, New Steel Casting Processes for Thin Slabs and Strip, Iron Steelmaker, Vol 15 (No 7), 1988, p 45-60 11 I Tamura, H Sekine, T Tanaka, and C Ouchi, Thermomechanical Processing of High-Strength Low-Alloy Steels, Butterworths, 1988 12 Thermomechanical Processing of Microalloyed Austenite, A.J DeArdo, G.A Ratz, and P.J Wray, Ed., The Metallurgical Society, 1982 13 Microalloyed HSLA Steels: Proceedings of Microalloying '88, ASM INTERNATIONAL, 1988 Processing: Ferritic Microstructure Large tonnages of hot-rolled steel strip are further processed to produce highly deformable sheet As shown in Fig 11, processing involves cold rolling to reduce thickness and improve surface quality, followed by annealing to produce microstructures consisting of ductile ferrite grains (Ref 14, 15) Cold-rolled and annealed sheet steels have low carbon contents, usually less than 0.10%, and therefore contain little pearlite in the slow-cooled hot-rolled condition However, if pearlite is present, it is deformed during cold rolling, and the cementite of the pearlite rapidly spheroidizes during annealing as the strained, cold-rolled ferrite recrystallizes to unstrained, equiaxed grains These microstructural changes are shown in Fig 12 for a 0.08% C steel containing 1.5% Mn and 0.21% Si (Ref 16) Fig 11 Temperature-time processing schedule for cold-rolled and annealed low-carbon sheet steel Continuous and batch annealing are schematically compared Source: Ref 14 Fig 12 Light micrographs of a 0.08C-1.5Mn-0.21Si steel (a) After cold rolling 50% (b) After cold rolling 50% and annealing at 700 °C (1290 °F) for 20 Source: Ref 16 A recently developed type of steel, made possible by the introduction of vacuum degassing into the steelmaking process, contains very low carbon, less than 0.008% (Ref 17) These steels are referred to as interstitial-free or ultralow-carbon steels and may contain small additions of niobium or titanium to tie up nitrogen and carbon residual from the steelmaking process Interstitial-free steels have excellent deep-drawing properties (see the article "High-Strength Structural and HighStrength Low-Alloy Steels" in this Volume) The carbon content of the interstitial-free steels is below that of the solubility limit of carbon in bcc ferrite (Fig and 13); therefore, no pearlite forms in these steels Also, the low interstitial content and the addition of stabilizing elements such as titanium or niobium eliminate strain aging and quench aging, as discussed below Fig 13 Iron-rich side of Fe-C diagram showing extent of ferrite-phase field and decrease of carbon solubility in ferrite with decreasing temperature Source: Ref The formability of cold-rolled and annealed sheet steels, especially in stamping operations that require deep drawing, is significantly improved by the development of crystallographic textures that defer necking and fracture in thin sheets to higher strains The preferred orientations, {111} planes parallel to the plane of the sheets and directions in the rolling direction (that is, {111} annealing textures), are promoted by aluminum deoxidation The aluminum-killed steels, if finish hot rolled at high temperatures and coiled at low temperatures, retain aluminum and nitrogen in solid solution in hot-rolled strip and through cold rolling (Ref 15) During batch annealing, aluminum nitride particles precipitate and suppress the nucleation and growth of recrystallized grains in orientations other than the preferred orientations for good formability The traditional method of annealing cold-rolled sheet has been to heat stacks of coils in a batch process (Fig 11) Batch annealing requires several days Recently, continuous annealing lines, in which the sheet is uncoiled and rapidly passed through high-temperature zones in continuous annealing furnaces, have been installed and used to anneal cold-rolled sheet steels (Ref 14) Continuous annealing requires only minutes to recrystallize a section of sheet as it passes through the hot zone of a furnace Although cold-rolled and annealed sheet steels have low carbon contents, during annealing at temperatures close to the A1 temperature, some carbon and nitrogen are always taken into solution (unless the steels are ultralow-carbon or interstitialfree steels) Figure 13 shows are carbon-rich side of the Fe-C diagram Carbon has its maximum solubility at the A1 temperature, and its solubility decreases with temperature to a negligible amount at room temperature Nitrogen shows a similar relationship Thus, if a steel is cooled from around A1 at a rate that prevents gradual relief of supersaturation by cementite formation during cooling, the ferrite at room temperature may be highly supersaturated with respect to carbon and nitrogen These interstitial elements then may segregate to dislocations in strained structures, a process referred to as strain aging, or they may precipitate out as fine carbide or nitride particles, a process referred to as quench aging (Ref 15) The aging processes may occur at room temperature or during heating at temperatures just above room temperature because of the high diffusivity of carbon and nitrogen in the bcc ferrite structure Strain aging and quench aging raise the yield strength of ferritic microstructures by pinning dislocations When yielding does occur, new dislocations are generated and the stress drops to a lower level at which localized plastic deformation propagates across a specimen The localized deformation is referred to as a Lüders band, and the process is described as discontinuous yielding In deformed sheet steels, Lüders bands are called stretcher strains and, if present, result in unacceptable surface appearance in formed parts In order to eliminate stretcher strains, cold-rolled and annealed sheet steels are temper rolled (Fig 11) The temper rolling introduces just enough strain to exceed the Lüders strain Beyond this point, sufficient dislocations are introduced, and all parts of a specimen or deformed sheet will strain uniformly or continuously Continuously annealed sheet steels are very susceptible to aging effects, because the thin sheet cools rapidly from the annealing temperature in contrast to batch annealing As a result, various types of overaging treatments are applied to continuously annealed steels, as shown in Fig 11 These treatments are designed to remove carbon and nitrogen from solid solution by the precipitation of relatively coarse carbide and nitride particles Most cold-rolled steels are subcritically annealed; that is, they are annealed below the A1 temperature However, continuous annealing lines have made possible intercritical heating into the ferrite-austenite field, with cooling that is rapid enough to cause the austenite to transform to martensite The martensite formation introduces a dislocation density exceeding that which can be pinned by the available carbon As a result, early yielding is continuous and occurs with high rates of strain hardening Intercritically annealed steels with ferrite-martensite microstructures are referred to as dualphase steels and offer another approach to producing high strength levels that range from 345 to 550 MPa (50 to 80 ksi) in low-carbon steels (Ref 18, 19, 20) Dual-phase steels are discussed in the article "High-Strength Structural and HighStrength Low-Alloy Steels" in this Volume References cited in this section G Krauss, Physical Metallurgy and Heat Treatment of Steel, in Metals Handbook Desk Edition, H.E Boyer and T.L Gall, Ed., American Society for Metals, 1985, p 28-2 to 28-10 14 P.R Mould, An Overview of Continuous-Annealing Technology, in Metallurgy of Continuous-Annealed Sheet Steel, B.L Bramfitt and D.L Mangonon, Jr., Ed., The Metallurgical Society, 1982, p 3-33 15 W.C Leslie, The Physical Metallurgy of Steels, McGraw-Hill, 1981 16 D.Z Yang, E.L Brown, D.K Matlock, and G Krauss, Ferrite Recrystallization and Austenite Formation in Cold-Rolled Intercritically Annealed Steel, Metall Trans A, Vol 11A, 1985, p 1385-1392 17 Metallurgy of Vacuum-Degassed Steel Products, R Pradhan, Ed., The Metallurgical Society, to be published in 1990 18 Structure and Properties of Dual-Phase Steels, R.A Kot and J.M Morris, Ed., The Metallurgical Society, 1979 19 Fundamentals of Dual-Phase Steels, R.A Kot and B.L Bramfitt, Ed., The Metallurgical Society, 1981 20 D.K Matlock, F Zia-Ebrahimi, and G Krauss, Structure, Properties and Strain Hardening of Dual-Phase Steels, in Deformation, Processing and Structure, G Krauss, Ed., ASM INTERNATIONAL, 1984 Martensite Martensite is the phase that produces the highest hardness and strength in steels (Fig 2) The martensitic transformation is diffusionless and occurs upon cooling at rates rapid enough to suppress the diffusion-controlled transformation of austenite to ferrite, pearlite, and bainite Neither the iron atoms nor the carbon atoms diffuse Therefore, the transformation occurs by shearing or the cooperative motion of large numbers of atoms Figure 14 shows schematically the formation of a martensitic crystal Macroscopically, the shears act parallel to a fixed crystallographic plane, termed the habit plane, and produce a uniformly tilted surface relief on a free surface Not only is the crystal structure change from austenite (fcc) to martensite (bcc) (referred to as the lattice deformation) accomplished by the transformation, but also the product martensite is simultaneously deformed because of the constraints created by maintaining an unrotated and undistorted habit plane within the bulk austenite (Ref 2, 21) The deformation of the martensite is referred to as the lattice invariant deformation, and it produces a high density of dislocations or twins in martensite This fine structure, together with the carbon atoms trapped within the octahedral interstitial sites of the body-centered tetragonal structure, produce the very high strength of as-quenched martensite (Ref 22) Fig 14 Schematic of shear and surface tilt associated with the formaton of a martensite plate Courtesy of M.D Geib Martensite begins to form at a critical temperature, defined as the martensite start (Ms) temperature The transformation is accomplished by the nucleation and growth of many crystals Because of the matrix constraints, the width of the martensitic units is limited, and the transformation proceeds primarily by the successive nucleation of new crystals This process occurs only upon cooling to lower temperatures and is therefore independent of time The latter type of transformation kinetics is termed athermal and is characterized by (Ref 23): f = - exp - [0.01] (Ms - Tq) (Eq 3) where f is the fraction of martensite formed after quenching to any temperature, Tq, below Ms Thus, the amount of martensite that forms at room temperature, for example, is a function only of Ms The Ms temperature is a function of the carbon and alloy content of steel, and a number of equations from which Ms can be calculated based on composition have been developed (Ref 2) Figure 15 shows that Ms decreases sharply with increasing carbon content in iron-carbon alloys Almost all other alloying elements also lower Ms A major effect of low Ms temperatures is incomplete martensite formation at room temperature Therefore, in all martensitic structures, some austenite is retained, the exact amount of which depends sensitively on composition Fig 15 Ms temperature as a function of carbon content in steels Composition ranges of lath and plate martensite in iron-carbon alloys are also shown Source: Ref 2; investigations referenced are identified by their number in that reference Two morphologies of martensitic microstructures form in iron-carbon alloys and steels (Fig 15) In low- and mediumcarbon alloys, lath martensite forms Lath martensite is characterized by parallel board or lath-shaped crystals The laths have an internal structure consisting of tangled dislocations, and the microstructure contains small amounts of retained austenite between the laths The groups of parallel laths are termed packets; many of the laths are too fine to be resolved in the light microscope Figure 16 shows an example of lath martensite in 4340 steel Fig 16 Light micrograph of lath martensite in 4340 steel quenched from 940 °C (1725 °F) and tempered at 350 °C (660 °F) Packets of parallel laths are below resolution of light microscope Source: Ref 24 In high-carbon steels, plate martensite forms The martensitic crystals have the shape of plates, and adjacent units tend to be nonparallel The fine structure associated with the plates often consists of fine transformation twins, and large amounts of retained austenite are present because of the low Ms temperatures Figure 17 shows plate martensite and austenite in an Fe-1.36C alloy Fig 17 Light micrograph of plate martensite and retained austenite in an Fe-1.39C alloy Source: Ref 25 Martensite can form only if the diffusion-controlled transformations of austenite can be suppressed On a practical level, this is accomplished by rapid quenching, for example, in water or brine baths However, such drastic cooling introduces high surface tensile residual stresses and may cause quench cracking Therefore, medium-carbon steels are alloyed with elements such as nickel, chromium, and molybdenum, which make it more difficult for the diffusion-controlled transformations to occur As a result, martensite can be formed with less drastic cooling, such as oil quenching The design of steels and cooling conditions to produce required amounts of martensite is the subject of the technology referred to as hardenability (Ref 26, 27) The application of hardenability concepts characterize not only the conditions that produce martensite but also those under which other microstructures form Thus, hardness gradients in bars of various diameters, cooled at various rates, can be estimated Continuous cooling diagrams, in which cooling conditions that produce various microstructures are defined for a given steel, are often related to hardness gradients measured on Jominy end-quenched specimens, as shown in Fig 18 Fig 18 Continuous transformation and isothermal transformation for steel containing nominally 0.4% C, 1.0% Cr, and 0.2% Mo Several cooling rates are related to positions and hardness on a Jominy end-quench specimen Source: Ref References cited in this section G Krauss, Physical Metallurgy and Heat Treatment of Steel, in Metals Handbook Desk Edition, H.E Boyer and T.L Gall, Ed., American Society for Metals, 1985, p 28-2 to 28-10 G Krauss, Steels: Heat Treatment and Processing Principles, ASM INTERNATIONAL, 1989 21 B.A Bilby and J.W Christian, The Crystallography of Martensite Transformations, Vol 197, 1961, p 122131 22 M Cohen, The Strengthening of Steel, Trans TMS-AIME, Vol 224, 1962, p 638-657 23 D.P Koistinen and R.E Marburger, A General Equation Prescribing the Extent of the Austenite-Martensite Transformation in Pure Iron-Carbon Alloys and Plain Carbon Steels, Acta Metall., Vol 7, 1959, p 59-60 24 J.P Materkowski and G Krauss, Tempered Martensite Embrittlement in SAE 4340 Steel, Metall Trans A, Vol 10A, 1979, p 1643-1651 25 A.R Marder, A.O Benscoter, and G Krauss, Microcracking Sensitivity in Fe-C Plate Martensite, Metall Trans , Vol 1, 1970, p 1545-1549 26 Hardenability Concepts with Applications to Steel, D.V Doane and J.S Kirkaldy, Ed., American Institute of Mining, Metallurgical, and Petroleum Engineers, 1978 27 C.A Siebert, D.V Doane, and D.H Breen, The Hardenability of Steels: Concepts, Metallurgical Influences, and Industrial Applications, American Society for Metals, 1977 Tempering of Martensite As-quenched martensite has very high strength, but has very low fracture resistance, or toughness Therefore, almost all steels that are quenched to martensite are also tempered, or heated, to some temperature below A1 in order to increase toughness Depending on time and temperature, tempering treatments can produce a wide variety of microstructures and properties As-quenched martensitic microstructures are supersaturated with respect to carbon, have high residual stresses, contain a high density of dislocations, have a very high lath or plate boundary area per unit volume, and contain retained austenite All these factors make martensitic microstructures very unstable and drive various phase transformations and microstructural changes during tempering Table lists the various reactions that develop during tempering The most important changes are a result of aging and precipitation phenomena, which are caused by the supersaturation of carbon, and range from carbon atom clustering to transition carbide precipitation to cementite formation and spheroidization Figure 19 shows a plate of high-carbon martensite tempered at 150 °C (300 °F) The dark basketweave structure is due to contrast associated with rows of very fine particles that are nm (0.08 μin.) in size, of the orthorhombic transition carbide, η The carbides must be imaged by other techniques (Ref 29, 30) Tempering at temperatures between 150 and 200 °C (300 and 390 °F) retains high hardness with increased toughness relative to as-quenched martensite Table Tempering reactions in steel Temperature range Reaction and symbol (if designated) Comments °C °F -40 to 100 -40 to 212 Clustering of two to four carbon atoms on octahedral sites of martensite (A1); segregation of carbon atoms to dislocations and boundaries Clustering is associated with diffuse spikes around fundamental electron diffraction spots of martensite 20 to 100 70 to 212 Modulated clusters of carbon atoms on (102) martensite planes (A2) Identified by satellite spots around electron diffraction spots of martensite 60 to 80 140 to 175 Long period ordered phase with ordered carbon atoms (A3) Identified by superstructure spots in electron diffraction patterns 100 to 200 212 to 390 Precipitation of transition carbide as aligned nm (0.08 μin.) diam particles (T1) Recent work identifies carbides as η(orthorhombic, Fe2C); earlier studies identified the carbides as ε(hexagonal, Fe2.4C) 200 to 350 390 to 660 Transformation of retained austenite to ferrite and cementite (T2) Associated with tempered-martensite embrittlement in low- and medium-carbon steels 250 to 700 480 to 1290 Formation of ferrite and cementite; eventual development of well-spheroidized carbides in a matrix of equiaxed ferrite grains (T3) This stage now appears to be initiated by χ-carbide formation in high-carbon Fe-C alloys 500 to 700 930 to 1290 Formation of alloy carbides in chromium-, molybdenum-, vanadium- and tungsten-containing steels The mix and composition of the carbides may change significantly with time (T4) The alloy carbides produce secondary hardening and pronounced retardation of softening during tempering or long-time service exposure around 500 °C (930 °F) 350 to 550 660 to 1020 Segregation and cosegregation of impurity and substitutional alloying elements Responsible for temper embrittlement Source: Ref 28 Fig 19 Transmission electron micrograph showing effects of transition carbide precipitation (dark contrast) in fine structure of a plate of martensite in an Fe-1.22C alloy tempered at 150 °C (300 °F) Source: Ref 29 Tempering at higher temperatures causes the formation of cementite, and, if strong carbide-forming elements are present, alloy carbides Concurrently, the laths or plates coarsen and the dislocation density is reduced by recovery mechanisms (Ref 2) In addition, retained austenite transforms to mixtures of cementite and ferrite between martensite laths and plates The carbides formed by high-temperature tempering are much coarser than the transition carbides and are present at residual martensite interfaces and dispersed within the ferrite of the tempered martensite (Fig 20) ... Shrinkage, % MPa ksi 1 0 -2 0 32 0-3 80 4 6-5 5 2- 5 5 0-5 2 1. 8 -2 .2 2 0-3 0 38 0-4 50 5 5-6 5 2- 6 4 8-5 0 2. 0 -2 .6 4 0-5 0 45 0-5 00 6 5-7 3 3-6 3 8-4 2 3. 2- 4 .6 Fig 11 Correlation between nodularity and elongation for ferritic... (21 2 °F) Spheroidal 400 °C (750 °F) 500 °C (930 °F) 3.8 50 .24 (29 . 02) 48.99 (28 .30) 45 .22 (26 . 12) 41.87 (24 .19) 38. 52 (22 .25 ) 53.39 (30.84) 50.66 (29 .27 ) 47.31 (27 .33) 43. 12 (24 .91) 38.94 (22 .49)... (22 .01) 41.0 (23 .69) 39.40 (22 .76) 37.30 (21 .55) 35 .20 (20 .34) 4.1 Compacted 300 °C (570 °F) 4.8 Flake 20 0 °C (390 °F) 43.54 (25 .15) 43. 12 (24 .91) 40.19 (23 .22 ) 37.68 (21 .77) 35.17 (20 . 32) 4.2