microstructures,processing and properties of steel.

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microstructures,processing and properties of steel.

Messaggioda Aldebaran » 05/05/2010, 15:28

Introduction
THE PERFORMANCE of steels depends on the properties associated with their microstructures, that is, on the
arrangements, volume fractions, sizes, and morphologies of the various phases constituting a macroscopic section of steel
with a given composition in a given processed condition. Because all the phases in steels are crystalline, steel
microstructures are made up of various crystals, sometimes as many as three or four different types, which are physically
blended by solidification, solid-state phase changes, hot deformation, cold deformation, and heat treatment. Each type of
microstructure and product is developed to characteristic property ranges by specific processing routes that control and
exploit microstructural changes. Thus, processing technologies not only depend on microstructure but are also used to
tailor final microstructures. For example, sheet steel formability depends on the single-phase ferritic microstructures of
low-carbon cold-rolled and annealed steel, while high strength and wear resistance are enhanced by carefully developed
microstructures of very fine carbides in fine martensite in fine-grain austenite of high-carbon hardened steels.
This article describes microstructures and microstructure-property relationships in steel. An important objective is to
relate microstructural evolution to the different processing schedules to which various types of steel are subjected. For
example, low-carbon sheet and plate steels are subjected to much different processing schedules and develop significantly
different microstructures and properties from those of medium-carbon forged and hardened steels. This article emphasizes
the correlation of microstructure and properties as a function of carbon content and processing in relatively low-alloy
steels. More highly alloyed steels, such as tool steels and stainless steels, are discussed in detail in the Section "Specialty
Steels and Heat-Resistant Alloys" in this Volume.
Iron-Carbon Phase Diagram
The major component of steel is iron, which exists in two crystal forms below its melting point. One is the body-centered
cubic (bcc) form, which is stable from below room temperature to 912 °C (1675 °F) and from 1394 °C (2540 °F) to the
melting point of 1530 °C (2785 °F). In the former temperature range, bcc iron is known as α-ferrite, while in the higher
temperature range, it is known δ-ferrite. The other crystal form, which is stable between 912 and 1394 °C (1675 and 2540
°F), is the face-centered cubic (fcc) form, known as austenite or γ-iron.
Steels also contain carbon in amounts ranging from very small, of the order of 0.005 wt% in ultralow-carbon, vacuumdegassed
sheet steels, to a maximum of 2.00 wt% in the highest-carbon tool steels. Carbon profoundly changes the phase
relationships, microstructure, and properties in steels. Generally, carbon content is kept low in steels that require high
ductility , high toughness, and good weldability, but is maintained at higher levels in steels that require high strength, high
hardness, fatigue resistance, and wear resistance.
Figure 1 shows the iron-carbon phase diagram and the changes that carbon induces in the phase equilibria of pure iron.
Carbon is an austenite stabilizer and expands the temperature range of stability of austenite. Its solubility is much higher
in austenite (a maximum of 2.11 wt% in equilibrium with cementite at 1148 °C, or 3000 °F) than in ferrite (a maximum of
0.0218 wt% in equilibrium with cementite at 727 °C, or 1340 °F). The solubility of carbon in ferrite and austenite is a
function of temperature; when the carbon atoms can no longer be accommodated in the octahedral interstitial sites
between the iron atoms, a new phase that can accommodate more carbon atoms in its crystal structure will form (Ref 2).
This phase is designated as cementite or iron carbide (Fe3C) and has an orthorhombic crystal structure. Cementite
formation and the temperature-dependent solubility of carbon in austenite and ferrite, as controlled by alloying and
processing, account for the great variety of microstructures and properties produced in steels.
Alloys of iron and carbon that contain up to 2.00 wt% C are classified as steels, while those containing over 2.00 wt% C
are classified as cast irons. Graphite is a more stable carbon-rich phase than cementite; its formation is promoted by a
high carbon concentration and the presence of large amounts of such elements as silicon. Therefore, graphite is an
important phase in cast irons but is rarely found in steels. When graphite does form, solubility limits and temperature
ranges of phase stability are changed slightly, as indicated by the dashed lines in Fig. 1.
The austenite phase field shown in Fig. 1 is the basis for the hot workability and heat treatability of carbon steels. Singlephase
austenite is readily hot worked; therefore, massive sections of steel can be hot reduced to smaller sections and
structural shapes. Austenite, upon cooling, must transform to other microstructures. If slow cooled under conditions
approximating equilibrium, austenite will change to mixtures of ferrite and cementite (Fig. 1); if cooled rapidly, it will
change to martensite. Such transformations provide the basis of the heat treatments applied to steels. Thus, the fortuitous
high-temperature stability of austenite in iron and iron-carbon alloys and the solid-state transformation of austenite upon
cooling create many opportunities to optimize shape, section size, microstructure, and properties for many different
applications.
Processing temperatures for the formation and transformation of austenite are set by the critical temperatures that mark
the boundaries between the various phase fields of Fig. 1. The critical temperatures as a function of carbon content, which
were initially identified by changes in slope or thermal arrests in heating and cooling curves, are given the designation
"A." If equilibrium conditions apply, the designations Ae1, Ae3 and Aecm, or simply A1, A3, and Acm, are used to indicate
the upper boundary of the ferrite-cementite phase field, the boundary between the ferrite-austenite and austenite phase
fields, and the boundary between the austenite and austenite-cementite phase fields (Fig. 1). If heating conditions apply
(which raise the critical temperatures relative to equilibrium), then Ac1, Ac3, and Accm are used for the critical
temperatures, with the c being derived from the French chauffant. If cooling conditions apply (which lower critical
temperatures relative to equilibrium), then the designations Ar1, Ar3, and Arcm are used, with the r being derived from the
French refroidisant. There is hysteresis in the transformation temperatures (that is, Ac temperatures are higher than Ae
temperatures and Ar temperatures are lower than Ae temperatures) because continuous heating and cooling leave
insufficient time for complete diffusion-controlled transformation at the true equilibrium temperatures.
In addition to iron and carbon, steels contain many other elements that shift the boundaries of the iron-carbon phase
diagram. Elements such as manganese and nickel are austenite stabilizers, which lower critical temperatures. Elements
such as silicon, chromium, and molybdenum are ferrite stabilizers and carbide formers, which raise critical temperatures
and shrink the austenite phase field (Ref 3). Other elements, such as titanium, niobium, and vanadium, may form
temperature-dependent dispersions of nitrides, carbides, or carbonitrides in the austenite. These effects must be taken into
account when setting processing temperature ranges for commercial alloys. For well-established grades, the optimum
temperature ranges for hot work are listed in Ref 4.
References cited in this section
1. 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
2. G. Krauss, Steels: Heat Treatment and Processing Principles, ASM INTERNATIONAL, 1989
3. J.S. Kirkaldy, B.A. Thompson, and E.A. Baganis, Prediction of Multicomponent Equilibrium and
Transformation Diagrams for Low Alloy Steels, in Hardenability Concepts with Applications to Steel , D.V.
Doane and J.S. Kirkaldy, Ed., The Metallurgical Society, 1978
4. Heat Treaters Guide, P.M. Unterweiser, H.E. Boyer, and J.J. Kubbs, Ed., American Society for Metals, 1982
Carbon Content and Properties
Figure 2 shows hardness as a function of carbon content for various types of microstructures. Hardness is readily
measured and generally is directly proportional to strength and inversely proportional to ductility and toughness.
Although the hardness-carbon relationships are shown as lines in Fig. 2, they are in fact better represented as bands
because many factors may cause variations in hardness in a given microstructure. For example, the strength of low-carbon
ferritic microstructures is very sensitive to grain size, while that of largely pearlitic microstructures is very sensitive to the
interlamellar spacing of cementite and ferrite.
All types of microstructures increase in strength with increasing carbon content, but martensitic microstructures show the
most dramatic increases. Because of the low solubility of carbon in ferrite (except for as-quenched martensite), the carbon
is primarily concentrated in carbide phases. Therefore, much of the higher strength of medium- and high-carbon steels is
due to higher volume fractions and finer dispersions of carbides in ferrite. Ferritic matrix grain sizes and morphology also
significantly affect mechanical behavior at any given carbon level.
Figure 2 shows that all types of microstructures could be produced in a steel of a given carbon content. There are,
however, practical limits to this observation. Low-carbon steels do not have sufficient hardenability to form martensite
except in the thinnest sections and are therefore produced primarily with ferritic microstructures, which have excellent
ductility for cold-working and forming operations. At the other extreme, medium- and high-carbon steels alloyed with
chromium, nickel, and/or molybdenum may have such high hardenability for the formation of martensite or bainite that
other microstructures are formed only by special annealing treatments.
The preceding comments should help explain why various types of steels have evolved based on the most readily
attainable microstructures and the property requirements that they satisfy for certain types of applications. Alloy design
and processing approaches have also evolved and are still evolving to exploit the best features of each type of steel. The
following sections in this article describe in detail several important ferrous microstructures and the steels and processing
methods used to produce them.
Reference cited in this section
1. 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
Pearlite and Bainite
Pearlite is the name given to the microstructure produced from austenite (A) during cooling of a steel by the following
solid-state reaction:
3 (0.77 % ) (0.02 % ) (6.67 % ) cooling
heating
A wt carbon ¬¾¾¾¾¾¾®F wt carbon + Fe C wt carbon (Eq 1)
Such a reaction, in which one solid phase transforms to two other solid phases, is generically termed a eutectoid reaction.
In the case of steels, the ferrite (F) and cementite (Fe3C) form as roughly parallel lamellae, or platelets, to produce a
composite lamellar two-phase structure. Figure 3 shows an example of pearlite formed in a rail steel that contains close to
the eutectoid carbon content, 0.77 wt%C. In this scanning electron micrograph the cementite lamellae appear light and
ferrite appears recessed because it has etched more deeply than the cementite.
The parallel lamellae act as diffraction gratings in the light microscope and diffract light of various wavelengths to
produce the colors and the luster characteristic of pearls. Thus, early metallographers used the name pearlite for the
unique lamellar structure they observed in steels. Within a given colony of pearlite, it has been shown that all the ferrite
and cementite have largely the same crystallographic orientations. Therefore pearlite may be described as two
interpenetrating single crystals (Ref 5).
Under equilibrium conditions, best approximated by very slow cooling, the pearlite reaction must occur at 727 °C (1340
°F), because austenite is not stable below that temperature . However, pearlite formation is accomplished by
diffusion, a time-dependent process. Carbon atoms diffuse away from regions that become ferrite to regions that become
cementite. Thus, rapid cooling, which gives less time for diffusion, depresses pearlite formation to lower temperatures.
Figure 4 shows an isothermal transformation diagram for 1080 steel, which contains 0.79 wt% C and therefore transforms
entirely to pearlite over a range of temperatures well below 727 °C (1340 °F). Figure 4 was produced by cooling the steel
rapidly to a series of temperatures below A1, holding at those temperatures, and then following, as a function of time, the
transformation of austenite to pearlite. An incubation time is required for the initiation of transformation, and this time, as
shown by the curve that marks the beginning of transformation, decreases with decreasing temperature. This acceleration
is associated with increased undercooling, which provides greater thermodynamic driving force, increasing nucleation
rates and decreasing interlamellar spacing. The latter, in turn, enhances growth rates because of reduced diffusion
distances.
The kinetics of pearlite formation, based on the nucleation and growth of spherical pearlite colonies, are described by
(Ref 6):
f(t) = 1 - 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 2 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.
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 5 and 6 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. 5 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. 6 Light micrograph showing lower bainite (dark plates or needles) formed in 4150 steel. Courtesy of F.A.
Jacobs.
References cited.
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.
shows a low-carbon steel ferrite-pearlite microstructure that formed during air cooling from the austenite phase
field . 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.

the ferrite appears white, and the boundaries between ferrite grains of different orientation appear as dark lines.
The pearlite, in contrast to, appears uniformly black because the interlamellar spacing of the pearlite in this
example is too fine to be resolved by the light microscope.
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.
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. 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.
The thermomechanical processing temperature ranges are shown in Fig. 9 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. 9. 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. 9. 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 Low-
Alloy 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. Similar controlled rolling schedules are used worldwide to produce fine-grain low-carbon steels. A processing parameter
introduced is the austenite recrystallization temperature (TR, which is primarily dependent on the amount of
deformation and alloying.
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.
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. for a 0.08% C steel containing 1.5% Mn and 0.21% Si (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 High-
Strength 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 ; 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.
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. For a 0.08% C steel containing 1.5% Mn and 0.21% Si (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 High-
Strength 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. 1 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.
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 <110> directions in the
rolling direction (that is, {111} <110> 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 High-
Strength Low-Alloy Steels" in this Volume.
References cited in this section
1. 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).
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 = 1 - 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, iin all martensitic structures, some
austenite is retained, the exact amount of which depends sensitively on composition.
Two morphologies of martensitic microstructures form in iron-carbon alloys and steels . 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.
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.
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.
References cited in this section
1. 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
2. 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 122-
131
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 1 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. The dark basketweave structure is due to
contrast associated with rows of very fine particles that are 2 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.
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 .
Toughness, or fracture resistance, generally increases with tempering temperature, but various types of enbrittlement or
reduced toughness can develop (Ref 2). Figure 22 shows impact toughness as a function of tempering temperature for
selected sets of steels with high and low levels of phosphorus. Carbon content has a major influence on toughness.
Medium-carbon tempered steels are quite tough, but high-carbon steels show very low impact toughness, which limits the
application of hardened and tempered high-carbon steels to conditions of compressive loading without impact, such as in
bearings. The effect of carbon on the toughness of low-temperature tempered specimens correlates with increasing
densities of transition carbides and associated high strain hardening rates as carbon content increasing (Ref 2).
Toughness reaches its peak in specimens tempered at 200 °C (390 °F); it drops to a minimum in specimens tempered
around 300 °C (570 °F). This drop is referred to as tempered martensite embrittlement and is associated with the
transformation of retained austenite to coarse carbide structures. Tempered martensite embrittlement is exacerbated by
phosphorus segregation to prior-austenite grain boundaries and carbide interfaces, but this effect appears to be constant
over the entire tempering range (Fig. 22). At higher tempering temperatures, between 350 and 550 °C (660 and 1020 °F),
another embrittlement phenomenon may develop in steels containing phosphorus, antimony, or tin (Ref 34). This
embrittlement is referred to as temper embrittlement, and requires long holding times or slow cooling through the
embrittling temperature range. Alloy steels are most susceptible, and the cosegregation of the alloying elements with the
impurities to prior austenite grain boundaries has been documented (Ref 35).
References cited in this section
2. G. Krauss, Steels: Heat Treatment and Processing Principles, ASM INTERNATIONAL, 1989
28. G. Krauss, Tempering and Structural Change in Ferrous Martensites, in Phase Transformations in Ferrous
Alloys, A.R. Marder and J.I. Goldstein, Ed., The Metallurgical Society, 1984
29. D.L. Williamson, K. Nakazawa, and G. Krauss, A Study of the Early Stages of Tempering in an Fe-1.22 pct
C Alloy, Metall. Trans. A, Vol 10A, 1979, p 1351-1363
30. Y. Hirotsu and S. Nagakura, Crystal Structure and Morphology of the Carbide Precipitated in Martensitic
High Carbon Steel During the First Stage of Tempering, Acta Metall., Vol 20, 1972, p 645-655
31. R.A. Grange, C.R. Hibral, and L.F. Porter, Hardness of Tempered Martensite in Carbon and Low Alloy
Steels, Metall. Trans. A, Vol 8A, 1977, p 1775-1785
32. D.L. Yaney, "The Effects of Phosphorus and Tempering on the Fracture of AISI 52100 Steel," M.S. thesis,
Colorado School of Mines, 1981
33. F. Zia-Ebrahimi and G. Krauss, Mechanisms of Tempered Martensite Embrittlement in Medium-Carbon
Steels, Acta Metall ., Vol 32, 1984, p 1767-1777
34. C.J. McMahon, Jr., Temper Brittleness: An Interpretive Review, in Temper Embrittlement in Steel, STP
407, American Society for Testing and Materials, 1968, p 127-167
35. M. Guttman, P. Dumonlin, and M. Wayman, The Thermodynamics of Interactive Co-Segregation of
Phosphorus and Alloying Elements in Iron and Temper-Brittle Steels, Metall. Trans. A, Vol 13A, 1982, p
1693-1711
Processing: Quenched and Tempered Microstructures
Hardened steels with tempered martensitic microstructures are most frequently used in machine components that require
high strength and excellent fatigue resistance under conditions of cyclic loading. Figure 23 shows a typical processing
sequence for these components. Hot-rolled bars are received and forged, generally at high temperatures where
deformation into complex shapes is readily accomplished. The forgings are air cooled, and ferrite-pearlite microstructures
develop upon cooling to room temperature. A normalizing treatment to refine the coarse microstructures that originated
because of high-temperature forging may be required, or a spheroidizing treatment to produce a microstructure of ferrite
and spheroidized cementite may be applied if extensive machining prior to hardening is required. The forgings are then
austenitized, quenched to martensite, and tempered to the properties described in the preceding section. Straightening and
stress relieving operations may be applied if required.
Processing: Direct-Cooled Forging Microstructures
To reduce the number of processing steps associated with producing quenched and tempered microstructures, new
alloying approaches have been developed to produce high-strength microstructures directly during cooling after forging.
Figure 24 shows a schematic of such a processing approach and an alternate processing sequence that cold finishes hotrolled
bars. Eliminating heat treatment processing steps by direct cooling relative to quenching and tempering has obvious
advantages.
One group of steels that has been developed for direct cooling is microalloyed medium-carbon steels (see Ref 36, 37 and
the article "High-Strength Low-Alloy Steel Forgings" in this Volume). These steels contain small amounts of vanadium
and niobium and transform to precipitation-hardened microstructures of ferrite and pearlite. The hardness produced by
rapid air cooling ranges from 25 to 30 HRC depending on the extent of precipitation and pearlite in the microstructure;
ultimate strength values are over 690 MPa (100 ksi). Thus, the hardness and strength levels are not as high as can be
produced by quenching and low-temperature tempering, but they are more than adequate for many automotive
applications that require intermediate strengths (Ref 36).
The fatigue resistance of direct-cooled microalloyed steels is comparable to that of quenched and tempered steels of the
same hardness, but the impact toughness is much lower. This reduced toughness is due to the well-known increase in the
ductile-to-brittle temperature in steels with ferrite-pearlite microstructures as pearlite content increases (Fig. 25). In order
to improve the toughness of direct-cooled forging steels, steels that transform to bainitic structures and forging steels with
lower carbon concentrations and finer ferrite-pearlite microstructures are being developed (Ref 38).
References cited in this section
1. 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
36. Fundamentals of Microalloying Forging Steels, G. Krauss and S.K. Banerji, Ed., The Metallurgical Society,
1987
37. G. Krauss, Microalloyed Bar and Forging Steels, in 29th Mechanical Working and Steel Processing
Conference Proceedings , Vol XXV, Iron and Steel Society, 1988, p 67-77
38. K. Grassl, S.W. Thompson, and G. Krauss, "New Options for Steel Selection for Automotive Applications,"
SAE Technical Paper 890508, Society of Automotive Engineers, 1989
Summary
This article has briefly described the major microstructures and the phase transformations by which these microstructures
are developed in carbon and low-alloy steels. Each type of microstructure and product is developed to characteristic
property ranges by specific processing routes that control and exploit microstructural changes. The incorporation of steel
carbon content into microstructure has a profound effect on microstructure and properties, and steels fall naturally into
low-strength/high ductility/high toughness or high-strength/high fatigue resistant/low toughness groups with increasing
carbon content. The use of new casting techniques, microalloying, and thermomechanical processing are being used
increasingly to reduce processing steps and to improve steel product microstructures and quality.
References
1. 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
2. G. Krauss, Steels: Heat Treatment and Processing Principles, ASM INTERNATIONAL, 1989
3. J.S. Kirkaldy, B.A. Thompson, and E.A. Baganis, Prediction of Multicomponent Equilibrium and
Transformation Diagrams for Low Alloy Steels, in Hardenability Concepts with Applications to Steel ,
D.V. Doane and J.S. Kirkaldy, Ed., The Metallurgical Society, 1978
4. Heat Treaters Guide, P.M. Unterweiser, H.E. Boyer, and J.J. Kubbs, Ed., American Society for Metals,
1982
5. 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
6. W.A. Johnson and R.F. Mehl, Reaction Kinetics in Processes of Nucleation and Growth, Trans. AIME,
Vol 135, 1939, p 416-458
7. 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
8. R.F. Hehemann, Ferrous and Nonferrous Bainite Structures, in Metals Handbook, 8th ed., Vol 8, American
Society for Metals, 1973, p 194-196
9. B.L. Bramfitt and J.G. Speer, A Perspective on the Morphology of Bainite, Metall. Trans. A, to be
published in 1990
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
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
21. B.A. Bilby and J.W. Christian, The Crystallography of Martensite Transformations, Vol 197, 1961, p 122-
131
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
28. G. Krauss, Tempering and Structural Change in Ferrous Martensites, in Phase Transformations in Ferrous
Alloys, A.R. Marder and J.I. Goldstein, Ed., The Metallurgical Society, 1984
29. D.L. Williamson, K. Nakazawa, and G. Krauss, A Study of the Early Stages of Tempering in an Fe-1.22
pct C Alloy, Metall. Trans. A, Vol 10A, 1979, p 1351-1363
30. Y. Hirotsu and S. Nagakura, Crystal Structure and Morphology of the Carbide Precipitated in Martensitic
High Carbon Steel During the First Stage of Tempering, Acta Metall., Vol 20, 1972, p 645-655
31. R.A. Grange, C.R. Hibral, and L.F. Porter, Hardness of Tempered Martensite in Carbon and Low Alloy
Steels, Metall. Trans. A, Vol 8A, 1977, p 1775-1785
32. D.L. Yaney, "The Effects of Phosphorus and Tempering on the Fracture of AISI 52100 Steel," M.S. thesis,
Colorado School of Mines, 1981
33. F. Zia-Ebrahimi and G. Krauss, Mechanisms of Tempered Martensite Embrittlement in Medium-Carbon
Steels, Acta Metall ., Vol 32, 1984, p 1767-1777
34. C.J. McMahon, Jr., Temper Brittleness: An Interpretive Review, in Temper Embrittlement in Steel, STP
407, American Society for Testing and Materials, 1968, p 127-167
35. M. Guttman, P. Dumonlin, and M. Wayman, The Thermodynamics of Interactive Co-Segregation of
Phosphorus and Alloying Elements in Iron and Temper-Brittle Steels, Metall. Trans. A, Vol 13A, 1982, p
1693-1711
36. Fundamentals of Microalloying Forging Steels, G. Krauss and S.K. Banerji, Ed., The Metallurgical
Society, 1987
37. G. Krauss, Microalloyed Bar and Forging Steels, in 29th Mechanical Working and Steel Processing
Conference Proceedings , Vol XXV, Iron and Steel Society, 1988, p 67-77
38. K. Grassl, S.W. Thompson, and G. Krauss, "New Options for Steel Selection for Automotive
Applications," SAE Technical Paper 890508, Society of Automotive Engineers, 1989
Introduction
STEELS constitute the most widely used category of metallic material, primarily because they can be manufactured
relatively inexpensively in large quantities to very precise specifications. They also provide a wide range of mechanical
properties, from moderate yield strength levels (200 to 300 MPa, or 30 to 40 ksi) with excellent ductility to yield strengths
exceeding 1400 MPa (200 ksi) with fracture toughness levels as high as 110 MPa m (100 ksi in ).
This article will review the various systems used to classify carbon and low-alloy steels*, describe the effects of alloying
elements on the properties and/or characteristics of steels, and provide extensive tabular data pertaining to designations of
steels (both domestic and international). More detailed information on the steel types and product forms discussed in this
article can be found in the articles that follow in this Section.
Note
* The term low-alloy steel rather than the more general term alloy steel is being used to differentiate the steels
covered in this article from high-alloy steels. High-alloy steels include steels with a high degree of fracture
toughness (Fe-9Ni-4Co), which are described in the article "Ultrahigh-Strength Steels" in this Section of the
Handbook. They also include maraging steels (Fe-18Ni-4Mo-8Co), austenitic manganese steels (Fe-1C-
12Mn), tool steels, and stainless steels, which are described in separate articles in the Section "Specialty
Steels and Heat-Resistant Alloys" in this Volume.
Classification of Steels
Steels can be classified by a variety of different systems depending on:
· The composition, such as carbon, low-alloy, or stainless steels
· The manufacturing methods, such as open hearth, basic oxygen process, or electric furnace methods
· The finishing method, such as hot rolling or cold rolling
· The product form, such as bar, plate, sheet, strip, tubing, or structural shape
· The deoxidation practice, such as killed, semikilled, capped, or rimmed steel
· The microstructure, such as ferritic, pearlitic, and martensitic (Fig. 1)
· The required strength level, as specified in ASTM standards
· The heat treatment, such as annealing, quenching and tempering, and thermomechanical processing
· Quality descriptors, such as forging quality and commercial quality
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