Bhadeshia H.K.D.H., Honeycombe R. Steels: Microstructure and Properties

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218 CHAPTER 10 THERMOMECHANICAL TREATMENT OF STEELS

stuck at about 1 µm. The reason is recalescence, which is the rise in tempera-

ture of the steel caused by release of the latent heat of transformation at a rate

which is so high that it cannot easily be dissipated by diffusion. It causes the

temperature of the steel to rise, thus reducing G

γα

and preventing the achieve-

ment of ultraﬁne grain structures. Large-scale thermomechanical processing is

therefore limited by recalescence and is unlikely to lead to grain sizes which are

uniformly less than about 1 µm.

10.2.4 Dispersion strengthening during controlled rolling

The solubility data imply that, in a micro-alloyed steel, carbides and carbo-

nitrides of Nb, Ti and V will precipitate progressively during controlled rolling

as the temperature falls. While the primary effect of these ﬁne dispersions is

to control grain size, dispersion strengthening will take place. The strength-

ening arising from this cause will depend both on the particle size r, and the

interparticle spacing  which is determined by the volume fraction of precipi-

tate (Equation (2.10)). These parameters will depend primarily on the type of

compound which is precipitating, and that is determined by the micro-alloying

content of the steel. However the maximum solution temperature reached and

the detailed schedule of the controlling rolling operation are also important

variables.

It is now known, not only that precipitation takes place in the austenite,

but that further precipitation occurs during the transformation to ferrite. The

precipitation of niobium, titanium and vanadium carbides has been shown to

take place progressively as the interphase boundaries move through the steel.

This is the interphase precipitation discussed in Section 4.4.3. As this precipi-

tation is normally on an extremely ﬁne scale occurring between 850

◦

C and

650

◦

C, it is likely to be the major contribution to the dispersion strengthening.

In view of the higher solubility of vanadium carbide in austenite, the effect

will be most pronounced in the presence of this element, with titanium and

niobium in decreasing order of effectiveness. If the rate of cooling through

the transformation is high, leading to the formation of supersaturated plates

of ferrite, the carbides will tend to precipitate within the grains, usually on the

dislocations which are numerous in this type of ferrite.

In arriving at optimum compositions of micro-alloyed steels, it should be

borne in mind that the maximum volume fraction of precipitate which can be

put into solid solution in austenite at high temperatures is achieved by use

of stoichiometric compositions. For example, if titanium (atomic weight 47.9)

is used, it will combine with approximately one quarter its weight of carbon

(atomic weight 12), so that for a 0.025 wt% C steel, 0.10 wt% of Ti will provide

carbide of the stoichiometric composition. In Fig. 10.9 the stoichiometric line

for TiC is shown superimposed on the solubility curves for titanium carbide at

1100

◦

C,1200

◦

C and 1300

◦

C. If the precipitation in steels with 0.10 wt% titanium

cooled from 1200

◦

C is considered, at low carbon contents, i.e. to the left of the

10.2 CONTROLLED ROLLING OF LOW-ALLOY STEELS 219

Fig. 10.9 Effect of stoichiometry on the precipitation ofTiC in a micro-alloyed steel (Gladman

et al., Micro-alloying 75, Union Carbide corporation, New York, USA, 1975).

stoichiometric line, the carbide fraction is limited by the carbon content, i.e.

zone A, lower diagram. For carbon contents between the stoichiometric line

and the solubility line at 1200

◦

C, the full potential volume fraction of ﬁne TiC

will form on cooling (zone B). When the carbon content exceeds the solubility

limit (>0.10 wt%),the titanium is progressively precipitated at 1200

◦

C as coarse

carbide, thus reducing the amount of titanium available to combine with carbon

to form ﬁne TiC during cooling. As coarse carbide particles are ineffective in

220 CHAPTER 10 THERMOMECHANICAL TREATMENT OF STEELS

controlling grain growth, it is highly desirable to have steel compositions which

avoid their formations. It also follows from Fig. 10.9 that high austenitizing

temperatures are essential to obtain full beneﬁt from the precipitation of ﬁnely

divided carbide phases.

10.2.5 Strength of micro-alloyed steels: an overall view

In modern control-rolled micro-alloyed steels, there are at least three strength-

ening mechanisms which contribute to the ﬁnal strength achieved. The relative

contribution from each is determined by the composition of the steel and,

equally important, the details of the thermomechanical treatment to which

the steel is subjected. The several strengthening contributions for steels with

0.2 wt% carbon, 0.2 wt% silicon, 0.15 wt% vanadium and 0.015 wt% nitrogen

as a function of increasing manganese content are shown schematically in

Fig. 10.10. Firstly, there are the solid solution strengthening increments from

manganese, silicon and uncombined nitrogen. Secondly, the grain size contribu-

tion to the yield stress is shown as a very substantial component, the magnitude

of which is very sensitive to the detailed thermomechanical history. Finally, a

typical increment for dispersion strengthening is shown. The total result is a

range of yield strengths between about 350 and 500 MN m

−2

. In this particular

example, the steel was normalized (air cooled) from 900

◦

C,but had it been con-

trol rolled down to 800

◦

C or even lower, the strength levels would have been

substantially raised.

The effect of the ﬁnishing temperature for rolling is important in determin-

ing the grain size and, therefore, strength level reached for a particular steel.

It is now becoming common to roll through the transformation into the com-

pletely ferritic condition, and so obtain ﬁne subgrain structures in the ferrite,

which provide an additional contribution to strength. Alternatively, the rolling

is ﬁnished above the γ/α transformation, and the nature of the transformation is

altered by increasing the cooling rate. Slow rates of cooling obtained by coiling

at a particular temperature will give lower strengths than rapid rates imposed

by water spray cooling following rolling. The latter route can change the ferrite

from equi-axed to Widmanstätten with a much higher dislocation density. The

result is a steel with improved mechanical properties and, in many cases, the

sharp yield point can be suppressed. This has practical advantages in fabrica-

tion of sheet steel, e.g. pipe manufacture, where a continuous stress–strain curve

is preferred.

10.3 DUAL-PHASE STEELS

The HSLA steels described in Section 10.2 give improved strength to weight

ratios over ordinary standard steels. However, they are not readily formed,

e.g. by cold pressing and related techniques. The worldwide demand for safety

10.3 DUAL-PHASE STEELS 221

Fig. 10.10 The contributions to strength in a 0.2C–0.15V wt% steel as a function of Mn

content (Gladman et al., Micro-alloying 75, Union Carbide Corporation, 1975).

and fuel economy in automobiles has led to the development of a number of

steel types which are not only strong, but at the same time have the formability

required for mass production of car bodies and components. A measure of

formability is the product of strength and uniform elongation.

222 CHAPTER 10 THERMOMECHANICAL TREATMENT OF STEELS

(a)

(b)

Fig. 10.11 (a)Typical microstructure of dual-phase steel, consisting of a mixture of martensite

(dark) and ferrite. (b) Schematic stress–strain curves comparing the behaviour of a conventional

automobile steel with that of a dual-phase steel.

The dual-phase steels are low-alloy steels which satisfy these requirements

by exploiting microstructures in which there are two major phases (Fig. 10.11),

one of which is soft and the other signiﬁcantly harder. The ferrite–martensite

dual-phase steels typically contain manganese and silicon, and are strong and

yet are formable. They exhibit continuous yielding, i.e. no sharp yield point,

and a relatively low proof strength (300–350 MN m

−2

). The simplest steels in

this category contain 0.08–0.2C, 0.5–1.5Mn wt%, but steels micro-alloyed with

vanadium are also suitable, while small additions of Cr (0.5 wt%) and Mo (0.2–

0.4 wt%) are frequently used to control the development of microstructure.

The simplest way of achievinga duplex structure is to use intercritical anneal-

ing in which the steel is heated into the (α +γ) region between the Ae

and Ae

and held, typically at 790

◦

C for several minutes to allow small regions of austen-

ite to form in the ferrite. As it is essential to transform these regions of austenite

into martensite, cooling to ambient temperature must be sufﬁciently rapid to

10.4 TRIP-ASSISTED STEELS 223

avoid other intervening transformations. Alternatively, the hardenability of the

austenite must be enhanced by adding between 0.2 and 0.4 wt% Mo to a steel

already containing 1.5 wt% manganese. The required microstructure can then

be obtained by air cooling after intercritical annealing.

To eliminate an extra heat treatment step, dual-phase steels have now been

developed which can be given the required structure during cooling after con-

trolled rolling. Typically, these steels have additions of 0.5Cr and 0.4Mo wt%.

After completion of hot rolling around 870

◦

C, the steel forms approximately

80% ferrite on the water cooled run-out table from the mill. The material is

then coiled in the metastable region (510–620

◦

C) below the pearlite/ferrite

transformation and, on subsequent cooling, the austenite regions transform to

martensite.

10.4 TRIP-ASSISTED STEELS

The steels developed to exploit the properties obtained when the martensite

reaction occurs duringplastic deformation areknown astransformation-induced

plasticity (TRIP) steels. They are strong and exhibit considerable uniform elon-

gation before failure. There are several varieties of such steels. Those which are

made fully austenitic by using large quantities of austenite-stabilizing solutes,

but transform to martensite when stressed, are simply called the TRIP steels

(discussed in Chapter 12). When the austenite is a minor phase in the overall

microstructure, but undergoes martensitic transformation during straining, the

steels are said to be TRIP assisted and are usually low alloy steels.

Martensitic transformation induced by local stress has the effect of relieving

stress concentrations,increasing the work-hardening rate, and promoting homo-

geneous deformation, with consequent improvements in the strength, ductility

and toughness of steels. TRIP-assisted steels are mass produced, made using a

complex heat treatment which is often completed within a short time during the

processing of steel strip. Their microstructure consists of allotriomorphic fer-

rite as the major phase together with a total of 30–40% of harder regions. The

latter consist of mixtures of bainite, martensite and carbon-enriched retained

austenite. The chemical composition is typically Fe–0.12C–1.5Si–1.5Mn wt%.

Some austenite is retained in spite of the low overall solute content because

when the bainite forms, the silicon prevents cementite precipitation, thereby

enriching the residual austenite with carbon (Chapter 6). The major application

of TRIP-assisted steels is in the automobile industries, both for painted surfaces

and for enhancing the safety of the passenger compartment in the event of a

crash.

There are two kinds of TRIP-assisted steels. In the ﬁrst case a cold-rolled

strip is heated rapidly from ambient temperature for intercritical treatment in

the α +γ phase ﬁeld between the Ac

and Ac

temperatures (Fig. 10.12). The

intercritical annealing induces partial transformation to austenite and at the

224 CHAPTER 10 THERMOMECHANICAL TREATMENT OF STEELS

Fig. 10.12 The two kinds of heat treatment used to generate the microstructures of

TRIP-assisted steels. The terms γ, α, α

and α

′

represent austenite, allotriomorphic ferrite,

bainite and martensite, respectively.

same time recrystallizes the residual ferrite. The strip is then cooled at a con-

trolled rate during which some of the austenite transforms into allotriomorphic

ferrite and at lower temperatures into bainitic ferrite. This latter reaction causes

the austenite to become enriched in carbon, allowing it to be retained to ambient

temperature (Fig. 10.13).

The details ofthe microstructureand mechanical propertiescan be alteredby

manipulating the cooling condition. For example, it is common practice to allow

more time in the bainite transformation range than at the higher temperatures

where allotriomorphic ferrite grows. Figure 10.14 shows the effect of holding in

the bainite transformation temperature range on the ﬁnal microstructure. An

inadequate amount of bainite leaves the austenite susceptible to martensitic

transformation. Similarly, because the carbon concentration in the austenite is

limited by the T

curve (Chapter 6), transformation to bainite at too high a

temperature also renders the austenite unstable.

The second kind of heat treatment starts from a hot-rolled strip which is

fully austenitic (Fig. 10.12) and forms both allotriomorphic ferrite and bainite

during the cooling part of the thermal cycle. This has the advantage that the

microstructure can be produced directly from the hot strip which has been rolled

to its ﬁnal dimensions. The process is cheap since the strip does not have to be

heated to the intercritical annealing temperature. However, hot-rolling mills

are restricted by rolling loads to strips thicker than about 3 mm, although there

are modern mills which can cope with 1.4 mm thickness. Cold-rolled strips can,

on the other hand, be made routinely into thinner gauges. Hot-rolled strips are

preferred for automobile applications where cost is a prime factor in the choice

of materials.

The transformation strain due to the formation of martensite does not

account for the observed uniform tensile elongation of some 15–30%. The

shape deformation due to martensite (Chapter 5) is at most equivalent to a

2% tensile strain because the amount of austenite available for transformation

10.4 TRIP-ASSISTED STEELS 225

(a)

(b)

(c)

Fig. 10.13 TRIP-assisted steel showing a mixed microstructure of allotriomorphic ferrite,

bainitic ferrite and retained austenite films. (a) Optical micrograph. (b) Bright field transmission

electron micrograph. (c) Dark field image of retained austenite.

is quite small in TRIP-assisted steels. The major contributions to uniform elon-

gation arise partly from the enhanced work-hardeningcoefﬁcient ofthe material

due to the progressive formation of hard martensite during deformation. There

is a further signiﬁcant contribution from dislocations induced into the ferrite

226 CHAPTER 10 THERMOMECHANICAL TREATMENT OF STEELS

Fig. 10.14 Evolution of room temperature microstructure as a function of the time dur-

ing isothermal transformation to bainite (Girault, E., Mertens, A., Jacques, P., Houbaert, Y.,

Verlinden, B. and van Humbeek, J., Scripta Materialia 44, 885, 2001).

by the strains associated with martensitic transformation.

These dislocations

strengthen the ferrite and are visible in Fig. 10.13.

The austenite also delays the necking process during a tensile test by trans-

forming to martensite at stress concentrations. It is therefore important to delay

the transformation of retained austenite to the late stages of deformation when

signiﬁcant damage accumulates in the steel. It is at this point that the TRIP

effect can be most beneﬁcial. It is useful therefore to examine further the

transformation of austenite as a function of plastic strain.

It is reasonable to assume that the change in the fraction ofmartensite (dV

′

)

obtained for a given increment of plastic strain (dǫ) should be proportional to

the fraction of remaining austenite:

′

dǫ

= k

, (10.1)

where k

is a function of the steel composition and test temperature and V

the fraction of austenite remaining untransformed. If the fraction of austenite at

zero strain is V

, then V

′

−V

, and integration of Equation (10.1) gives:

ln{V

}−ln{V

}=k

ǫ.

The form of this equation is illustrated in Fig. 10.15.

10.4.1 Low- or zero-silicon TRIP-assisted steels

The substantial silicon addition to TRIP-assisted steels leads to the formation

of hard, adherent oxide (Fe

SiO

) which is difﬁcult to remove prior to hot

Jacques, P., Furnemont, Q., Mertens, A. and Delannay, F. Philosophical Magazine A 81,

1789, 2001.

10.4 TRIP-ASSISTED STEELS 227

Fig. 10.15 Martensitic transformation of retained austenite in a TRIP-assisted steel as a func-

tion of deformation temperature and plastic strain (Sherif, M., Garcia-Mateo, C., Sourmail, T.

and Bhadeshia, H. K. D. H. Materials Science and Technology 20, 319, 2004).

rolling, resulting in a poor surface ﬁnish. This restricts automobile applications

to components which are hidden from view. Low-silicon TRIP steels do exist,

as illustrated in Fig. 10.16. The austenite in such alloys is more difﬁcult to retain

because of the tendency to precipitate cementite, but this can be minimized with

careful heat treatment. When this is done, strengths in excess of 600 MPa with

a uniform ductility of 15% have been achieved.

10.4.2 Galvanizing of TRIP-assisted steels

There are two basic methods of galvanizing, by dipping the steel in liquid zinc

or by electrolytically depositing the zinc. The typical concentrations of silicon

and manganese in TRIP-assisted steels lead to a stable Mn

SiO

oxide ﬁlm on

the surface during the heat treatment that leads to the desired microstructure.

This makes it difﬁcult for the zinc to wet the steel surface, making it necessary

to electrolytically galvanize such alloys.

The problem can be alleviated by increasing the humidity in the annealing

furnace. The oxide coverage of the surface is then reduced, giving better wet-

ting by zinc. The higher humidity causes the internal oxidation of Mn and Si

below the steel surface, thus reducing their availability to form Mn

SiO

at the

surface.

An alternative approach is to eliminate the silicon and add aluminium to

retain the cementite-free microstructure. Aluminium oxidizes easily to form

alumina by internal oxidation near the surface, again limiting the amount of

FeAl

that can form at the free surface when the humidity in the annealing

furnace is low. Such a steel can easily be hot-dip galvanized. At high humidity,