Bhadeshia H.K.D.H., Honeycombe R. Steels: Microstructure and Properties
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208 CHAPTER 9 THE TEMPERING OF MARTENSITE
Irvine, K. J. and Pickering, F. B., High strength 12% chromium steels, in Iron and Steel Institute
Special Report No. 86, London, UK, p. 34, 1964.
Krauss, G., Phase Transformation in Ferrous Alloys (eds Marder, A. R. and Goldstein, J. I.),
TMS-AIME,Warrendale, PA, pp. 101–1123, 1984.
Krauss, G., Steels: Heat Treatment and Processing Principles, ASM International, Ohio, USA,
1990.
Kurdjumov, G., Journal of the Iron and Steel Institute 195, 26, 1960.
Leslie,W.C.,The Physical Metallurgy of Steels, McGraw-Hill, Tokyo, Japan, 1981.
Nutting, J., Physical metallurgy of alloy steels, Journal of the Iron and Steel Institute 207, 872,
1969.
Pickering,F. B.,Physical Metallurgy and the Design of Steels,Applied Science Publishers, 1978.
Roberts,C.S.,Averbach,B. L. and Cohen,M.,Transactions of theAmerican Society of Materials
45, 576, 1953.
Speich, G. R.,Tempering of plain carbon martensite,Transactions of the American Institute of
Mining, Metallurgical and Petroleum Engineers 245, 2553, 1969.
Speich, G. R., Symposium Iron and Steel Soc., AIME, Warrendale, PA, 1992.
Speich, G. R. and Clark, J. B. (eds), Precipitation from Iron-base Alloys, Gordon and Breach,
London, UK, 1965.
Stewart, J., Thomson R. C. and Bhadeshia, H. K. D. H., Cementite precipitation during
tempering of martensite under stress, Journal of Materials Science 29, 6079, 1994.
Thomas, G., Retained austenite and tempered martensite embrittlement. Met. Trans. 9, 439,
1978.
Winchell, G., Symposium on ‘The Tempering of Steel’, Metallurgical Transactions A, 991–1145,
1983.
Winchell, P. G. and Cohen, M.,Trans. ASM 55, 347, 1962.
Woodhead, J. H. and Quarrell,A. G.,The Role of Carbides in Low Alloy Creep-resisting Steels,
Climax Molybdenum Co., Michigan, USA, 1965.
10
THERMOMECHANICAL TREATMENT OF
STEELS
10.1 INTRODUCTION
Thermomechanical treatment involves the simultaneous application of heat and
a deformation process to an alloy, in order to change its shape and refine the
microstructure. Thus, hot rolling of metals, a well-established industrial pro-
cess, is a thermomechanical treatment which plays an important part in the
processing of many steels, particularly those produced in very large quantities.
Continuously cast segments of steel, ranging from 1 to 50 tonnes in weight,
are introduced into the rolling sequence at a temperature typically in the range
1200–1300
◦
C.Theyare then progressively rolled into a variety of shapes depend-
ing on application. The deformation leads to a breaking down of the original
coarse microstructure by repeated recrystallization of the steel while in the
austenitic condition, and by the gradual reduction in the chemical segrega-
tion introduced during casting. Also, the inevitable non-metallic inclusions, i.e.
oxides, sulphides and silicates, are broken up, some deformed and distributed
throughout the steel in a more refined and uniform manner.
The hot rolling process is now a highly controlled operation in which a bil-
lion tonnes of steel is produced annually using computer-controlled arrays of
equipment, resulting in impressive levels of productivity and reproducibility.
The compositions of the low-alloy steels are carefully chosen to provide opti-
mum mechanical properties when the hot deformation and subsequent cooling
is complete. This process, in which the rolling parameters (temperature, strain,
number of rolling passes, finishing temperature, etc.) are predetermined and
precisely defined, is called controlled rolling. It is now of the greatest import-
ance in obtaining reliable mechanical properties in steels for pipelines, bridges,
buildings and a huge variety of other products.
209
210 CHAPTER 10 THERMOMECHANICAL TREATMENT OF STEELS
10.2 CONTROLLED ROLLING OF LOW-ALLOY STEELS
10.2.1 General
Before the Second World War, strength in hot-rolled low-alloy steels was
achieved by the addition of carbon up to 0.4 wt% and manganese up to 1.5 wt%,
giving yield stresses of 350–400 MN m
−2
. However, such steels are essentially
ferrite–pearlite aggregates, which do not possess adequate toughness for many
modern applications. Indeed, the toughness, as measured by the ductile/brittle
transition, decreases dramatically with carbon content, i.e. with increasing vol-
ume of pearlite in the steel (Fig. 10.1). Furthermore, with the introduction
of welding as the main fabrication technique, the high carbon contents led
to serious cracking problems, which could only be eliminated by the use of
lower-carbon steels. The great advantage of producing in these steels a fine fer-
rite grain size soon became apparent (see Section 2.5), so controlled rolling
in the austenitic condition was gradually introduced to achieve this. Fine fer-
rite grain sizes in the finished steel were found to be greatly expedited by the
addition of small concentrations (<0.1 wt%) of grain refining elements such as
niobium, titanium and vanadium, and also aluminium. On adding such elements
to steels with 0.03–0.008%C and up to 1.5 wt% Mn, it became possible to pro-
duce fine-grained material with yield strengths between 450 and 550 MN m
−2
,
and with ductile/brittle transition temperatures as low as −70
◦
C. Such steels
are now referred to as high-strength low-alloy (HSLA) steels, or micro-alloyed
Fig. 10.1 Effect of carbon content on the impact transition temperature curves of ferrite/
pearlite steels (Pickering, in Micro-alloying 75, Union Carbide Corporation, New York, USA,
1975).
10.2 CONTROLLED ROLLING OF LOW-ALLOY STEELS 211
steels. This progress, from the relatively low strength of ordinary mild steel
(220–250 MN m
−2
) in a period of 20 years represents a major metallurgical
development, the importance of which, in engineering applications, cannot be
overstated.
The general features of controlled rolling are summarized in Fig. 10.2.
10.2.2 Grain size control during controlled rolling
The primary grain refinement mechanism in controlled rolling is the recrystal-
lizationof austeniteduringhot deformation,known as dynamic recrystallization.
This process is clearly influenced by the temperature and the degree of defor-
mation which takes place during each pass through the rolls. However, in
austenite devoid of second-phase particles, the high temperatures involved in
Fig. 10.2 The variety of thermomechanical processing routes. Controlled rolling, followed by
accelerated cooling is often designated thermomechanical controlled processing or TMCP.
212 CHAPTER 10 THERMOMECHANICAL TREATMENT OF STEELS
hot rolling lead to marked grain growth, with the result that grain refinement
during subsequent working is limited.
The situation is greatly improved if fine particles are introduced into the
austenitic matrix. The particles are usually found on grain boundaries, because
an interaction takes place between the particles and the boundary. A short
length of grain boundary is replaced by the particle and the interfacial energy
ensures a stable configuration. When the grain boundary attempts to migrate
away from the particles, the local energy increases and thus a drag is exerted on
the boundary by the particles.
The theory of boundary pinning by particles has already been referred to in
Chapter 9. Equation (9.1) defines the critical size of particle belowwhich pinning
is effective. Clearly, the control of grain size at high austenitizing temperatures
requires as fine a grain boundary precipitate as possible, and one which will not
dissolve completely in the austenite, even at the highest working temperatures
(1200–1300
◦
C). The best grain refining elements are very strong carbide and
nitride formers, such as niobium, titanium and vanadium, also aluminium which
forms only a nitride. As both carbon and nitrogen are present in control-rolled
steels, and as the nitrides are even more stable than the carbides (see Fig. 4.5,
Chapter 4), it is likely that the most effective grain refining compounds are the
respective carbo-nitrides, except in the case of aluminium nitride.
Equally important is the degree of solubility that such stable compounds
have in austenite. It is essential that there is sufficient solid solubility at the
highest austenitizing (soaking) temperatures to allow fine precipitation to occur
during controlled rolling at temperatures which decrease as rolling proceeds.
The solubility products (in atomic per cent) of several relevant carbides and
nitrides have been shown in Fig. 4.12 as a function of the reciprocal of the tem-
perature. All of these compounds have a small but increasing solubility in the
critical temperature range (900–1300
◦
C) (Fig. 10.3). In contrast, the carbides of
chromium and molybdenum have much higher solubilities, which ensure that
they will normally go completely into solution in the austenite, if the tempera-
ture is high enough, and will not precipitate until the temperature is well below
the critical range for grain growth. Data from another source
1
have provided the
following equations for solubilities expressed in weight per cent as a function
of absolute temperature:
log
10
[Al][N] =−6770/T + 1.03,
log
10
[V][N] =−8330/T + 3.46,
log
10
[Nb][C] =−6770/T + 2.26,
log
10
[Ti][C] =−7000/T + 2.75.
1
Irvine, K. J. et al., Journal of the Iron and Steel Institute 205, 161, 1967.
10.2 CONTROLLED ROLLING OF LOW-ALLOY STEELS 213
Fig. 10.3 Solubility curve for NbC in a steel with 0.15C–1.14Mn–0.04Nb wt% (Hoogendoorn
and Spanraft, Micro-alloying 75, Union Carbide Corporation, New York, USA, 1975).
The compositional changes possible are many, so discussion will be limited
to general principles which apply equally, whichever compound is the effective
grain refiner in a given steel. While grain growth at the highest austenitizing
temperatures may be restricted to some extent by a residual dispersion, the
main refinement is achieved during rolling as the temperature progressively
falls, and fine carbo-nitrides are precipitated from the austenite. These new
precipitates will:
1. Increase the strain, for a given temperature, at which recrystallization will
commence (Fig. 10.4).
2. Restrict the movement of recrystallized grain boundaries.
It should be borne in mind that the austenite may recrystallize several times
during a controlled-rolling schedule and the total effect of this will be a marked
austenite grain refinement by the time the steel reaches the γ/α transformation
temperature (Fig. 10.4). In the later stages of austenite deformation,at the lower
temperatures, recrystallization may not occur, with the result that deformed
austenite grains elongated and flattened by rolling may transform directly to
ferrite. In the final stages of controlled rolling, austenite grain growth can be
further suppressed byrapid coolingfrom the finishingtemperature, which allows
214 CHAPTER 10 THERMOMECHANICAL TREATMENT OF STEELS
Fig. 10.4 Critical strain needed to complete recrystallization of austenite as a function of
deformation temperature and grain size. Comparison of Nb steel with plain carbon steel
(Tanaka et al., Micro-alloying 75, Union Carbide Corporation, NewYork, USA, 1975).
the γ/α transformation to take place sub-critically, i.e. below Ar
1
, in austenite
which is still deformed. It is becoming common practice now to continue rolling
through the γ/α transformation and even into the fully ferritic region. Such
treatments lead to finer grain sizes, and higher yield stresses in the finished
product (Fig. 10.5), but impose much higher load factors on the rolling mills.
As a result of the combined use of controlled rolling and fine dispersions of
carbo-nitrides in low-alloy steels, it has been possible to obtain ferrite grain
sizes between 5 and 10 µm, in commercial practice. Laboratory tests have
achieved grain sizes approaching 1 µm, which would appear to be a practical
limit using this approach. The Hall–Petch relationship between grain size and
yield strength, which was discussed in Chapter 2, is very relevant to micro-
alloyed steels and, in fact, linear plots are obtained for the yield stress against
d
−1/2
(Fig. 10.6).Addition of0.05–0.09 wt% Nb toa plaincarbon steel refinesthe
ferrite grain size, allowing it to be reduced to below 5 µm(d
−1/2
=14 mm
−1/2
),
with a consequent substantial increase in yield strength. The displacement of
∼100 MN m
−2
between the C–Mn and C–Mn–Nb curves arises from dispersion
strengthening due to NbC. This is further illustrated in the two lower curves of
Fig. 10.7, which were obtained from specimens austenitized at 950
◦
C prior to air
cooling. If, however, progressively higher austenitizing temperatures are used,
e.g. 1100
◦
C and 1250
◦
C followed by air cooling, the resulting curves, although
still linear, have much steeper slopes, indicating a marked increase in yield
10.2 CONTROLLED ROLLING OF LOW-ALLOY STEELS 215
Fig. 10.5 Effect of finish rolling temperature on final ferrite grain size of a micro-alloyed steel
(after McKenzie, Proceedings of the Rosenhain Centenary Conference, Royal Society, London, UK,
1976).
Fig. 10.6 Effect of grain size on yield stress of a carbon–manganese–niobium steel (Le Bon
and Saint Martin, Micro-alloying 75, Union Carbide Corporation, NewYork, USA, 1975).
216 CHAPTER 10 THERMOMECHANICAL TREATMENT OF STEELS
Fig. 10.7 Effect of austenitizing temperature on the yield strength of a 0.1C–
0.6Mn–0.09Nb wt% steel (Gladman et al., Micro-alloying 75, Union Carbide Corporation,
NewYork, USA, 1975).
strength for a particular grain size. This large increment in strength is due to
the precipitation of NbC during cooling, following its solution at the higher
austenitizing temperatures.
10.2.3 Minimum achievable grain size
It is interesting to consider what might be the smallest grain size achievable
using large-scale thermomechanical processing. A reduction in ferrite grain size
(linear intercept d) is equivalent to an increase in the amount of grain boundary
surface per unit volume (S
V
) since d =2/S
V
. Grain boundaries have an energy
σ per unit area, so that the interfacial energy stored per unit volume of steel is:
σ × S
V
≡ 2σ/d.
10.2 CONTROLLED ROLLING OF LOW-ALLOY STEELS 217
Fig. 10.8 Plot of the logarithm of ferrite grain size versus the free energy change avail-
able to generate grain boundaries. The curve represents the values of d
min
α
. The points are
experimental data.
This stored energy cannot exceed the magnitude of the free energy change when
austenite transforms to ferrite, i.e. |G
γα
V
|:
|G
γα
V
|≥
2σ
α
d
α
.
It follows that the smallest ferrite grain size that can be achieved is when all of
G
γα
V
is used up in creating α/α grain boundaries, so that:
d
min
α
=
2σ
α
|G
γα
V
|
.
Figure 10.8 shows the variation in this limiting ferrite grain size as a function
of G
γα
V
, together with a compilation of experimental data on the smallest
grain size achieved commercially, using thermomechanical processing.
2
The
curve indicates that at large grain sizes, d
min
α
is sensitive to G
γα
V
and hence
to the undercooling below the equilibrium transformation temperature. How-
ever, reductions in grain size in the sub-micrometre range require huge values
of |G
γα
V
|, meaning that the transformations would have to be suppressed to
large undercoolings to achieve fine grain size.
Comparison of the industrial data against the calculated curve indicates that
in spite of tremendous efforts in developing processing routes, the smallest
ferrite grain size obtained commercially using thermomechanical processing is
2
Yokota,T., Garcia-Mateo, C. and Bhadeshia, H. K. D. H., Scripta Materialia 51, 767, 2004.