Keywords: ferrite, grain refinement, transformation, austenite, texture, C-Mn steel, hot rolling Abstract. Ferrite grain refinement by hot rolling mostly above Ae3 being followed by an ultra-fast
cooling has been investigated. An emphasis has been paid on the interval, Δt, between the finish of rolling and the start of water spray cooling of which cooling rate is more than 1000 °C·s -1. When Δt
was nearly equal to zero, ultra-fine ferrite of about 1 or 2 μm in grain diameter was obtained for 1.2 to
1.3 mm thick 0.1%C-1%Mn steel sheets near sheet surfaces or in thickness center regions
respectively, although the grain size at Δt of 0.5 s was about 3 μm in both regions. The ferrite grains
were almost equiaxed and surrounded by high angle boundaries. This grain refinement is likely to be
caused by an increased number of nucleation sites for the transformation from austenite to ferrite due
to the ultra-fast cooling. Such a grain refinement mechanism is discussed in terms of prior-austenite
structures deduced by the misorientation distribution function analysis.
IntroductionTMCP (Thermomechanical Controlled Process) has been a key technology to produce low-cost
high strength steel. In this process, hot rolling, which is mostly carried out at stable austenite ( γ)
temperatures above Ae3, and subsequent controlled cooling cause fine structures of ferrite (α) in final
products. Finer grain size of α leads to beneficial properties of steel such as higher tensile strength
and better ductility. Therefore, much effort has been paid to reduce the grain size of α by TMCP in
the past 30 years. However, the progress in reducing α grain size has been sluggish, and it has been
thought by many that the lower bound of α grain size for plain C-Mn steel appears to exist around 3 to
5 μm regardless of the amount of strain imparted by hot rolling and the rate of subsequent cooling [1].
This lower bound is only about 1 μm below the smallest obtainable grain size in commercial TMCP
mills.To reduce the lower limit of α grain size, many researchers have investigated methods based on
rolling at temperatures well below Ae3 or even below Ar3, and alternative methods capable of
reducing grain size down to 1μm have been proposed in the last decade [2-6]. Hurley and Hodgson
[4] have reported that as an undercooled γ from a high temperature to below Ae3 is deformed by
single-pass rolling of 30 to 40% in reduction, the intergranular nucleation of α is activated and α
grains of less than 1μm in diameter appear in a region near sheet surfaces. Similar grain refinement
phenomena have been observed by other investigators [1], and this type of phenomenon has been
called as the "dynamic" strain-induced transformation [4]. Other proposed methods incorporate either
heavy warm rolling in α temperature region [5] or cold rolling and subsequent annealing [6] to cause
continuous or discontinuous recrystallization of α, respectively. Although these methods are of considerable scientific interest, there are still difficulties for them to be commercially applied due to a
large deformation resistance at the required low temperatures and complexity of processes.
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The question that arises in the course of these investigations is whether or not the lower bound of α
grain size known for TMCP is inevitable as hot rolling is carried out in a stable γ temperature region.
The ultra-fine α formation by the strain-induced transformation [4] implies that, as a sufficient
amount of strain is kept remained in γ, the subsequent transformation may lead to a dramatic
reduction in grain size regardless of rolling temperature. Therefore, the aim of this research was to
investigate a method that employed an extremely fast cooling after multi-pass rolling with an attemptto suppress the annihilation of dislocations in γ and to achieve ultra-fine grains in C-Mn steel.
Experimental Procedure
The plates of plain 0.1%C-1%Mn steel of 30 or 35 mm in thickness were reheated to the
temperatures of 1000 °C and hot-rolled to about 1.2 to 1.3 mm in thickness by 6 passes using a
three-stand experimental rolling mill. The first three rolling passes were carried out in a reversing
way using one of the three stands and the last three rolling passes were made in an advancing way
using all the stands. The reduction of each pass of the last three passes and the time periods between
them were 40 to 50 % and about 1 s respectively. The end temperature of rolling was controlled to be
800 to 820 °C; the calculated paraequilibrium Ae3 for 0.1%C-1%Mn steel was 818 °C.
After rolling, the samples were subjected to a high pressure water spray cooling. The time period
during which the sample ran from the roll gap exit position to the position at which a jet of water
made contact with the sample, “Δt”, was controlled to be from 0.05 to 1 s. The cooling rate, estimated
by monitoring temperatures at the entry and exit positions of the cooling zone, was more than
1000 °C·s-1
. This cooling has been continued until the sample temperature has decreased to from 650
to 700 °C, at which the transformation from γ to α is known to become active, or to the room
temperature to investigate the structure of prior- γ. In addition, the plates of a 70%Ni-30%Fe alloy
were hot rolled and then cooled mostly in the same fashion as described above. This Ni-Fe alloy has
been comparatively used in the texture analyses of prior- γ, in which fcc to bcc transformation never
occurs.The microstructures of these steels were examined via transmission electron microscopy (TEM),
scanning electron microscopy (SEM), and electron back scattering diffraction analysis (EBSD), and
the textures were analyzed by X-ray diffraction techniques and orientation distribution function
(ODF) analysis. The positions in the samples at which above measurements were made were either
0.1 mm below the sample surfaces or in the center region in the thickness direction.
Misorientation distribution function method (MODF)
The MODF is a calculation method developed based on the ODF representation of textures of
materials, which relates the texture of a parent phase to that of its product phase with a set of linear
equations [7]. Let the crystal orientation distribution function of γ, f γ, and that of α, f α, be representedby the following series expansions,
),()()()(0
)(
1
)(
1
.:
0
)(
1
)(
1
.:
gT C g f and gT C g f M N M N
∑ ∑ ∑∑ ∑ ∑∞
= = =
∞
= = =
==λ
λ
μ
λ
ν
μν
λ
μν
λ
α
α λ
λ
μ
λ
ν
μν
λ
μν
λ
γ
γ
where.:
T λμν
(g)’s are the symmetrically invariant functions defined by Bunge [7] for cubic crystals in
the sample with orthorhombic symmetry with g being a set of Euler angles that represent crystal
orientations. Then, given the crystal axis rotation Δg caused by the transformation from γ to α and
assuming no variant selection rule, the series expansion coefficients, γCλ
T λμμ'(g)’s are the symmetrically invariant functions defined for cubic crystals in the sample
with cubic symmetry [7]. The above equations are solvable for γCλ
μν, if
αCλ
μνand Δg are known. In
the present investigation, the prior- γ textures were thereby deduced using the determinedαCλ
μν’s by
the ODF analysis and assuming the K-S orientation relation between γ and α [8]. The calculation was
truncated at the order of 20th.
Experimental Results
Microstructures of C-Mn steels. The change in average grain size of α with changing Δt is
presented in Fig.1; these samples have been subjected to the ultra-fast cooling to about 700 °C being
followed by a slow cooling. The grain sizes at Δt of 0.5 s are about 3.3 μm near sheet surfaces and 2.5
μm in thickness center regions, which are not much different from the lower bound of grain size
reported for C-Mn steel, although the rate of cooling employed in this experiment is extremely large
being larger than 1000 °C·s-1
. However, as Δt decreased, the grain size was markedly reduced, in
particular near sheet surfaces. As Δt decreased from 0.5 to 0.2 s, grain size decreased to about 2 μm
near the sheet surface and to about 2.5 μm at the thickness center. Then, near the sheet surface, grain
size dramatically decreased to 1.3 μm as Δt decreased to 0.05 s, whereas it scarcely decreased at the
thickness center.
Figures 2 and 3 show the typical examples of SEM and TEM micrographs for the samples
ultra-fast cooled to about 650 °C at Δt being nearly equal to zero, respectively. The α grains were
almost equiaxed with a low internal dislocation density. This is an evidence that the ultra-fine α
scarcely underwent deformation once formed. The second phase observed in near surface areas was
mainly granular cementite existing at α grain boundaries. As presented in Figs. 4 and 5, EBSP
analysis revealed that the most of the grains were surrounded by high angle boundaries with the
misorientation angles greater than 15°. The ratios of such high angle boundaries to all the boundaries
were 88% and 71%, and the average grain diameters calculated with only such high angle boundarieswere 1.0 and 1.8 μm near sheet surfaces and in thickness center regions, respectively.
In order to investigate the structure of prior- γ just before the α transformation in relation with the
grain refinement mechanism, the microstructures of steels that were subjected to the ultra-fast cooling
to the room temperature were examined. In Fig. 6, the SEM micrographs of these samples are shown.
0 0.1 0.2 0.3 0.4 0.5
Δt (sec)
1
1.5
2
2.5
3
3.5
F e
r r i t e D i a m e t e r (μ m ) surface
center
Fig. 1: Effect of Δt on a grain diameter.
Surface Center
Fig. 2: SEM micrographs of C-Mn steel with ultra-fine
grains (Δt = 0.05s).
Surface Center
Fig. 4: Mapping of grain boundaries with misorientation
i.e., of a deformed state, whereas the Ni-Fe cooled at Δt of 1 s has been almost completely
recrystallized. The main components of textures lie around {001}<110> to {225}<110> and
{111}<110> for the deformed sample, and around {013}<031> to {013}<331> and {112}<021> for
the recrystallized sample.
Discussion
The mechanisms that have been proposed in relation with the α grain refinement by hot rolling of
C-Mn steel are (1) dynamic or (2) static transformation from γ to α, (3) the grain refinement of prior- γ
by dynamic recrystallization of γ, and (4) dynamic recrystallization of α [1−6]. Among them, (1) and
(4) can be omitted for the mechanism of the grain refinement in this experiment. The reason is that the
structure of C-Mn steel in this experiment is thought to be γ just after the final rolling, since the C-Mn
steel can be quenched into martensite by the ultra-fast cooling to the room temperature except for the
near surface region at Δt being nearly equal zero (see Fig 6). Although the quenched microstructure
near the surface at Δt of 0.05 s has been ultra-fine α, it is not conceivable that the reversetransformation of the ultra-fine ferrite to γ can occur around Ae3. This ultra-fine α is thus thought to
have been formed during the ultra-fast cooling.
Then the question is if the prior- γ has been recrystallized before transformation. For the thickness
center region, it is clear from Fig.6 that the prior- γ has undergone little recrystallization. However, for
the near sheet surface region, an implication exists in Fig.6 that the prior- γ might have undergone
recrystallization as mentioned earlier. Therefore, MODF calculation was performed to deduce the
texture of the prior- γ in order to clarify it.
The MODF results for near sheet surface regions ultra-fast cooled to the room temperature are
shown in Fig. 9. At Δt of 0.05 s, the main components of the texture of prior- γ lie around {001}<110>
to {112}<110>, and a subcomponent is also observed in the proximity of {233}<331>. This texture isvery similar to that of the deformed Ni-Fe alloy shown in Fig. 8. On the other hand, at Δt of 0.3 s, the
main components of the texture of prior- γ lie around {013}<031> to {013}<331> and {001}<110>,
although the {001}<110> components split into two. This texture can be recognized as the mixture of
the two textures of deformed and recrystallized Ni-Fe alloys shown in Fig.8; some inconsistency is
probably due to the truncation and the omission of variant selection rules in MODF calculation.
Therefore, the prior- γ near sheet surfaces at Δt of 0.05 s is considered to be deformed, while that at Δt
of 0.3 s is partly recrystallized.
Thus, the mechanism (3) is also omitted for Δt being nearly equal to zero, and the most likely
mechanism operating for the grain refinement in this experiment is the one based on the mechanism
(2). Perhaps, a large amount of strain imparted in the deformed prior- γ is kept remained by the
ultra-fast cooling until the temperature of samples decreases to the range in which α transformation
becomes active, then the increased number of nucleation sites caused by the strain may result in the
observed ultra-fine grain structure. Moreover, it is likely that this transformation to ultra-fine ferrite is
greatly enhanced by the presence of shear strain, which is known to be imparted mainly in near
Ultra-fast cooled at Δt of 0.05 s Cooled at Δt of 1 s
0 30 60 90
ϕ1
90
60
30
0
φ
2
2
4
4
{001}<110>
ϕ 2=0° 0 30 60 90
ϕ1
90
60
30
0
φ
1
1
1
2
2
2
3
3
3
44
ϕ 2=45°
{111}<110>
{001}<110>
{225}<110>
0 30 60 90
ϕ1
90
60
30
0
φ
0.8 0.8
0.8 0.8 0.8
0.80.8
1.21.2
1.2
1.2
1.6
1.6
ϕ 2=0°
{013}<331>
{013}<031>
0 30 60 90
ϕ1
90
60
30
0
φ
0.8
0.8
0.8
1.2
1.21.2
1.6
ϕ 2=45°
{112}<021>
Fig. 8: Textures for near sheet surface regions of 70%Ni-30%Fe alloys.
surface regions due to the friction between sample and roll surfaces, since the greater grain
refinement tendency has been observed near sheet surfaces. The presence of shear strain can also
enhance the recovery and crystallization of γ, and, therefore, the greater grain refinement tendency
near the sheet surface quickly disappears as Δt increases (see Figs. 1 and 9). It is also thought that an
internal nucleation of α in γ grains should play an important role in the grain refinement in this
experiment.
Δt = 0.05 s Δt = 0.3 s
0 30 60 90
ϕ1
90
60
30
0
φ
2
2
2
2
22
6
6
10
10
ϕ 2=0°
{001}<110>
0 30 60 90
ϕ1
90
60
30
0
φ
ϕ 2=45°
{001}<110>{112}<110>
{233}<331>
5
6
10 10
4
4
2
2
2
0 30 60 90
ϕ1
90
60
30
0
φ
ϕ 2=0°
{013}<331>
{013}<031>
2.5
4.5
2
1
4.5
4.5 4.5
2 2.5
1
0 30 60 90
ϕ1
90
60
30
0
φ
2 2
2
2
2
2
4 4
ϕ 2=45°
3
Fig. 9: Textures of prior- γ calculated from α textures of C-Mn steel by MODF.
Conclusions
A new approach to breakthrough in grain refinement of C-Mn steel has been described, which
allow polygonal α of about 1 and 2 μm in grain size to be achieved near sheet surfaces and in
thickness center regions respectively.
It has been found that, as 0.1%C-1%Mn steel is hot-rolled mostly in a stable γ region being
followed by an ultra-fast cooling of over 1000 °C·s-1
in cooling rates without any duration, a dramatic
effect in grain refinement appears. In near surface regions, the grain size gradually decreases fromabout 2.5 to 2 μm with decreasing Δt from 0.5 to 0.2 s, and then the reduction in grain size toward 1
μm occurs as Δt is reduced nearly to zero. In thickness center regions, the grain size gradually
decreases from over 3 μm to about 2 μm, as Δt is reduced from 0.5 to 0.05 s. These ultra-fine grains
are almost equiaxed and surrounded by high angle boundaries. The analysis on prior- γ structures
employing MODF suggests that this grain refinement should be caused by the almost complete
suppression of the recovery and recrystallization of deformed γ by the ultra-fast cooling.
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[4] P.J. Hurley and P.D. Hodgson: Mater. Sci. Eng. Vol. A302 (2001), p. 206
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