Microstructural Features of Austenite Formation in C35 and C45
The microstructural evolution during continuous heating experiments has been studied for two
C-Mn steels with carbon contents in the range 0.35 to 0.45 wt pct using optical microscopy,
scanning electron microscopy (SEM), and electron probe microanalysis (EPMA). It is shown
that the formation of the austenitic phase is possible in pearlite as well as in ferrite regions.
Thus, a considerable overlap in time of ferrite-to-austenite and pearlite-to-austenite transformations is likely to occur. Another observation that was made during the experiments is that,
depending on the heating rate, the pearlite-to-austenite transformation can proceed in either one
or two steps. At low heating rates (0.05 C/s), ferrite and cementite plates transform simultaneously. At higher heating rates (20 C/s), it is a two-step process: first ferrite within pearlite
grains transforms to austenite and then the dissolution of the cementite lamellae takes place.
Several types of growth morphologies were observed during the experiments. The formation of
a finger-type austenite morphology was noticed only for low and intermediate heating rates
(0.05 C/s and 20 C/s), but not for the heating rate of 300 C/s. The formation of this fingertype austenite occurs on pearlite-ferrite grain boundaries and coincides with the direction of
cementite plates. The carbon inhomogeneities in the microstructure affect the formation of
martensitic/bainitic structures on cooling.
DOI: 10.1007/s11661-007-9128-3
The Minerals, Metals & Materials Society and ASM International 2007
HEAT treatment of steels is an everyday routine to
obtain materials with desired properties and structures.
The first step in the heat-treatment process for the vast
majority of commercial steels is austenitization—formation of austenite upon heating from a variety of initial
phases. In spite of its importance, there has been little
work carried out on the formation of austenite, compared to the huge effort put into studying its decomposition. The major interest toward austenitization was
drawn after dual-phase (DP) steels were developed. The
DP steels are most commonly used in structural applications where they have replaced more conventional
high strength low alloy (HSLA) steels. They were
developed to provide high strength formable alloys
and offered a significant weight reduction of the final
Speich et al.,[1] who studied the intercritical annealing
of DP steels, distinguished several stages in the ferriteto-austenite transformation. According to Reference 1,
V.I. SAVRAN, Postdoctoral Student, C. KWAKERNAAK, Technical Staff Member, and W.G. SLOOF and J. SIETSMA, Associate
Professor, are with the Department of Materials Science and
Engineering, Delft University of Technology, 2628 CD, Delft,
The Netherlands. Contact e-mail: [email protected] Y. VAN
LEEUWEN, Scientist, is with the Nuclear Safety Department, P.O.
Box 16191, 2500 BD, The Hague, The Netherlands. D.N. HANLON,
Metals Researcher, is with the Corus Research, Development &
Technology, P.O. Box 10000, 1970 CA, IJmuiden, The Netherlands.
Manuscript submitted July 26, 2006.
Article published online May 1, 2007.
946—VOLUME 38A, MAY 2007
the first step of ferrite-to-austenite transformation
consists of the nucleation of austenite (c) at the ferritepearlite interfaces and growth of austenite into pearlite
(a + h) until the pearlite dissolution is complete. The
nucleation of austenite is argued to occur instantaneously, with essentially no nucleation barrier. The rate
of growth in this stage is controlled primarily by the rate
of carbon diffusion in austenite between adjacent
pearlitic cementite (h) lamellae, but may also be influenced by diffusion of substitutional elements at low
temperatures.[2] At the end of this first step, a highcarbon austenite has been formed, which is not in
equilibrium with ferrite (a). The second step of the
transformation consists of the growth of this austenite
into ferrite to achieve partial equilibrium with ferrite.
The lower growth rate of austenite in this step is
controlled either by the carbon diffusion in austenite
over larger distances or by the manganese diffusion in
ferrite. In the final step, very slow final equilibration of
ferrite and austenite is achieved by manganese diffusion
through austenite. Jayaswal and Gupta,[3] who studied
in detail the second and third stages of transformation in
HSLA steel, observed that in addition to the growth of
austenite from regions of prior pearlite, austenite was
also observed to form at the a-a grain boundaries. They
were not able to give explanations for this phenomenon,
but indicated that the possible reason could be the
presence of retained austenite in the starting microstructure. On the other side, Garcia de Andres et al.[4] in their
study of the pearlite dissolution in DP steel reported a
clear differentiation between the pearlite dissolution
process and the a-c transformation.
An interesting observation that is often reported in
relation to the ferrite-to-austenite transformation is the
formation of acicular structures. Zel’dovich et al.[5]
distinguished three different mechanisms of austenite
formation depending on the heating rate. At a very low
heating rate (a few degrees per minute) or at a very rapid
heating (thousands degrees per second), newly formed
austenite grains have an acicular structure, and a
structural heredity (the original austenite grain is recovered both in size and crystallographic orientation) is
present. It is said that the phenomenon of structural
heredity must indicate an ordered mechanism of austenite formation, that is, diffusionless during rapid
heating and what is known as the homogeneous
mechanism of diffusional transformation during slow
heating. Heating at a certain intermediate range results
in loss of ordering and in grain refinement. The newly
formed austenite grains have more of a rounded
The formation of acicular structures during the
ferrite-to-austenite transformation was also observed
by Jayaswal and Gupta.[3] They noticed that the
austenite phase, instead of growing with a planar or
nearly planar front, changed into a Widmanstätten
structure on both the ferrite-ferrite grain boundaries and
on well-advanced ferrite-pearlite (now austenite) interfaces. Law and Edmonds[6] performed studies of the
morphology and crystallography of austenite precipitates in Fe-0.2 pct C-1 pct V alloy. They noticed that
austenite formed on grain boundaries is idiomorphic or
allotriomorphic, while that formed on lath boundaries
can become acicular by inheriting the lath dimensions.
Grain boundary austenite was proven to nucleate in low
carbon ferrite with the Kurdjumow–Sachs orientation
relationship with one ferrite grain, and to grow predominantly into an adjacent grain with which it was not
related. Based on this observation and the general
absence of planar facets or sideplate morphologies, they
proposed that austenite grows by migration of incoherent interfaces.
This article presents experimental observations obtained during continuous heating experiments of C-Mn
steels with 0.35 or 0.45 wt pct carbon, using optical
microscopy, scanning electron microscopy (SEM), and
electron probe microanalysis (EPMA). The alloys were
heated with two different heating rates, 0.05 C/s and 20
C/s, to different temperatures of the intercritical region
and were directly quenched. The effect of the extreme
heating rates on the microstructure evolution was
studied by heating the samples with 300 C/s to different
temperatures within and above the intercritical region.
The results of experimental studies for different heating
rates on the development of the microstructure during
the ferrite-to-austenite transformation focusing on austenite nucleation and growth morphologies are presented and analyzed in this article. Experimental
observations to support the idea of probable overlapping of pearlite-to-austenite and ferrite-to-austenite
transformations, as observed by Jayaswal and Gupta,[3]
are presented together with a possible explanation for
this phenomenon.
A. Driving Force for Nucleation in Hypoeutectoid Steel
On heating a hypoeutectoid steel from room temperature to a single-phase region, a phase transformation
occurs, which consists of two stages, namely, nucleation
and growth. The essential driving force behind this
transformation is the difference in the Gibbs free energy,
DG, between the initial and final states.[7] For a
transition from phase i to phase j to occur, the condition
DG ¼ Gj Gi <0
must be satisfied (G and G are the free energies of the
parent and the new phase, respectively).
A schematic representation of the Gibbs free energy G
as a function of the carbon concentration is shown in
Figure 1(a) at a temperature above the eutectoid temperature of the Fe-C system. In this temperature range,
two phases, a (ferrite) with composition Ca/c (Figure 1(b), point 1) and c (austenite) with composition
Cc/a (point 2), are in equilibrium. The formation of the
c phase with the mentioned equilibrium composition
leads to a maximum gain in Gibbs free energy (DGmax).
In this case, the newly formed austenite grains have a
different composition from the original ferrite phase,
and a significant enrichment in carbon must take place.
Fig. 1—Schematic representation of the (a) Gibbs free energy G as a function of the carbon concentration in ferrite (a), austenite (c), and
cementite (h) at a temperature above A1; and (b) metastable Fe-C phase diagram, indicating notations for the carbon atomic fractions used
throughout the text. Numbers in brackets correspond to figurative points (1) through (5) in Figures 1 and 2. The solid thin lines in (a) represent
the common tangent lines between a-h and a-c.
VOLUME 38A, MAY 2007—947
From Figure 1(a), it also follows that even though the
maximum gain in free energy is achieved for a large
carbon enrichment of the c phase (equilibrium condition), some decrease in the Gibbs free energy, even
though of a smaller value, is also realized with the
formation of the c phase with lower carbon content (for
example, DG1, a situation in which the system departs
from the thermodynamic equilibrium). Thus, even
though with the formation of low-carbon austenite the
gain in Gibbs free energy is smaller compared to the
equilibrium value (DG1 < DGmax), this process is nevertheless thermodynamically possible. By low-carbon
austenite, the austenite with a carbon content less than
equilibrium according to the phase diagram is understood, and not a carbon-free austenite. Some degree of
enrichment does have to take place and the nucleation
and growth will be stimulated in carbon-rich areas or in
their vicinity.
B. Temperature Range 1: A1 < T < A3
The changes in the microstructure of steel on heating
can in part be understood in terms of the Fe-C phase
diagram (Figure 1(b)). At room temperature and normal pressure, the microstructure of carbon hypoeutectoid steel after slow cooling consists of ferrite and
pearlite. Upon heating the steel from room temperature
to the A3 temperature, two different situations can be
distinguished. The first one is related to the formation of
austenite within the pearlite phase at the a/h interface, as
the one schematically represented in Figure 2(b), and is
described as
aCa=h þ h ! aCa=c þ cCc=a
where C
and C are the equilibrium carbon concentrations in ferrite, changing with temperature
according to lines QP and PG, respectively (Figure 1(b)); Cc/a is the equilibrium carbon concentration
in austenite, changing according to line SG; and h is
cementite and is considered to be of a constant composition.
It is known that the velocity of the phase boundary
can be considered in first approximation inversely
related to the carbon concentration difference on it.[8]
For the value of this difference, the concentration should
change to form a new phase. The carbon difference on
the a-c grain boundaries is much less than on the c-h
grain boundary; thus, austenite can be expected to grow
much faster in the ferrite phase than in the cementite
Figure 2(b) is a one-dimensional representation of the
planar geometry. In this case, the interface movement is
controlled by the diffusion of carbon through the
austenite phase. The situation at a triple line between
cementite, ferrite, and austenite is not considered as the
diffusion distances become negligible, and it will not be a
limiting factor for the transformation.
The second situation is related to the possible
formation of austenite on ferrite-ferrite grain boundaries
(Figure 2(c)) and can be described as
aCa=h ! aCa=c þ cCc=a
From the Fe-C phase diagram (as the one shown in
Figure 1(b)), it is seen that the carbon concentration in
ferrite decreases with increasing temperature. This will
lead to austenite nucleation at the a-a grain boundaries.
Thus, not only does the carbon of cementite play a role
in the formation of austenite, but so does the carbon
rejected from the a-solid solution. On the other hand,
due to the difference in carbon solubility at the a/h and
a/c grain boundaries, a concentration gradient within
the ferrite phase is present (Ca/h > Ca/c, Figure 2(c)).
This creates the driving force for carbon diffusion
toward the a-c grain boundary.
The subsequent growth of the austenite nuclei
involves the removal of carbon from cementite with its
Fig. 2—Schematic view of the microstructure (a) and variation of the carbon content across (b) the cementite-austenite-ferrite and (c) the ferriteaustenite-ferrite boundaries. F = ferrite, A = austenite, and P = pearlite.
948—VOLUME 38A, MAY 2007
diffusion into the not-transformed ferrite in order to
‘‘feed’’ carbon to austenite at the a/c interface. The
velocity of the a-c grain boundary movement depends
on how fast the carbon is supplied to it and, consequently, on the diffusion path length. This leads to the
notion that ferrite within the pearlitic phase will
transform much faster than proeutectoid ferrite.
Taking all of the aforementioned into account, it is
possible to assume that the nucleation of austenite
grains in pearlite and in proeutectoid ferrite can both
occur; however, the transformation will proceed at a
much higher rate in the pearlite phase than in proeutectoid ferrite due to the shorter diffusion distances and
the surplus of carbon that is available from the
dissolving cementite plates.
C. Temperature Range 2: T > A3
The transformation that takes place at temperature
T > A3 can be described as
aCa=c þ cCc=a ! cC0
where C is the average carbon concentration in the
alloy, and therefore also the austenite composition under equilibrium conditions after the transformation is
completed (above A3).
Upon further heating of the sample in the temperature range A1 < T < A3, the a fi c transformation at
conditions close to equilibrium proceeds most probably
by diffusional growth. The transformation completes
above the A3 temperature. In alloys with low carbon
content and thus a low amount of pearlitic phase
present, nucleation on the ferrite-ferrite grain boundaries can take place. For the alloys with high pearlitic
volume fractions, transformation proceeds via the
growth from the already existing austenitic areas
(mainly former pearlite grains) into proeutectoid ferrite.
In order to examine the evolution of the microstructure during continuous heating in C35 and C45 steels, a
set of interrupted heating experiments was performed
using a Bähr 805A/D dilatometer (Hüllhorst, Germany).
Typical micrographs of the initial microstructures for
C35 and C45 alloys are shown in Figure 3. The micro-
structures consist of ferrite and pearlite mixture with
different phase volume fractions. Chemical compositions
of the experimental alloys are shown in Table I. The Ac1
and Ac3 temperatures for studied alloys are presented in
Table II. For the results that will be presented in this
article, the differences in chemical compositions between
studied alloys do not play a significant role.
The samples, with a diameter of 5 mm and a length of
10 mm, were heated using a high-frequency induction
coil with heating rates 0.05 C/s and 20 C/s to different
temperatures within the ferrite + pearlite-to-austenite
transformation region and directly quenched with cooling rates in the range of 700 C/s to 750 C/s. A
thermocouple, spot welded on the sample, was used to
control the temperature during the test. Samples for
metallographic examinations were ground, polished,
and etched with a 2 pct Nital cleaning agent and
examined under an optical microscope.
The secondary electron images were produced with a
LEO 438VP scanning electron microscope fitted with a
tungsten filament. Samples for SEM measurements were
first mechanically polished to 1 lm and then electrolytically polished using electrolyte that contained 400 mL
99 pct 2-bytoxyethanol with 20 mL HClO4 (hydrogen
peroxide 30 pct).
The composition profiles of the samples were determined using EPMA. On the cross sections prepared from
the samples, 0.5-lm equidistant points were selected
along lines defined in backscattered electron images. The
measurements were performed with a JEOL* JXA
*JEOL is a trademark of Japan Electron Optics Ltd., Tokyo.
8900R wavelength-dispersive/energy-dispersive combined microanalyzer, operated with a focused electron
beam of 15 keV and 25 nA. These electron beam
conditions were a compromise between sufficient spatial
resolution and intensity of C Karadiation. Wavelength
dispersive spectrometry was employed to record the C
Ka, Si Ka, Cr Ka, Mn Ka, and Cu Ka intensities
simultaneously. A W/Si X-ray reflective multilayer with
a 2d spacing of 9.80 nm was used for selecting the C Ka
radiation, a (100)-TAP was used for selecting Si Ka, a
(002)-PET for Cr Ka, and a (200)-LiF crystal for Mn Ka
and Cu Ka radiation. The peak intensity for a single spot
on the specimen was determined from measuring the
Fig. 3—Initial microstructures for (a) C35 and (b) C45. F = ferrite (white) and P = pearlite (dark).
VOLUME 38A, MAY 2007—949
Table I.
Composition of the Experimental Alloys in Weight Percentages
Table II. Ac1 and Ac3 Temperatures for C35 and C45
Alloys for Two Different Heating Rates 0.05 and 20 C/s
Ac1, C
0.05 C/s
Ac1, C
20 C/s
Ac3, C
0.05 C/s
Ac3, C
20 C/s
number of counts during 4 minutes. The background
intensities of C Ka, Si Ka, Cr Ka, Mn Ka, and Cu Ka were
determined similarly at the same spot. The background
intensity of C Ka was measured separately on a pure a-Fe
reference. The surface of the specimen was decontaminated 30 seconds prior to and during each measurement
using an air jet. This procedure removes any carbonaceous surface contamination at the measurement location. The composition at each analysis location of the
sample was determined using the X-ray intensities of the
constituent elements after background correction relative to the corresponding intensities of reference materials, i.e., h-Fe3C[9] for C, and the pure elements for Si,
Cr, Mn, and Cu, respectively. The thus obtained
intensity ratios were processed with a matrix correction
program CITZAF based on the F(qz) method[10] to
compute the composition with the matrix element Fe
taken as balance. The carbon concentrations determined
are accurate within 0.03 wt pct, including the background error.[9]
A. Nucleation
Figures 4 and 5 show typical optical and SEM images
from the interrupted heating experiments. In these
down-quenched samples, martensite islands reveal the
locations of the austenite grains. Two interesting
observations can be made. The first observation is
related to the nucleation of the new austenite grains,
which takes place predominantly in pearlite areas that
are rich in carbon (Figures 4(a) and (b)). Depending on
the heating rate, this process can take place in one or
two steps. At a very low heating rate, 0.05 C/s, there is
no essential delay between the ferrite-to-austenite transformation and the cementite dissolution within the
pearlite grain. Figures 4(a) and 5(b) show that no
inhomogeneities are present in the martensite phase that
was austenite prior to cooling. Figure 5(a) (area marked
with oval) shows an austenite grain nucleated on the
ferrite-pearlite grain boundary. The growth of the newly
formed austenite grain is not planar and the position of
austenite ‘‘fingers’’ coincides with the direction and
position of the cementite lamellae, which are rich in
A completely different situation was observed in the
case when the heating rate was 20 C/s. A clear delay in
the cementite dissolution in comparison to the ferrite-toaustenite transformation resulted in a time-step difference between the two processes.[4] A closer look using
SEM (Figure 4(b)) reveals a partially transformed
pearlite grain in which the cementite is not completely
dissolved in martensite, which was austenite at high
temperatures prior to quenching.
The second interesting observation is related to the
nucleation of austenite on the ferrite-ferrite grain
boundaries: at a triple point (arrow in Figure 5(a))
and a grain boundary (Figure 5(b)). The austenite nuclei
have a classical cuplike shape and appear at the very
early stages of the transformation. Thus, two transformations, pearlite-to-austenite and ferrite-to-austenite,
appear to overlap. The degree to which the two
Fig. 4—Typical micrographs from the interrupted heating experiments. (a) Optical micrograph of the C35 alloy heated with the heating rate of
0.05 C/s to 745 C (close to the end of the pearlite-to-austenite transformation) and (b) SEM micrograph of the C45 alloy heated with the heating rate of 20 C/s to 765 C (middle of a pearlite-to-austenite transformation). The arrow indicates a pearlite grain that was transformed into
austenite on heating and subsequently into martensite on cooling. In this grain, the cementite plates are still visible and, in some cases, partially
dissolved. F = ferrite, M = martensite, and P = pearlite.
950—VOLUME 38A, MAY 2007
Fig. 5—SEM micrographs of the C35 alloy, heated with 0.05 C/s to (a) 740 C (start of the pearlite-to-austenite transformation) and (b) 745 C
(close to the end of the pearlite-to-austenite transformation). The arrows indicate the nucleation of austenite at a triple point (a) and a ferriteferrite grain boundary (b). F = ferrite, M = martensite, and P = pearlite.
processes overlap cannot be established from the present
Figure 6 shows the distribution of the alloying elements across the austenite nuclei on the ferrite-ferrite
grain boundary similar to the one shown in Figure 5(b).
The carbon concentration across the austenite nucleus
region varies significantly: from approximately
0.01 wt pct in the ferrite phase and up to 0.27 pct in
the austenite nuclei. Other alloying elements (Si, Cr,
Mn, and Cu) do not show any significant variations in
B. Growth Morphologies
Several types of growth morphologies were observed
during the experiments. Acicular (finger) type growth is
spotted on the pearlite-ferrite grain boundaries (Figure 7). New austenite grains nucleate on the grain
boundary and grow into the neighboring ferrite grain,
most likely inheriting the lath dimensions. This type of
growth is only detected on pearlite laths being perpendicular to the ferrite/pearlite grain boundary and is not
present if the laths are parallel to it. In the latter case,
the formation of a bainitic structure on cooling takes
place (Figures 7 and 8). This bainite was austenite at
Fig. 6—Alloying element distribution across austenite nuclei formed
on ferrite-ferrite grain boundary (C35 alloy). The maximum of the
carbon concentration corresponds to the middle of austenite nuclei.
The low level of the carbon concentration corresponds to the ferrite
higher temperatures prior to quenching. The bainitic
phase is clearly visible in the vicinity of former pearlite
grains (martensite after quenching), is black in color,
and lies along the grain boundaries.
Figure 9 shows the distribution of the alloying elements across the fingers, as determined by EPMA.
Similar to Figure 6, the only diffusing element is carbon;
the other elements (Si, Cr, Mn, and Cu) show no or
negligible variations in concentrations. The carbon
content varies from approximately 0.2 wt pct between
the fingers to approximately 0.8 wt pct inside the finger.
The micrographs showing the effect of the extreme
heating rate (300 C/s) and different holding times on
the microstructure evolution are shown in Figure 10. At
770 C and holding time 1 second, the microstructure is
highly inhomogeneous and consists of ferrite (white
areas), bainite (black areas), and martensite matrix (gray
areas). The black areas form a continuous network and
reproduce the original grain size. The ferritic phase lies
along these black lines, presumably the former grain
Figure 11 shows the distribution of alloying elements
along the line T-T (Figure 10(a)). Alloying elements
(Si, Cr, Mn, and Cu) do not show significant variations
in concentrations. In contrast, the carbon concentration
in the black phase is remarkably low: around 0.2 wt pct.
In the rest of the sample, the carbon content is
approximately 0.4 wt pct, which is about the average
carbon content in the sample.
With increasing the holding time to 10 seconds
(Figure 10(b)), the structure tends to become more
homogeneous and the amount of ferritic phase first
decreases and eventually almost disappears at a holding
time of 60 seconds (Figure 10(c)). The temperature of
770 C corresponds to the austenite area in the Fe-C
phase diagram. At higher overheating (900 C) and
short holding time (1 second), the microstructure is as
well highly inhomogeneous and consists of a martensitic
matrix (gray areas) and a bainitic phase (black areas)
forming a continuous network (Figure 10(d)).
It has been known for a long time that the situation
during forming of austenite is much different from the
VOLUME 38A, MAY 2007—951
Fig. 7—Examples of acicular growth morphologies as found in C35 and C45 alloys. The arrow (S-S) indicates the direction of EPMA measurement (Fig. 9).
energy of the product, austenite, and the starting
microstructure, such as pearlite or ferrite). Also, with
increasing temperature, the atomic mobility increases.
Thus, the rate of austenite formation increases with
increasing temperature. Because both the thermodynamic driving force for the formation of austenite
and the atomic mobility become larger with higher
temperatures, both the rate of nucleation and the
rate of growth continually increase with increasing
A. Nucleation
Fig. 8—SEM micrograph of the bainite structure.
Fig. 9—Alloying element distribution across the fingers in the C45
alloy (Fig. 7). The maximum of the carbon concentration corresponds to the fingers and the minimum between them.
transformation upon cooling. As the temperature is
raised into the austenite region, the driving force for
transformation increases (this is the difference in the free
952—VOLUME 38A, MAY 2007
Above the A1 temperature, the ferrite + pearlite
phase mixture, which was stable at lower temperatures,
becomes unstable. The system will try to decrease its free
energy by creating austenite, and consequently, an
austenite-ferrite interface is formed. The essential driving force behind this transition is the difference in the
Gibbs free energy between the initial and final states. As
was already shown in Section II, maximum gain in free
energy is achieved for a large carbon enrichment of the c
phase. Some decrease in the Gibbs free energy, even
though of a smaller value, is also possible with the
formation of the c phase with a lower than equilibrium
carbon content (for example, DG1 in Figure 1(a)). In
both cases, nevertheless, the nucleation and growth is
stimulated in carbon-rich areas or their vicinity, because
some enrichment of austenite phase has to take place.
There are two possible sources of carbon. First of all,
these are the pearlitic areas, which have lamellar
structures, consisting of alternating ferrite (low in
carbon) and cementite (high in carbon) plates (Figures 2(a) and (b)). The second is the proeutectoid ferrite
itself, because Ca/c(1) < Ca/h (5) (Figure 2(c)). Hence, it
follows that, for the austenitic nucleus formed on the
boundary of ferrite grains, contact with cementite is not
necessary. The decreasing equilibrium C content with
increasing temperature in the ferritic phase is in itself a
direct source of carbon. After supplying carbon to
Fig. 10—Typical micrographs from the interrupted heating experiments of the C45 alloy heated with 300 C/s to (a) 770 C, 1 s; (b) 770 C,
10 s; (c) 770 C, 60 s; and (d) 900 C, 1 s. The arrow (T-T) indicates the direction of EPMA measurement (Fig. 11). Ferrite = white, martensite = gray, and bainite = black.
Fig. 11—Alloying element distribution along the line T-T in the C45
alloy (Fig. 10(a)). The minimum of the carbon concentration corresponds to the position of the black phase (bainite).
austenite, ferrite is, in its turn, resupplied with carbon by
the dissolution of cementite plates, thus creating a
constant driving force for the carbon diffusion toward
the austenite-ferrite interface (Figure 2).
During the experiments, both types of nucleation were
observed for the 0.05 C/s heating rate and only one
type of nucleation (within the pearlitic areas) for the
20 C/s heating rate. The EPMA measurements show a
significant variation in carbon content in austenite
nucleated on the ferrite-ferrite grain boundaries (about
0.27 wt pct, low-carbon austenite) and the one nucleated on pearlite-ferrite boundaries (around 0.8 wt pct,
high-carbon austenite, which is close to the equilibrium
value at temperature just above A1). The possible
presence of carbides on ferrite-ferrite grain boundaries
prior to transformation, which could stimulate the
austenite nucleation, is doubtful. First of all, they were
not observed in the initial microstructures, and second,
their presence on ferrite-ferrite grain boundaries would
lead to austenite nucleation on them for all heating
rates, which was not the case. This observation indicates
the importance of carbon diffusion during the transformation to the growing austenite phase. At the higher
heating rates, there is not enough time for carbon to
diffuse, so the austenite nucleation takes place only on
ferrite-pearlite boundaries in the vicinity of the carbon
Obviously, the diffusion of carbon is not an infinitely
fast process. The indication of the delay in cementite
dissolution compared to the austenite formation is
clearly seen in Figure 4(b) (the grain indicated by an
arrow). The pearlitic ferrite undergoes the transformation to austenite at higher temperatures, and upon
further quenching, it transforms into martensite. At the
same time, the cementite plates are still visible within
martensite. Similar observations were made in References 1, 2, and 11 through 14. They noticed that new
grains of austenite grow along the plates of ferrite in a
pearlite colony and expand to replace the ferrite in the
colony. The cementite plates do dissolve in this austenite, but the austenite grains grow into ferrite at a faster
rate than that at which the cementite plates dissolve.
This residual cementite is eventually first thin and then
spheroidizes and dissolves completely in austenite,
depending on the carbon content and temperature.
This points out the time difference between austenite
formation and cementite dissolution. Thus, the main
VOLUME 38A, MAY 2007—953
difference between the pearlite formation on cooling and
pearlite transformation on heating is that, on cooling,
the ferrite and cementite plates grow together, whereas
on heating, it is a two-step process and thus pearlite
should be considered as consisting from two different
phases: ferrite and cementite.
B. Growth Morphologies
The carbon diffusion process, the rate of which
increases with temperature, plays an important role in
the occurrence of growth morphologies. As was indicated in Section IV.B, several types of growth morphologies were observed. The finger-type austenite growth
morphologies (Figure 7) were spotted growing from the
former pearlitic areas into neighboring proeutectoid
ferrite grains. The average carbon concentration within
the fingers is around 0.8 wt pct, indicating that the
fingers grow from the carbon-rich areas. Thus, it can be
argued that the formation of the fingers coincides with
the position of former cementite plates that were
perpendicular to the grain boundary. After pearlite
transformation to austenite is completed, the newly
formed ‘‘pearlitic’’ austenite is inhomogeneous in composition, the carbon content being highest at the
location of pre-existing cementite plates.[11,15] Thus,
the driving force at these grain boundaries is highly
inhomogeneous. At low heating rates, which correspond
to a low degree of overheating, the driving force for
ferrite-to-austenite transformation is not high, and it is
energetically more favorable for austenite nuclei to grow
with minimum surface energy. Minimal surface energy is
secured by an orientation relationship between austenite
and ferrite. In the case of finger-type growth, the
coherent broad sides should have a relatively low
interfacial energy, whereas the incoherent tip would
have a much higher interfacial energy. The presence of
an orientation relationship between ferrite and austenite
was shown by Law and Edmonds,[6] who determined
that the austenite was always within 15 deg of a
Kurdjumow–Sachs orientation relationship:
11bcc ==½0
(110)bcc //(111)fcc , [1
For the bcc-fcc phase combination, these are the only
planes that are more or less identical in each crystal, and
by choosing the correct orientation relationship, it is
possible for a low energy coherent or semicoherent
interface to be formed. There are, however, no other sets
of planes of good matching and the austenite plate is
thus bounded by the incoherent interface. It is known
that an incoherent interface has a much higher mobility
than a coherent one.[16] The incoherent interface will
move as fast as diffusion allows and the growth will take
place under diffusion-controlled mode. At the very fast
heating rates, as the ones shown in Figure 10, ‘‘fingers’’
of austenite were not formed. This can be explained by
the lack of time for the diffusion to proceed because of
the relatively high driving forces for ferrite-to-austenite
954—VOLUME 38A, MAY 2007
In the places where the cementite plate direction is not
perpendicular to the grain boundary, the formation of a
‘‘black etching’’ structure on cooling was observed (for
example, Figures 7 and 10). The EPMA measurements
(Figure 10(a), line T-T, and Figure 11) show that the
black phase has a lower carbon content (around
0.2 wt pct) compared to the rest of the sample, which
is around 0.4 wt pct and is bainite (Figure 8). It is well
known that steels with lower carbon content require
higher cooling rates in order to obtain a martensitic
structure. Thus, in the structure with the carbon inhomogeneities, for the same cooling rate during the
quenching of the sample, it is possible to obtain
martensite in the areas rich in carbon and bainite in the
areas with lower carbon content. The formation of the
‘‘black’’ phase was noticed only on the former pearlite
(which is austenite oversaturated with carbon) and
proeutectoid ferrite (low in carbon) grain boundaries.
Hence, the formation of bainite on cooling is possible.
1. Formation of the austenitic phase upon heating is
possible in pearlite as well as in ferrite areas; however, in the first one, it proceeds at a much faster
rate due to the shorter diffusion distances. The carbon content of the austenite nuclei formed on the
ferrite-ferrite grain boundary is about 0.27 wt pct,
which is much lower than the equilibrium value
determined by the metastable phase diagram. A
possible explanation is proposed based on thermodynamic considerations for the formation of lowcarbon austenite (on ferrite-ferrite grain boundaries)
and high-carbon austenite (on pearlite-ferrite grain
2. Depending on the heating rate, the pearlite-to-austenite transformation can proceed in either one or
two steps. At low heating rates (0.05 C/s), the ferrite and cementite plates transform simultaneously.
At higher heating rates (20 C/s), a two-step process
is observed: first ferrite within the pearlite grain
transforms into austenite and then the dissolution
of the cementite lamellae takes place.
3. Carbon inhomogeneities give rise to specific phenomena in the a/c structure. The formation of finger-type austenite occurs on pearlite-ferrite grain
boundaries and coincides with the position of
cementite plates. In places where the direction of
cementite lamellae is not perpendicular to the grain
boundary, the formation of the ‘‘black’’ phase,
which is believed to be bainite, takes place.
This work is done within the framework of the
research program of the Netherlands Institute for Metals Research (, Project No. MC5.03171.
The authors thank Dr. Arjan Rijkenberg (Corus, the
Netherlands) for performing the SEM measurements.
1. G.R. Speich, V.A. Demarest, and R.L. Miller: Metall. Trans. A,
1981, vol. 12A, pp. 1419–28.
2. M. Hillert, K. Nilsson, and L.E. Torndahl: J. Iron Steel Inst.,
London, 1971, vol. 209, pp. 49–66.
3. S.K. Jayaswal and S.P. Gupta: Metallkd., 1992, vol. 83, pp. 809–
4. C. Garcia de Andres, F.G. Caballero, and C. Capdevila: Scripta
Mater., 1998, vol. 38 (12), pp. 1835–42.
5. V.I. Zel’dovich, I.V. Khomskaya, and O.S. Rinkevich: Phys. Met.
Metallogr., 1992, vol. 73 (3), pp. 250–65.
6. N.C. Law and D.V. Edmonds: Metall. Mater. Trans. A, 1980,
vol. 11A, pp. 33–46.
7. J.W. Christian: The Theory of Transformations in Metals and
Alloys, Pergamon, Amsterdam, The Netherlands, 2002, pp. 422–79.
8. I.I. Novikov: Teoria Termicheskoi Obrabotki Metallov {Theory of
Heat Treatment of Metals}, Metallurgia, Moskow, 1986, pp. 154–
9. S. Saunders, P. Karduck, and W.G. Sloof: Microchim. Acta, 2004,
vol. 145, pp. 209–13.
10. J.T. Armstrong: in Electron Probe Quantitation, K.F.J. Heinrich
and D.E. Newbury, eds., Plenum Press, New York, NY, 1991,
pp. 261–315.
11. G.A. Roberts and R.F. Mehl: Trans. ASM, 1943, Sept., pp. 613–
12. G.R. Speich, A. Szirmae, and M.J. Richards: Trans. TMS-AIME,
1969, vol. 245, pp. 1063–74.
13. C.R. Brooks: Principles of the Austenitization of Steels, Elsevier
Applied Science, London, 1992, pp. 81–144.
14. A. Jacot, M. Rappaz, and R.C. Reed: Acta Mater., 1998, vol. 46,
pp. 3949–62.
15. L.E. Samuels: Optical Microscopy of Carbon Steels, ASM, Metals
Park, OH, 1980, pp. 101–498.
16. D.A. Porter and K.E. Easterling: Phase Transformations in Metals
and Alloys, 2nd ed., Nelson Thornes Ltd., Cheltenham, United
Kingdom, 2001, pp. 171–80.
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