Formation of ferritic-bainitic structure with retained austenite in

Formation of ferritic-bainitic structure with retained austenite in

ADAM GOŁASZEWSKI, JERZY SZAWŁOWSKI, WIESŁAW ŚWIĄTNICKI
Formation of ferritic-bainitic structure
with retained austenite in hypoeutectoid steel
INTRODUCTION
EXPERIMENTAL STUDY
The very important aim of developing new grades of steel is to
obtain a specific structure which guaranties a combination of high
strength and high ductility. A newly developed family of
Advanced High Strength Steels (AHSS) meets such requirements
to a large extent [1÷4]. These include different multiphase steels
such as: Dual Phase (DP), Complex Phase (CP), and
Transformation Induced Plasticity (TRIP) steels [1], which give
the possibility of obtaining good mechanical properties through
manipulating with the type of phase, its volume fraction and
spatial arrangement. DP steels are ferritic-martensitic steels,
whereas CP steels consist of at least three phases amongst: ferrite,
bainite, martensite and austenite. Annealing in the intercritical
region plays an important role in the process of forming phase
structure in CP and DP steels, since it enables to manipulate with
the volume fractions of austenite and ferrite. Another interesting
subgroup of AHSS family is steels with TRIP effect, in which
carbon-enriched austenite is formed as a result of precise control
of phase transformations [2]. The aforementioned phase is
transformed into martensite during plastic deformation, which
prevents necking deformation and premature rupture [3]. The
presence of martensite in addition to ferrite (in the DP steels)
results in limited formability when plastic forming methods are
used [4]. Replacement of martensite with bainite in such grades of
steels may increase this property [4].
In this paper, an attempt to produce a complex-phase steel with
ferritic-bainitic structure with retained austenite was described. To
form such a structure the process of annealing in intercritical
region was used in order to obtain defined fractions of austenite
and ferrite, followed by quenching with an isothermal holding in
the bottom range of temperatures of bainitic transformation. It was
assumed that the presence of retained austenite would increase
ductility of steel due to the TRIP effect, whereas low-temperature
bainite would provide a high strength due to the significant grain
size reduction and limited carbides precipitation during transformation [5]. Such a type of bainite is obtained from austenite
containing the adequate amount of carbon and silicon [5], during
the process of annealing in the temperatures just above the MS
temperature. The research was performed on 35CrSiMn5-5-4 steel
which contained adequate amount of silicon, whereas the desired
concentration of carbon was expected to be obtained through the
intercritical annealing [6].
Steel was subjected to the appropriate heat treatment process,
which would allow obtaining desired microstructure. After heat
treatment the microstructure was carefully examined by means of
light microscopy and transmission electron microscopy (TEM).
Moreover, various mechanical tests were performed in order to
characterise hardness, strength, ductility and impact strength.
Chemical composition of the tested steel is shown in Table 1.
The JMatPro [7] computer program was used to select the
conditions for intercritical annealing in order to obtain specific
phase composition and to achieve the desired carbon concentration
in the austenitic phase.
The studies of phase transformations occurring in steel and the
preheat treatments were conducted with the use of Baehr DIL805L
dilatometer on samples with the diameter of 3 mm and the length of
10 mm. The annealing time in the two-phase region was constant at
all temperatures and was of 1 hour. The cooling was performed at
the rate of 50°C/s with the use of compressed helium. Ac1 and Ac3
temperatures of 35CrSiMn5-5-4 steel were revealed during the
heating process at 2°C/min (up to the standard) [8] and 0.28°C/min.
The obtained microstructure was examined with the use of
Nikon Eclipse MA200 light microscope and transmission electron
microscopy (TEM – JEOL JEM1200). The stereological analysis
of metallographic specimen was performed by the line secant
method [9] in order to characterize the microstructure and the
volume fraction of phases. The volume fraction of Vv() austenite
was estimated on the basis of the measurement of martensite
volume fraction in steel samples after thermal treatment.
The samples used in mechanical tests (static tensile test and
impact strength test) were annealed in the critical region in a gastight furnace in the nitrogen atmosphere, whereas their isothermal
quenching was conducted in the Sn bath. The hardness was
examined by Vickers method, with load of 2 kg. Tensile tests on
the fivefold samples of 6 mm in diameter were carried out on the
Zwick/Roell Z250 testing machine, using an extensometer with
25 mm base. Impact tests were performed on the samples with
a "V" notch with the use of the Charpy method.
Mgr Adam Gołaszewski ([email protected]), dr hab. inż. Jerzy
Szawłowski, dr hab. inż. Wiesław A. Świątnicki – Wydział Inżynierii Materiałowej,
Politechnika Warszawska
MODELLING OF PHASE COMPOSITION
IN THE INTERCRITICAL REGION
BY COMPUTER SIMULATIONS
In order to obtain a ferritic structure with a low-temperature
bainite and retained austenite, a ferritic-austenitic structure with an
expected carbon concentration in austenite must be formed at the
initial stage of the heat treatment. According to JMatPro
(Fig. 1a), the range of temperature for the intercritical region
of 35CrSiMn5-5-4 steel is 738÷805°C. Changes in volume
fraction of austenite and ferrite as a function of temperature are
presented in Figure 1a and the changes in volume fraction of the
phases occurring in steel are shown in Figure 1b.
Table 1. Chemical composition of 35CrSiMn5-5-4 steel, wt %
Tabela 1. Skład chemiczny stali 35CrSiMn5-5-4. % mas.
C
Mn
Si
Cr
Ni
Mo
0.35
0.95
1.3
1.31
0.14
0.018
Al
Cu
P
S
Ti
V
0.04
0.15
0.01
0.01
0
0.006
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Temperature in intercritical region, °C
The changes in carbon concentration in austenite in the
intercritical region are presented in Figure 1c. At the temperature
of 756°C the volume fraction of ferrite is equal to volume fraction
of austenite (~49.5%), the rest are carbides.
Cementite present in steel dissolves entirely at 757°C, whereas
MC and M7C3 carbides dissolve at 755°C and 773°C respectively
(Fig. 1b). The only non-metallic inclusions that remain in the
whole intercritical region are: manganese sulfide (MnS) and
a small amount of titanium carbon sulfide Ti4C2S2. Their presence
results from the occurrence of impurities. Complete dissolution of
carbides and carbon saturation of austenite occurs therefore at
773°C. The carbon concentration reaches a maximum value of
0.51 wt. % (Fig. 1c) at the temperature of 757°C, in which cementite dissolves completely and the austenite volume fraction is equal
to 50.4%. At the temperature of 773°C, in which all the carbides
are dissolved, austenite volume fraction equals 76% and carbon
concentration is still high (0.46%) as compared to the average
carbon concentration in steel (0.35%). Thus, both temperatures:
757ºC – in which the cementite is completely dissolved and 773°C
at which all the carbides are dissolved, may be considered to be
used in order to obtain a mixture of ferrite and low-temperature
bainite in 35CrSiMn5-5-4 steel. At these temperature a high
amount of austenite with high carbon concentration may be
obtained (Fig. 1c).
EXPERIMENTAL DETERMINATION OF Ac1, Ac3
AND MS TEMPERATURES
Temperature in intercritical region, °C
Temperature in intercritical region, °C
Fig. 1. JMatPro simulations for 35CrSiMn5-5-4 steel in the α + γ
temperature range: a) phase composition, b) carbide volume fraction,
c) carbon concentration in the austenite
Rys. 1. Symulacje z programu JMatPro dla stali 35CrSiMn5-5-4
w zakresie temperatury współistnienia α + γ: a) skład fazowy, b) udział
węglików, c) stężenie węgla w austenicie
Temperature of annealing in the intercritical region is the vital
parameter affecting austenite volume fraction in the structure of
multi-phase steels. The intercritical region occurs between Ac1 and
Ac3 temperatures. Therefore, it is crucial to determine such
temperatures adequately. The graph (Fig. 2) presents the values of
Ac1 and Ac3 temperatures for 35CrSiMn5-5-4 steel taken from
different sources: literature data [10], dilatometric tests and
computer simulations. According to the literature and standard
dilatometric studies (with the heating rate equal to 2°C/min), Ac1
temperature is between 770 and 780°C. However, the dilatometric
tests conducted at lower rate (0.28°C/min) than the standard rate
and JMatPro simulations show that the Ac1 temperature is lower.
Such decrease in Ac1 temperature may signify that the standard
heating rate is too high and does not guarantee equilibrium
conditions. The effect of heating rate on Ac1 and Ac3 temperatures
was noticed also by Pawłowski [11].
Fig. 2. Comparison of Ac1 and Ac3 temperatures obtained from
various sources. Dashed line indicates the annealing temperature
before isothermal quenching process, selected for structural and
mechanical testing
Rys. 2. Zestawienie temperatur Ac1 i Ac3 uzyskanych z różnych źródeł.
Przerywaną linią zaznaczono temperaturę wygrzewania przed procesem
hartowania izotermicznego wybraną do badań strukturalnych i mechanicznych
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Figure 2 also reveals the differences between the values of Ac3
temperature. The results obtained from dilatometric tests are
higher than these presented in the literature or obtained from
computer simulations. The increase in Ac3 temperature at low
heating rates observed during dilatometric tests was surprising,
since the drop was rather expected as in the case of Ac1
temperature.
Such an increase may result from the fact that some subtle
changes in length of the sample were observed in the dilatometric
curve for the lower heating rate and they might have been
unregistered at higher speed.
The dashed line presented in Figure 2 indicates the annealing
temperature (770°C) in the intercritical region, in which 56% of
austenite was obtained, after one hour of annealing. According to
the literature data and a standard procedure for determining the
characteristic temperatures, that temperature is below or exactly at
the same level as Ac1. It means that it would be impossible to
obtain such amount of austenite.
On the basis of dilatometric tests, the temperatures of the
beginnings of martensitic transformation (MS) were selected.
As expected, MS temperature decreases with the diminishing of
annealing temperature in the intercritical region. Such drop results
from an increase in the carbon concentration in austenite, and that
is confirmed by the simulations (Fig. 1c and Fig. 3). The carbon
concentration in austenite is the result of two parallel processes:
carbides dissolution and austenite formation. Carbides dissolution
for 35CrSiMn5-5-4 steel initiates at the beginning of austenite
formation, so that a small amount of newly formed austenite is
enriched in carbon from dissolved carbides (Fig. 1c). At higher
temperature, in which the carbides are no longer present, the
carbon concentration in austenite begins to decrease steadily until
reaching the average concentration level for the investigated steel.
This is only the consequence of an increase in austenite volume
fraction in the intercritical region.
This result confirms the validity of the adopted assumptions
and the correctness of the simulations, which showed that the
austenite was enriched in carbon to a level higher than the average
one during the annealing in the intercritical region.
Fig. 3. A change in MS temperature and carbon concentration in the
austenite in the intercritical region, in relation to annealing
temperature according to the JMatPro simulations
Rys. 3. Zmiana temperatury MS i stężenia węgla w austenicie według
JMatPro w funkcji temperatury wyżarzania w zakresie międzykrytycznym
AUSTENITE VOLUME FRACTION IN STEEL
In order to verify experimentally the austenite volume fraction in
steel as a function of temperature, the measurements of the phase
composition in the samples after the heat treatment were
conducted, with the use of a light microscope. It was assumed that
the austenite volume fraction is equal to the fraction of martensite
areas visible in steel structure after cooling to a room temperature.
The austenite volume fraction values obtained from the stereological analysis differ from the results obtained from the simulation, but their changes as a function of temperature exhibit the
same tendency as in the case of simulations (Fig. 4). Only at the
temperature of 760°C the values of austenite volume fraction are
significantly different from that obtained by simulation.
Such differences in the results obtained from the measurements
and simulations may be caused by the fact that during the onehour annealing in the intercritical region the equilibrium state was
not achieved.
At the temperature of 760°C, only 5% of austenite in steel
structure was obtained after annealing in the intercritical region.
This temperature is close to Ac1 temperature, in which a large
amount of carbides is present. However, at 770°C the amount of
obtained austenite (56%) was similar to the one determined by the
simulation for 760°C. Therefore, the temperature of 770°C was
selected as the optimal one for the further studies.
It can be remarked that even a slight change in the temperature
such as five degrees may result in a large change in the austenite
volume fraction in steel. This may cause some difficulties in
control of the phase composition during the heat treatment.
HEAT TREATMENT DESIGN
The next step in the study was to design and perform a complete
heat treatment, which would consist of the annealing in the
intercritical region and of austempering. On the basis of the
dilatometric tests, the temperature of isothermal annealing after
quenching was chosen to be equal to 300°C which was higher of
about 35 degrees than the MS temperature. In order to reduce the
time of the whole process, the isothermal holding lasted 1 hour,
which would be long enough to enable the transformation in 95%.
Dilatometric tests confirmed that during cooling to a room
temperature after isothermal quenching, martensitic transformation
did not occur, which signified that the retained austenite enriched
with carbon during the bainitic transformation was stable.
Fig. 4. The amount of austenite in 35CrSiMn5-5-4 steel in the selected
temperatures of intercritical annealing determined on the basis of the
JMatPro measurements and simulations
Rys. 4. Ilość austenitu w stali 35CrSiMn5-5-4 w wybranych temperaturach wygrzewania międzykrytycznego wyznaczona na podstawie
pomiarów i symulacji z programu JMatPro
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MICROSTRUCTURE OF STEEL
AFTER ISOTHERMAL TREATMENT
In order to verify whether the designed heat treatment would lead
to the desired phase structure, the microstructure observations
were carried out with the use of light microscope and TEM.
35CrSiMn5-5-4 steel after the incomplete austenitization and
isothermal quenching is homogenous and consists of: bainitic
grains embedded in a ferritic matrix containing carbides (Fig. 5a).
Carbides were observed only in ferrite, which means that they
precipitate during the soft annealing process, and that they did not
dissolve during the annealing in the intercritical region. It might
have resulted from too low temperature of intercritical annealing
for the carbides dissolution or from too short time of the
annealing.
The TEM observations showed that bainite grains were
composed of the ferrite laths separated by thin austenite layers and
are carbide-free (Fig. 5b). In fact, the diffraction patterns of bainite
grains contained reflections coming only from ferrite and
austenite. The carbide reflections were not present in diffraction
patterns of bainite. These results confirm that in steel submitted to
intercritical annealing followed by low temperature austempering
it is possible to obtain a carbide free bainite containing retained
austenite.
The images of the microstructures obtained by light microscope
(Fig. 6a and b) are alike. The differences may result from the
etching process of such a complex structure.
Fig. 6. Microstructure of steel after incomplete austenitization at
770°C for 1 hour and after isothermal quenching at 300°C for 1 h:
a) microstructure of a dilatometric sample, b) microstructure of
a sample for mechanical testing; LM
Rys. 6. Mikrostruktura stali po niepełnej austenityzacji w 770°C przez
1 h i hartowaniu izotermicznym w 300°C przez 1 h: a) mikrostruktura
z próbki dylatometrycznej, b) mikrostruktura próbki do badań mechanicznych; LM
MECHANICAL PROPERTIES
Fig. 5. Microstructure of steel after incomplete austenitization at
770°C for 1 hour and isothermal quenching at 300°C for 1 h:
a) microstructure of a dilatometric sample, b) bainitic microstructure;
TEM
Rys. 5. Mikrostruktura stali po niepełnej austenityzacji w 770°C przez
1 h i hartowaniu izotermicznym w 300°C przez 1 h: a) mikrostruktura
z próbki dylatometrycznej, b) mikrostruktura bainitu; TEM
The mechanical properties of the examined steel after the
intercritical annealing and quenching with isothermal holding were
compared to the properties of steel samples obtained after two
kinds of heat treatment:
(A) full austenitization at 900°C for 1 h followed by quenching
with isothermal holding at 310°C for 2 h,
(B) conventional quenching and tempering, which was carried out
in two stages:
(1st) stage-quenching from 900°C in the oil followed by tempering
at 700°C,
(2nd) stage-quenching from 890oC in the oil and then tempering at
230°C [9].
The results of mechanical tests (Fig. 7) conducted on steel
samples after the full austenitization and after austempering
(treatment A) are similar to those obtained after quenching and
tempering (treatment B).
These results are significantly different from those obtained for
steel after the treatment presented in this study with the annealing
in the intercritical region (incomplete austenitization) and
isothermal quenching at 300ºC for 1 h. The greatest differences are
visible in terms of hardness (HV) and the yield strength (R0.2).
Hardness and yield strength decrease about twice in comparison to
the sample isothermally quenched from full austenitization and
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Fig. 7. Comparison of the mechanical properties after isothermal quenching from incomplete and complete austenitization (A) and after
quenching and tempering (B)
Rys. 7. Zestawienie właściwości mechanicznych po hartowaniu izotermicznym z niepełnej i pełnej austenityzacji (A) oraz po ulepszaniu cieplnym (B)
after quenching and tempering. Low yield ratio (R0.2/Rm – 0.56), as
compared to the corresponding value of that parameter (0.79) in
steels after the reference treatments (A and B), shows that the
applied treatment should improve the formability.
After such a treatment, the mechanical strength (Rm) is lower
than the one for the isothermal quenching after the full
austenitization and lower than the one for quenching and
tempering. Nevertheless, it still remains at a high level
(1200 MPa). Ultimate elongation (A5) for all three treatments is
similar and equal to about 10%.
The stress-strain curves are shown in the diagram (Fig. 8) for
the samples after isothermal quenching with the full (a) and the
incomplete (b) austenitization. The curves overlap for each of the
treatments, which means that every one of them would lead to
a homogeneous structure and to the similar properties.
SUMMARY
Fig. 8. Stress-strain curve after incomplete austenitization (a) and
after full austenitization (b)
Rys. 8. Krzywe rozciągania po niepełnej austenityzacji (a) oraz po pełnej
austenityzacji (b)
The ability to determine the phase composition is vital in terms of
designing the heat treatments for the multiphase steels. The
JMatPro computer simulations are an effective and useful tool for
modeling the phase composition. By means of such simulations, it
was possible to determine the conditions of the heat treatment for
35CrSiMn5-5-4 steel, what allowed to obtain the phase composition as it was planned. Further experimental studies confirmed
the results of computer simulations to a large extent.
The critical temperatures of the phase transformations as well
as the thermal stability of carbides in steel are necessary for
designing the heat treatments of complex phase steel. In order to
design the annealing treatments in the intercritical region, the
initial and final temperatures of austenite formation should be
precisely determined. The results presented in this study proved
that the standard measurements of Ac1 and Ac3 temperatures are
insufficient. The results obtained from the JMatPro simulations
and from the dilatometric tests carried out with a low heating rate
(lower than recommended) turned out more realistic. These
temperatures are only slightly different from the values obtained
under standard conditions, which are sufficient for designing
classic heat treatments after full austenitization, where the
austenitization is always performed at the temperature at least
several degrees above the Ac3 temperature. However, selection of
the annealing temperature in the intercritical region requires much
more precision.
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The investigations have shown that in 35CrSiMn5-5-4 steel
a microstructure composed of bainitic grains, embedded in ferritic
matrix, is formed during a heat treatment consisting of an
intercritical annealing followed by the low temperature
austempering. The bainite formed is carbide-free and is composed
of thin ferritic lath separated by the retained austenite layers.
The mechanical properties of investigated steel significantly
vary in relation to the applied heat treatment. The steel after the
isothermal quenching from the incomplete austenitization has low
hardness and low yield ratio with the high mechanical strength in
comparison to steel after isothermal quenching from the full
austenitization and to steel after quenching and tempering. The
application of newly-developed heat treatment leads to the
interesting properties of steel different from the properties of steel
after conventional thermal treatments.
ACKNOWLEDGMENT
The study was accomplished within the Structural Project
“Nanocrystalline structure formation in steels using phase
transformation” No. POIG.01.01.02-14-100/09 co-financed by
EU, within the funds of the Operational Programme Innovative
Economy, 2007÷2013 and supported by Warsaw University of
Technology.
Special thanks to Prof. Adam Grajcar for his valuable comments and to MA Dorota Pietrzyk for her help with translation.
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