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Kovove Mater. 58 2020 129–140
DOI: 10.4149/km 2020 2129
129
Mechanical properties of carbide-free lower bainite
in complex-alloyed constructional steel:
Effect of bainitizing treatment parameters
V. I. Zurnadzhy1,V.G.Efremenko
1*, I. Petryshynets2, K. Shimizu3,M.N.Brykov
4,
I. V. Kushchenko1,V.V.Kudin
4
1Pryazovskyi State Technical University, Mariupol, Ukraine
2Institute of Materials Research, Slovak Academy of Sciences, Kosice, Slovak Republic
3Muroran Institute of Technology, Muroran, Japan
4”Zaporizhzhia Polytechnic” National University, Zaporizhzhia, Ukraine
Received 10 December 2019, received in revised form 19 January 2020, accepted 21 January 2020
Abs tract
Tensile/impact behaviour of lower bainite obtained in high-Si steel 55Si3Mn2CrMoVNb
was studied using SEM, TEM, and XRD. Specimens were austenitized at 900◦
C and isother-
mally treated at 250, 270, and 300◦
C with holding up to 600 min. The heat treatment results
in the formation of cementite-free lower bainite/retained austenite structure, where retained
austenite was found as blocky “islands” and interlaths “films”. The width of bainitic ferrite
laths decreases from 170–240 µm to 45–80 µm with holding temperature decreasing. This re-
sults in increasing UTS (to 1700 MPa) and hardness (to 52 HRC). The optimal combination of
mechanical properties (UTS 1397–1522 MPa, hardness 45–47 HRC, total elongation 18–21 %,
U-notched impact toughness 105–139 J cm−2) refers to holding at 300◦
C to be associated
with higher amount of retained austenite (30–33 %). With prolonging the bainitizing dura-
tion the hardness and ductility decreases while impact toughness increases. Prolonged holding
at 300◦
C leads to a continuation of bainite transformation and precipitation of transitional
carbides within ferrite laths.
K e y w o r d s : carbide-free lower bainite, retained austenite, phase transformations, mi-
crostructure, mechanical properties
1. Introduction
Currently, various heat treatment technologies are
actively used to improve mechanical properties and ex-
ploitation durability of steel machine parts. One of the
most effective approaches is the isothermal quench-
ing (austempering) consisting in austenitization fol-
lowed by the holding (bainitizing) at a temperature
slightly higher than Ms point with eventual final cool-
ing [1]. Isothermal quenching results in lower bainite
structure, which provides an improved combination of
strength, ductility, and impact toughness [2]. In ad-
dition, isothermal quenching leads to lower residual
stresses with no distortion or cracking in the products,
which is associated with temperature equalization in
the products during the isothermal holding [3].
*Corresponding author: e-mail address: vgefremenko@gmail.com
Beginning from the pioneering research of H.
Bhadeshia et al. in the 90s and to the present, the
approach to develop the high-strength steels with
nanostructured carbide-free bainite (nanobainite) has
been actively studied [4–6]. The main peculiarity of
this structure is nanoscaled (30–60 nm width) laths
of bainitic ferrite with interlayers of carbon-enriched
retained austenite (RA) with no cementite carbides
[7, 8]. Nanosized transition ε-carbide may precipitate
within ferritic laths, providing additional strength-
ening effect [9, 10]. The formation of carbide-free
nanobainite is ensured by alloying with increased (1.5–
3.0 wt.%) amount of silicon or aluminium to sup-
press the cementite precipitation during bainite trans-
formation [5, 11, 12]. Preventing cementite precipita-
tion eventually results in austenite enrichment with
130 V. I. Zurnadzhy et al. / Kovove Mater. 58 2020 129–140
carbon leading to increase in RA volume fraction
[13]. Nanoscaled ferritic laths impart higher strength
(reaching 2000–2400 MPa), while the absence of ce-
mentite carbides and the presence of retained austen-
ite films provide the increased ductility (elongation
of 10 % or higher) and impact toughness [5, 14, 15].
Moreover retained austenite may significantly con-
tribute to mechanical properties due to TRIP-effect
[16]. With that in mind, carbide-free lower bainite is
considered as a desirable microstructural component
of multiphase steels such as TRIP-aided steels [17] and
Q&P-treated (Quenching and Partitioning) steels [18,
19].
The 0.3–0.4 wt.% C steels with bainite were suc-
cessfully used in the production of rails for particularly
difficult operating conditions, such as heavily loaded
curved sections of the railroad [20–22]. Nanobainite
steels are known to be applied for the armours con-
struction [23]. They are also considered as perspec-
tive materials with increased fatigue durability and
wear resistance [24–27]. Improved wear behaviour of
carbide-free lower bainite is associated with the trans-
formation of retained austenite into martensite under
wear, providing an increase in surface hardness [28]
and stress relaxation [29]. The suitability of nanobai-
nite steels as elements of welded constructions is in-
vestigated in [30] focusing on the properties in the
heat-affected zone.
The production of nanobainite steels involves long-
lasting isothermal holding at a temperature below
250◦
C reaching up a week or more [4, 5, 31] which
is detractive for steels manufacturers. This refers to
a very low rate of bainite transformation at tem-
peratures below Ms point [32]. Attempts to accel-
erate the low-temperature transformation led to the
development of nanobainitic steels alloyed with alu-
minium and cobalt [33]. Another approach is obtain-
ing a certain fraction of athermal martensite in the
structure that drastically reduces the incubation pe-
riod of bainitic transformation [34]. The first approach
increases the cost of steel, and the second one leads to
reducing ductility because of the appearance of brittle
martensite in the structure.
The alternative approach to nanobainite is lower
carbide-free bainite obtained at the temperatures
above Ms point; its formation does not involve too
long isothermal holding [14]. Appropriate design of
chemical composition could allow obtaining carbide-
free lower bainite even at temperatures above Ms [12].
However, tensile/impact behaviour of this microstruc-
ture should be studied additionally to prove its com-
petitiveness to nanobainite. Besides, the effect of pro-
longed isothermal holding (longer than it needs to
complete bainite transformation) on mechanical prop-
erties of lower bainite has not been studied so well
to choose the optimal duration of isothermal holding
[33]. Based on above the object of this work was to
evaluate the mechanical properties of newly designed
Ni-free steel 55Si3Mn2CrMoVNb after heat treatment
for the structure of carbide-free lower bainite. The
specific tasks of the research were: (a) to investigate
the effect of bainitizing treatment parameters on ten-
sile and impact behaviour of steel focusing on pro-
longed isothermal holding; (b) to explore the relation-
ship “microstructure/mechanical properties” concern-
ing lower bainite microstructure.
2. Methods
The study material was the steel 55Si3Mn2CrMo
VNb containing: 0.56 wt.% C; 2.50 wt.% Si; 1.70 wt.%
Mn; 0.50 wt.% Cr; 0.21 wt.% Mo; 0.12 wt.% V;
0.05 wt.% Nb; 0.006 wt.% S; 0.015 wt.% P. The chem-
ical composition of steel was specifically designed
to gain an advantageous Q&P treatment to form
complex microstructure “martensite/carbide-free bai-
nite/RA” [35]. The steel was smelted in laboratory
induction furnace of 60-kg capacity and rolled into
a strip 15 mm thick, from which the specimens were
manufactured for research. Heat treatment was ful-
filled in an electric muffle furnace (austenitization)
and in liquid 60 wt.% Sn-40 wt.% Pb bath (bainitiz-
ing holding). The used modes of heat treatment are
describedinsection3.1.
Optical mirror dilatometer was used to determine
the steel critical temperatures (Ac1,Ac3,Ms)and
the kinetics of austenite/bainite transformation. The
specimen of 2 mm in diameter and 20 mm in length
was used for the dilatometric study. The specimen was
heated with the rate of 1 K s−1to 900 ◦
Candthen
cooled at still air. The heating/cooling curves were
recorded by light beam deflection fixed on the screen,
located at a distance of 600 mm from the mirror. The
critical points were found by the inflexion of heat-
ing/cooling curves connected with phase transforma-
tions. In the case of the kinetics of bainite transforma-
tion preliminarily austenitized at 900 ◦
C specimen was
transferred to bainitizing dilatometer furnace where it
was held with dilatometric curve recording.
Tensile properties were determined by a tensile test
using the samples of 5 mm diameter and 30 mm gauge.
Impact toughness was measured by the Charpy test
at room temperature using U-notched specimens of
7×10 ×55 mm3size. Hardness was measured by
Rockwell according to C scale. The microstructures
and fracture surfaces were observed using Ultra-55
(Carl Zeiss) scanning electron microscope (SEM). The
fine microstructure was examined using JEM-100-C-
-XII (JEOL) transmission electron microscope (TEM)
at an accelerating voltage of 100 kV. The phase sta-
tus of the steel was determined by the XRD method
with diffractometer Pro-IV (Rigaku) in CuKαradia-
tion. The volume fraction of retained austenite and
V. I. Zurnadzhy et al. / Kovove Mater. 58 2020 129–140 131
carbon content in RA were calculated as described in
[36, 37].
3. Results
3.1. Critical points determination and
transformation kinetics study
The selection of the parameters of bainitizing treat-
ment was preceded by a determination of the critical
points (Ac1,Ac3,andMs) and by studying the ki-
netics of austenite transformation in the bainite tem-
perature range (225–350◦
C). As follows from inflec-
tions on dilatometric heating/cooling curves (Fig. 1)
the critical points of steel 55Si3Mn2CrMoVNb were
determined as 800◦
C(Ac1), 840 ◦
C(Ac3), and 240 ◦
C
(Ms). Taking into account Ac3position the tempera-
ture of austenitization before bainitizing holding was
chosen as 900◦
C with holding of 10 min.
Figure 2 depicts the kinetics of austenite trans-
formation in steel 55Si3Mn2CrMoVNb at 250, 270,
300, and 350◦
C after austenitization at 900◦
C. As
follows from the analysis of kinetic curves (Fig. 2a),
the minimum incubation period corresponds to 350 ◦
C
(55 s). With holding temperature decrease the incuba-
tion period gradually increases to 420 s at 250◦
C.
As seen, kinetics curves approach the horizontal
line at different beam deflection. Beam deflection on
dilatogram (dilatation effect) is influenced by different
transformation features such as rate of γFe →αFe
transition, carbon partitioning from bainitic ferrite
(BF) to austenite, carbide precipitation, etc. Thus,
light beam deflection is a superposition of mentioned
factors; therefore, it does not directly correlate with
the transformation rate only. The minimum beam de-
flection was noted for holding at 300◦
C which reflects
a very low rate of transformation at this temperature.
Supposedly, during holding at 300◦
C, the transforma-
tion proceeded very slowly due to active enrichment
of austenite by carbon leading to transformation sup-
pressing on its earlier stage. The carbon enrichment
at 300◦
C presumably was facilitated by higher car-
bon diffusivity and by presence of 2.5 wt.% Si which
suppresses cementite precipitation from austenite [5].
Increasing holding temperature to 350 ◦
C stimulated
the transformation due to an increase in carbon dif-
fusion and cementite precipitation that resulted in
higher transformation rate with much more signifi-
cant dilatation effect. At 270◦
C the transformation de-
veloped more thoroughly as well in comparison with
300◦
C due to less austenite enrichment with carbon.
Moreover, the dilatation effect at 270 ◦
C exceeds the
one for 350◦
C that can be explained by increasing the
carbon content in BF. Accordingtopreviousresearch
based on XRD study [11, 38, 39], the carbon content
Fig. 1. Dilatometric curves for (a) heating and (b) cooling
to determine critical temperatures of steel 55Si3Mn2CrMo
VNb.
in BF increases with transformation temperature de-
crease that results in increasing ferrite specific volume.
This conclusion was confirmed lately by Rementeria
et al. [40]. Using atom probe tomography, the authors
[40] revealed in BF lath the carbon-clustered regions
which are responsible for an increase in total tetrago-
nality with transformation temperature decrease [41].
Isothermal holding at 250◦
C unexpectedly resulted in
lower curve height as compared with 270◦
Cdespite
lower content of retained austenite. This behaviour
may occur because of lower transformation rate at
250◦
C when the specimen extending (as a result of
γFe →αFe phase transition) was partially compen-
sated with specimen shortening due to carbon parti-
tioning from BF to austenite and ε-carbide precipita-
tion.
Using the data derived from kinetics curves, the
bainite domain of TTT-diagram was constructed
(Fig. 2b). Taking into account the data of critical
points and transformation kinetics the parameters of
132 V. I. Zurnadzhy et al. / Kovove Mater. 58 2020 129–140
Fig. 2. (a) Kinetics curves of austenite transformation and (b) TTT-diagram of steel 55Si3Mn2CrMoVNb with schedules
of bainitizing heat treatment used in work.
bainitizing heat treatment were determined as graph-
ically shown in Fig. 2b. The specimens were austen-
itized at 900◦
C for 10 min with further fast trans-
fer to liquid Sn-Pb bath for isothermal holding at
temperatures slightly above the Ms point which are
specifically 300, 270, and 250◦
C. Holding duration var-
ied in the range of 45–600 min in order to: (a) reach
the moment of bainite transformation accomplish-
ment according to TTT-diagram (which are 45 min
for 300◦
C, 160 min for 270 ◦
C, 400 min for 250 ◦
C)
and (b) study the effect of further prolonged hold-
ing on bainite’s mechanical properties. After hold-
ing completion, the specimens were cooled in calm
air.
3.2. Mechanical properties evaluation
Table 1 presents the mechanical properties of
studied steel after bainitizing heat treatment. As
seen, the treatment resultedinhighstrength(UTS:
1397–1733 MPa, YTS: 1162–1496 MPa) while ductil-
ity retained at an increased level (Total Elongation
V. I. Zurnadzhy et al. / Kovove Mater. 58 2020 129–140 133
Ta b l e 1. Mechanical properties of steel 55Si3Mn2CrMoVNb after bainitizing treatment
YTS UTS TEL PSE KCU
Heat treatment Hardness, HRC
(MPa) (%) (GPa%) (J cm−2)
300◦
C (45 min) 1210 1522 18 27.5 47 105
300◦
C (120 min) 1162 1397 21 28.8 45 116
300◦
C (240 min) 1171 1397 12 16.9 43 139
270◦
C (160 min) 1353 1585 10 15.1 50 126
270◦
C (240 min) 1337 1542 12 18.4 49 133
270◦
C (300 min) 1447 1604 11 17.5 49 136
250◦
C (400 min) 1456 1697 12 20.9 52 107
250◦
C (600 min) 1438 1733 10 17.0 48 114
Fig. 3. Effect of holding duration at different bainitizing temperature on mechanical properties of steel 55Si3Mn2CrMoVNb.
(TEL): 10–21 %). PSE index (Product of Strength
and Elongation) varied from 15.1 to 28.8 GPa%.
The impact toughness values were in the range of
105–139 J cm−2. These data reveal that middle car-
bon steel 55Si3Mn2CrMoVNb acquired an advanced
“strength/ductility/impact toughness” combination
wherein its ductility/toughness level is rather char-
acteristic for low-carbon steels.
The maximum strength (UTS 1733 MPa, YTS
1697 MPa) was achieved after bainitizing at 250 ◦
Cfor
600 min. With such strength, steel performed accept-
able ductility (TEL 10 %) and higher impact tough-
ness (KCU 114 J cm−2). The optimal complex of me-
chanical properties was ensured by holding at 300◦
C
for 45–120 min. In this case, the steel possessed a
tensile strength of 1397–1522 MPa and hardness of
45–47 HRC with total elongation of 18–21 %. Notably,
impact toughness values were 105–139 J cm−2while
134 V. I. Zurnadzhy et al. / Kovove Mater. 58 2020 129–140
PSE index reached its maximum values (27.5–28.8
GPa%).
The effect of bainitizing parameters on the mecha-
nical properties of steel is shown graphically in Fig. 3.
There is pronounced trend ofincreasingofstrength
and hardness as the holding temperature decreases
(Fig. 3a). Conversely, ductility and impact toughness
tend to decrease with bainitizing temperature decrease
(Figs. 3b,e).
As seen in Fig. 3a, holding duration slightly de-
creases UTS at 300◦
C while it has no significant ef-
fect on UTS at 270 and 250◦
C. In contrast, hardness
dropped with holding duration at any bainitizing tem-
perature (Fig. 3d). Holding at 250◦
C for 400 min gives
the maximum hardness (52 HRC) though it tends to
decrease to 48 HRC after prolonged holding (600 min).
The minimum hardness value (43 HRC) corresponds
to bainitizing at 300 ◦
C for 240 min.
For ductility other trend was observed. The hold-
ings at 270 and 250◦
C provided no effect on TEL,
resulting in approximately the same total elonga-
tion (10–12 %). In contrast, at 300 ◦
CTELvariedby
the curve with maximum corresponded to 120 min
(Fig. 3b).
Impact toughness increases monotonically with in-
creasing holding duration at all temperatures. This
trend is more pronounced at 300◦
C,andtoalesser
extent, at 250◦
C (Fig. 3e). Analyzing the data shown
in Fig. 3e, it should be noted that for each bainitizing
temperature impact toughness reaches a rather high
level (over 100 J cm−2), which is usually observed in
steels alloyed with 2–4 wt.% Ni [42]. Holding at 270–
300◦
C ensures a higher level of impact toughness com-
pared with holding at 250◦
C.
3.3. Microstructure and fracture surface
characterization
The microstructure of bainitized steel is presented
in Fig. 4. The characteristic feature of all heat treat-
ment regimes is a fine grain (number 10–10.5) due
to alloying by strong carbide-forming elements (V,
Nb). As seen in Fig. 4a, bainitizing at 300 ◦
Cresults
in a structure consisting of packages of ferritic α-
phase laths with retained austenite areas distributed
throughout the structure. The decrease in bainitizing
temperature to 250◦
C led to decrease in the bainitic
laths width and RA volume fraction, which is in agree-
ment with the data presented in [4, 5] (Fig. 4c).
The features of the fine structure of lower bainite
were revealed using transmission electron microscopy
(Fig. 5). The lower bainite is composed of the packets
consisting of parallel α-phase laths divided by film-
like retained austenite; blocky “islands” of austen-
ite are also observed at the junction of the packets
(Fig. 5a,b). The presence of austenite films is con-
firmed by dark-field observation in the austenite re-
Fig. 4. Microstructure (SEM) of steel 55Si3Mn2CrMoVNb
after holding at (a) 300 ◦
C for 240 min, (b) 270◦
Cfor
300 min, (c) 250 ◦
C for 400 min (A, B – austenite, bainite,
accordingly).
flection (Fig. 5c). The thickness of the bainitic α-
phase laths decreases from 170–240 nm at 300 ◦
Cto
45–80 nm at 250◦
C; thus, at the latter case, bainite
is nanostructured. The α-phase laths have a higher
density of dislocation, induced by accommodative mi-
V. I. Zurnadzhy et al. / Kovove Mater. 58 2020 129–140 135
Fig. 5. TEM microstructure of steel 55Si3Mn2CrMoVNb after holding at (a) 250◦
C for 400 min and (b–e) 300◦
Cfor
240 min; (a, b) bainitic α-phase laths with “film” and blocky “island” of retained austenite; (c) dark-field image of
austenite reflection from the axis of the [111]γzone; (d) dislocation structure of α-phase laths; (e) transitional carbides
inside α-phase lath (A – retained austenite).
croplastic deformation during displacive γ→αtrans-
formation (Fig. 5d). The analysis of TEM images and
diffraction patterns revealed no presence of cementite
carbide, which confirms the carbide-free nature of ob-
tained lower bainite. Instead, the transitional (ε)car-
bides were revealed to be precipitated within ferrite
laths (Fig. 5e).
Evaluation of the phase status of bainitized steel
was performed using XRD analysis (Fig. 6). Diffrac-
tion peaks related to α-Fe ((110)α, (200)α, (211)α,
(220)α) and to austenite ((111)γ, (200)γ, (220)γ,
(311)γ) were revealed in all XRD patterns. It was
found that the integral intensity of the austenite peaks
increases with holding temperature increase. The vol-
136 V. I. Zurnadzhy et al. / Kovove Mater. 58 2020 129–140
Fig. 6. XRD patterns of heat-treated specimens.
Fig. 7. The effect of holding duration on (a) RA volume fraction and (b) carbon content in RA.
ume fraction of retained austenite and carbon content
in RA are presented in Fig. 7 as a function of holding
duration. The highest RA volume fraction (32.6 vol.%)
was measured for holding at 300◦
C for 45 min, while
the carbon concentration in austenite was 0.94 wt.%,
which is almost two-fold of the average carbon con-
tent in the steel. With an increase in holding duration
to 240 min, the RA volume fraction slightly decreases
to 30.6 vol.% with an increase in carbon concentra-
tion to 1.08 wt.%. After holding at 270 and 250◦
C,
RA volume fractions are much less to be 18–19 vol.%
and 18.1 vol.% accordingly, while the carbon concen-
tration in RA reaches about 1.05 and 0.94 vol.%, re-
spectively. At these temperatures the effect of holding
duration on RA volume fraction/C content is not so
pronounced as for 300◦
C. XRD data are in accordance
with the above finding of sharp suppression of bainite
transformation at 300◦
C (Fig. 2a) because of strong
chemical stabilization of austenite (with C) resulting
in increased RA volume fraction.
The question arises if some of alpha-phase is
fresh martensite formed under the final cooling af-
ter isothermal holding. This martensite obviously may
form from austenite remaining to the end of holding. It
V. I. Zurnadzhy et al. / Kovove Mater. 58 2020 129–140 137
Fig. 8. Fractured surface of the specimens after impact
loading. Bainitizing at 300 ◦
C for (a) 45 min and (b)
240 min.
depends on Ms temperature, which can be calculated
by using the empirical equation [38]:
Ms (◦
C) = 525−350(C−0.05)−45Mn −35V −30Cr−
20Ni −16Mo −5Si −16Al + 6Co,(1)
where the symbols present the concentration of the
elements (wt.%).
Assuming that carbon content in retained austen-
ite is 0.94–1.08 wt.% (as revealed from XRD measure-
ments) and taking the elements (Si, Mn, Cr, Mo) con-
tent in steel 55Si3Mn2CrMoVNb as austenite chem-
ical composition, then calculated Ms temperature is
53–102◦
C. That means that fresh martensite may ap-
pear after isothermal holding at 250–300◦
C, contribut-
ing to steel strength. However, its volume fraction
should be negligible since Ms is close to room tempe-
rature. Furthermore, fresh martensite should appear
within small austenite areas making its observation
difficult even using TEM technique.
U-notched fractured specimens were characterized
with SEM to reveal fracture behaviour of carbide-free
lower bainite under impact loading. In general, all the
specimens performed rather ductile rupture mecha-
nism. Figure 8 depicts the fractured specimens baini-
tized at 300◦
C. The specimen held for 45 min has the
fractured surface consisting of trans-crystalline quasi-
cleavage facets (shown by arrows), surrounded by ar-
eas of copious dimples (Fig. 8a). The characteristic
brittle “river-like” pattern was not found. With the
prolongation of holding to 240 min the area fraction
of the ductile component (dimples) increases, which
is consistent with the increase in impact toughness
(Fig. 8b).
4. Discussion
As follows from Fig. 3, two main trends connected
with bainitizing treatment regime were observed. The
first is that the strength/hardness of lower bainite in
steel 55Si3Mn2CrMoVNb is inversely proportional to
holding temperature. Based on microstructural char-
acterization this behaviour can be explained by a de-
crease in the thickness of bainitic ferrite plates [4, 5].
Another probable reason is an increase in carbon con-
tent in bainitic α-phase because of lower carbon dif-
fusivity [11, 42].
The second trend is connected with an effect of
prolonged holding on bainite’s mechanical properties
that is of great interest. As seen in Fig. 3, holding du-
ration at 250 and 270◦
C very slightly affects strength
and ductility. On the contrary, holding at 300◦
Cleads
to decrease in strength and to non-monotonous vari-
ation of TEL with a maximum at 120 min. As to
hardness and impact toughness, they were signifi-
cantly influenced by holding duration at any bainitiz-
ing temperature. Specifically, hardness decreased and
impact toughness rose with holding duration increase.
The latter reached an excellent level (>130 J cm−2)
which is beneficial for Ni-free steel with increased
strength/hardness.
As was proven by TEM and XRD study, trend
mentioned above is caused by superposition of struc-
tural processes taking place under prolonged holding,
namely: (a) precipitation of transitional carbide in-
side ferrite laths, (b) gradual prolongation of bainite
transformation with eventual decrease in RA volume
fraction, and (c) enrichment of retained austenite with
carbon. Precipitation of ε-carbide led to a decrease in
carbon content in ferrite resulting in a decrease in the
micro distortion of its crystal lattice (i.e., softening
effect). This consequently reduced the hardness while
improved impact toughness due to promoting dislo-
cation movement. Besides, copious ε-carbide precipi-
tates inside bainitic ferrite acted as nuclei for dimples
formation which ensured the ductile mechanism of im-
pact fracture.
138 V. I. Zurnadzhy et al. / Kovove Mater. 58 2020 129–140
Notably, the carbon content in RA at 300◦
Cin-
creased with holding duration increase (Fig. 7b). This
can be considered as a proof of bainite transformation
continuation under prolonged holding since the trans-
formation is accompanied by carbon partitioning to
austenite. The higher transformation rate the higher
C content in remaining austenite that is promoted by
suppression of cementite precipitation due to alloying
with 2.5 wt.% Si.
It should be noted that prolongation of duration
at 300◦
C to 240 min leads to decrease in ductility
while impact toughness tends to be increased. This
means that total elongation is more sensitive to RA
volume fraction while impact toughness is somewhat
influenced by the strength of ferritic laths. Thus, the
decrease in retained austenite is detrimental for TEL
while the softening of ferrite (due to ε-carbide precip-
itation) is beneficial for impact toughness.
The results obtained show that the structure of
carbide-free lower bainite provides an advanced com-
plex of strength, ductility, and impact toughness in
medium-carbon Ni-free steel (55Si3Mn2CrMoVNb),
which makes it promising as high-strength material for
use in various operating conditions combining static
and dynamic loading. It is essential that this structure
is obtained after significantly shorter holding as com-
pared with nanobainitic steels thus providing lower
production costs.
6. Conclusions
Based on the results obtained, the major conclu-
sions were drawn from this research:
(a) Bainitizing at 250–300◦
C with holding up to
600 min results in formation in steel 55Si3Mn2CrMo
VNb the microstructure of carbide-free lower bainite
consisting of bainitic ferrite laths and retained austen-
ite of blocky or interlath film morphologies. Cementite
carbides were not detected in the structure; instead,
transitional carbides precipitated within ferritic laths.
Lower bainite obtained at 250 ◦
C is nanoscaled having
ferrite laths of 45–80 nm width.
(b) Carbide-free lower bainite structure ensured
an advanced complex of strength/ductility/impact
toughness in steel 55Si3Mn2CrMoVNb. The optimal
properties combination (UTS = 1397–1522 MPa, TEL
= 18–21 %, KCU20 = 105–116 J cm−2, PSE = 27.5–
28.8 GPa%) is achieved by holding at 300◦
Cfor
45–120 min. This combination refers to an increased
amount of retained austenite (30–33 %), stabilized by
carbon (up to 1.08 wt.%) during bainitic transforma-
tion. The bainitized steel is fractured under impact
loading by a predominantly ductile mechanism com-
bining trans-crystalline quasi-cleavage facets and dim-
ple relief.
(c) Prolonged holding at bainitizing temperature
leads to a decrease in hardness and to increase in im-
pact toughness as compared with the moment of trans-
formation completion. Prolonged holding at 300◦
Cre-
sults in the continuation of bainite transformation
with correspondent enrichment of remaining austen-
ite with carbon.
Acknowledgements
This work was financially supported by SAIA (Slovak
Academic Information Agency).
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