Effect of reheating rate on microstructure and
properties of high-strength-toughness steel
Z. J. Xie1†, Y. P. Fang1†, Y. Cui2, X. M. Wang1, C. J. Shang∗ 1 and R. D. K. Misra3
Effects of reheating rate and holding time during tempering on microstructure and properties of a
960-MPa-grade low-alloy steel were investigated. Different reheating rates from 1 to 300 K s−1 and
holding time from 5 to 1200 s on tempering at 823 K were carried out. The combination of high
reheating rate and short holding time is obviously beneficial forgetting more uniformly distributed
and refined carbides, which are believed to contribute to the high strength and high toughness.
Yield strength of 960 MPa with elongation of ∼16% and impact toughness of 76 J (half size
specimen) at 233 K was obtained by induction reheating tempering at 823 K for 5 s at a rate of
373 K s−1.
Keywords: High-strength low-alloy steel, Nano-scale carbides, Reheating rate, Induction tempering
Introduction
High-strength low-alloy (HSLA) steels are widely used in
energy field, such as off-shore drilling platforms, pipelines
and other fossil fuel industries, because of their outstanding mechanical properties, good weldability and low cost.
To meet the requirements, especially for the severe service
environment, a superior combination of high-strength
and low-temperature toughness steel is the primary aim
in engineering.1 In consideration of the weldability and
toughness, the carbon content of HSLA steels is usually
controlled at a lower level. As a result, adjusting the
microstructure by heat treatment becomes a main route
for obtaining desired strength and toughness combination
in a given steel. One of the most effective heat treatments
is quenching and tempering. During quenching and tempering, matrix recovery and carbide precipitation usually
occur. However, the extent of the recovery and the size
and morphology of carbides, especially cementite carbides, depend significantly on the appropriate selection
of heat treatment parameters, such as reheating rate, holding time and tempering temperature, which are important
to obtain desired properties.2
Quenching process inclusive of on-line quenching
directly after finishing rolling and off-line quenching
after reheating the hot-rolled steel plates to re-austenitisation temperature is the common practical way to
obtain high strength in HSLA steel.3 Toughness could
be improved by the subsequent tempering with the cost
of strength, because of the reciprocal relationship
between high-strength and low-temperature toughness.
†
Contributed equally to this work.
School of Materials Science and Engineering, University of Science and
Technology Beijing, Beijing, China
2
Shougang Research Institute of Technology, Beijing, China
3
Department of Metallurgical and Materials Engineering, University of
Texas at El Paso, El Paso, TX 79968, USA
1
∗
Corresponding author, email cjshang@ustb.edu.cn
© 2016 Institute of Materials, Minerals and Mining
Received 1 October 2015; accepted 30 December 2015
DOI 10.1080/02670836.2015.1138044
It is widely reported that the coarse carbide precipitates
and brittle particles result in cleavage fracture and are
harmful for low-temperature toughness.4–6 In the case
of HSLA steels, the conventional tempering (CT) process
with slow reheating rate and long holding time would
create coarse carbides and particles inhomogeneous distribution, which do harm to toughness.7 However, with
the advantage of rapid reheating rate and high efficiency,
induction reheating tempering (IT) is favourable to
obtain fine microstructure and dispersed precipitates.8
Thus, selection of appropriate tempering parameters is
the key factor in controlling carbide precipitation and
in obtaining good combination of high-strength and
low-temperature toughness.9,10
The aim of the present study is to elucidate the relationship between microstructure and mechanical properties of
960-MPa-grade HSLA steel by altering tempering parameters, with particular attention to reheating rate and
holding time on carbide size and distribution, and their
relationship to mechanical properties.
Experimental methods
The chemical composition of experimental steel in weight
percent (wt-%) was Fe–0.12C–0.22Si–1.44Mn, <1.0(Cr
+ Mo + Nb + Ti + V+B). The experimental steel plates
were melted in a vacuum induction furnace, cast and
were hot rolled to a strip of 8-mm thickness. The experimental steel plates were austenitised at 1173 K for 30
min and water quenched to room temperature at a cooling rate of ∼40 K s−1.Two aspects of experimental test
were carried out to investigate the influence of tempering
parameter on the mechanical properties and carbide
evaluation.
One part of the quenched specimens were machined to
cylindrical specimens that were 10 mm in length and 4
mm in diameter with the axis parallel to the rolling direction, and heat treated as illustrated in Fig. 1 using a DIL
805a-type dilatometer for tempering simulation. The
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Xie et al. Reheating rate and high-strength-toughness steel
Results and discussion
Influence of tempering parameters on the
Vickers hardness
1 Schematic illustration of the quenching and tempering
cycle treatment
cylindrical specimens were tempered at 823 K for periods
of 5, 40, 150 and 1200 s at reheating rates of 1, 20, 100 and
300 K s−1, respectively. After tempering, the specimens
were cooled at ∼40 K s−1 to room temperature. The simulated specimens were prepared for the Vickers hardness
and microstructure evaluation. The average Vickers hardness was obtained from 12 test points, with a load of 9.8 N
using HVS-1000ZDT. The microstructure of specimens
were characterised via combination of a scanning electron
microscope (ZEISS ULTRA-55 field SEM) and a transmission electron microscope (TEM; JEM-2100F) at 200
kV. Transmission electron microscope foils were prepared
by mechanical polishing and twin-jet polishing in a 5%
perchloric acid solution using a twin-jet polisher.
The second part of the quenched samples were induction reheated at a heating rate of ∼100 K s−1 to 823 K
using an induction heating furnace, followed by air cooling to room temperature. These samples were designated
as Q-IT. For comparing with the CT process, the
quenched specimens were reheated at a reheating rate of
∼1 K s−1 to 823 K and held for 20 min by conventional
furnace, followed by air cooling to room temperature.
These samples were designated as Q-CT (Fig. 2). The
tempered Q-IT and Q-CT samples were machined to the
tensile and the Charpy impact tests according to the
ASTM standard. Tensile tests were carried out at a strain
rate of 2.5 × 10–3 s−1 at room temperature to determine
the yield strength, tensile strength and elongation. Impact
toughness was determined using the Charpy V-notch
specimen of dimensions 5 mm × 10 mm × 55 mm prepared from the longitudinal orientation and tested at
233 K.
2 Schematic illustration of the tempering treatment of
quenched specimens with different reheating rates
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With regard to tempering, the reheating rate and holding
time are the key factors to attain good combination of
high-strength and low-temperature toughness.8 By
induction tempering, a reheating rate range of 1–300
K s−1 can be obtained and the holding time varied
from 5 to 1200 s. The correlation of tempering parameter, microstructure and hardness can be revealed
from Fig. 3. It shows the mean hardness obtained
after tempering at 823 K by different reheating rates
and holding times. The hardness of the as-quenched
specimen was 388 HV. Compared with the as-quenched
sample, tempering process led to a decrease in the hardness. For holding of specimens for 5 s, the Vickers hardness of tempered specimens increased from 330 to 347
HV with an increase in the reheating rate from 1 to
300 K s−1. There was no significant difference in hardness between the rapid reheating rate and the slow
reheating rate on tempering for 1200 s. In addition,
when the reheating rate was 300 K s−1, the Vickers hardness decreased from 347 to 309 HV on holding from 5 to
1200 s. It can be concluded that increasing reheating rate
and decreasing holding time significantly contributed to
higher Vickers hardness after tempering. Meanwhile,
prolonging the holding time decreased the contribution
of the rapid heating rate on the hardness.
Microstructure characterisation
Bright-field TEM micrographs showing carbides in
quenched and tempered samples with different reheating
rates are presented in Fig. 4. It can be seen that the
quenched microstructure comprised lath martensite and
lower bainite. Owing to the diffusionless, shear mechanism of martensite transformation, there was no carbide
precipitation within martensitic lath.11 However, needleshaped carbides were observed in the bainitic ferrite
matrix. The average length of the long axis of carbide
was approximately 200 nm (Fig. 4b). The previous result
confirmed the carbides as cementite precipitation on the
basis of selected area diffraction pattern.12
Figure 4c–h shows the bright-field TEM micrographs
of as-quenched samples tempered at 823 K for 5 s at a
3 Evolution of the Vickers hardness as a function of heating
rate and holding time at the 823 K tempering treatment
Xie et al. Reheating rate and high-strength-toughness steel
reheating rate of 1–300 K s−1. There are two types of carbides forming in the microstructure: one is needle-like
within the bainitic ferrite lath, and the other one is spherical along the lath boundary or prior austenite boundary.
With an increase in the reheating rate, the distribution
of carbides became homogenous and the carbides were
finer (Fig. 4c–f). With an increase in the reheating rate
from 1 to 300 K s−1, the morphology of carbides
precipitated within bainitic ferrite lath changed from needle to rod-like, and the dimension of the long axis was
decreased from 200 to 80 nm.
Figure 4g–h reveals the changes in spherical-type carbides at different reheating rates. The diameter of spherical-type carbides was ∼200 nm at a reheating rate of 1 K
s−1, while their size decreased to less than 100 nm at a
reheating rate of 100 K s−1. This clearly gives evidence
4 Bright-field TEM micrographs showing cementite carbides of as-quenched specimens and tempered with different reheating
rates a and b as-quenched, c and g reheating rate 1 K s−1, d 20 K s−1, e and h 100 K s−1 and f 300 K s−1, respectively
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Xie et al. Reheating rate and high-strength-toughness steel
to support that increasing the reheating rate significantly
contributes to refining carbides.
Influence of tempering parameter on
refinement of carbides’ size
Generally, during the precipitation process, carbides preferentially choose the prior austenite boundaries, packet
boundaries (high-angle boundaries), lath boundaries
and dislocation within laths (low-angle boundaries) to
nucleate in the microstructure.8 Considering the difference between prior austenite boundaries and martensitic
or bainitic lath boundaries, it is important to analyse
the evolution of carbides. As mentioned above, needle-
type carbides precipitated primarily within the lath and
spherical carbides precipitated along the lath boundary
or prior austenite boundaries. Figure 5 shows the effect
of tempering parameters on the size distribution of carbides during tempering. The results indicate that after
holding for 5 s(short time), the length of long and short
axes of the carbides was significantly decreased irrespective of the locations of the precipitates for reheating
rates of 1–300 K s−1 (Fig. 5a–b). By increasing the reheating rate from 1 to 300 K s−1 during tempering, the length
of the long axis of carbides precipitated at the prior austenite boundaries decreased from 300 to 100 nm, while,
for carbides those precipitated within lath, the length of
the long axis decreased from 210 to 80 nm. On the
5 The relationship between long axis, short axis and axis ratio of the carbides with different tempering parameters at 823 K with
different reheating rates and holding times a and b 1 to 300 K s−1 for 5 s; c and d 1 K s−1 for 5 to 1200 s; e and f 100 K s−1 for 5
to 1200 s at prior austenite boundaries a, c and e and within the lath b, d and f
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Xie et al. Reheating rate and high-strength-toughness steel
holding times, it can be found that the holding time
makes a significant effect on the size of carbides coarsening at a rapid reheating rate.
During tempering process, rapid reheating rate retards
the recovery of dislocations.8,13 On the other hand, the
time–temperature–precipitation diagram for high reheating rate is expected to be different from the normal reheating rate.14–17 The rapid reheating rate reduces the
incubation time of carbides, resulting in acceleration of
precipitation kinetics of carbides. From this aspect, the
rapid nucleation rate of carbides and short holding time
are basic parameters to obtain fine carbides. Figure 5
further confirmed that tempering parameters such as
reheating rate and holding time are important to regulate
the size and distribution of secondary precipitation. When
the tempering temperature was 823 K, increasing the
reheating rate to 100 K s−1 and decreasing the holding
time is an ideal approach to obtain nano-scale carbides
in the experimental steel.
Relationship between the size of carbides and
the Vickers hardness
6 Relationship between mean carbides’ size and hardness
in 1 and 100 K s−1 with different holding times a at prior
austenite boundaries and b within the lath
other hand, the morphology of carbides changed from
needle-like to short rod-like with a decrease in the long/
short axis and axis ratio.
To analyse the effect of holding time on the size and distribution of carbides, Fig. 5c–f summarises carbides’ size
distribution obtained for specimens reheated at 1 and
100 K s−1, held for 5–1200 s, respectively. When the
reheating rate was 1 K s−1, prolonging the holding time
led to coarsening of carbides, consequently, the length
of long and short axes was increased dramatically (Fig.
5c and d ). The ratio of long/short axis of carbides was
decreased from ∼5.0 to 4.0, implying that the morphology
of carbides was changed to spherical type. As the reheating rate increased to 100 K s−1, with an increase in the
holding time, the length of long and short axes behaved
in the same manner as that at a low reheating rate (Fig.
5e and f ). However, comparing the size distribution of
carbides at a constant reheating rate with different
Table 1
Figure 6 shows the relationship between the mean size of
carbides and the Vickers hardness. It was clearly noted
that tempering parameters make a big difference in the
Vickers hardness. Irrespective of precipitation at
prior austenite boundaries (Fig. 6a) or within the lath
(Fig. 6b), the general behaviour of plot is similar. The hardness decreased with the increasing holding time from 5 to
1200 s. In other words, the coarsening of carbide is the
main reason for the lower hardness. While combining
rapid reheating rate with short holding time, fine precipitates were obtained, the hardness was higher. This is
because the fine and dispersed precipitates contribute
more to strengthening the matrix.18,19 The lower hardness
can be explained from two aspects: the first one is the coarsening of carbides caused by the slow reheating rate and
long holding time; the other one is the dramatical decrease
in the dislocation density due to the recovery process.
Mechanical properties
Table 1 shows mechanical properties of experimental steel
quenched and tempered by two different tempering procedures, namely conventional furnace reheating and
induction reheating. After IT and CT at 823 K as
described in the experimental procedure, the yield
strength of both samples was greater than 960 MPa,
and the tensile strength was greater than 1000 MPa.
The Charpy impact energy of the conventional tempered
specimen (half thickness) was 35 J at 233 K, and that of
the induction tempered sample is 76 J at 233 K. The latter
is twice as much as the former. This suggests that rapid
reheating rate significantly contributes to improve the
low-temperature impact toughness. As mentioned
above, Fig. 4e and h shows short rod-like carbides
Mechanical properties of the studied steel after tempering at 823 K with different reheating rates
Processing
Q-CT550 (1 K s−1)
Q-IT550 (100 K s−1)
Yield strength/MPa
Tensile strength/MPa
Total elongation/%
Charpy impact energy at 233 K
(half thickness specimen)/J
988
997
1000
1026
15.5
16.7
35
76
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Xie et al. Reheating rate and high-strength-toughness steel
precipitated within the lath and finer spherical carbides
located at the prior austenite boundaries or lath boundaries during the induction rapid reheating tempering, consistent with the underlying reason for improvement in
toughness reported in previous literatures.20,21
Comparing the mechanical properties and microstructure of the experiment steel tempered by different reheating methods, it can be concluded that IT is an excellent
approach to achieve good combination of high strength
and excellent toughness. Rapid reheating tempering is
an effective heat treatment to obtain high strength and
toughness in HSLA steel.
Conclusions
This study focused on the heat treatment of a low-carbon
HSLA steel with particular emphasis on the effect of
reheating rate and holding time of tempering process on
the size distribution of carbides. It was observed that
increasing reheating rate and decreasing holding time
are conducive to reduce the size of carbides and alter
their morphology to short rod-like or spherical type.
Prolonging the holding time led to coarsening of carbides
whether the reheating rate is rapid or slow. Dispersed fine
carbides and high density of dislocations were the reasons
for yield strength greater than 960 MPa and low-temperature toughness (half thickness) at 233 K of ∼76 J, which
was two times as high as that by the CT process.
Acknowledgements
This work was financially supported by the National
Basic Research Program of China (973 Program) through
the Contract No. 2010CB630801 and the Natural Science
Foundation of China through the Contract No.
51371001. Z. J.X.thanks the China Scholarship Council
for the award of a scholarship for studying at McMaster
University, Canada. R.D.K.M gratefully acknowledges
the support from the University of Texas at El Paso, USA.
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