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Materials Science & Engineering A 734 (2018) 398–407
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Materials Science & Engineering A
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Microstructure characterization and mechanical behavior analysis in a high
strength steel with different proportions of constituent phases
Fei Penga, Yunbo Xua, , Xingli Gua, Yuan Wanga, Jianping Lia, Hua Zhanb
State Key Laboratory of Rolling and Automation, Northeastern University, Shenyang 110819, China
Maanshan Iron & Steel Company Limited, Maanshan 243003, China
Isothermal bainite holding
Quenching and partitioning
Austenite retention
Constituent phases
Mechanical behavior
A high strength steel with Nb-V-Ti addition was heat treated by a series of isothermal bainite holding (IBH)
processes and quenching and partitioning (Q&P) processes with different intercritical annealing temperatures.
The microstructure observation showed that some martensite was formed during quenching in IBH condition
with annealing temperature above 850 °C, mainly ascribed to heterogeneous carbon content distribution in
parent austenite with relatively large size. Meanwhile, film-like RA in martensite, relatively coarser lath-like RA
in bainite and blocky RA located in grain boundaries or phase boundaries were observed in both IBH and Q&P
treatments. Moreover, Q&P process was proved to be more beneficial to retain austenite than IBH process, while
the ability of austenite retention was similar in IBH conditions. An excellent combination of strength and
ductility was obtained in Q&P process annealed at 880 °C with tensile strength of 1126 MPa and total elongation
of ~ 18%, attributing to TRIP effect mainly occurred in the latter part of strain and fine microstructure with
homogeneous distribution.
1. Introduction
In order to meet the requirements of improving fuel efficiency and
occupant security in automotive industry, variable microstructure
strategies are applied to obtain good combination of strength and
ductility [1]. As one of the most promising approach, the addition of
retained austenite can effectively improve the strength/ductility combination due to the transformation induced plasticity (TRIP) effect
[2,3]. For the sake of obtaining abundant volume fraction of retained
austenite (RA), isothermal bainite holding (IBH) treatment is applied as
a traditional method in TRIP steel and the final microstructure comprises an aggregate of ferrite, carbide-free bainite including carbonenriched austenite and a small fraction of martensite [4,5]. However,
due to the limited strength grade of IBH treatment, a novel heat
treatment called quenching and partitioning (Q&P) is proposed and
shows attractive mechanical properties at high strength levels [6,7]. In
Q&P process, austenite after annealing is partially transformed into
initial martensite after quenching to a controlled temperature (referred
to as QT) between the martensite-start (Ms) and martensite-finish (Mf)
temperatures. Subsequently, partitioning process is applied at QT (onestep Q&P) or a higher temperature (two-step Q&P). Simultaneously, the
initial martensite tempers during partitioning process and the carbon
atoms in supersaturated martensite transfers into neighbored austenite
to achieve austenite retention [8]. Instead of bainite ferrite, tempered
martensite is introduced into Q&P steel to act as main hard phase and Q
&P steel generally shows a more excellent combination of high strength
and good elongation as compared to that of TRIP steel with similar
chemical compositions [9].
Moreover, as a critical constituent phase, intercritical ferrite mainly
acts as a soft phase to change strength level. It is well known that the
volume fraction of ferrite could be precisely regulated by intercritical
annealing temperature to change the chemical composition and phase
transformation behavior of austenite during subsequent process [10].
As a consequence,different phase constituents are obtained and hence
the final mechanical properties are controlled by annealing temperature. Actually, the martensite and bainite transformation kinetics were
also affected by prior ferrite formation [11]. Furthermore, as reported
in some articles, heterogeneous carbon distribution existed in austenite
after intercritical annealing, which showed a relatively higher carbon
content in the exterior and carbon-depleted in the interior of austenite
[12,13]. Hence, different phase transformations would happen depending on the corresponding position in parent austenite [12]. However, as for high strength steels, few articles were focused on the effect
of intercritical annealing temperature on phase transformation, especially phase transformation types and final phase constituents, which
influence the final mechanical behaviors.
Corresponding author.
E-mail address: (Y. Xu).
Received 24 June 2018; Received in revised form 5 August 2018; Accepted 6 August 2018
Available online 08 August 2018
0921-5093/ © 2018 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 734 (2018) 398–407
F. Peng et al.
electro-polished using an 875 ml CH3CH2OH + 125 ml HClO4 solution
at 20 °C and 22 V.
Characterization of microstructure was obtained by a FEI G2 F20
transmission electron microscope (TEM) with an operating voltage of
200 kV. In addition, the precession electron diffraction (PED) technique
was performed using digital precession unit from NanoMEGAS
(Digistar/ASTAR) equipped in TEM to obtain the orientation and phase
distribution maps with nanometer-scale spatial resolution. The precession angle of 0.5° was applied and camera length was selected to be
71 mm [14]. The scanning step size was 10 nm. A ϕ3 thin foil with
40 µm thickness was twin-jet polished using a solution of 12.5 vol%
perchloric acid alcohol at − 25 °C and 32 V.
In order to investigate the retained austenite fraction and corresponding carbon content in austenite, a D/max 2400 X-ray diffractometer (56 kV, 182 mA) was applied with Cu Kα radiation and a 2θ
range of 40–100° (step size: 2°/min). The measured samples were
prepared the same as EBSD samples. The amount of retained austenite
was calculated based on the measured peaks of (200)α, (211)α, and
(200)γ, (220)γ, (311)γ. Austenite carbon concentration can be obtained
based on the peak of (220)γ using an empirical equation. The detailed
calculation method can be obtained in the authors' prior [15].
Table 1
Chemical composition of tested steel (wt%).
In this article, IBH process with different intercritical annealing
temperatures and two typical Q&P treatments were applied in a high
strength steel with micro-alloyed addition. This paper aims to provide
an insight into the influence of phase constituents under different heat
treatment conditions on final mechanical behavior. Meanwhile, the
characteristics of retained austenite were also observed and compared
in detail.
2. Experimental procedures
The chemical composition of tested steel is listed in Table 1, which
is a low-carbon TRIP composition with Nb, V and Ti additions. The steel
was melt in a vacuum induction furnace and then forged into a billet
with transverse dimension of 60 mm × 60 mm. After homogenization
at 1200 °C for 5 h, the billet was hot rolled to 4 mm with the finish
rolling temperature of 830 °C. Then, the sheet was ultra-fast cooling to
650 °C and covered by asbestos to simulate coiling procedure. The final
microstructure was composed of ferrite and pearlite. Subsequently, the
sheet was cold rolled to 1 mm after pickling in 20 vol% hydrochloric
acid. The critical temperatures of Ac1 and Ac3 were measured by dilatometry as 676 °C and 938 °C, respectively. Two kinds of heat treatments were applied in this article: isothermal bainite holding at 350 °C
for 200 s with different intercritical annealing temperatures ranged
from 750 °C to 900 °C (hereafter referred to as IBH-750–900) and typical two-step Q&P process with annealing temperatures of 850 °C and
880 °C (hereafter referred to as Q&P-850/880). The Q&P process consists of quenching to 250 °C for 10 s after intercritical annealing, followed by partitioning at 380 °C for 100 s and finally water quenched to
ambient temperature. The schematic diagrams of IBH and Q&P heat
treatments are presented in Fig. 1. The tensile specimens were machined along the rolling direction with gauge length of 50 mm and
width of 12.5 mm.
The secondary electron (SE) images were obtained using JEOL JXA8530F electron probe micro-analyzer (EPMA) and corresponding samples were etched by 4% nital solution after mechanical polishing.
Electron backscatter diffraction (EBSD) technique (step size: 0.05 µm or
0.1 µm, tilt angle: 70°) was conducted in a Zeiss Ultra-55 field emission
scanning electron microscope (FE-SEM) at 20 kV and the result data
was post-processed by Channel 5 software. The specimens of EBSD were
3. Results and discussion
3.1. Characterization and comparison of microstructure in different
annealing treatments
The microstructures of different heat treatments were characterized
by secondary electron images, as presented in Fig. 2. Meanwhile, metallographic method was used to obtain ferrite fraction and mean diameter of secondary phase region (referred to as SPR), which indicates
the retaining austenite before IBH or partitioning process [16]. In
particular, the martensite packet with high-angle grain boundary,
which represents the actual dimension of martensite, was also regarded
as a special SPR during statistical process. The volume fraction of
martensite formed during first quenching process was calculated using
the Koistinen-Marburger equation [17] and the corresponding Ms
temperature was obtained by the following empirical equation [18]:
Ms(°C) = 539 − 423C − 30.4Mn − 7.5Si + 30Al (in wt. %)
Where the austenite composition was calculated by Thermal-calc based
on practical ferrite fraction. The volume fraction of retained austenite
was measured by XRD technique. All of the statistical results are summarized in Fig. 3.
In the cases of IBH-750 and IBH-780 (Fig. 2a), the deformed ferrite
(DF) band is observed along the rolling direction and abundant amount
Fig. 1. Schematic diagrams of different heat treatments (a) IBH condition; (b) Q&P condition.
Materials Science & Engineering A 734 (2018) 398–407
F. Peng et al.
Fig. 2. Secondary electron images of different heat treatments (a) IBH-750; (b) IBH-800; (c) IBH-850; (d) IBH-880; (e) IBH-900; (f) Q&P-880 F: ferrite; DF: deformed
ferrite; B: bainite; M: martensite; A: austenite; M1: initial martensite.
our previous article [15]. Meanwhile, the mean diameter of SPR increased gradually, i.e. a coarser matrix was obtained at a higher annealing temperature. It should be emphasized that initial martensite
(referred to as M1), i.e. martensite formed during quenching process
and tempered in the following IBH procedure, was observed when annealing temperature was 880 °C and 900 °C in IBH conditions. Both of
the inhomogeneous carbon distribution in SPR and change of phase
transformation during IBH process are taken into account to explain
this phenomenon. On one hand, it can be explained that the carbon
distribution in SPR was inhomogeneous, which means that carbon-enrichment was existed in the exterior of austenite as a result of ferrite
formation, while relatively lower carbon content was existed in the
interior of austenite [21]. Subsequently, the carbon-depleted region
transformed into martensite and the exterior of austenite retained after
final quenching. Actually, this martensite transformation was not observed when annealing temperature was below 880 °C due to the relatively small size of SPR, in which the gradient of carbon distribution
was not dominant. On the other hand, the carbon-enrichment from
ferrite to austenite reduced with increasing annealing temperature, and
a higher Ms temperature was obtained at relatively high annealing
temperature. In particular, the Ms temperature of IBH-900 was calculated to be 356.6 °C, which is higher than the isothermal temperature of
350 °C. Hence, martensite transformation would take place in theory
of cementite is distributed along the DF boundary. Actually, both of the
DF band and cementite were derived from the cold rolling microstructure, i.e. deformed ferrite and pearlite, and were recrystallized
incompletely during annealing process. Moreover, some of the cementite particles have transformed into austenite during annealing and
have a relatively high stability as a result of small size and high element
content [19]. In addition, the parent austenite after intercritical annealing was partially transformed into carbide free bainite during IBH
process. Meanwhile, austenite between bainite ferrite laths was carbonenriched to enhance corresponding thermal stabilization. It should be
noted that partial austenite was still transformed into martensite during
the final quenching due to insufficient stabilization. Hence, the retained
part of austenite after IBH process was finally retained as M/A island. It
is reliable to conclude that the austenite of M/A island adjacent to
ferrite is retained primarily attributing to carbon-enrichment from
ferrite to austenite [20].
With annealing temperature increases to 800 °C and above, the
volume fraction of ferrite decreases gradually (Fig. 3a). Actually, the
microstructure was evolved from ferrite matrix to M/A matrix with
increasing annealing temperature. Additionally, the cold-rolling microstructure was recrystallized completely and all of the cementite was
dissolved. In particular, some particle carbides were also observed in
the ferrite matrix and have determined to be microalloyed carbide in
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3.2. Retained austenite characterization and austenite stabilization analysis
In order to investigate the relationship between austenite stabilization during tensile test and corresponding location distribution,
austenite of IBH-900 before and after tensile test was characterized
using EBSD technique and the results are presented in Fig. 5. It reveals
that retained austenite is mainly located in packet boundary (PB) or
parent austenite grain boundary (PAGB) (enveloped by dotted line)
with a blocky morphology (Fig. 5a), which is consistent with the result
reported in Ref. [22]. As a comparison, the fraction of retained austenite after tensile test is substantially declined and most of the blocky
austenite in PB and PAGB is consumed (Fig. 5b). In particular, some
film-like austenite between martensite laths can also be observed in
specimen after tensile test. Based on misorientation maps calculated by
EBSD results, the stress status among martensite could be revealed and
it shows a negligible change after tensile test, which implies a slight
deformation of martensite during tensile test (Fig. 5c and d).
Considering the detect limit of EBSD technique, a more accurate
microstructure characterization of IBH-900 was performed by TEM and
the results are shown in Fig. 6. It is clearly that some film-like retained
austenite with size of 40–60 nm is located between the martensite laths
(Fig. 6a and b), which means that a significant carbon partitioning
behavior takes place in IBH-900. In addition, some relatively coarser
lath-like RA with more than 200 nm thickness is also observed in bainite (Fig. 6c). Meanwhile, the part of RA neighbored to grain boundary
shows a blocky morphology, which corresponds to the blocky RA observed by EBSD (Fig. 5a). As for the IBH-900 specimen after tensile test,
the film-like RA in martensite is still retained (Fig. 6d), which means a
negative TRIP effect due to excessive stabilization of RA. As a comparison, the coarser lath-like or blocky RA in bainite was not observed
after tensile test. Moreover, some blocky regions of twin martensite are
also found after tensile test (Fig. 6e). The twin martensite was derived
from unstable austenite that transformed during tensile test or austenite
insufficient carbon-enriched and transformed into twin martensite in
the final quenching procedure. Some cementite among the martensite
matrix are also observed (Fig. 6f), which can be ascribed to the temper
behavior of martensite during isothermal holding at 350 °C [23].
As a comparison, a similar observation was obtained for IBH-850, as
presented in Fig. 7. The relatively coarser lath-like RA and blocky RA in
bainite are also observed. Meanwhile, the blocky RA of bainite is
mainly located in the grain boundary of SPR (Fig. 7a and b). In particular, film-like RA in martensite was also observed in IBH-850 due to
the carbon gradient as mentioned above. Identically, twin martensite
and film-like RA in martensite are also found in specimen after tensile
test, without the existence of RA in bainite (Fig. 7e). Hence, it is reliable
to conclude that RA in bainite was mainly transformed into twin martensite during tensile test, which means a positive TRIP effect. Meanwhile, the film-like RA in martensite was still retained after tensile test,
due to the stress shield of martensite lath and relatively higher thermal
stabilization compared with blocky RA and lath-like RA in bainite
The typical morphology and distribution of retained austenite in Q&
P-880 was characterized by precession electron diffraction (PED)
technique and traditional TEM observation, respectively. As a novel
technique to obtain the information of phase and orientation, the PED
technique can overcome the detect limit of EBSD and clearly reveal the
distribution and crystallographic orientation of film-like RA between
martensite laths with nanometer scale (Fig. 8a~c) and those in bainite
with sub-micrometer scale (Fig. 8d~f). It is apparent that the film-like
RA in martensite shows a thickness of less than 200 nm, which is relatively larger than the scale of filmy RA in IBH-900. Moreover, the
existence of epitaxial ferrite is also found (Fig. 8c), which shows a small
angle grain boundary with neighbored intercritical ferrite [26]. The
morphology and distribution of RA in bainite under Q&P condition
reveal a similar characteristic compared with those under IBH conditions. Hence, the RA in bainite under Q&P condition would reveal a
Fig. 3. The measured or calculated volume fraction of different phases (a)
volume fraction of ferrite (metallographic method); (b) volume fraction of initial martensite (calculated); (c) volume fraction of RA (measured by XRD); (d)
mean diameter of secondary phase region(metallographic method). F: ferrite;
M: martensite; RA: retained austenite; SPR: secondary phase region.
during the quenching process in IBH-900, i.e. the IBH-900 is actually a
one-step Q&P process, not a traditional IBH process. In order to describe
conveniently, this process is still indicated as IBH-900. As a consequence, the M1 formation in IBH-880/900 could act as a carbon
source in the following IBH process to achieve carbon partitioning from
M1 to adjacent austenite. However, it should be mentioned that the
temperature discrepancy over the volume of sample, which is
about ± 5 °C measured by thermocouple, influenced the phase transformation behavior and phase constituents, as well as corresponding Ms
temperature of retained austenite. As for region with lower Ms temperature, the bainite transformation accelerated due to larger driving
force at 350 °C. Hence, more bainite was obtained and larger volume
fraction of austenite was retained simultaneously.
The microstructure of Q&P condition consists of ferrite, M1, M/A
island and slight volume fraction of bainite (Fig. 2f). Compared with the
microstructure of equivalent annealing temperature in IBH treatment,
the microstructure of Q&P treatment shows a relatively small size
(Fig. 3d), which is to say, Q&P treatment can refine the microstructure
significantly. Actually, the refinement of microstructure in IBH conditions was mainly ascribed to bainite formation during IBH process.
However, as for Q&P condition, the quenching temperature was selected to give rise to severe martensite transformation in whole region
of SPR, which can effectively refine the microstructure by various high
angle grain boundary, like grain boundary and packet boundary. Subsequently, some bainite was also formed to refine microstructure further. Therefore, Q&P treatment can obtain a relatively finer microstructure compared with IBH treatment. Consistently, as revealed by
the inverse pole figure (IPF) of EBSD results (Fig. 4), the size of martensite block, which shows identical orientation in a specific block region, is relatively finer in Q&P-880 as compared with that of IBH-880.
Materials Science & Engineering A 734 (2018) 398–407
F. Peng et al.
Fig. 4. IPF obtained by EBSD results (step size: 0.1 µm) (a) IBH-880; (b) Q&P-880.
Fig. 5. EBSD results of IBH-900 (step size: 0.05 µm) Before tensile test: (a) combined map of band contrast (BC) and IPF; (c) Misorientation map; After tensile test: (b)
combined map of BC and IPF; (d) Misorientation map.
reality, the precipitation of ε carbide inside initial martensite will reduce the carbon content that can partition from M1 to austenite and
hence reduce the final amount of retained austenite.
The volume fractions of retained austenite (VRA) and corresponding
austenite carbon contents (X CRA) of different heat treatments were obtained by XRD and the results are presented in Fig. 10. With increasing
annealing temperature in IBH conditions, the volume fraction of austenite increases from ~ 6% at 750 °C to a saturated value of 9.6% at
850 °C, after which a distinct reduction at 880 °C and 900 °C is observed. Furthermore, most of the austenite has consumed after tensile
test to introduce a positive TRIP effect except for that of IBH-900, of
which the austenite fraction is just reduced slightly after tensile test.
This discrepant result is supposed to be related to the limited ductility
in IBH-900 and will describe in detail in Section 3.3. As a comparison,
the volume fraction of Q&P treatment shows an identical increase trend
with annealing temperature and both of the evaluated Q&P treatments
show a positive TRIP effect. Fig. 10b shows the corresponding austenite
carbon content before tensile test. It is clearly that the value of X CRA
increases slightly from 1.03 wt% to 1.28 wt% in IBH conditions, whilst
X CRA keeps a constant value of ~ 1.2 wt% in the range of 780–880 °C.
similar transformation behavior during tensile test, as compared with
those of IBH treatments.
In addition, some particular characteristics of microstructure were
found by traditional TEM observation, as presented in Fig. 9. Some
blocky RA is located in triple grain boundary (Fig. 10a and b), which is
ascribed to the accelerated effect of grain boundary on elements diffusion [27], especially carbon element in this article. Fig. 9c presents
film-like RA in M1 with thickness ranged from 80 to 170 nm, which is
identical to the results obtained by PED technique (Fig. 8b). Meanwhile,
some RA particles with size of more than 500 nm are also found in the Q
&P-880 specimen, which is derived from parent austenite that retained
after ferrite transformation [28]. In particular, this kind of particle RA
has been confirmed to transform into martensite in the latter part of
tensile process [29], i.e. is beneficial to delaying necking process and
enhance the whole ductility. Furthermore, some lath-like carbide is
observed in M1 (Fig. 9f), which has been confirmed as ε carbide that
derived from auto-tempering process of M1 [30]. Compared with IBH900, the initial martensite in Q&P-880 formed during quenching to
250 °C possessed a wider temperature range to achieve auto-temper and
hence Q&P process was more beneficial to precipitate ε carbide. In
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F. Peng et al.
Fig. 6. TEM micrographs of specimen (IBH-900) before tensile test (a) bright field image of film-like RA and corresponding (b) dark field image of film-like RA, (c) RA
in grain boundary; after tensile test (d) film-like RA, (e) twin martensite (TM), and (f) cementite. Note that the scales of images are not the same.
Meanwhile, the variation of X CRA in Q&P conditions shows a similar
increasing trend at an equivalent carbon level compared to that of IBH
With increasing annealing temperature, the extent of austenitization
increased and hence the volume fraction of bainite showed an increasing trend in IBH conditions. Consequently, the increasing trend of
VRA with annealing temperature is mainly attributed to bainite formation, which brings in obvious carbon enrichment in austenite to enhance its thermal stability. Simultaneously, as mentioned above, some
M1 was formed in IBH conditions with relatively high annealing temperatures (850 °C and above), which is mainly ascribed to inhomogeneous distribution of carbon in relatively large austenite region.
Actually, carbon partitioning behavior the same as Q&P treatment takes
place during subsequent IBH process and is conducive to austenite retention as well. However, some cementite precipitated inside tempered
martensite during IBH process in IBH-900 (Fig. 6f). Hence, it is reasonable to conclude that the distinct reduction in RA fraction is ascribed
to the formation of cementite, while the austenite carbon content is
insensitive to constituent phases, irrespective of annealing temperature.
As for Q&P conditions, carbon partitioning from M1 to austenite was
dominant in austenite retention as abundant M1 of 50–60 vol% formed
in Q&P-850/880. Moreover, as austenite fraction before partitioning
process, i.e. the volume fraction of secondary phase region (VSPR), was
just ~ 10 vol% in Q&P-850 (Fig. 11), which is much less than that in
IBH-850, hence a relatively less VRA was obtained in Q&P-850. In particular, although ε carbide was precipitated inside M1 in Q&P-880
(Fig. 9f), it does not reduce corresponding VRA significantly and an
equivalent level of VRA was obtained as 9.4 vol% with corresponding
austenite carbon content of 1.2 wt%, as compared to that of IBH-880.
Furthermore, the ratio of volume fraction of RA and corresponding
volume fraction of SPR was calculated in this article to represent the
ability of different phase proportions in austenite retention. As presented in Fig. 11, the volume fraction of SPR increases significantly
with increasing annealing temperature, while the corresponding ratio
of VRA/VSPR keep a constant value of ~ 0.18 in the temperature range of
750–850 °C followed by a slight decline to ~ 0.07 at 880 °C and 900 °C.
It can be concluded that phase proportions of constituent phases in IBH
condition do not influence the ability of austenite retention, while the
formation of cementite in IBH-880/900 slightly deteriorates the corresponding austenite retention ability. Moreover, the values of VRA/
VSPR in Q&P conditions are much larger than those of IBH specimens,
i.e. 0.55 in Q&P-850 and 0.77 in Q&P-880, which means that the introduction of abundant fraction of M1 in Q&P treatment is much more
beneficial to austenite retention than IBH treatment, which is consistent
with the reported result of Speer [31].
3.3. Mechanical properties and work hardening behavior
The engineering stress-strain curves and corresponding mechanical
properties of different heat treatments are shown in Fig. 12. It is obvious that all of the curves present continuous yielding except for that
of IBH-750, which has a yielding plateau due to the existence of
abundant undissolved cementite (Fig. 2a). With increasing annealing
temperature in IBH conditions, both of yielding strength and ultimate
tensile strength show an ascending trend due to increasing fractions of
hard phases (Fig. 12b and c), i.e. bainite and martensite. Meanwhile,
the variation of total elongation presents a reverse trend. Actually, as
the ferrite fraction decreased to a critical extent, here is 25.2 vol% in
IBH-880, the matrix phase was varied from ferrite to martensite/bainite
and the stress could not transfer among ferrite during tensile test.
Hence, local stress concentration will take place significantly at the
interface of ferrite and martensite/bainite, where microvoids formed
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Fig. 7. TEM micrographs of specimen (IBH-850) before tensile test (a) bright field image of blocky RA in bainite (B) and (b) corresponding dark field image, (c) SAED
pattern of RA in (a); after tensile test (d) twin martensite (TW), (e) film-like RA in martensite and (f) corresponding SAED pattern of RA in (e).
Fig. 8. Phase and orientation distribution of Q&P-880 by PED (step size: 10 nm) (a) (d) Phase reliability; (b) (e) BCC (red)+FCC (green); (c) (f) IPF + high angle
grain boundary (10°) M: martensite; B: bainite; F; ferrite; EF: epitaxial ferrite. (For interpretation of the references to color in this figure legend, the reader is referred
to the web version of this article.)
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Fig. 9. TEM micrographs of Q&P-880 (a) bright field image of blocky RA in triple grain boundary and (b) corresponding dark field image, (c) film-like RA in M1; (d)
bright field image of particle RA and (e) corresponding dark field image; (f) ε carbide in martensite.
and coalesced to result in final fracture [32]. Meanwhile, the deformation of hard phase was not significant (Fig. 5d) and therefore, the
total elongation of IBH-880/900 decreased distinctly compared to that
of samples annealed at relatively lower temperature (Fig. 12d). As for Q
&P treatments, Q&P-850 shows a similar combination of strength and
ductility as that of IBH-850. However, the mechanical properties of Q&
P-880 show a remarkable PSE value of ~ 20000 MPa%, while YS and
UTS show a negligible discrepancy compared with those of IBH-880.
The work hardening behavior of different heat treatments in IBH
conditions and corresponding comparison of IBH-850/880 and Q&P850/880 are shown in Fig. 13, respectively. As shown in Fig. 13a, it is
clearly that the whole work hardening exponent n decreases with increasing annealing temperature. Meanwhile, obvious TRIP effect is
significantly observed at relatively low annealing temperature, i.e. IBH750/780. In reality, a large amount of carbon-enriched martensite,
which is mainly twin martensite (Fig. 6e and Fig. 7d), was formed in
IBH specimens due to relatively weak ability of austenite retention
(Fig. 11) and hence enhances the value of n. However, most of the IBH
heat treatments show analogous work hardening ability, therefore, the
discrepancy of different treatments is ascribed to other factors. Actually, the relatively higher work hardening ability of IBH-750 and IBH780 is attributed to undissolved cementite (Fig. 2a) and deformed
Fig. 10. XRD results of different heat treatments (a) volume fraction of RA; (b) mean carbon concentration in RA.
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like RA in Q&P conditions was tempered at a relatively higher temperature, resulting in a weaker stress shield effect on the film-like RA,
which also possessed a larger size than that in IBH conditions due to
more adequate carbon partitioning behavior from M1 to austenite
(Fig. 9c). Hence, the lath-like RA in Q&P specimens could give rise to
positive TRIP effect and further avoid stress concentration [24,33].
Furthermore, the relatively finer microstructure of Q&P specimens was
also beneficial to enhance the corresponding work hardening ability.
Especially, Q&P-880 showed an excellent combination of mechanical
properties with tensile strength of 1126 MPa and total elongation of
~ 18%, which was attributed to TRIP effect mainly occurred in the
latter part of strain and fine microstructure with homogeneous distribution.
4. Conclusions
In this article, a high strength steel with lean microalloyed addition
was heat treated by isothermal bainite holding (IBH) and quenching
and partitioning (Q&P) processes with a series of intercritical annealing
temperature. All of the conclusions are shown as follows:
Fig. 11. Volume fraction of secondary phase region (SPR) and corresponding
ratio of VRA/VSPR.
(1) Some tempered martensite was observed in IBH conditions with
annealing temperature above 850 °C, mainly ascribed to heterogeneous carbon content distribution in parent austenite with relatively large size. Meanwhile, Q&P treatment could obtain a relatively finer microstructure, which was divided by high angle
boundaries deriving from severe martensite transformation.
(2) Three kinds of retained austenite were mainly observed in both IBH
and Q&P treatments: film-like RA between martensite laths, relatively coarser lath-like RA in bainite and blocky RA located in high
angle grain boundaries, like parent austenite grain boundary,
martensite packet boundary and bainite/ferrite interface.
Furthermore, some film-like RA in martensite was still retained
after tensile test, while blocky and lath-like RA were mostly transformed into twin martensite.
(3) Both of IBH and Q&P processes showed an increasing trend of RA
fraction (5–10 vol%) with decline in ferrite fraction, while a significant decrease was observed as a result of cementite formation in
IBH condition annealed at 880 °C and 900 °C. Moreover, Q&P process was more beneficial to retain austenite than IBH process, while
the ability of austenite retention was nearly similar in IBH
ferrite, which was recrystallized incompletely and easily formed dislocation tangle to enhance work hardening ability. Moreover, as
abundant volume fraction of ferrite was formed in IBH-750/780, a large
amount of blocky RA with relatively low stability was retained [20] and
hence gave rise to a significantly obvious TRIP effect. Furthermore, due
to the initial martensite formed and subsequently tempered in IBH-880/
900 (Fig. 2d and e), a relatively softer matrix was obtained, which
explains the weaker work hardening ability at higher annealing temperature. It should also be emphasized that the abnormal work hardening behavior of IBH-900 is mainly ascribed to severely stress concentration at the interface of ferrite and bainite/martensite, as
mentioned above.
At the onset of strain in Fig. 13b (part Ⅰ), the IBH specimens show
analogous or higher n values than those of Q&P samples annealed at
equivalent temperatures, while the n values of Q&P treatments increase
rapidly with true strain and presents relatively higher values in the
latter part of strain (part Ⅱ). It is clearly that the critical true strain of
these two parts increased from 0.0182 at 850 °C to 0.0633 at 880 °C.
Compared with IBH treatments, the corresponding M1 matrix of film-
Fig. 12. Mechanical properties of uniaxial tensile test (a) Engineering stress-strain curves; (b)yielding strength (YS); (c) ultimate tensile strength (UTS); (d) total
elongation (TEL); (e) product of UTS and TEL (PSE).
Materials Science & Engineering A 734 (2018) 398–407
F. Peng et al.
Fig. 13. Work hardening behavior of different heat treatments (a) different annealing temperatures in IBH conditions; (b) comparison of IBH and Q&P processes.
conditions, irrespective of different phase proportions.
(4) An excellent combination of strength and ductility was obtained in
Q&P-880 with tensile strength of 1126 MPa and a PSE value of
~ 20000 MPa%, as a result of TRIP effect mainly occurred in the
latter part of strain and fine microstructure with homogeneous
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This work was financially supported by the National Key R&D
Program of China (2017YFB0304105)
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