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Cold rolling and intercritical annealing of C-Mn steel sheets with different initial microstructures
Author’s Accepted Manuscript Cold rolling and intercritical annealing of C-Mn steel sheets with different initial microstructures Y. Karimi, S. Hossein Nedjad, H. Shirazi, M. Nili Ahmadabadi, H. Hamed Zargari, K. Ito www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(18)31196-1 https://doi.org/10.1016/j.msea.2018.09.008 MSA36889
To appear in: Materials Science & Engineering A Received date: 22 May 2018 Revised date: 3 September 2018 Accepted date: 4 September 2018 Cite this article as: Y. Karimi, S. Hossein Nedjad, H. Shirazi, M. Nili Ahmadabadi, H. Hamed Zargari and K. Ito, Cold rolling and intercritical annealing of C-Mn steel sheets with different initial microstructures, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2018.09.008 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Cold rolling and intercritical annealing of C-Mn steel sheets with different initial microstructures
Y. Karimia, S. Hossein Nedjada*, H. Shirazib, M. Nili Ahmadabadib, H. Hamed Zargaric, K. Itoc
a
Faculty of Materials Engineering, Sahand University of Technology, P.O. Box: 51335-1996, Tabriz, Iran.
b
School of Metallurgy and Materials Engineering, University of Tehran, P.O. Box: 14395-731, Tehran, Iran.
c
Joining and Welding Research Institute, Osaka University, 11-1 Mihogaoka, Ibaraki, Osaka 567-0047, Japan.
*
Corresponding author. Tel.: +98-41- 3345 9449; fax: +98-41-3344 4333. [email protected] (Syamak Hossein Nedjad).
Abstract The microstructure and mechanical properties of cold-rolled dual-phase steels produced from different initial microstructures, i.e., complex ferrite-pearlite, martensite, and bainite, were studied for a C-Mn steel contained 0.15C, 2.50Mn, and 0.30Si (wt. %). Sheets of 4 mm thickness, and with the different initial microstructures were cold rolled for 50% thickness 1
reduction, intercritically annealed at 710 and 740 ˚C for various times (10 – 1800 s) in a salt bath, and then water cooled. Different amounts and distributions of martensite were produced from the different initial microstructures by the same annealing treatment. Bainite was overwhelming which produced slightly higher amounts of martensite and the net-arrayed martensite distribution. Electron backscattering diffraction revealed recrystallization of coldrolled bainite before austenite formation at 710 ˚C; and that the ferrite bound near martensite has remarkable transformation induced effects. Tensile strength and ductility of the steels annealed for 1200 s at 710 ˚C were almost the same. The effect of the initial microstructure on the tensile properties is remarkable for steels annealed for the same time at 740 ˚C. The netarrayed martensite produced by using bainite represented higher initial strain hardening at both temperatures, which is attributed to the transformation effects in the ferrite bound around martensite.
Keywords: Dual-phase steel; Cold rolling; Intercritical annealing; Recrystallization; Austenite formation.
1. Introduction The automotive industries use up the advanced high strength steels (AHSS) because of better crashworthiness and weight reduction [1]. Dual-phase (DP) steels are the most prominent category of AHSS, which have earned widespread applications in the fabrication of chassis, pillars, shields, rocker rail and cross members; owing to good formability, high strength and affordability [2, 3]. The microstructure of DP steels mainly consists of a ferrite-martensite mixture. They are produced by annealing or thermo-mechanical processing in the austenite2
ferrite intercritical region [4]. Mechanical properties of DP steels depend mainly on the mean size of ferrite grains and the volume fraction of martensite [5-8]. Martensite morphology (shape, size, and distribution) also influences the mechanical properties, and effect of different heat treatments, producing different martensite morphologies, has been well studied [9-23]. Three kinds of heat treatments have frequently been compared [9, 10, 12, 16]: (i) intercritical annealing; (ii) intermediate quenching; and (iii) step quenching. The intermediate quenching provides martensite before the intercritical annealing treatment and gives fine and fibrous microstructures with superior strength-ductility balance. Step quenching results in coarse and blocky microstructures with poor tensile properties. The intercritical annealing yields globular microstructures with intermediate tensile properties. Indeed, it is the initial microstructure before intercritical annealing treatment, which is modified by using these kinds of heat treatments. A part of DP steels is produced by additional cold rolling of available steel products, followed by annealing in the austenite-ferrite intercritical region, and subsequent cooling to transform austenite grains into martensite [24-26]. A feature of the heat treatment of coldrolled DP steels is the recrystallization of cold-rolled microstructure and its interaction with austenite formation. Complete recrystallization of the deformed structure before austenite formation provides a random distribution of austenite while incomplete recrystallization gives rise to the formation of banded austenite grains [27, 28]. Austenite formation in the deformed structure retards the growth of subsequently recrystallized ferrite grains and results in grain refinement. Recrystallization precedes austenite formation at low intercritical annealing temperatures and low heating rates [29-31]. It has also been reported that the cooling rate
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after intercritical annealing influences carbon content of austenite and thus the martensite start temperature (Ms) in the cold-rolled DP steels. When subsequent isothermal treatment (aging) is carried out at a temperature lower than Ms, tempered martensite and ferrite appear; Otherwise bainite formation becomes possible [32]. In addition to the heating and cooling rate effects, the influence of the initial microstructure of cold-rolled DP steels has to be taken into account. Hot-rolled steel plates are often used for cold rolling, which usually consists of the ferrite-pearlite microstructure. During intercritical annealing of a cold-rolled ferrite-pearlite microstructure; recrystallization of the deformed ferrite, spheroidization of fragmented cementite and austenite formation take place concomitantly [33, 34]. Nucleation of austenite firstly occurs at ferrite-ferrite grain boundaries, and then on the cementite-ferrite interfaces. However, advantages of making the use of bainite and martensite as the initial microstructure for cold-rolled DP steels have been reported. For example, austenite formation in the deformed bainite has resulted in remarkable microstructure refinement [35]. Recrystallization of the deformed martensite during intercritical annealing has led to a refined dual-phase microstructure [36]. It has also been reported that cold-rolled bainite-ferrite-pearlite and martensite have higher rates of recrystallization and austenite formation than cold rolled ferrite-pearlite [37]. A DP steel produced by cold rolling and intercritical annealing of bainite-ferrite initial microstructure has shown net-arrayed martensite distribution which resulted in a marked strain hardening, increased tensile strength and fatigue endurance [38]. A ferrite-fibrous martensite initial microstructure has shown more refined microstructure and superior tensile properties than ferrite-pearlite after cold rolling and intercritical annealing [39].
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The efficacy of step-quenched ferrite-martensite and bainite initial microstructures for grain refinement during cold rolling and annealing treatment, along with the banding effects has been studied for a C-Mn steel contained about 0.15C and 2.50Mn (wt. %) [40, 41]. For the same steel, this paper is to compare the microstructure and mechanical properties of coldrolled DP steels produced from different initial microstructures, i.e., complex ferrite-pearlite, martensite, and bainite.
2. Experimental Procedure An ingot weighing 35 kg was prepared using induction melting under inert argon protection and subsequently refined by electroslag remelting (ESR). The ingot contained 0.15C, 2.50Mn, and 0.30Si (in wt. %). It was hot rolled to a sheet of about 4 mm thickness in several passes then air cooled to room temperature. The AC1 and AC3 temperatures of the air-cooled steel were determined about 690 ˚C and 820 ˚C, respectively, using the Adamel dilatometer. The hot-rolled sheet was used to create different initial microstructures. For this purpose, strips (4 mm × 100 mm × 150 mm) cut from the hot-rolled sheet were austenitized in a resistance furnace at 900 ˚C for 3600 s, then cooled at different conditions, as shown in Figure 1. The strips were subsequently rolled by 50% thickness reduction at room temperature. The next step involved heating of the cold-rolled microstructures in a salt bath to intercritical annealing temperatures (710 and 740 ˚C) and soaking for different times (10 – 1800 s). Finally, the samples were quenched in water to produce DP microstructures. Specimens were ground and polished using conventional metallography techniques. The Nital reagent (2 vol. % of nitric acid in ethyl alcohol) and sodium metabisulfite reagent (1 g in 100 ml distilled water) were used to reveal 5
the microstructure for optical and scanning electron microscopy (SEM). Microstructural analyses were carried out on the rolling direction (RD) of the sheet. The volume fraction of martensite was measured using universal image analyzing software. Electron backscattering diffraction (EBSD) analyses were conducted using a focused ion & electron beam system nanoDUE’T NB5000 equipped with TSL-solutions orientation imaging microscopy. The EBSD was undertaken in the RD and the normal direction (ND) with step sizes of 0.3 and 0.05 μm. Vickers hardness measurement was performed using a load of 30 kg. The uniaxial tensile test was carried out at room temperature and a crosshead speed of 1 mm/min on tensile test pieces with a gauge length of 25 mm and a width 6.25 mm by ASTM E 8M-04.
3. Results Figure 2 shows optical micrographs of the present initial microstructures. The furnace-cooled steel (Fig. 2a) has a complex ferrite-pearlite (CFP) microstructure consisting of polygonal ferrite grains (pF), massive ferrite (mF) and pearlite (P). The water-cooled steel (Fig. 2b) shows martensite (M), and the steel isothermally transformed for 300 s at 450 ˚C (Fig. 2c) represents bainite (B). Steel sheets with the aforementioned microstructures were cold rolled and intercritically annealed. Hereafter, we will refer to the cold rolled and annealed steels by the corresponding CFP, M and B designations. Figure 3 shows changes in the hardness of the coldrolled initial microstructures with annealing time at 710 and 740 ˚C. It shows hardness decreasing at initial stages, followed by a remarkable increment at the later stages. Decreasing of hardness at the early stage is attributed to the recrystallization of the deformed structures mainly. After that, hardness increases with increasing of annealing time which demonstrates 6
ongoing austenite formation and its transformation to martensite after quenching. CFP shows slightly rapid hardening at 710 ˚C than M and B. Increasing of annealing temperature accelerates the hardness change. Figure 4 shows SEM micrographs of the steels annealed for 1200 s at 710 and 740 ˚C. In these micrographs, ferrite appears in dark gray, martensite in light gray and carbide precipitates in white color. The martensite grains have evolved by transformation of intercritical austenite grains during subsequent water cooling. In the CFP (Fig. 4a), martensite grains have formed on the deformed mF+P bands. In the M (Fig. 4b), martensite grains have formed at grain boundaries. There are innumerous carbide precipitates, and a mild banding of martensite grains. In the B (Fig. 4c), net-arrayed martensite grains has uniformly covered grain boundaries. Annealing at high temperature increases martensite volume fraction and changes the distribution of martensite. In the CFP (Fig. 4d), martensite grains have formed on the earlier mF+P bands. In the M (Fig. 4e), fibrous martensite distribution has been realized. In the B (Fig. 4f), the net-arrayed martensite distribution can be seen. The ferrite grains have been recrystallized. Table 1 gives corresponding volume fractions of martensite (VM) for the microstructures annealed for 1200 s at 710 and 740 ˚C. The volume fraction of martensite is slightly higher in the B. Figure 5 shows EBSD micrographs from the ND of the CFP, M and B after cold rolling and intercritical annealing for 120 s at 710 ˚C. The inverse pole figure (IPF) maps indicate that the CFP (Fig. 5a) has not been recrystallized, but M (Fig. 5b) and B (Fig. 5c) have been recrystallized completely. The Kernel average misorientation (KAM) maps indicate high misorientation in the CFP (Fig. 5d), addressing the non-recrystallized microstructure; but the average misorientation has been decreased in M (Fig. 5e) and B (Fig. 5f) due to recrystallization. Phase mapping identified that the volume fraction of retained austenite is less than 2% in these
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microstructures. A high-resolution EBSD-KAM map obtained from B, annealed for 500 s at 710 ˚C, is shown in Figure 6. This map highlights a ferrite bound around the net-arrayed martensite (see green color around the red). It is assumed that the ferrite bound has been affected by volume expansion during the austenite-martensite transformation. Figure 7 shows EBSD micrographs from the RD of the CFP annealed for 1200 s at 710 ˚C. It is found out that the microstructure has been partially recrystallized. The engineering stress-strain curves of the steels annealed for 1200 s at 710, and 740 ˚C are presented in Figure 8. Corresponding tensile properties are given in Table 2. The initial microstructure marginally affects ultimate tensile strength (UTS) and ductility (EL. %) of the steels annealed at 710 ˚C. However, these properties are slightly more influenced by the initial microstructure after annealing at 740 ˚C. Moreover, it is unlikely to influence the ratio of yield strength to ultimate tensile strength (YS/UTS) at both temperatures. The strength-elongation balance (UTS.EL%) is the same for microstructures annealed at 710 ˚C but changes slightly more for microstructures annealed at 740 ˚C. Strain hardening of each type of the tensile test pieces was evaluated by fitting the stress-strain data into the Holloman equation [42]: σ = Kεn
(1)
where σ is the logarithmic stress, ε is the logarithmic strain, K is the strength coefficient, and n is the strain hardening exponent. A plot of ln (σ) versus ln (ε) was drawn for each type of tensile test pieces as in Figure 9. Corresponding values of n were determined from the slope of the trend lines as given in Table 3. The first exponent (n1) changes with the initial microstructure remarkably, but the second exponent (n2) is almost the same for different initial microstructures.
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4. Discussion The results of this study show that the initial microstructure influences martensite distribution in the cold-rolled DP steels considerably. The amount of martensite realized by the same annealing treatment depends to the initial microstructure too. The martensite distribution manifests the distribution of austenite nucleation sites during intercritical annealing. It becomes net-arrayed when austenite nucleates heterogeneously at grain boundaries; otherwise the banded distribution is realized. Cold-rolled bainite recrystallized well before austenite formation at 710 ˚C; austenite nucleated at boundaries of the recrystallized grains subsequently, and thus net-arrayed martensite distribution was realized. The same process seems to take place at 740 ˚C, though the austenite (martensite) volume fraction increases. Cold-rolled martensite similarly recrystallized before austenite formation at 710 ˚C and produced the net-arrayed distribution; but a fibrous martensite distribution was obtained at 740 ˚C, which indicates austenite nucleation in the deformed martensite before recrystallization. In the present complex ferrite-pearlite, austenite nucleation always occurs in the pearlite bands before recrystallization, and leads to the banded structure. The deformed polygonal ferrite grains are unlikely to recrystallize, but recrystallization of the deformed massive ferrite occurs, which consumes the polygonal grains at the later stages of annealing. A combined effect of high strain accumulation and lower solute carbon in the polygonal ferrite grains probably leads to rapid recovery and thus inhibits recrystallization, as pointed out in Reference [43]. These results are in agreement with the report that deformed ferrite-pearlite recrystallizes slower than deformed bainite and martensite [36]. In addition to the austenite 9
nucleation sites, recrystallization influences the rate of subsequent austenite formation in the different initial microstructures during the same annealing treatment. Hardness change indicates that the austenite forms more rapidly in the deformed ferrite-pearlite than recrystallized martensite and bainite during annealing at 710 ˚C. It has been reported that austenite forms fast in the deformed structures due to the increased nucleation sites and enhanced diffusion [44]. The effect of the initial microstructure on the tensile strength and ductility is remarkable at the higher annealing temperature which provides higher volume fraction of martensite. The net-arrayed microstructure produced by using bainite shows higher strength and reduced ductility. However, there are different amounts, and distributions of martensite, ferrite, and carbide precipitates in the microstructures produced from different initial microstructures. Therefore, it is difficult to correlate the tensile strength and ductility with the different phase combinations formed in the different initial microstructures. Nevertheless, the initial strain hardening exponent (initial flow) seems to systematically change with the microstructure than the tensile strength and ductility. For instance, the net-arrayed martensite produced by using either bainite or martensite shows higher initial strain hardening than the partially recrystallized microstructure realized by using ferrite-pearlite. The net-arrayed martensite produced by annealing of cold-rolled bainite at 740 ˚C shows higher initial strain hardening than the others too. Micromechanical simulation has shown that the initial flow is affected by the ferrite bound having residual stresses and plastic strains, and that the skeleton type (netarrayed) martensite gives higher initial strain hardening [45, 46]. The ferrite bound near martensite is thought to prohibit crack propagation into the ferrite and thus improves tensile
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properties [36]. Cold-rolled bainite provides the net-arrayed martensite by stimulated austenite nucleation and growth at grain boundaries due to the enhanced recrystallization.
Conclusion 1. Cold-rolled martensite and bainite recrystallized, but cold-rolled ferrite-pearlite was unlikely to recrystallize before austenite formation at 710 ˚C. The Hardness of the recrystallized microstructures increases similarly, but in a lower rate than the nonrecrystallized microstructure. Increasing of annealing temperature stimulates austenite formation against recrystallization and accelerates the hardness change. 2. If recrystallization takes place before austenite formation, it will result in net-arrayed martensite distribution; otherwise banded or fibrous distribution of martensite is realized. Intercritical annealing at low temperature and using bainite microstructure act in favor of enhanced recrystallization, and thus producing the net-arrayed martensite distribution. 3. Effect of initial microstructure on the strength and ductility is marginal at the lower volume fraction of martensite, but becomes appreciable at the higher volume fraction. 4. The net-arrayed martensite gives higher strength and strain hardening at higher martensite volume fraction. It shows higher initial strain hardening presumably due to the transformation induced effects in ferrite bound near martensite.
Acknowledgment
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This research was collaboratively supported by Iranian Mines & Mining Industries Development & Renovation Organization through Grant No. 30.13727, the Sahand University of Technology through Grant in Aid for Ph. D. Students (No. 30.7650). A part of this study is financially supported by the Project to Create Research and Educational Hubs for Innovative Manufacturing in Asia supported by the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan.
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Fig 1. Schematic representation of cold rolling and heat treatments. Fig 2. Optical micrographs of the initial microstructures: (a) furnace cooled steel; (b) water cooled steel; (c) isothermally transformed steel. pF: Polygonal ferrite; mF: massive ferrite; P: pearlite. Fig 3. Hardness changes of the cold-rolled initial microstructures with intercritical annealing time at (a) 710 ˚C; (b) 740 ˚C. Fig 4. SEM micrographs of the cold-rolled initial microstructures intercritically annealed for 1200 s at 710 (left column) and 740 ˚C (right column): (a, d) CFP; (b, e) M; (c, f) B. Fig 5. EBSD IPF (a-c) and KAM (d-f) maps of the cold-rolled initial microstructures annealed for 120 s at 710 ˚C. (a,d) CFP; (b, e) M; (c, f) B. Fig 6. EBSD KAM map of cold-rolled bainite (B) annealed for 500 s at 710 ˚C. Fig 7. EBSD IPF (a) and KAM (b) maps of the CFP annealed for 1200 s at 710 ˚C. Fig 8. (a) Engineering tensile stress-strain curves of the cold rolled initial microstructures intercritically annealed for 1200 s at (a) 710 ˚C; (b) 740 ˚C. Fig 9. Plots of ln (σ) versus ln (ε): (a) cold-rolled initial microstructures annealed for 1200 s at 710 ˚C; (b) cold-rolled initial microstructures annealed for 1200 s at 740 ˚C.
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Figure(1)
Austenitization Treatment 900 °C – 3600 s
Water Cooling à M
Isothermal Transformation* à B
Furnace Cooling à CFP (pF+mF+P)
Cold Rolling (50%)
Ae1 = 690 °C
Ae3 = 820 °C
740 °C / 10 – 1800 s 710 °C / 10 – 1800 s
Inter-Critical Annealing (F+A)
Designations: Polygonal Ferrite (pF) –Massive Ferrite (mF) Pearlite (P) - Complex Ferrite-Pearlite (CFP) Austenite (A) – Bainite (B) – Martensite (M) * 450 ° C – 300 s
Figure(2)
Figure(3)
500 CFP
M
B
400 300 200
Hardness HV30
Hardness HV30
500
M
B
400 300 200 100
100
a
CFP
10
100 1000 10000 Annealing time (s)
b
10 100 1000 10000 Annealing time (s)
Figure(4)
Figure(5)
Figure(6)
Figure(7)
700
1400
600
1200
500
1000
Stress (MPa)
Stress (MPa)
Figure(8)
400 300
CP
200
B
100
M
CP
600
B
400
M
200
0
0 0
a
800
10 20 Strain (%)
30
0
b
5 10 Strain (%)
15
Fig. 9
7
7.25 7 ln (σ)
ln (σ)
6.5
6.75
6 6.5 CFP
M
CFP
B
5.5
a
M
B
6.25 -6
-5
-4 -3 ln (ε)
-2
-1
-6
b
-5
-4 -3 ln (ε)
-2
-1
Table 1. Volume fraction of martensite (Vm) in the cold-rolled microstructures intercritically annealed for 1200 s steel
annealed at
annealed at 740
710 ˚C
˚C
CFP
25
59
M
23
57
B
29
62
Table 2. Tensile properties of the cold-rolled and intercritically annealed microstructures
steel
YS (MPa)
UTS (MPa)
EL%
YS/UTS
(UTS.EL%)×10 0
annealed for 1200 s at 710 ˚C CFP 282 653
22.0
0.43
144
M
281
644
22.1
0.44
142
B
269
634
22.9
0.43
145
annealed for 1200 s at 740 ˚C CFP 603 1186
9.1
0.51
107
M
615
1207
8.5
0.51
103
B
606
1218
7.3
0.50
89
1
Table 3. Strain hardening exponents (n) of the cold-rolled and intercritically annealed microstructures
steel
n1
n2
annealed for 1200 s at 710 ˚C CFP
0.14
0.29
M
0.27
0.30
B
0.27
0.30
annealed for 1200 s at 740 ˚C CFP
0.45
0.20
M
0.30
0.19
B
0.53
0.16
2
PII: DOI: Reference:
S0921-5093(18)31196-1 https://doi.org/10.1016/j.msea.2018.09.008 MSA36889
To appear in: Materials Science & Engineering A Received date: 22 May 2018 Revised date: 3 September 2018 Accepted date: 4 September 2018 Cite this article as: Y. Karimi, S. Hossein Nedjad, H. Shirazi, M. Nili Ahmadabadi, H. Hamed Zargari and K. Ito, Cold rolling and intercritical annealing of C-Mn steel sheets with different initial microstructures, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2018.09.008 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Cold rolling and intercritical annealing of C-Mn steel sheets with different initial microstructures
Y. Karimia, S. Hossein Nedjada*, H. Shirazib, M. Nili Ahmadabadib, H. Hamed Zargaric, K. Itoc
a
Faculty of Materials Engineering, Sahand University of Technology, P.O. Box: 51335-1996, Tabriz, Iran.
b
School of Metallurgy and Materials Engineering, University of Tehran, P.O. Box: 14395-731, Tehran, Iran.
c
Joining and Welding Research Institute, Osaka University, 11-1 Mihogaoka, Ibaraki, Osaka 567-0047, Japan.
*
Corresponding author. Tel.: +98-41- 3345 9449; fax: +98-41-3344 4333. [email protected] (Syamak Hossein Nedjad).
Abstract The microstructure and mechanical properties of cold-rolled dual-phase steels produced from different initial microstructures, i.e., complex ferrite-pearlite, martensite, and bainite, were studied for a C-Mn steel contained 0.15C, 2.50Mn, and 0.30Si (wt. %). Sheets of 4 mm thickness, and with the different initial microstructures were cold rolled for 50% thickness 1
reduction, intercritically annealed at 710 and 740 ˚C for various times (10 – 1800 s) in a salt bath, and then water cooled. Different amounts and distributions of martensite were produced from the different initial microstructures by the same annealing treatment. Bainite was overwhelming which produced slightly higher amounts of martensite and the net-arrayed martensite distribution. Electron backscattering diffraction revealed recrystallization of coldrolled bainite before austenite formation at 710 ˚C; and that the ferrite bound near martensite has remarkable transformation induced effects. Tensile strength and ductility of the steels annealed for 1200 s at 710 ˚C were almost the same. The effect of the initial microstructure on the tensile properties is remarkable for steels annealed for the same time at 740 ˚C. The netarrayed martensite produced by using bainite represented higher initial strain hardening at both temperatures, which is attributed to the transformation effects in the ferrite bound around martensite.
Keywords: Dual-phase steel; Cold rolling; Intercritical annealing; Recrystallization; Austenite formation.
1. Introduction The automotive industries use up the advanced high strength steels (AHSS) because of better crashworthiness and weight reduction [1]. Dual-phase (DP) steels are the most prominent category of AHSS, which have earned widespread applications in the fabrication of chassis, pillars, shields, rocker rail and cross members; owing to good formability, high strength and affordability [2, 3]. The microstructure of DP steels mainly consists of a ferrite-martensite mixture. They are produced by annealing or thermo-mechanical processing in the austenite2
ferrite intercritical region [4]. Mechanical properties of DP steels depend mainly on the mean size of ferrite grains and the volume fraction of martensite [5-8]. Martensite morphology (shape, size, and distribution) also influences the mechanical properties, and effect of different heat treatments, producing different martensite morphologies, has been well studied [9-23]. Three kinds of heat treatments have frequently been compared [9, 10, 12, 16]: (i) intercritical annealing; (ii) intermediate quenching; and (iii) step quenching. The intermediate quenching provides martensite before the intercritical annealing treatment and gives fine and fibrous microstructures with superior strength-ductility balance. Step quenching results in coarse and blocky microstructures with poor tensile properties. The intercritical annealing yields globular microstructures with intermediate tensile properties. Indeed, it is the initial microstructure before intercritical annealing treatment, which is modified by using these kinds of heat treatments. A part of DP steels is produced by additional cold rolling of available steel products, followed by annealing in the austenite-ferrite intercritical region, and subsequent cooling to transform austenite grains into martensite [24-26]. A feature of the heat treatment of coldrolled DP steels is the recrystallization of cold-rolled microstructure and its interaction with austenite formation. Complete recrystallization of the deformed structure before austenite formation provides a random distribution of austenite while incomplete recrystallization gives rise to the formation of banded austenite grains [27, 28]. Austenite formation in the deformed structure retards the growth of subsequently recrystallized ferrite grains and results in grain refinement. Recrystallization precedes austenite formation at low intercritical annealing temperatures and low heating rates [29-31]. It has also been reported that the cooling rate
3
after intercritical annealing influences carbon content of austenite and thus the martensite start temperature (Ms) in the cold-rolled DP steels. When subsequent isothermal treatment (aging) is carried out at a temperature lower than Ms, tempered martensite and ferrite appear; Otherwise bainite formation becomes possible [32]. In addition to the heating and cooling rate effects, the influence of the initial microstructure of cold-rolled DP steels has to be taken into account. Hot-rolled steel plates are often used for cold rolling, which usually consists of the ferrite-pearlite microstructure. During intercritical annealing of a cold-rolled ferrite-pearlite microstructure; recrystallization of the deformed ferrite, spheroidization of fragmented cementite and austenite formation take place concomitantly [33, 34]. Nucleation of austenite firstly occurs at ferrite-ferrite grain boundaries, and then on the cementite-ferrite interfaces. However, advantages of making the use of bainite and martensite as the initial microstructure for cold-rolled DP steels have been reported. For example, austenite formation in the deformed bainite has resulted in remarkable microstructure refinement [35]. Recrystallization of the deformed martensite during intercritical annealing has led to a refined dual-phase microstructure [36]. It has also been reported that cold-rolled bainite-ferrite-pearlite and martensite have higher rates of recrystallization and austenite formation than cold rolled ferrite-pearlite [37]. A DP steel produced by cold rolling and intercritical annealing of bainite-ferrite initial microstructure has shown net-arrayed martensite distribution which resulted in a marked strain hardening, increased tensile strength and fatigue endurance [38]. A ferrite-fibrous martensite initial microstructure has shown more refined microstructure and superior tensile properties than ferrite-pearlite after cold rolling and intercritical annealing [39].
4
The efficacy of step-quenched ferrite-martensite and bainite initial microstructures for grain refinement during cold rolling and annealing treatment, along with the banding effects has been studied for a C-Mn steel contained about 0.15C and 2.50Mn (wt. %) [40, 41]. For the same steel, this paper is to compare the microstructure and mechanical properties of coldrolled DP steels produced from different initial microstructures, i.e., complex ferrite-pearlite, martensite, and bainite.
2. Experimental Procedure An ingot weighing 35 kg was prepared using induction melting under inert argon protection and subsequently refined by electroslag remelting (ESR). The ingot contained 0.15C, 2.50Mn, and 0.30Si (in wt. %). It was hot rolled to a sheet of about 4 mm thickness in several passes then air cooled to room temperature. The AC1 and AC3 temperatures of the air-cooled steel were determined about 690 ˚C and 820 ˚C, respectively, using the Adamel dilatometer. The hot-rolled sheet was used to create different initial microstructures. For this purpose, strips (4 mm × 100 mm × 150 mm) cut from the hot-rolled sheet were austenitized in a resistance furnace at 900 ˚C for 3600 s, then cooled at different conditions, as shown in Figure 1. The strips were subsequently rolled by 50% thickness reduction at room temperature. The next step involved heating of the cold-rolled microstructures in a salt bath to intercritical annealing temperatures (710 and 740 ˚C) and soaking for different times (10 – 1800 s). Finally, the samples were quenched in water to produce DP microstructures. Specimens were ground and polished using conventional metallography techniques. The Nital reagent (2 vol. % of nitric acid in ethyl alcohol) and sodium metabisulfite reagent (1 g in 100 ml distilled water) were used to reveal 5
the microstructure for optical and scanning electron microscopy (SEM). Microstructural analyses were carried out on the rolling direction (RD) of the sheet. The volume fraction of martensite was measured using universal image analyzing software. Electron backscattering diffraction (EBSD) analyses were conducted using a focused ion & electron beam system nanoDUE’T NB5000 equipped with TSL-solutions orientation imaging microscopy. The EBSD was undertaken in the RD and the normal direction (ND) with step sizes of 0.3 and 0.05 μm. Vickers hardness measurement was performed using a load of 30 kg. The uniaxial tensile test was carried out at room temperature and a crosshead speed of 1 mm/min on tensile test pieces with a gauge length of 25 mm and a width 6.25 mm by ASTM E 8M-04.
3. Results Figure 2 shows optical micrographs of the present initial microstructures. The furnace-cooled steel (Fig. 2a) has a complex ferrite-pearlite (CFP) microstructure consisting of polygonal ferrite grains (pF), massive ferrite (mF) and pearlite (P). The water-cooled steel (Fig. 2b) shows martensite (M), and the steel isothermally transformed for 300 s at 450 ˚C (Fig. 2c) represents bainite (B). Steel sheets with the aforementioned microstructures were cold rolled and intercritically annealed. Hereafter, we will refer to the cold rolled and annealed steels by the corresponding CFP, M and B designations. Figure 3 shows changes in the hardness of the coldrolled initial microstructures with annealing time at 710 and 740 ˚C. It shows hardness decreasing at initial stages, followed by a remarkable increment at the later stages. Decreasing of hardness at the early stage is attributed to the recrystallization of the deformed structures mainly. After that, hardness increases with increasing of annealing time which demonstrates 6
ongoing austenite formation and its transformation to martensite after quenching. CFP shows slightly rapid hardening at 710 ˚C than M and B. Increasing of annealing temperature accelerates the hardness change. Figure 4 shows SEM micrographs of the steels annealed for 1200 s at 710 and 740 ˚C. In these micrographs, ferrite appears in dark gray, martensite in light gray and carbide precipitates in white color. The martensite grains have evolved by transformation of intercritical austenite grains during subsequent water cooling. In the CFP (Fig. 4a), martensite grains have formed on the deformed mF+P bands. In the M (Fig. 4b), martensite grains have formed at grain boundaries. There are innumerous carbide precipitates, and a mild banding of martensite grains. In the B (Fig. 4c), net-arrayed martensite grains has uniformly covered grain boundaries. Annealing at high temperature increases martensite volume fraction and changes the distribution of martensite. In the CFP (Fig. 4d), martensite grains have formed on the earlier mF+P bands. In the M (Fig. 4e), fibrous martensite distribution has been realized. In the B (Fig. 4f), the net-arrayed martensite distribution can be seen. The ferrite grains have been recrystallized. Table 1 gives corresponding volume fractions of martensite (VM) for the microstructures annealed for 1200 s at 710 and 740 ˚C. The volume fraction of martensite is slightly higher in the B. Figure 5 shows EBSD micrographs from the ND of the CFP, M and B after cold rolling and intercritical annealing for 120 s at 710 ˚C. The inverse pole figure (IPF) maps indicate that the CFP (Fig. 5a) has not been recrystallized, but M (Fig. 5b) and B (Fig. 5c) have been recrystallized completely. The Kernel average misorientation (KAM) maps indicate high misorientation in the CFP (Fig. 5d), addressing the non-recrystallized microstructure; but the average misorientation has been decreased in M (Fig. 5e) and B (Fig. 5f) due to recrystallization. Phase mapping identified that the volume fraction of retained austenite is less than 2% in these
7
microstructures. A high-resolution EBSD-KAM map obtained from B, annealed for 500 s at 710 ˚C, is shown in Figure 6. This map highlights a ferrite bound around the net-arrayed martensite (see green color around the red). It is assumed that the ferrite bound has been affected by volume expansion during the austenite-martensite transformation. Figure 7 shows EBSD micrographs from the RD of the CFP annealed for 1200 s at 710 ˚C. It is found out that the microstructure has been partially recrystallized. The engineering stress-strain curves of the steels annealed for 1200 s at 710, and 740 ˚C are presented in Figure 8. Corresponding tensile properties are given in Table 2. The initial microstructure marginally affects ultimate tensile strength (UTS) and ductility (EL. %) of the steels annealed at 710 ˚C. However, these properties are slightly more influenced by the initial microstructure after annealing at 740 ˚C. Moreover, it is unlikely to influence the ratio of yield strength to ultimate tensile strength (YS/UTS) at both temperatures. The strength-elongation balance (UTS.EL%) is the same for microstructures annealed at 710 ˚C but changes slightly more for microstructures annealed at 740 ˚C. Strain hardening of each type of the tensile test pieces was evaluated by fitting the stress-strain data into the Holloman equation [42]: σ = Kεn
(1)
where σ is the logarithmic stress, ε is the logarithmic strain, K is the strength coefficient, and n is the strain hardening exponent. A plot of ln (σ) versus ln (ε) was drawn for each type of tensile test pieces as in Figure 9. Corresponding values of n were determined from the slope of the trend lines as given in Table 3. The first exponent (n1) changes with the initial microstructure remarkably, but the second exponent (n2) is almost the same for different initial microstructures.
8
4. Discussion The results of this study show that the initial microstructure influences martensite distribution in the cold-rolled DP steels considerably. The amount of martensite realized by the same annealing treatment depends to the initial microstructure too. The martensite distribution manifests the distribution of austenite nucleation sites during intercritical annealing. It becomes net-arrayed when austenite nucleates heterogeneously at grain boundaries; otherwise the banded distribution is realized. Cold-rolled bainite recrystallized well before austenite formation at 710 ˚C; austenite nucleated at boundaries of the recrystallized grains subsequently, and thus net-arrayed martensite distribution was realized. The same process seems to take place at 740 ˚C, though the austenite (martensite) volume fraction increases. Cold-rolled martensite similarly recrystallized before austenite formation at 710 ˚C and produced the net-arrayed distribution; but a fibrous martensite distribution was obtained at 740 ˚C, which indicates austenite nucleation in the deformed martensite before recrystallization. In the present complex ferrite-pearlite, austenite nucleation always occurs in the pearlite bands before recrystallization, and leads to the banded structure. The deformed polygonal ferrite grains are unlikely to recrystallize, but recrystallization of the deformed massive ferrite occurs, which consumes the polygonal grains at the later stages of annealing. A combined effect of high strain accumulation and lower solute carbon in the polygonal ferrite grains probably leads to rapid recovery and thus inhibits recrystallization, as pointed out in Reference [43]. These results are in agreement with the report that deformed ferrite-pearlite recrystallizes slower than deformed bainite and martensite [36]. In addition to the austenite 9
nucleation sites, recrystallization influences the rate of subsequent austenite formation in the different initial microstructures during the same annealing treatment. Hardness change indicates that the austenite forms more rapidly in the deformed ferrite-pearlite than recrystallized martensite and bainite during annealing at 710 ˚C. It has been reported that austenite forms fast in the deformed structures due to the increased nucleation sites and enhanced diffusion [44]. The effect of the initial microstructure on the tensile strength and ductility is remarkable at the higher annealing temperature which provides higher volume fraction of martensite. The net-arrayed microstructure produced by using bainite shows higher strength and reduced ductility. However, there are different amounts, and distributions of martensite, ferrite, and carbide precipitates in the microstructures produced from different initial microstructures. Therefore, it is difficult to correlate the tensile strength and ductility with the different phase combinations formed in the different initial microstructures. Nevertheless, the initial strain hardening exponent (initial flow) seems to systematically change with the microstructure than the tensile strength and ductility. For instance, the net-arrayed martensite produced by using either bainite or martensite shows higher initial strain hardening than the partially recrystallized microstructure realized by using ferrite-pearlite. The net-arrayed martensite produced by annealing of cold-rolled bainite at 740 ˚C shows higher initial strain hardening than the others too. Micromechanical simulation has shown that the initial flow is affected by the ferrite bound having residual stresses and plastic strains, and that the skeleton type (netarrayed) martensite gives higher initial strain hardening [45, 46]. The ferrite bound near martensite is thought to prohibit crack propagation into the ferrite and thus improves tensile
10
properties [36]. Cold-rolled bainite provides the net-arrayed martensite by stimulated austenite nucleation and growth at grain boundaries due to the enhanced recrystallization.
Conclusion 1. Cold-rolled martensite and bainite recrystallized, but cold-rolled ferrite-pearlite was unlikely to recrystallize before austenite formation at 710 ˚C. The Hardness of the recrystallized microstructures increases similarly, but in a lower rate than the nonrecrystallized microstructure. Increasing of annealing temperature stimulates austenite formation against recrystallization and accelerates the hardness change. 2. If recrystallization takes place before austenite formation, it will result in net-arrayed martensite distribution; otherwise banded or fibrous distribution of martensite is realized. Intercritical annealing at low temperature and using bainite microstructure act in favor of enhanced recrystallization, and thus producing the net-arrayed martensite distribution. 3. Effect of initial microstructure on the strength and ductility is marginal at the lower volume fraction of martensite, but becomes appreciable at the higher volume fraction. 4. The net-arrayed martensite gives higher strength and strain hardening at higher martensite volume fraction. It shows higher initial strain hardening presumably due to the transformation induced effects in ferrite bound near martensite.
Acknowledgment
11
This research was collaboratively supported by Iranian Mines & Mining Industries Development & Renovation Organization through Grant No. 30.13727, the Sahand University of Technology through Grant in Aid for Ph. D. Students (No. 30.7650). A part of this study is financially supported by the Project to Create Research and Educational Hubs for Innovative Manufacturing in Asia supported by the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan.
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[41] Y. Karimi, S. Hossein Nedjad, G. Miyamoto, H. Shirazi, T. Furuhara, Banding effects on the process of grain refinement by cold deformation and recrystallization of acicular C-Mn steel, Mater. Sci. Eng. A, 697 (2017) 1-7. [42] J.H. Holloman, Tensile deformation, Trans. AIME 162 (1945) 268–290. [43] A. Karmakar, M. Mandal, A. Mandal, M. Basiruddin, S. Mukherjee, D. Chakrabarti, Effect of starting microstructure on the grain refinement in cold-rolled low-carbon steel during annealing at two different heating rates, Metall. Mater. Trans. A. 47 (2016) 268- 281. [44] M. Ollat, V. Massardier, D. Fabregue, E. Buscarlet, F. Keovilay, M. Perez, Modeling of the recrystallization and austenite formation overlapping in cold-rolled dual-phase steels during intercritical treatments, Metall. Mater. Trans. A. 48A (2017) 4486-4499. [45] A. Fillafer, E. Werner, C. Krempaszky, On phase transformation induced effects controlling the initial flow behavior of ferritic-martensitic dual-phase steels, Mater. Sci. Eng. A 708 (2017) 556–562. [46] U. Liedl, S. Traint, E.A. Werner, An unexpected feature of the stress–strain diagram of dualphase steel, Comp. Mater. Sci. 25 (2002) 122–128.
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Fig 1. Schematic representation of cold rolling and heat treatments. Fig 2. Optical micrographs of the initial microstructures: (a) furnace cooled steel; (b) water cooled steel; (c) isothermally transformed steel. pF: Polygonal ferrite; mF: massive ferrite; P: pearlite. Fig 3. Hardness changes of the cold-rolled initial microstructures with intercritical annealing time at (a) 710 ˚C; (b) 740 ˚C. Fig 4. SEM micrographs of the cold-rolled initial microstructures intercritically annealed for 1200 s at 710 (left column) and 740 ˚C (right column): (a, d) CFP; (b, e) M; (c, f) B. Fig 5. EBSD IPF (a-c) and KAM (d-f) maps of the cold-rolled initial microstructures annealed for 120 s at 710 ˚C. (a,d) CFP; (b, e) M; (c, f) B. Fig 6. EBSD KAM map of cold-rolled bainite (B) annealed for 500 s at 710 ˚C. Fig 7. EBSD IPF (a) and KAM (b) maps of the CFP annealed for 1200 s at 710 ˚C. Fig 8. (a) Engineering tensile stress-strain curves of the cold rolled initial microstructures intercritically annealed for 1200 s at (a) 710 ˚C; (b) 740 ˚C. Fig 9. Plots of ln (σ) versus ln (ε): (a) cold-rolled initial microstructures annealed for 1200 s at 710 ˚C; (b) cold-rolled initial microstructures annealed for 1200 s at 740 ˚C.
18
Figure(1)
Austenitization Treatment 900 °C – 3600 s
Water Cooling à M
Isothermal Transformation* à B
Furnace Cooling à CFP (pF+mF+P)
Cold Rolling (50%)
Ae1 = 690 °C
Ae3 = 820 °C
740 °C / 10 – 1800 s 710 °C / 10 – 1800 s
Inter-Critical Annealing (F+A)
Designations: Polygonal Ferrite (pF) –Massive Ferrite (mF) Pearlite (P) - Complex Ferrite-Pearlite (CFP) Austenite (A) – Bainite (B) – Martensite (M) * 450 ° C – 300 s
Figure(2)
Figure(3)
500 CFP
M
B
400 300 200
Hardness HV30
Hardness HV30
500
M
B
400 300 200 100
100
a
CFP
10
100 1000 10000 Annealing time (s)
b
10 100 1000 10000 Annealing time (s)
Figure(4)
Figure(5)
Figure(6)
Figure(7)
700
1400
600
1200
500
1000
Stress (MPa)
Stress (MPa)
Figure(8)
400 300
CP
200
B
100
M
CP
600
B
400
M
200
0
0 0
a
800
10 20 Strain (%)
30
0
b
5 10 Strain (%)
15
Fig. 9
7
7.25 7 ln (σ)
ln (σ)
6.5
6.75
6 6.5 CFP
M
CFP
B
5.5
a
M
B
6.25 -6
-5
-4 -3 ln (ε)
-2
-1
-6
b
-5
-4 -3 ln (ε)
-2
-1
Table 1. Volume fraction of martensite (Vm) in the cold-rolled microstructures intercritically annealed for 1200 s steel
annealed at
annealed at 740
710 ˚C
˚C
CFP
25
59
M
23
57
B
29
62
Table 2. Tensile properties of the cold-rolled and intercritically annealed microstructures
steel
YS (MPa)
UTS (MPa)
EL%
YS/UTS
(UTS.EL%)×10 0
annealed for 1200 s at 710 ˚C CFP 282 653
22.0
0.43
144
M
281
644
22.1
0.44
142
B
269
634
22.9
0.43
145
annealed for 1200 s at 740 ˚C CFP 603 1186
9.1
0.51
107
M
615
1207
8.5
0.51
103
B
606
1218
7.3
0.50
89
1
Table 3. Strain hardening exponents (n) of the cold-rolled and intercritically annealed microstructures
steel
n1
n2
annealed for 1200 s at 710 ˚C CFP
0.14
0.29
M
0.27
0.30
B
0.27
0.30
annealed for 1200 s at 740 ˚C CFP
0.45
0.20
M
0.30
0.19
B
0.53
0.16
2