effects of distributions of constituent phases on

ISIJ International, Vol. 61 (2021), No. 1

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ISIJ International, Vol. 61 (2021), No. 1, pp. 1–11

https://doi.org/10.2355/isijinternational.ISIJINT-2020-435

* Corresponding author: E-mail: [email protected]

© 2021 The Iron and Steel Institute of Japan. This is an open access article under the terms of the Creative Commons Attribution-NonCommercial-NoDerivs license (https://creativecommons.org/licenses/by-nc-nd/4.0/).CCBYNCND

1. Introduction

High strength (HS) steels are increasingly used in the automotive industry in order to improve automobile safety and solve environmental problems associated with carbon dioxide emission. For such HS steels, however, strengthen-ing may cause deterioration of formability. Compared with conventional HS steels, multiphase HS steels like dual phase (DP) and low alloyed transformation induced plasticity (TRIP) steels display superior combination of strength and formability. Thus, multiphase HS steels are also widely used in the automotive industry. The formability of DP and low

Effects of Distributions of Constituent Phases on Mechanical Properties of C–Si–Mn Dual-phase Steel

Tatsuya NAKAGAITO,1)* Takako YAMASHITA,2) Takeshi YOKOTA,2) Yoshihiro TERADA3) and Masanori KAJIHARA3)

1) Steel Research Laboratory, JFE Steel Corporation, 1 Kokan-cho, Fukuyama, Hiroshima, 721–8510 Japan.2) Steel Research Laboratory, JFE Steel Corporation, 1 Kawasaki-cho, Chuo-ku, Chiba, 260-0835 Japan.3) Department of Materials Science and Engineering, Tokyo Institution of Technology, Nagatsuta-chou 4259, Midori-ku, Yokohama, 226-8502 Japan.

(Received on July 20, 2020; accepted on September 11, 2020; J-STAGE Advance published date: October 31, 2020)

Two kinds of ferrite were produced due to slow cooling after intercritical annealing, one being intercriti-cally-annealed ferrite (αa phase), and the other transformed ferrite (αc phase). The effects of distributions of the αa and αc phases on the mechanical properties of a Fe–0.17C–1.5Si–1.7Mn dual-phase steel were examined experimentally at room temperature. Two types of intercritical annealing (IA) were conducted to control the distribution. The IA temperature and time are Ta = 800°C (1 073 K) and ta = 0.5 h (1.8 ks), respec-tively, for the first type, and Ta = 740°C (1 013 K) and ta = 4 h (14.4 ks), respectively, for the second one. For both types of IA, the steel was slowly cooled to 400°C (673 K) at a cooling rate of 10°C/s, followed by water quenching. While the total volume fraction fα of the αa and αc phases is close to 0.68–0.69 for both Ta = 800 and 740°C (1 073 and 1 013 K), the combination of the volume fraction of each α phase is different for Ta =800 and 740°C. The volume fraction fαa of the αa phase and fαc of the αc phase is 0.33 and 0.36 for Ta = 800°C, respectively, and 0.68 and 0 for Ta = 740°C, respectively. The ultimate tensile strength su is about 970 MPa for fα = 0.68–0.69 independent of combination of fαa and fαc values. Thus, the effect of the α phase on su is close to each other between the αa and αc phases. In contrast, both the uniform strain eu and the local strain el increase with increasing volume fraction fαc of the αc phase. Such increase in eu is attributed to larger values of the strain hardening rate dσ/dε in the large strain region. Misfit strain stored at the boundary between the αa and αc phases causes to the larger values of dσ/dε. On the other hand, suppression of void formation deduces the increase in el. Acceleration of dynamic recovery of the α phase adjacent to martensite (M phase) and decrease in the hardness of the M phase adjacent to the α phase due to formation of the αc phase can suppress void formation. Consequently, formation of the αc phase contributes to improvement of the mechanical properties of the dual-phase steel.

KEY WORDS: phase transformation; ferrite; martensite; dual phase; intercritical annealing; partitioning; local equilibrium; paraequilibrium; elongation; ductility; uniform elongation; strain hardening; void.

alloyed TRIP steels has been extensively investigated from the viewpoint of distributions of the constituent phases, such as ferrite, martensite, retained austenite, etc.1–8) Among these phases, ferrite is relatively soft and ductile, and thus contributes to improvement of formability. Usually, multi-phase HS steels are manufactured by intercritical annealing after cold rolling, followed by controlled cooling. Two kinds of ferrite form owing to these heat treatments. One is the intercritically-annealed ferrite (αa phase) produced during intercritical annealing, and the other one is the transformed ferrite (αc phase) formed during controlled cooling. It is reported by some researchers9–11) that formation of the αc phase effectively improves ductility of DP and low alloyed TRIP steels. Since it is difficult to distinguish the αc phase

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from the αa phase in the microstructure, however, the mechanism of the improvement is not well understood yet.

Recently, microstructure evolution of a Fe–0.17C–1.5Si–1.7Mn DP steel during intercritical annealing and subsequent slow cooling was experimentally studied by Nakagaito et al.12) In that experiment, a new field emission electron probe microanalysis (FE-EPMA) technique13,14) was used to observe the microstructure of the DP steel. This technique enables highly accurate measurement of two-dimensional (2D) mapping not only for substitutional solute elements such as Mn and Si but also for interstitutional solute elements such as C. Due to the accurate 2D mapping of C, Mn and Si in the microstructure, the αc phase was clearly distinguished from the αa phase. Here, the content of C is close to each other between the αa and αc phases. In contrast, the content of Mn is greater for the αc phase than for the αa phase, but that of Si is smaller for the αc phase than for the αa phase. Such difference in the chemi-cal composition between the αa and αc phases is attributed to partitioning behavior of the substitutional solute element with the austenite formation during intercritical annealing and the ferrite transformation during slow cooling.

In the present study, effects of distributions of the αa and αc phases on mechanical properties of a DP steel were experimentally examined using a Fe-0.17C-1.5Si-1.7Mn steel. The DP steel was intercritically annealed at 800°C (1 073 K) for 0.5 h (1.8 ks) or at 740°C (1 013 K) for 4 h (14.4 ks), and then slowly cooled to various temperatures down to 400°C (673 K) at a cooling rate of 10°C/s, followed by water quenching. An Instron type testing machine and a nano-indentation hardness testing machine were used to measure mechanical properties of the heat treated DP steel. In contrast, characteristic features of the microstructure for the DP steel were observed by the new FE-EPMA tech-nique. The effects of distributions of the αa and αc phases on the mechanical properties were extensively discussed on the basis of the observation.

2. Experimental Procedures

A Fe–0.17C–1.5Si–1.7Mn steel was used in the present study. For this steel, the Ae1 and Ae3 temperatures are calcu-lated as 708 and 857°C (981 and 1 130 K), respectively, by Thermo-Calc® software. Here, Ae1 is the highest temperature at which cementite (θ phase) completely disappears, and Ae3 is the lowest temperature where the microstructure is com-posed of only austenite (γ phase). The chemical composi-tion and the Ae1 and Ae3 temperatures of the steel are shown in Table 1. A slab of the steel with a thickness of 27 mm prepared by the technique reported in a previous study12) was heated at 1 250°C (1 523 K) and then hot-rolled to a thickness of 4.0 mm with a finishing temperature of 900°C (1 173 K), followed by cooling and holding at 600°C (873 K) for 1 h (3.6 ks). The surface of the hot-rolled sheet was removed by grinding to a thickness of 3.2 mm. After cold rolling to a thickness of 1.2 mm, heat treatments were carried out with a salt bath furnace. The heat treatment is schematically depicted in Fig. 1. As shown in Fig. 1, the cold-rolled specimen was intercritically annealed at Ta = 800°C (1 073 K) for ta = 0.5 h (1.8 ks) or at Ta = 740°C (1 013 K) for ta = 4 h (14.4 ks). Here, Ta and ta are the

intercritical annealing temperature and time, respectively. After the intercritical annealing, the specimen was slowly cooled to the temperature Tq at a cooling rate of 10°C/s, followed by water quenching. Here, Tq is the quenching temperature taking the minimum value of 400°C (673 K).

Tensile tests were conducted on the specimen of JIS5 type at room temperature using an Instron type testing machine. For the tensile test, the gauge length was 50 mm, the parallel-part width was 25 mm, and the cross-head speed was 10 mm/min. After obtaining the nominal stress s versus the nominal strain e (s–e) curve, the maximum value of s was used as the ultimate tensile strength su. The total strain et was evaluated from the measured length of the fractured pieces divided by the gauge length. Hardness tests were also made on the specimen at room temperature using a nano-indentation hardness testing machine with an iMicro nano indenter (KLA Co.) under load of 250 μN. The hardness dis-tribution was measured at 900 points on the cross-sectional surface of the specimen, where the surface was polished with alumina and colloidal silica solutions.

The microstructure of each specimen was observed by scanning electron microscopy (SEM). For the SEM observation, a cross sectional surface of the specimen was mechanically polished and then chemically etched with Nital containing 1 vol% of nitric acid and 99 vol% of etha-nol. On the SEM image, the volume fraction fϕ of phase ϕ was determined by a quantitative image analysis. In the case of retained austenite (γ r phase), the volume fraction fγ r was measured by an X-ray diffraction (XRD) technique

with Co–Kα. For the XRD measurement, a surface of the specimen was mechanically polished to a thickness of 3/4 and then chemically polished with oxalic acid. The value of fγ r was estimated by analyzing the integrated intensities for

the (200), (220) and (311) peaks of the γ r phase and those for the (200), (211) and (220) peaks of ferrite (α phase).

The new FE-EPMA technique13,14) reported in a previous study12) was used to measure two-dimensional (2D) map-ping of C, Mn and Si for the microstructure of the specimen. To prevent influence of contamination on the accuracy, the mapping of C was measured first. For the C measurement,

Table 1. Chemical composition of steel used in experiment and calculated Ae1 and Ae3 temperatures.

Chemical composition (mass%)Ae1 Ae3

C Si Mn P S Al

0.17 1.5 1.7 0.010 0.001 0.03 708°C (981 K) 857°C (1 130 K)

Fig. 1. Schematic diagram of heat treatment patterns.

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the accelerating voltage Ve and the probe current Ie of the primary electron beam were 7 kV and 5 nA, respectively. For the Mn and Si measurements after the C measurement, Ve and Ie were changed to 9 kV and 10 nA, respectively, to obtain sufficiently large values of signal-to-noise (S/N) ratio. A calibration technique was adopted to evaluate the C content of each phase using pure Fe and reference steels with 0.089–0.680 mass% C as standard specimens. The Mn and Si contents of each phase were evaluated by the ZAF correction method15) using pure Mn and Si, respectively, as standard specimens. Crystal orientation of each phase was measure by an electron back scattered diffraction (EBSD) technique with Ve = 20 kV and a step size of 50 nm.

3. Results

3.1. MicrostructureThe volume fraction fα of the α phase is plotted against the

quenching temperature Tq as solid circles in Fig. 2. Figures 2(a) and 2(b) shows the result of Ta = 800°C (1 073 K) and ta = 0.5 h (1.8 ks) and that of Ta = 740°C (1 013 K) and ta = 4 h (14.4 ks), respectively. As can be seen in Fig. 2(a), fα takes a value of 0.33 at Tq = 800°C (1 073 K) and gradu-ally increases from 0.33 to 0.69 with decreasing quenching temperature from Tq = 800°C (1 073 K) to Tq = 400°C (673 K). The microstructure consists of the α phase and martens-ite (M phase) for all the cases in Fig. 2. Here, the M phase corresponds to the γ phase before water quenching. Thus, the change in fα in Fig. 2 means that the γ phase is partially transformed into the α phase during the slow cooling from 800°C (1 073 K) to 400°C (673 K). Hereafter, the transfor-mation of the γ phase into the α phase is denoted by the α transformation. In contrast, the α phase formed during the intercritical annealing is called the αa phase, and that pro-duced during the slow cooling is designated the αc phase. According to the result in Fig. 2(a), the volume fraction fα c of the αc phase is 0 and 0.36 for Tq = 800 and 400°C (1 073 and 673 K), respectively. In contrast, as shown in Fig. 2(b), fα is rather insensitive to Tq and takes the mean value of 0.68 at Tq = 400–740°C (673–1 013 K). Consequently, for the specimens in Fig. 2(b), the α transformation does not occur during the slow cooling, and thus the αc phase does not form. This is because that the growth of the α phase is extremely

sluggish due to large concentration of Mn into the γ phase during intercritical annealing at Ta = 740°C (1 073 K) for ta = 4 h (14.4 ks) as reported in a previous study.12) As a result, fα c = 0 for Tq = 400–740°C (673–1 013 K).

Since the α transformation hardy occurs during the water quenching with an extremely fast cooling rate, the value of fα for Tq = Ta corresponds to the volume fraction fαa of the αa phase. In contrast, the increase in fα for Tq < Ta correlates with the volume fraction fα c of the αc phase. As a consequence, fαa = 0.33 and 0.68 for the results in Figs. 2(a) and 2(b), respectively. Although fα c increases from 0 to 0.36 with decreasing quenching temperature from Tq = 800°C (1 073 K) to Tq = 400°C (673 K) in Fig. 2(a), it is close to 0 at Tq = 400–740°C (673–1 013 K) in Fig. 2(b). As a result, the α phase consists of only the αa phase for the specimens with Ta = 740°C (1 013 K) and ta = 4 h (14.4 ks), though it is composed of the αa and αc phases for those with Ta = 800°C (1 073 K) and ta = 0.5 h (1.8 ks).

3.2. Tensile PropertiesThe tensile test was conducted on the specimens water-

quenched from Tq = 400°C (573 K) after intercritical annealing at Ta = 800°C (1 073 K) for ta = 0.5 h (1.8 ks) or at Ta = 740°C (1 013 K) for ta = 4 h (14.4 ks). Hereafter, the specimens intercritically annealed under the former and latter conditions are called Specimens A and B, respectively. The nominal stress s versus the nominal strain e (s–e) curves of Specimens A and B are shown as black and gray curves, respectively, in Fig. 3. As can be seen, both the black and gray curves show continuous yielding, and hence give no clear yield points. Therefore, the 0.2% proof stress s0.2% is defined as the yield stress sy. Furthermore, the maximum value of s was used as the ultimate tensile strength su, and the total strain et was calculated from the measured length of the broken pieces divided by the gauge length. In contrast, the value of e corresponding to s = su was considered as the uniform strain eu, and the local strain el was calculated from the following relationship: el = et − eu. According to the s–e curves in Fig. 3, sy = 460 MPa and su = 969 MPa for Specimen A, and sy = 488 MPa and su = 975 MPa for Specimen B. Thus, sy and su are not so dissimilar to each other between Specimens A and B. In contrast, et = 0.217, eu = 0.139 and el = 0.078 for Specimen A, and et = 0.158,

Fig. 2. The volume fraction fα of the α phase versus the quenching temperature Tq under the intercritical annealing conditions of (a) Ta = 800°C (1 073 K) and ta = 0.5 h (1.8 ks) and (b) Ta = 740°C (1 013 K) and ta = 4 h (14.4 ks).



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eu = 0.105 and el = 0.053 for Specimen B. The microstruc-ture and mechanical properties are summarized in Tables 2 and 3, respectively. Both eu and el of Specimen A are higher than those of Specimen B. Hence, Specimen A is more duc-tile than Specimen B.

The dependencies of su and et on fα are shown in Figs. 4(a)

and 4(b), respectively. In Fig. 4, solid circles indicate the results obtained under the conditions of Ta = 800°C (1 073 K) and ta = 0.5 h (1.8 ks), and open squares represent those taken under the conditions of Ta = 740°C (1 013 K) and ta = 4 h (14.4 ks). Furthermore, the horizontal axis on the upper side shows the volume fraction fM of the M phase. Since the α + γ two-phase microstructure is realized before water quenching for all the heat treatments as previously men-tioned, the summation of fα and fγ is equal to unity for all the specimens. In contrast, the γ phase was not necessarily completely transformed into the M phase under the present quenching conditions. For Specimens A and B, the volume fraction fγ r of the retained austenite (γ r phase) is 0.02 as listed in Table 2. Since fγ r = 0.02 is negligible, however, we consider that fM is equivalent to fγ and hence the summa-tion of fα and fM is equal to unity. In Fig. 4(a), su = 1 416, 1 227, 1 046, 984 and 969 MPa at fα = 0.33, 0.35, 0.60, 0.65 and 0.69, respectively, for the solid circles, and su is close to 1 000 MPa at fα = 0.68 for the open squares. Thus, for the solid circles, su monotonically decreases with increas-ing value of fα. Furthermore, the values of su for the open squares are close to that of su for the solid circle at fα = 0.69. In contrast, in Fig. 4(b), et = 0.11, 0.13, 0.18, 0.20 and 0.22 at fα = 0.33, 0.35, 0.60, 0.65 and 0.69, respectively, for the solid circles. Hence, for the solid circles, et monotonically increases with increasing value of fα. This is because strength is greater for the M phase than for the α phase, but ductility is lower for the M phase than for the α phase. Unlike Fig. 4(a), however, the values of et around 0.15 for the open squares are much smaller than that of et = 0.22 for the solid circle at fα = 0.69 in Fig. 4(b). It is reported that grain size or distributions of the α and M phases affect tensile proper-ties of DP steel.16,17) Grain sizes of the α and M phases in Specimen B are coarser than those in Specimen A. However, the differences of grain sizes or distributions of the α and M phases in Specimens A and B are not so large compared with previous study, the effects on the mechanical properties are considered to be small. On the other hand, according to the results in Fig. 2, the volume fraction of the αc phase is fα c = 0.36 and 0 for the solid circle and the open square,

respectively, with fα = 0.68–0.69 at Tq = 400°C (673 K) in Fig. 4. Thus, we may consider that the effect of the α phase

Fig. 3. The nominal stress s versus the nominal strain e at room temperature for Specimens A and B shown as black and gray curves, respectively.

Table 2. Volume fraction and mean grain size of each phase for Specimens A and B.

SpecimenVolume fraction Mean grain size

αa phase, fαa

αc phase, fα c

M phase, fM

γr phase, fγ r

α phase, dα (μm)

M phase, dM (μm)

A 0.33 0.36 0.29 0.02 4.6 3.1

B 0.68 0 0.30 0.02 5.0 4.2

Fig. 4. Dependencies of tensile properties on the volume fraction fα of the α phase: (a) ultimate tensile strength su, and (b) total strain et.

Table 3. Tensile properties for Specimens A and B.

Specimen sy/MPa su/MPa et eu el

A 460 969 0.217 0.139 0.078

B 488 975 0.158 0.105 0.053



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on su is close to each other between the αa and αc phases but the contribution to improvement of et is greater for the αc phase than for the αa phase. The difference between the dependencies of su and et on fα is attributed to such mechani-cal properties of the αa and αc phases.

The true strain ε and the true stress σ are calculated from the nominal strain e and the nominal stress s as follows.18)

� � �ln( )e 1 ................................ (1)

� � �s e( )1 ................................. (2)

Using Eqs. (1) and (2), the s–e curves are converted into the true stress σ versus the true strain ε (σ–ε) curves. σ is expressed as a power function of ε as follows.19)

� � �� 0 k n ................................ (3)

Here, σ0 is the material constant, k is the proportionality coefficient, and n is the strain hardening exponent. Since ε and n are dimensionless, k has the same dimension as σ and σ0. The limit of uniform deformation corresponds to the start of necking, which is associated with localized thinning and deformation of the specimen. Necking occurs according to the plastic instability criterion defined by the equation20)

d d� � �/ � ................................ (4)

The differential dσ/dε is the slope of the σ–ε curve, and usu-ally called the strain hardening rate. For the dependence of σ on ε in Eq. (3), we obtain the strain hardening rate as follows.

d d n� � �/ � �kn 1 ............................. (5)

Changes in the value of dσ/dε with ε of Specimens A and B are shown as black and gray curves, respectively, in Fig. 5. In the range of ε > 0.02, the strain hardening rate is typi-cally greater for the black curve than for the gray curve. This means that strain hardening is more remarkable for Speci-men A than for Specimen B. The black and gray curves for σ of Specimens A and B are also shown in Fig. 5. As can be seen, the black curves for dσ/dε and σ are intersected each other at ε = 0.13, and the gray curves for dσ/dε and σ are intersected each other at ε = 0.10. The intersection of dσ/dε and σ indicates start of necking due to the plastic instability as described above. Therefore, we can conclude that high strain hardening rate of Specimen A contributes high uniform strain eu of it.

3.3. Nano-Indentation HardnessAs mentioned in Section 2, the nano-indentation hard-

ness HNI was measured at 900 points on the cross sectional surface of the specimen. The sizes of projection area in nano-indentation hardness test were approximately from 0.045 to 0.060 mm2 and from 0.025 to 0.040 mm2 for the α and M phases, respectively. They were small enough for the grain size of each phase to measure the hardness accurately. According to the measurement, the values of HNI were dis-tributed between 3 and 14 GPa. Thus, Hi was calculated by the following equation.

H H i H ii � � � �0 0 29� ( ), ..................... (6)

where ΔH is defined as follows.

�HH H

N�

�max min ........................... (7)

Here, Hmin = 3 GPa, Hmax = 14 GPa, and N = 30. From Eqs. (6) and (7), the following relationships hold.

H H0 = min .................................. (8)

H H H29 � �max � ............................ (9)

Furthermore, the frequency Qi is defined as follows.

QM

Mi

i≡ ................................. (10)

Here, M = 900, and Mi is the number of the data points for Hi ≤ HNI < Hi+1. Typical results are shown in Fig. 6. Figure 6(a) indicates the results for Ta = 800°C (1 073 K) and ta = 0.5 h (1.8 ks), and Fig. 6(b) represents those for Ta = 740°C (1 013 K) and ta = 4 h (14.4 ks). Solid and dashed curves show the results for Tq = 400 and 800°C (673 and 1 073 K), respectively, in Fig. 6(a), and the former and latter ones indicate those for Tq = 400 and 740°C (673 and 1 013 K), respectively, in Fig. 6(b). For the dashed curve in Fig. 6(a), Qi takes peak values at HNI = 5.2 and 8.5 GPa. The peaks of Qi at HNI = 5.2 and 8.5 GPa correspond to the αa and M phases, respectively. In contrast, for the solid curve in Fig. 6(a), a peak of Qi appears only at HNI = 4.8 GPa. The disappearance of the Qi peak at HNI = 8.5 GPa is attributed to the transformation of the γ (M) phase into the αc phase during the slow cooling and the broadening of the hardness distribution of the M phase. The broadening can be caused by the large dispersion of the C and Mn contents in the γ phase before water quenching.12) The mean value of HNI of the M phase for Tq = 400°C is 10.3 GPa, and it is higher than that for Tq = 800°C. This increase in the mean hardness of the M phase is attributed to the C concentration in the γ phase before water quenching by the α transformation. In addition, the results in Fig. 6(a) indicate that the αc phase is slightly softer than the αa phase. For the dashed curve in Fig. 6(b), peaks of Qi exist at HNI = 5.2 and 10.7 GPa. Also in this case, the Qi peaks at HNI = 5.2 and 10.7 GPa correlate to the αa and M phases, respectively. Unlike Fig. 6(a), however, the two Qi peaks appear also for the solid curve in Fig. 6(b). This means that the transformation of the γ (M) phase to the αc phase does not occur during the slow cooling. Mean value of HNI of the M phase for Tq = 400°C is approximately equal between Figs. 6(a) and 6(b).Fig. 5. Dependencies of σ and dσ/dε on ε for Specimens A and B

shown as black and gray curves, respectively.



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3.4. Microstructural Change during Tensile Deforma-tion

Typical SEM images of Specimen A for e = 0.04, 0.10 and 0.22 (fractured) are shown in Figs. 7(a1), 7(a2) and 7(a3), respectively. Figures 7(b1), 7(b2) and 7(b3) indicates image quality (IQ) mapping constructed from the EBSD data for Figs. 7(a1), 7(a2) and 7(a3), respectively, and Figs. 7(c1), 7(c2) and 7(c3) represents the corresponding inverse pole figure (IPF) mapping for Figs. 7(a1), 7(a2) and 7(a3), respectively. The IQ value shows the sharpness of the Kikuchi pattern at the measurement point. Lattice distor-tion due to crystalline defects, such as dislocation, grain boundary, etc., misshapes the Kikuchi pattern, and thus provides lower IQ values. A large number of dislocations are included in the M phase but not in the α phase. Hence, bright and dark crystal grains in Fig. 7(b1) correspond to the α and M phases, respectively. As the strain e increases, the IQ value gradually decreases, and thus the α phase becomes darker as shown in Figs. 7(b2) and 7(b3). The increase in the strain e also yields inclination of each grain of the α phase. Such inclination is recognized in Figs. 7(c1)–7(c3).

Similar SEM images of Specimen B for e = 0.04, 0.10 and 0.16 (fractured) are shown in Figs. 8(a1), 8(a2) and 8(a3), respectively. Figures 8(b1), 8(b2) and 8(b3) indicates IQ mapping for Figs. 8(a1), 8(a2) and 8(a3), respectively, and Figs. 8(c1), 8(c2) and 8(c3) represents IPF mapping for Figs. 8(a1), 8(a2) and 8(a3), respectively. The relationships in Fig. 7 are also recognized in Fig. 8. However, the influ-ence of e on the IQ and IPF mapping is less remarkable in Fig. 8 than in Fig. 7.

4. Discussion

4.1. Strain HardeningCalculating the common logarithm of Eq. (5), we obtain

the following equation.21–24)

log d d log log10 10 101( ) ( ) ( )/ ( )� � �� kn n ...... (11)

From Eq. (11), the value of log10(dσ/dε) was calculated as a function of log10(ε) from σ–ε curves. The calculations of Specimens A and B are shown as black and gray curves, respectively, in Fig. 9. As in previous reports,10,23) the black curve is divided into Stages I, II and III. Here, Stages I, II

and III correspond to the ranges of ε = 0.001–0.003, ε = 0.003–0.02 and ε > 0.02, respectively. The black curve is convex upward at Stages I and III but convex downward at Stage II. As described in Section 3.2, the strain hardening rate is greater for Specimen A than for Specimen B at ε > 0.02. This relationship corresponds to Stage III in Fig. 9. A similar relationship holds also at Stage I. At Stage II, how-ever, the strain hardening rate becomes smaller for Speci-men A than for Specimen B. As shown in Figs. 3 and 5, at Stage III, the strain hardening rate is greater for Specimen A than for Specimen B, but the flow stress σ of Specimen A is slightly lower than that of Specimen B. As a result, eu is greater for Specimen A than for Specimen B.

Kernel average misorientation (KAM) mapping for Figs. 7(a1)–7(a3) and 8(a1)–8(a3) is shown in Figs. 10(a1)–10(a3) and 10(b1)–10(b3), respectively. Here, KAM mapping of only the α phase is shown by rejecting low IQ regions witch correspond to the M phase, and the mean KAM value is shown in each mapping. The KAM value is defined as the average misorientation between the measurement point and the neighboring points.25) The mean KAM value increases with increasing strain e. Nakada et al. clarified that the KAM value closely related to strain in the microstructure by a combinational technique of EBSD and Digital Image Correlation (DIC) methods.26) Thus, the Fig. 10 indicates that the strain in the α phase increases with increasing the nominal strain e. Even for the same value of e = 0.10, regions with larger KAM values in the α grains is slightly larger for Fig. 10(a2) than for Fig. 10(b2). It is also reported that the misorientation angle was related to the dislocation density ρgnd of geometrically necessary dislocation (GND) by Calcagnotto et al.27) This means that ρgnd is lightly greater for α grains in Fig. 10(a2) than for those in Fig. 10(b2). As previously mentioned, the α phase consists of the αa and αc phases for Specimen A, but it is composed of only the αa phase for Specimen B. Hence, the increase in ρgnd seems to be caused by formation of the αc phase. At Stage III in Fig. 5, the strain hardening rate is greater for Specimen A than for Specimen B. This relationship can be attributed to larger values of ρgnd in the α phase of Specimen A.

Evolution of dislocation substructure due to plastic deformation in a DP steel was experimentally observed by Korzekwa et al.23) In their experiment, a Fe–0.063C–

Fig. 6. The frequency Qi versus the nano-indentation hardness HNI for Tq = 400°C (673 K) and Tq = Ta shown as solid and dashed curves, respectively: (a) Ta = 800°C (1 073 K) and ta = 0.5 h (1.8 ks), and (b) Ta = 740°C (1 013 K) and ta = 4 h (14.4 ks).



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1.29Mn–0.24Si steel was annealed at Ta = 810°C (1 083 K) for ta = 10 min (600 s), and then cooled at an average cooling rate of 60°C/s. Tensile tests were carried out on

the steel at room temperature, and dislocation substructure was observed for various degrees of plastic deformation by transmission electron microscopy (TEM). According to their

Fig. 7. (ai) SEM image, (bi) IQ mapping and (ci) IPF mapping for Specimen A, where i = 1, 2 and 3 correspond to εn = 0.04, 0.10 and 0.22 (fractured), respectively.

Fig. 8. (ai) SEM image, (bi) IQ mapping and (ci) IPF mapping for Specimen B, where i = 1, 2 and 3 correspond to e = 0.04, 0.10 and 0.16 (fractured), respectively.



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observation,23) the dislocation density ��d in the α phase is

greater at the region near by the M phase than at that far from the M phase. Thus, ��

d is inhomogeneous and varies depending on the location.

For Specimen A with e = 0.10, SEM image and the 2D mapping of C, Mn and Si measured by FE-EPMA are shown in Figs. 11(a), 11(b) ,11(c) and 11(d), respectively, and the corresponding IQ and KAM mappings are indicated in Figs. 11(e) and 11(f), respectively. The SEM observation and EBSD analysis were performed after FE-EPMA analysis. As the result, the microstructure in SEM observation and EBSD analysis was slightly changed for that in FE-EPMA analysis by polishing after FE-EPMA analysis. On the basis of the information in Figs. 11(b)–11(d), each phase is iden-tified as shown in Fig. 11(g). Here, Crystal grains of the αa and αc phases are identified as follows. The C and Mn contents are larger in the M (γ ) phase than in the αa phase,

Fig. 10. KAM mapping and mean KAM value for (ai) Specimen A and (bi) Specimen B, where (ai) i = 1, 2 and 3 cor-respond to e = 0.04, 0.10 and 0.22 (fractured), respectively, and (bi) i = 1, 2 and 3 correspond to e = 0.04, 0.10 and 0.16 (fractured), respectively.

Fig. 11. (a) SEM image, 2D mapping of (b) C, (c) Mn and (d) Si, (e) IQ mapping and (f) KAM mapping for Specimen A with e = 0.10. Each phase is depicted in (g).

Fig. 9. Dependencies of log10(dσ/dε) on log10(ε) for Specimens A and B shown as black and gray curves, respectively.



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but the Si content is smaller in the M (γ ) phase than in the αa phase. This means that the αa phase forms owing to the reverse transformation in the partitioning local equilibrium (PLE) mode.12) In contrast, the C content is larger in the M (γ ) phase than in the αa and αc phases, but the Mn and Si in the αc phase are close to those, respectively, in the M (γ ) phase. This is because the αc phase is produced by the α transformation in the negligible partitioning local equilib-rium (NPLE) mode.12) The regions with high KAM values in Fig. 11(f) are represented as green lines in Fig. 11(g). As can be seen in Fig. 11(g), the green lines are rather prefer-entially distributed at the boundary between the αa and αc phases. The boundary is not gain boundary but boundary between αc and αa phases in an α grain. As can be seen in Figs. 11(b)–11(d), the Mn content is greater for the αc phase than for the αa phase, but the Si content is smaller for the αc phase than for the αa phase. Thus, there exists a certain difference in the lattice parameter between the αa and αc phases, and misfit strain can be stored at the boundary of those phases during straining. This may be the reason why the green lines are preferentially located at the boundary between the αa and αc phases in Fig. 11(g).

4.2. Influence of Ferrite on Void Formation and Local Strain

Typical SEM images of the specimen fractured by the tensile test are shown in Fig. 12. Figures 12(ai) and 12(bi) indicates the results of Specimens A and B, respectively, where i = 1, 2 and 3 correspond to d = 0, 0.2 and 5 mm, respectively. Here, d is the distance from the fractured sur-face of the specimen. Furthermore, arrows represent voids. All the positions in Fig. 12 were in the necking region where continuous thickness change from the fractured surface occurred. As can be seen in Figs. 12(b1)–12(b3), many voids are observed in Specimen B. The size of the void is greater for d = 0 mm than for d = 0.2–5 mm. Although voids are recognized also for Specimen A in Figs. 12(a1)–

12(a3), the number of voids is smaller for Specimen A than for Specimen B.

The void number density ρvn is plotted against the dis-tance d as solid circles and open squares for Specimens A and B, respectively, in Fig. 13. Number of voids were counted in backscattered electron image followed by the observation in secondary electron image with etched speci-men to confirm the position of voids in the microstructure. Void number density ρvn was obtained by dividing the number of voids by the area of the observation with each specimen. At d = 0 mm, ρvn = 10 × 103 and 11 × 103 mm −2 for Specimens A and B, respectively. For Specimen A, ρvn decreases from 10 × 103 mm −2 to 5 × 103 mm −2 with increasing distance from d = 0 mm to d = 0.5 mm, and becomes almost constant at d = 0.5–5 mm. In contrast, for Specimen B, ρvn increases from 11 × 103 mm −2 to 14 × 103 mm −2 with increasing distance from d = 0 mm to d = 2 mm, and becomes mostly constant at d = 2–5 mm. Thus,

Fig. 12. SEM images of fractured pieces for (ai) Specimen A and (bi) Specimen B. Here, i = 1, 2 and 3 correspond to d = 0, 0.2 and 5 mm, respectively, where d is the distance from fractured surface of specimen. (Online version in color.)

Fig. 13. The void number density ρvn versus the distance d from fractured surface for Specimens A and B shown as solid circles and open squares, respectively.



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at d = 2–5 mm, ρvn is almost three times greater for Speci-men B than for Specimen A. Furthermore, for Specimen B, the size of voids is larger for d = 0 mm than for d = 2–5 mm, so the decrease in ρvn at d = 0 mm for Specimen B can be caused by coalescence of voids. The void formation and coalescence during plastic deformation will promote the fracture. The larger value of el for Specimen A than for Specimen B may be attributed to the lower value of ρvn for Specimen A than for Specimen B.

Other SEM images at d = 0 mm for Specimens A and B are shown in Figs. 14(a1) and 14(b1), respectively, where voids are indicated by arrows. Figures 14(a2), 14(a3) and 14(a4) indicates the 2D mapping of C, Mn and Si, respec-tively, measured by FE-EPMA for Figs. 14(a1), and 14(b2), 14(b3) and 14(b4) represents that of C, Mn and Si, respec-tively, for Fig. 14(b1). The distribution of the constituent phases is depicted in Figs. 15(a) and 15(b) for Figs. 14(a1)–14(a4) and Figs. 14(b1)–14(b4), respectively. As can be seen in Fig. 15, voids tend to form at the concave part of the M phase adjacent to the αa phase. Hereafter, the M/α inter-face at such a concave part is called the concave interface. Table 4 summarizes the number of the voids in an area of 35 × 35 μm2 observed in the M phase adjacent to various α phases. As shown in Table 4, most voids form adjacent to

the αa phase not only in Specimen B but also in Specimen A. Here, αa + αc means that a void is in contact with both the αa and αc phases at the M/α interface like the void in the upper right in Fig. 15(a). During plastic deformation, dislocations pile up at the interface between the α and M phases, and thus plastic strain concentrates in the interface. When the strain reaches to a certain critical value, it induces interface decohesion and then void formation.

It is reported by Azuma et al.28) that the misorientation ϕ in the α phase near the M/α interface is three times greater for the concave interface than for the flat interface. Such a larger value of ϕ results in a larger value of the disloca-tion density and thus promotes void formation. They also reported that softening of the M phase due to tempering increased the critical strain and decreased the topical strain at the concave part. This causes retardation of void forma-

Fig. 14. (a1 and b1) SEM images and (ai and bi) 2D mapping of C, Mn and Si for i = 2, 3, and 4, respectively: at d = 0 mm for (ai) Specimen A, and (bi) Specimen B.

Fig. 15. Constituent phases in Figs. 14(ai) and 14(bi) are shown in (a) and (b), respectively. Voids are indicated as arrows.

Table 4. Number of voids in area of 35 × 35 μm2 observed in the M phase adjacent to various α phases.

αa αa + αc αc

Specimen A 14 16 2

Specimen B 42 0 0



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tion and improves local strain. This means that void forma-tion can be suppressed if the concentration of topical strain at the M/α interface is relaxed by dynamic recovery in the α phase. It is reported by some researchers29–31) that addition of Si to the α phase lowers the stacking fault energy and inhibits the cross slip of dislocation. Thus, dynamic recov-ery of the α phase is suppressed by the addition of Si. As shown in Fig. 11, the Si content is greater for the αa phase than for the αc phase. The suppression of dynamic recovery in the αa phase near the M/α interface may finally promote void formation. In addition, a hardness difference between the M phase adjacent to the αa and that adjacent to the αc may also affect the void formation behavior. In a previous study,12) the Mn content in the γ phase during intercritical annealing tended to be larger near the grain boundary than at the center of grain. Based on this result, when the αc phase grows into the γ phase during the slow cooling, the Mn content in the γ phase adjacent to the αc phase is considered to be smaller than that adjacent to the αa phase. As a result, the M phase adjacent to the αc phase may be softer than that adjacent to the αa phase due to decreasing the content of Mn which is an element for improving hardenability. The decrease in hardness of the M phase can suppress void formation as reported in previous studies.26,32) These two factors can affect the suppression of void formation by for-mation of the αc phase. In summary, when the total volume fraction of the αa and αc phases is constant, increase in the volume fraction of the αc phase is effective for suppressing void formation and for achieving large local strain.

5. Conclusions

The effects of intercritically-annealed ferrite (αa phase) and transformed ferrite (αc phase) on the mechanical proper-ties of the Fe–0.17C–1.5Si–1.7Mn DP steel were experimen-tally examined using the tensile-test and nano-indentation hardness-test techniques. The new FE-EPMA technique was also used to characterize the microstructure of the steel. The main conclusions are summarized as follows.

(1) The ferrite (α) transformation progresses during the slow cooling at a cooling rate of 10°C/s after intercritical annealing at 800°C (1 073 K) for 0.5 h (1.8 ks). Since Mn concentrates in austenite (γ phase) owing to intercritical annealing at 740°C (1 013 K) for 4 h (14.4 ks), however, the α transformation hardly occurs during the slow cooling after the intercritical annealing.

(2) The tensile strength su is rather determined only by the total volume fraction fα of the αa and αc phases. On the other hand, for a constant value of fα, both the uniform strain eu and the local strain el increase with increasing volume fraction fα c of the αc phase.

(3) Such increase in eu is attributed to larger values of

the strain hardening rate dσ/dε in the large strain region. Misfit strain stored at the boundary between the αa and αc phases causes to the larger values of dσ/dε.

(4) Suppression of void formation deduces the increase in el. Since the Si content is smaller for the αc phase than for the αa phase, dynamic recovery in the α phase adjacent to martensite (M phase) may occur more preferentially in the αc phase than in the αa phase. This dynamic recovery can relax topical strain at the M/α interface and hence suppresses void formation. In addition, the decrease in the hardness of the M phase adjacent to the αc phase was also considered to contribute the suppression of void formation.

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