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Microstructure and mechanical properties of dual phase strip steelin the overaging process of continuous annealing

Chang-sheng Li a,n, Zhen-xing Li a, Yi-ming Cen a,b, Biao Ma a, Gang Huo a,c

a State Key Laboratory of Rolling and Automation, Northeastern University, Shenyang 110819, Chinab Ningbo Branch of China Academy of Ordnance Science, Ningbo 315103, Chinac Pohang Cold Strip Rolling Co., Ltd., Bensteel, Benxi 117000, China

a r t i c l e i n f o

Article history:Received 5 November 2014

Received in revised form

23 December 2014

Accepted 24 December 2014Available online 5 January 2015

Keywords:

Dual phase steel

Cold rolled strip steel

Overaging temperature

Microstructure

Mechanical properties

a b s t r a c t

Microstructure and properties of DP590 steel in overaging process of continuous annealing were investigatedin the laboratory. Effects of the overaging temperature (OT) in conventional process (Process A) and reheated

overaging temperature (ROT) of the proposed process (Process B) on microstructure and mechanical properties

were analyzed. The results showed that, in process A, with the increase of the OT from 280 1C to 400 1C, the

amount of granular retained austenite increases, the tensile strength decreases, the yield strength and

elongation both increase. In process B, with the increase of the ROT from 310 1C to 400 1C, the tensile strength

decreases, whereas the yield strength changes little. Compared with the conventional process, the tensile

strength of the experimental steel increases obviously and the ratio of yield-to-tensile strength decreases in

process B. The variation of instantaneous work hardening index (nn) with tensile strain can be divided into

three stages. The increase in the OT or ROT leads to the continuous increase ofnn in the third stage.

&2015 Elsevier B.V. All rights reserved.

1. Introduction

As a type of advanced high strength steel, dual phase (DP)

steels are increasingly used in automobile industry. The micro-

structure of DP steel is usually composed of soft ferrite matrix,

hard martensite and small amounts of retained austenite [1,2].

This steel has many advantages, such as continuous yielding, high

initial work hardening rate, high bake hardening ability and

relatively high formability [3,4]. These mechanical characteristics

of DP steels essentially originate from the interactions of soft

ferrite matrix and hard martensite[5].

Generally, DP steels are produced by intercritical annealing of

cold rolled low carbon steels, or by hot rolling in the austenite

region followed by step cooling (i.e., cooling to the region and

then quenching)[68]. The nal mechanical behavior of DP steels

depends on the microstructure characteristics, such as the grainsize, volume fraction and morphology of martensite, which are

strongly inuenced by the heat treatment process [911]. There-

fore, it is of great importance to determine the relationship

between annealing parameters and microstructure of DP steels.

Li et al. [12] studied the effect of heating rate on ferrite recrys-

tallization and austenite formation in DP steels, it was found that

with increasing heating rate, the ferrite recrystallization would be

retarded. As a result, before ferrite recrystallization, some auste-

nite starts to form at the deformed ferrite grain boundary, whichcan rene the nal grain size. Benoit et al.[13]studied the effect of

intercritical annealing temperature on banded structure in DP

steels, they found that the banded structure was more likely to be

broken at higher annealing temperature. Movahed et al. [14]and

Bag et al. [15] reported that with the increase of intercritical

annealing temperature, the martensite volume fraction increases,

which in turn leads to decrease of carbon content of martensite.

Therefore, with the increase of annealing temperature, tensile

strength increases rst and then decreases. Abouei et al. [16]and

Kelestemur and Yildiz[17]found that with the increase of holding

time at intercritical annealing temperature, the martensite volume

fraction increases, the wear properties are improved, but the

corrosion resistance decline. Chang [18] and Calcagnotto et al.

[19] have studied the effect of tempering and aging on themechanical properties of DP steels that were directly quenched

to room temperature from the region, they found that

tempering will delay the necking and enhance the ductility

obviously. It should be noted that, in industry production, DP

steels are usually subjected to intercritical annealing by contin-

uous annealing furnace, in which the DP steels are seldom directly

quenched to room temperature, and online overaging is an

essential part. However, few works focused on the effect of

overaging on microstructure and mechanical behavior of DP steels.

In this paper, the intercritical annealing process of DP590 steels

was simulated by the CAS-300II continuous annealing simulator in

Contents lists available at ScienceDirect

jo ur nal ho mep ag e: www.elsevier.com/locate/msea

Materials Science & Engineering A

http://dx.doi.org/10.1016/j.msea.2014.12.109

0921-5093/&2015 Elsevier B.V. All rights reserved.

n Corresponding author. Tel.: 86 24 83687749; fax: 86 24 23906472.

E-mail address:[email protected](C.-s. Li).

Materials Science& Engineering A 627 (2015) 281289
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the laboratory. The effect of overaging temperature on microstruc-

ture and mechanical properties was studied in the conventional

overaging and a proposed overaging process of continuous anneal-

ing. The corresponding microstructure and mechanical properties

of the experimental steel were analyzed. This research can provide

a theory basis for optimizing the intercritical annealing process of

cold rolled DP steel.

2. Experimental material and methods

The experimental material was cold rolled DP590 strip steel

with 1.2 mm thickness from Bensteel in China. The chemical

composition (wt%) of the experimental steel was 0.08 C, 0.479 Si,

1.81 Mn, 0.014 P, 0.004 S, 0.162 Cr, 0.004 N, 0.04 Al and bal. Fe. Hot

rolling process parameters were described as following. The

continuous casing slab with thickness of 220 mm was heated to

11901200 1C and rolled by 2300 mm hot strip rolling mill. The

slab was rolled from the thickness of 220 mm to 2025 mm by R1

and R2 rough rolling mill. Subsequently the slab was rolled from

2025 mm to 46 mm by F1F7nising rolling mill. Rough rolling

temperature was 11201170 1C. Entrance temperature ofnishing

rolling was 9501030 1C. Delivery temperature ofnishing rollingwas 850870 1C. Coiling temperature was controlled 600620 1C.

After acid pickling, the hot rolled strip was cold rolled to the

thickness of 1.2 mm by ve stands 6-high tandem cold rolling mill.

Fig. 1shows the schematic illustration of continuous annealing

process used in this work. Apart from the conventional annealing

process (Process A), a new annealing process (Process B) is also

depicted in Fig. 1. In process A, the tested steels were directly

cooled to various temperatures for overaging. However, in process

B, the tested steels were cooled to a lower temperature rstly and

held for a short time, and then they were rapidly reheated to a

higher temperature, i.e., reheated overaging temperature start

point. From this point the overaging temperature decreases slowly

with time in OAS1 stage of process B. HS, SS, SCS, RCS,

OAS

and FCS

correspondingly denote the stage of heating,

soaking, slow cooling, rapid cooling, overaging and nal cooling,

while LHS, RHS and OAS1 represent the low temperature

holding stage after rapid cooling, the reheating stage and the

overaging stage, respectively. The detailed annealing parameters

are given inTable 1.

In order to investigate the effect of overaging temperature on

the microstructure and mechanical properties, the experimental

steels were rapidly cooled to various overaging temperatures of

280 1C, 320 1C, 350 1C and 400 1C in process A. However, in process

B the experimental steels were rapidly cooled to 250 1C and held

for 50 s, and then they were reheated at a rate of 50 1C/s to various

overaging temperatures of 310 1C, 330 1C, 350 1C and 400 1C. For

the convenience of discussion, the overaging temperature in

process A is represented by the OT, the reheated overaging

temperature start point in process B is represented by the ROT

in the following text.

The continuous annealing experiment was carried out by CAS-

300II continuous annealing simulator in the laboratory. The

microstructure of experimental steels was observed by scanningelectron microscopy (QUANTA 600) and transmission electron

microscopy (FEI Tecnai G2 F20). The volume fraction of martensite

phase was measured and analyzed by Image-pro Plus software.

TEM foils were prepared by electrolytic polishing with a solution

composed of 10 vol% perchloric acid and 90 vol% ethanol. The

tensile tests were performed with the aid of Inston 4206-006

tensile testing machine at room temperature. The tensile speci-

mens with a gauge length of 50 mm were prepared by wire

electrical discharge machine. In order to observe the microstruc-

ture adjacent to the fracture surface, fractured tensile specimens

were sectioned through-thickness along the mid-width plane in

the longitudinal direction.

3. Results and discussion

3.1. Microstructure of experimental steels in conventional process

Fig. 2 shows the microstructure of experimental steels

annealed by process A. The dark color regions are ferrite and the

bright color regions are martensite or retained austenite. It can be

seen from Fig. 2 that, in the specimen with the OT of 280 1C,

martensite are less tempered. As the OT exceeds 320 1C, substan-

tial tempered martensite was observed, and the boundary of some

tempered martensite becomes indistinct. At the OT of 400 1C,

martensite volume fraction is quite small and the amount of

granular retained austenite is larger. In addition, in the specimen

with the OT of 400 1C, lots of granular retained austenite, havingsize of less than 0.8 m, is observed within the ferrite grain.

However, at the lower OT, most of the granular retained austenite

or martensite islands, having size of less than 1 m, are distributed

along the ferrite grain boundary.

According toFig. 2, it can be concluded that with the increase

of the OT from 280 1C to 400 1C, the amount of granular retained

austenite increases and martensite volume fraction decreases. It is

because that when the experimental steels were rapidly cooled

from the region to a certain temperature for overaging, there

must exist diffusion of carbon from ferrite to the untransformed

Fig. 1. Schematic illustration of conventional overaging (a) and a proposed overaging process (b) of cold rolled strip continuous annealing.

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austenite. The carbon enrichment will lead to the stabilization of

austenite and the suppression of Ms temperature [20,21]. When

the OT increases within the range ofMfMs, the volume fraction of

untransformed austenite will increase. The stabilization of theseuntransformed austenite will lead to the increase in the content of

retained austenite in the nal microstructure. If the OT is larger

thanMs, no austenite transforms into martensite during overaging.

In addition, some metastable austenite may transform into ferrite.

As some austenite nucleates on carbides during the heating or

holding at the intercritical temperature, they will inherit alloy

elements from these carbides, i.e., this kind of austenite is so stable

that most of them will be retained in the nal microstructure, as

shown inFig. 2d[2,22].

The microstructure of experimental steels annealed by process A

was further observed by TEM as shown in Fig. 3. In the bright eld

images (i.e.,Fig. 3ac), the bright color regions are ferrite and the dark

color regions are martensite or retained austenite. Fig. 3d is the dark

eld image obtained from the same area asFig. 3c. It can be seen from

Fig. 3a that there are lots of geometrically necessary dislocations

(GNDs) in the ferrite adjacent to martensite at the OT of 280 1C. As the

OT increases, the dislocation density decreases. Furthermore, at the OT

of 400 1C, lots of granular retained austenite can be found in the ferritematrix. The austenite-to-martensite transformation inevitability

involves volume expansion, which leads to the plastic deformation

of ferrite grains[23]. Therefore, there exists residual stress and GNDs

in the ferrite adjacent to martensite [24]. With the increase of the OT,

the amount of granular retained austenite increases and martensite

volume fraction decreases, which in turn results in the decreasing of

GNDs density. In addition, at the higher OT, the recovery of ferrite also

leads to the decrease of GNDs density.

3.2. Microstructure of experimental steels in the proposed process

Fig. 4shows the microstructure of the tested steels which were

annealed by process B. It can be clearly seen that the microstruc-

ture is composed of ferrite matrix, blocky martensite and granular

Table 1

Parameters of the experimental steel during continuous annealing process A and B.

Sections Process A Process B

HS (A-B) 30 1C/s, 20-800 1C 30 1C/s, 20-800 1C

SS (B-C) 800 1C, 110 s 800 1C, 110 s

SCS (C-D) 2 1C/s, 800-680 1 C 2 1C/s, 800-680 1C

RCS (D-E) 30 1C/s; 680-280, 320, 350, 400 1C 30 1C/s, 680-250 1C

OAS (E-F) 280, 320, 350, 400 1C; 420 s

LHS (E-E1) 2501

C, 50 sRHS(E1-E2) 50 1C/s; 250 1C-310, 330, 350, 400 1C

OAS1(E2-F) 310, 330, 350, 400 1C-280 1C; 400 s

FCS (F-G) 280 1C-20 1C, 80 s 280 1C-20 1C, 80 s

Fig. 2. SEM morphology of the tested steels with the overaging temperature of (a) 280 1C, (b) 320 1 C, (c) 350 1C and (d) 400 1C in process A. TM is tempered martensite,

UM is untempered martensite, RA/UM is retained austenite or untempered martensite, RA is retained austenite.

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retained austenite. It can be found that when the ROT exceeds

330 1C, some of the martensite particles are tempered obviously.

At the ROT of 400 1C, some of the martensite particles are almost

decomposed completely, which leads to the decrease of martensite

volume fraction. Furthermore, approximately 5 m martensitegrain size can be frequently observed in the specimens with the

ROT of lower than 350 1C. However, the martensite is less than

4 m at the ROT of 400 1C, which is attributed to the decomposi-

tion of martensite. According toFig. 4, it can also be observed that

with the increase in ROT, the amount of granular retained

austenite increases. The granular retained austenite is mostly

distributed along the grain boundary, yet some retained austenite

is also observed within the ferrite grain, as shown inFig. 4b. The

TEM morphology and diffraction patterns of granular retained

austenite are illustrated in Fig. 5. It can be seen that the size of

granular retained austenite is about 0.45 m.

In process B, the experimental steels were rapidly cooled to

250 1C at rst, which is certainly below the Ms temperature. As a

result, some austenite transformed into martensite. However, forlow carbon and low alloy steels, the martensite transformation

belongs to the athermal transformation, i.e., the volume fraction of

transformed martensite is a function of undercooling degree[25].

Therefore, there will exist a certain amount of untransformed

austenite after isothermal holding at 250 1C for 50 s. As the

experimental steels were reheated to different temperatures for

overaging, the carbon would diffuse from the martensite or

supersaturated ferrite to the surrounding austenite. When the

ROT increases from 310 1C to 400 1C, the diffusivity of carbon

increases and larger amounts of untransformed austenite will be

stabilized. In addition, at higher overaging temperature, due to the

diffusion of carbon, the ferriteaustenite interface will move

towards the austenite, i.e., some metastable austenite may trans-

form into ferrite, which further increases the carbon content of

untransformed austenite, and the size of untransformed austenite

is also decreased. However, due to the lower carbon diffusivity, the

stabilization of untransformed austenite is suppressed at the lower

ROT. Subsequently, the austenite will transform into martensite

during the nal cooling. Therefore, the content of granularretained austenite increases with the increasing of the ROT.

Fig. 6shows the TEM morphology of martensite in the speci-

mens annealed by process B. It is evident that the martensite

belongs to lath martensite[26]. It should be noted that cementite

was found in the martensite. This is because that some martensite

was formed during the isothermal holding at 250 1C. The marten-

site was tempered, and cementite began to precipitate when the

experimental steels were reheated to a higher temperature for

overaging. The tempering of martensite and precipitation of

cementite are benecial to decrease the strength mismatch

between martensite and the ferrite matrix. As a result, during

deformation, the martensite starts to deform simultaneously with

ferrite at a lower strain, and the strain distribution is more

uniform.

3.3. Mechanical properties

3.3.1. Yield and tensile strength

Yield strength, tensile strength and the ratio of yield-to-tensile

strength (YS/TS) for the experimental steels were shown in Fig. 7.

In this gure, the horizontal ordinate means the OT for process A

and the ROT for process B. When the OT, in process A, increases

from 280 1C to 400 1C, tensile strength of experimental steels

decreases from 661.8 MPa to 584.6 MPa and the yield strength

increases from 313.6 MPa to 415.2 MPa. In process B, as the ROT

increases from 310 1C to 400 1C, the tensile strength decreases

from 739.0 MPa to 660.3 MPa, whereas the yield strength changes

little. At the beginning, the yield strength in process B is larger.

Fig. 3. TEM micrograph of specimens with the overaging temperature of (a) 280 1C, (b) 320 1C, (c) 400 1C and (d) 400 1C (dark eld image) in process A. M indicates

martensite, RA indicates retained austenite, GND indicates geometrically necessary dislocation.


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Fig. 4. SEM morphology of the tested steels with the reheated overaging temperature of (a) 310 1C, (b) 330 1C, (c) 350 1C and (d) 400 1C in process B. TM is tempered

martensite, RA is retained austenite, RA/UM is retained austenite or untempered martensite, DM is the completely decomposed martensite.

Fig. 5. TEM micrograph and diffraction patterns of retained austenite in the specimens with the reheated overaging temperature of (a) 310 1C and (b) 400 1C.

Fig. 6. TEM morphology of martensite in the experimental steels with different reheated overaging temperatures, (a) 3101

C and (b) 4001

C.


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However, as the overaging temperature increases, due to the

return of yield point and appearance of yielding terrace, the yield

strength in process A increases signicantly and exceeds the yield

strength in process B. In addition, it is found that with the increase

in the OT for process A, the ratio of YS/TS increases from 0.47 to

0.71. However, with the increases in the ROT for process B, the

ratio of YS/TS increases from 0.46 to 0.54.

According to Fig. 7, in the range of research work, compared

with the conventional annealing process (Process A), the new

annealing process (Process B) increases the tensile strength by

76107 MPa. When the OT or the ROT is lower than 330 1C, the

ratio of YS/TS is similar for both annealing process. As the OT or

ROT exceeds 350 1C, compared with process A, the ratio of YS/TS in

process B decreases by 0.160.17.The higher tensile strength in process B should be related with

the larger carbon content of martensite. It was reported that

increasing carbon content of martensite would lead to the increase

in the tensile strength of DP steel [27]. For process B, most of the

martensite is formed at the temperature of 250 1C, at which the

carbon diffusivity is lower. However, most of the martensite in

process A is formed at the temperature higher than 280 1C, at

which the carbon is more likely to diffuse towards the neighboring

untransformed austenite or ferrite. Furthermore, the overaging

temperature in process B decreases with time slowly, which

means the solubility of carbon in ferrite will decrease. As a result,

during overaging, for process B, the carbon in ferrite tends to

diffuse towards the untransformed austenite. During the subse-

quent nal cooling, some of the untransformed austenite will

transform into martensite, which has higher carbon content than

that formed in process A.

Furthermore, the relation between tensile strength and volume

fraction of the martensite in the steels of process A and B was

further analyzed. When the OT was 280 1C, 320 1C, 350 1C and

400 1C, the volume fraction of martensite in the steels was 19.2%,

17.8%, 17.6% and 11.2% respectively in process A. While the volume

fraction of martensite was 21.1%, 20.3%, 19.0% and 13.1% under the

ROT of 310 1C, 330 1C, 350 1C and 400 1C in process B. It was noted

that the above measured data of the volume fraction was the

average value of 5 statistic SEM morphology under each condition.

It was indicated that the volume fraction of martensite in process B

is slightly larger than that of process A. This was also the reason

why the tensile strength of the steel increased in process B.In addition to the higher carbon content and volume fraction of

martensite in process B, the formation of nanoscale cementite in

ferrite may also contributes to the higher tensile strength, as

shown inFig. 5.Fig. 8further shows the cementite precipitates in

ferrite. It is found that, for process A, the size of cementite

precipitates in the specimen with the OT of 400 1C is larger than

30 nm. However, for process B, the cementite size is approximately

10 nm in the specimen with the ROT of 400 1C, and the cementite

is distributed more uniform or disperse. The formation of nanos-

cale cementite could be due to the decrease of overaging tem-

perature with time. Because, in process B, the solubility of carbon

in ferrite decreases gradually with time, which facilitates the

precipitation of cementite during overaging. In addition, due to

the precipitation of cementite or diffusion of carbon towards the

Fig. 7. Yield strength, tensile strength and ratio of YS/TS for the experimental steels.

Fig. 8. Cementite precipitates in ferrite of the specimen with overaging temperature of 400 1C of process A (a) and reheated overaging temperature of 400 1C of process B (b).


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untransformed austenite, the formation of Cottrell atmosphere

would be suppressed. The Cottrell atmosphere affects the onset of

yielding by restricting the dislocation motion and thus increases

yield strength[28].Therefore, at the higher OT or ROT, the ratio of

YS/TS of experimental steels in process B is lower.

3.3.2. Elongation

Fig. 9 shows the variation of elongation with overaging tem-

perature. Note that, in this gure, the horizontal ordinate indicatesthe OT for process A and the ROT for process B. It is evident that

with the increase of the OT, the elongation increases continuously

in process A. However, in process B, the elongation changes little at

rst and then increases signicantly when the ROT exceeds 350 1C.

This is because with the increase of the OT or ROT, the martensite

plasticity increases and the strength mismatch between the softer

ferrite matrix and harder martensite phase decreases.

In order to analyze the variation of elongation, the microstruc-

ture in the vicinity of fracture surface was further examined, as

illustrated in Fig. 10. It is found from Fig. 10a and b that the

microvoids density decreases evidently with the increase of the OT

in process A. Similarly, with the increase of the ROT, the micro-

voids density also decreases in process B, as shown in Fig. 10c

and d. During the tensile deformation, the strain incompatibility ofthese two phases leads to the inhomogeneous deformation or

deformation localization of ferrite[29]. As a result, microvoids or

microcracks begin to nucleate at the severely deformed region

adjacent to the ferritemartensite interfaces. Kang et al. [30]

studied the microscopic strain distribution in DP steels during

tensile deformation. They found that, for the tempered DP steels,

the local strain for initiation of damage is higher, and the

martensite deforms much earlier. In other words, if the martensite

plasticity or the strength mismatch is improved, the nucleation

site of microvoids would certainly decrease and then the elonga-

tion of DP steels would increase.

The fracture morphology is illustrated inFig. 11. It can be seen

that the fracture surface is lled with dimples, which means all

specimens fractured in a ductile manner. However, it should benoted that the dimples are relative small or shallow in the

specimens with lower overaging temperature. The ductile fracture

of DP steels is closely related with the nucleation, growth and

coalescence of microvoids[31]. If the number of nucleation sites of

microvoids was quite large, the microvoids growth would be

inhibited due to the intersecting effect or linking up of neighbor-

ing voids[32]. As a result, the dimples on the nal fracture surface

are very small. On the contrary, if few nucleation sites of micro-

voids were present in the matrix, the nal dimples would be quite

large. According toFig. 10, the microvoids density decreases with

the increase of overaging temperature. Therefore, there exist large

and deep dimples on the fracture surface of specimens with higher

overaging temperature. In addition, the ductile fracture surface of

specimens further demonstrates that the new proposed overaging

process (process B) is an effective method to increase the strength

and guarantee the ductility.

3.3.3. Work hardening

The plastic ow behavior of DP steels can be described by

Hollomon relation as follows[33]:

Kn 1

where is true stress, is true strain, n is the work hardening

index, Kis strength coefcient. The value ofn signies the work

hardening characteristic, that is, the material with a higher value

ofn will work hardens at a higher rate during plastic deformation.

Furthermore, the larger the value ofn, the more the material can

deform before instability. Thus, the work hardening behavior is

directly associated with formability[34]. However, previous work

indicated that the DP steels usually show two or three stagehardening[3537]. This means the only one n cannot describe the

work hardening behavior of DP steels in detail. In this paper, the

instantaneous work hardening index (nn) was calculated according

to the Eq.(2), which is derived by Hollomon relation.

nn dln

dln 2

The plots of instantaneous work hardening index as a function

of true strain for process A and B are presented inFig. 12. Owing to

the reoccurrence of yield terrace, for the specimen with the OT of

400 1C in process A, the value of nn remains approximately zero

within the strain range of 0.0030.02. The sharp decrease/increase

in the nn value at the strain of 0.003 and 0.02 corresponds to the

lower/upper yield point. Apart from the specimen with the OT of400 1C in process A, the variation trend ofnn with strain in process

A is similar to that of process B. On the whole, in the strain range

studied, the variation trend ofnn can be divided in three stages. In

stage I, the nn value decreases signicantly, which is attributed to

the glide and annihilation of GND in the ferrite matrix. This can be

demonstrated by the phenomenon that the ferrite adjacent to the

ferrite/martensite interface softens after the DP steel was

deformed in uniaxial tension[4]. In stage II, thenn value increases,

which is due to the activation of new dislocation sources and the

consequent increase of statistically stored dislocation density[23].

In addition, the strain induced martensite transformation also

contributes to the increase in the nn value[2,38]. In stage III, the

variation trend ofnn is relative complex. This is because that with

the increase in strain, the increasing trend of dislocation densitybecomes slower, which leads to the decrease ofnn value. Mean-

while, with the increase in strain, the strain induced martensite

transformation may continue and the strength difference between

deformed ferrite and martensite decreases, i.e., the martensite

begins to deform with the ferrite simultaneously, which results in

the increase of nn value. At lower overaging temperature, the

strength difference between ferrite and martensite is larger, which

means that that martensite dose not deform at lower strain. As a

result,nn value decreases rst and then increases gradually. On the

contrary, at higher overaging temperature, the martensite is easier

to deform due to the tempering effect, and the amount of retained

austenite seems to be larger, which means that the strain induced

martensite transformation could continue at larger strain. Thus,nn

value increases continuously.

Fig. 9. Elongation as a function of temperature of the experimental steel in process A

and B.


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4. Conclusion

(1) In conventional annealing process, the tested steels were

directly cooled to various overaging temperatures from

region. With the increase of overaging temperature from

280 1C to 400 1C, some ferritemartensite interface becomes

indistinct, the amount of granular retained austenite increases,

the GND density decreases. With the increasing of overaging

Fig. 10. Microvoids pattern in mid-width plane of tensile specimens after fracture. (a) 280 1C in process A, (b) 400 1C in process A, (c) 310 1C in process B, (d) 400 1C in

process B.

Fig. 11. Fracture morphology. (a) 280 1 C and (b) 400 1C in process A, (c) 310 1C and (d) 400 1C in process B.


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temperature, the tensile strength decreases, whereas the yield

strength, ratio of YS/TS and elongation increase.

(2) In the proposed annealing process, at rst, the tested steels

were rapidly cooled to 250 1C and held for 50 s, and then they

were reheated to various overaing temperatures. At the

reheated overaging temperature of 400 1C, the amount of

retained austenite seems to be larger, and some martensite

was decomposed. With the increase in reheated overaging

temperature from 310 1C to 400 1C, the tensile strength

decreases gradually, the yield strength changes little. The

elongation changes little at rst and then increases signicantly

when the reheated overaging temperature exceeds 350 1C.

(3) The variation of instantaneous work hardening index (nn) with

strain can be divided into three stage. In stage I, the nn value

decreases signicantly due to the glide of geometrically

necessary dislocation. In stage II, the nn value increases due

to the increase in the density of statistically stored dislocation

or the occurence of strain induced martensite transformation.

In stage III, for specimens with lower overaging temperature,

the nn value decreases rst and then increases gradually, for

specimes with higher overing temperature, the nn valueincreases coninuously, which is attributed to the decrease of

strength mismatch between two phases and the increase in

the amount of granular retained austenite.

(4) At lower overaging or reheated overaging temperature, the

microvoids were easy to to nucleate. As a result, the dimples

on fracture surface was relative small and shallow.

(5) Compared with the conventional overaging process, the pro-

posed overaging process can increase tensile strength by 76

107 MPa. When the overaging or reheated overaging tempera-

ture is lower than 330 1C, the ratio of YS/TS is similar for both

annealing process. As the overaging or reheated overaging

temperature exceeds 350 1C, compared with process A, the

ratio of YS/TS in process B decreases by 0.160.17. This could be

attributed to the increase in carbon content of martensite andthe precipitation of nanoscale cementite in process B.

Acknowledgments

The authors are very grateful to the nancial support of the

National Natural Science Foundation of China (51174057, 51274062);

the National High Technology Research and Development Program of

China (2012AA03A503) and Research Fund for the Doctoral Program

of Higher Education of China (20130042110040).

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Fig. 12. Plots of instantaneous work hardening index as a function of true strain.

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