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Materials Science and Engineering A 553 (2012) 96– 104

Contents lists available at SciVerse ScienceDirect

Materials Science and Engineering A

journa l h o me pa ge: www.elsev ier .com/ locate /msea

icrostructural characteristics and impact fracture behavior of a high-strengthow-alloy steel treated by intercritical heat treatment

. Kang ∗, C. Wang, G.D. Wangtate Key Laboratory of Rolling and Automation, Northeastern University, Shenyang 110819, PR China

r t i c l e i n f o

rticle history:eceived 24 November 2011eceived in revised form 10 May 2012ccepted 24 May 2012vailable online 7 June 2012

eywords:

a b s t r a c t

Effects of the intercritical heat treatment (IHT) on microstructural evolution and Charpy impact fracturebehavior of a high-strength low-alloy (HSLA) steel were investigated. The toughening mechanism wasclarified by analyzing microstructural characteristics and crack propagation paths. Results showed thata composite microstructure of ferrite phase separated by globular, rod and irregular shape martensitewas obtained by adding the intercritical quenching to the conventional heat treatment of quenching andtempering. And 3.6% retained austenite was detectable in the microstructure. The percentage content

◦

igh-strength low-alloy steelicrostructure

rack propagationntercritical heat treatment

of high-angle (15 or more) boundaries reached 78.5%. It was also found that the steel had a high ratioof propagation energy (average: 152 J) to the total absorbed energy (average: 212 J) during impacting at−40 ◦C. Two crack propagation path models were observed: along the long axis direction of banded fer-rite, and across the grains and corresponding interfaces. The improvement of impact toughness wasattributed mainly to the retained austenite, the interlocking arrangement of banded ferrite and theferrite–martensite interfaces with high-angle misorientation, which exhibited effective resistance to the

crack propagation.

. Introduction

High-strength low-alloy (HSLA) steel plates are widely usedhroughout the world for various structural applications becausef the significant mechanical properties of high strength and goodoughness [1–5]. Due to possible applications in severe service envi-onments, the low-temperature impact toughness has been givenore attentions and becomes significant to the application of HSLA

teels. However, it becomes more difficult to achieve excellentoughness when the strength is very high, especially at a certain lowemperature, due to the conflict of the commonly used mechanismsor improving strength and toughness.

The most common heat treatment that is applied to improveechanical properties of HSLA steels consists of quenching and

empering. The intercritical heat treatment (IHT) is a new approacho enhance HSLA steels to dual phase microstructure with combina-ion of high strength, excellent toughness, low yield ratio and good

niform elongation [5–8]. The IHT is performed by quenching from

suitable temperature between the lower transformation temper-ture (Ac1) and the upper transformation temperature (Ac3). In thisntercritical quenching the austenite transforms to martensite, and

∗ Corresponding author at: No. 11, Lane 3, WenHua Road, HePing District,henyang 110819, PR China. Tel.: +86 24 83686415; fax: +86 24 23906472.

E-mail address: wwwkangjian@126.com (J. Kang).

921-5093/$ – see front matter © 2012 Elsevier B.V. All rights reserved.ttp://dx.doi.org/10.1016/j.msea.2012.05.098

© 2012 Elsevier B.V. All rights reserved.

finally a ferrite–martensite structure can be obtained [9–12]. Pre-vious investigations [3,13] have shown that the impact toughnessof dual phase steel is obviously influenced by martensite charac-teristics. It is not yet, however, clear exactly how ferrite phaseaffects the impact fracture behavior of dual phase steel, in spiteof its well-known importance in improving the impact toughness.

The aim of this work is to obtain high toughness microstruc-ture in HSLA steels through adding the IHT to the conventionalquenching and tempering treatment. The effects of the IHT onmicrostructural evolution and Charpy impact fracture behaviorwere investigated. And the mechanism of improving toughness byIHT was elucidated through the microstructural considerations andthe observation of microcracks in the fractured specimen.

2. Experimental procedure

The steel used in this work was manufactured in a 150 kg vac-uum induction furnace, then forged and cut into square ingot with120 mm thickness, 130 mm width and 200 mm length. The chemi-cal composition of the steel is shown in Table 1.

The rolling conditions and heat treatments used in this work

are shown in Fig. 1. The steel was austenitized at 1200 ◦C for 2 hand then rolled to 15 mm in a series of steps between 1150 ◦C and820 ◦C. The controlled rolling (CR) process was conducted with passreduction ratio of 20–24% in the austenite recrystallization regionand total reduction of 68% in the austenite non-recrystallization

J. Kang et al. / Materials Science and Engineering A 553 (2012) 96– 104 97

Table 1Chemical composition of the investigated steel (wt.%).

rit

Fatmtsfbatcsu

dsDcwmmac1Tbpc−c(pa(

lraw

Fr

The substructure of the steel was further investigated by TEM.TEM observations indicate that tempered lath martensite coex-

C Si Mn P S Nb

0.13 0.25 1.37 0.01 0.002 0.026

egion. The hot rolling process was carried out in a two-high revers-ng hot-rolling mill. And then the steel plate was cooled down tohe room temperature in the air.

In this study, Ac1 and Ac3 temperatures were determined by aormastor-FII dilatometer. Cylinder specimens with 10 mm lengthnd 3 mm diameter were used in the dilatation test. Results showedhat Ac1 and Ac3 were 727 ◦C and 831 ◦C respectively. The key

odification to the heat treatment process was the introduc-ion of IHT between the conventional quenching and temperingteps. Specimens were quenched after austenization at 930 ◦Cor 25 min, intercritically quenched from varying temperaturesetween 740 ◦C and 820 ◦C holding for 30 min, and then temperedt 500 ◦C for 50 min. It was found that the austenite volume frac-ion increased with increasing intercritical temperature, which wasonsistent with the previous investigations [14,15]. Among thepecimens above, the one intercritically quenched from 780 ◦C wassed for further investigation in this study.

Metallographic specimens were prepared according to the stan-ard procedure from the steel plates and etched with 4% nitalolution. Microstructural examinations were carried out by LEICAMIRM optical microscope (OM). In order to identify the martensiteharacteristics of the as-received steel, metallographic specimensere also etched with LePera solution (4% picral + 1% sodiumetabisulfite solution). Volume fraction of martensite was deter-ined by automatic areal analysis with the help of an image

nalyzer. And the mean linear intercept method was used to cal-ulate the size of the phases. Measurements were carried out at5 different areas of the specimen and average values were taken.he thin foil specimens for further microstructural investigationsy TECNAI G220 transmission electron microscope (TEM) wererepared on a twin-jet electropolisher at 28 V using the solutionomposed of 8% perchloric acid and 92% ethanol maintained below30 ◦C. Electron back scattering diffraction (EBSD) analysis was

onducted on a FEI QUANTA 600 scanning electron microscopeSEM) with the step size of 0.5 �m, and then the data were inter-reted by the orientation imaging microscope (OIM) system. Themount of retained austenite was estimated by X-ray diffractionXRD) analysis.

Tensile specimens with 8 mm diameter and 40 mm gageength were prepared in the transverse direction and tested atoom temperature on an INSTRON universal testing machinet a cross-head speed of 3 mm/min. Instrumented impact testsere performed on Charpy V-notch (CVN) specimens with the

ig. 1. Diagram of the processing route of the steel under study (CR: controlledolling, AC: air cooling, and WQ: water quenching).

Ti Ni Cr Mo B Al

0.016 0.27 0.36 0.33 0.0013 0.033

10 mm × 10 mm × 55 mm size in the longitudinal orientation at−40 ◦C using an INSTRON 9250HV impact tester.

The fracture surfaces of CVN specimens were subsequentlyexamined by SEM. And in order to determine the crack propagationpaths, a metallographic section was prepared from the fracturedCVN specimen, as shown in Fig. 2. Then the observed section wasexamined by SEM.

3. Results and discussion

3.1. Microstructural evolution

Fig. 3 presents the optical micrographs showing microstructuralevolution during heat treatments. The hot-rolling microstructure(before quenching) is composed of martensite–austenite (M–A)islands, bainite ferrite (BF) and degenerated pearlite (DP) as shownin Fig. 3(a). The microstructure after conventional quenching is typ-ical lath martensite as shown in Fig. 3(b). Fig. 3(c) reveals thatin the specimen treated by quenching and IHT, the ferrite phase(white) appears to be separated by globular, rod and irregular shapemartensite (gray). During reheating and isothermal holding at theintercritical temperature, the quenched martensite dissociates intoferrite plus carbide, then the austenite forms at grain boundariesand martensite lath (or lath packets) boundaries, and finally themicrostructure reverts to ferrite and austenite [12,16]. In the sub-sequent water quenching, the austenite transforms to martensite.Fig. 3(d) shows the tempered microstructure of Fig. 3(c), which isthe composite structure of tempered martensite and ferrite.

3.2. Microstructural characteristics

The intercritical temperature and holding time are the key con-trolling parameters for microstructural characteristics of the steel.The volume fraction of martensite at the selected intercritical tem-perature (780 ◦C) and holding time (30 min) reached 60.4%. Themicrostructural parameters are presented in Table 2.

ists with banded ferrite, as shown in Fig. 4(a), and ferrite bandswith 0.31–1.16 �m width and 6–15 �m length were frequently

Fig. 2. Schematic of the observed section used in the investigations of microvoidsand microcracks.

98 J. Kang et al. / Materials Science and Engineering A 553 (2012) 96– 104

F quencq

odamidcof

swTwa(ci[Kdpids

TM

ig. 3. Typical optical micrographs showing microstructural evolution: (a) before

uenching, IHT and tempering.

bserved. Fig. 4(b) reveals that tempered martensite consists ofislocation cells in the lath interiors as we know. The tanglednd zigzagged dislocations were found in the vicinity of temperedartensite and within banded ferrite (Fig. 4(c) and (d)). Fine precip-

tates on dislocation lines produce strong pinning effect, stabilizingislocation structure. And the interaction of dislocations and pre-ipitates remarkably increases the resistance to cross slip and climbf dislocation, which delays the recovery process and strengthenserrite phase.

The retained austenite was found in the microstructure of theteel. Fig. 5(a) and (b) shows the retained austenite in matrix,hich was examined respectively by bright field and dark field of

EM micrograph. The film-like retained austenite with 70–150 nmidth is mainly distributed between the martensite laths and

long the grain boundaries. Fig. 5(c) is selected area diffractionSAD) patterns of the retained austenite and the matrix and theirorresponding analyzing results. It is seen that the retained austen-te has the coherent orientation relationship with the matrix:1 1 1]M // [ 1 0 1 ] � . Apparently, the result is in agreement with–S relationship. And 3.6% retained austenite was quantitativelyetected by means of X-ray. The retained austenite consists of two

arts. One part is that remaining directly from intercritical quench-

ng process. And the other part is the reversed austenite formeduring tempering, which nucleates and grows in the fresh marten-ite laths obtained from intercritical quenching process [17,18].

able 2icrostructural parameters of the steel processed by quenching and IHT.

Volumefraction ofmartensite(%)

Average size ofmartensite(�m)

Width range ofbanded ferrite(�m)

60.4 2.14 0.31–1.16

hing, (b) after conventional quenching, (c) after quenching and IHT, and (d) after

The misorientation map of grain boundaries is shown in Fig. 6.Grain boundaries showing misorientations of 2–5◦ and 5–15◦ arecolored in red and green respectively. The percentage content ofhigh-angle (15◦ or more) boundaries colored in blue reaches 78.5%.It is found that boundaries of neighboring ferrite bands are mostlylow-angle ones of 2–15◦. The tempered martensite phases are high-angled each other, and boundaries between tempered martensiteand ferrite are also mostly high-angle ones.

3.3. Mechanical properties

Yield and tensile strengths of the specimens were found tobe in the range of 688–702 MPa and 807–820 MPa, respectively.The average values of mechanical properties are listed in Table 3.The low yield ratio (approximate 0.85) was obtained. This isclosely related to the composite microstructure constituted bysoft ferrite and hard martensite, where the soft phase ensuresgood deformability and the hard phase provides high tensilestrength [8].

Continuous yielding behavior in the tensile test is usually con-sidered to be one of the characteristics of dual phase steels [19–21].

However, this phenomenon was not observed for as-received spec-imens. As can be seen from Fig. 7, tensile stress–strain curves ofspecimens exhibit well-defined yield point phenomenon, whichcan be related to the followings:

Table 3Mechanical properties of the specimens.

Yieldstrength(MPa)

Tensilestrength(MPa)

Yield ratio(%)

Elongationin 40 mm(%)

Impactenergy (J)

693 813 85.2 21.8 212

J. Kang et al. / Materials Science and Engineering A 553 (2012) 96– 104 99

F a) band in ban

(

(

3

avsimcotat

unstable crack propagation (the sharp decrease in load carryingcapacity). Thus, the energy for crack propagation (Ep) consists ofthe stable and unstable crack propagation energy. And the ratioof propagation energy to the total absorbed energy is between

Table 4Impact fracture energies of three specimens determined by instrumented impacttest.

Specimen Total absorbed Initiation Propagation Propagation to

ig. 4. TEM micrographs of the steel processed by quenching, IHT and tempering: (islocations around phase boundaries, and (d) high-density zigzagged dislocations

1) Although high-density dislocations were created in ferrite viaplastic deformation during the austenite to martensite transfor-mation, many of them with enough thermal activation energywould interact with each other and annihilate during temper-ing at 500 ◦C [22]. Simultaneously, fine iron carbide precipitatesproduced strong pinning effects on dislocation movement (asshown in Fig. 4(c)). Consequently, the mobile dislocations werestrongly reduced.

2) Even if the dislocations breaking away from the pinning pre-cipitates could move, other immobile dislocations and bandedferrite boundaries would block them immediately (Fig. 4(d)).Thus, the mean free path (MFP) of slipping dislocations wasrestricted.

.4. Charpy impact energy

The average Charpy impact energy of specimens reached 212 Jt −40 ◦C, as seen in Table 3. Curves of load and impact energyersus displacement obtained by instrumented impact test arehown in Fig. 8. The load–displacement curve indicates that thempact process of this steel is characterized by initial elastic defor-

ation, the occurrence of the general yield, the initiation of fracture

rack at the peak load and crack propagation till the occurrencef fracture. It is worth mentioning, however, that crack initia-ion generally occurs somewhere between the general yield pointnd the peak load point. And it is a common practice to considerhat cracks roughly initiate at peak load [23,24]. Thus, the same

ded ferrite (F) and tempered martensite (M), (b) tempered martensite, (c) tangledded ferrite.

assumption was used in this investigation to calculate the crackinitiation energy.

The total absorbed energy (the area under theload–displacement curve) is composed of crack initiation energy(Ei) and crack propagation energy (Ep). The energy absorbed up tothe peak load can be regarded as the energy for crack initiationEi, which includes elastic bending and plastic deformation energy.The total absorbed energy values and their contributions (Ei andEp) determined by the test data for three specimens are listed inTable 4.

As already observed in Fig. 8, the stable crack propagation occursafter crack nucleation (the peak load) and before the onset of

energy (J) energy Ei (J) energy Ep (J) total absorbedenergy (%)

1 210 58.2 151.8 72.32 218 64.1 153.9 70.63 208 57.8 150.2 72.2

100 J. Kang et al. / Materials Science and Engineering A 553 (2012) 96– 104

F , IHT ar

7abpaen

Fa

ig. 5. TEM micrographs of retained austenite in the steel processed by quenchingetained austenite and the matrix as indicated by the circle in (a).

0.6 and 72.3% for tested specimens (Table 4). Obviously, it can beppreciated that the crack initiation energy is a small contri-ution to the total absorbed energy and the energy for crack

ropagation plays a major role for this steel. This can bettributed mainly to the dual phase microstructure, which exhibitsffective resistance to the crack propagation once the crack isucleated.

ig. 6. EBSD analysis of the microstructure of the steel processed by quenching, IHTnd tempering (F: ferrite and M: tempered martensite).

nd tempering: (a) bright field image, (b) dark field image, and (c) SAD patterns of

3.5. Microfracture mechanism

Fig. 9(a) and (b) shows the fracture surface near the notch rootof a fractured CVN specimen. The morphology exhibits typical dim-

ple rupture. During impact deformation, ferrite (the soft phase)deforms first, and the specimen begins yielding on the tensile side(at the notch root). Then plasticity spreads and the plastic strainsubsequently reaches the critical point at which ductile fracture isinitiated [25]. Moreover, this steel exhibits good ductility at even

Fig. 7. Stress–strain curves obtained from the room-temperature tensile test of thespecimens.

J. Kang et al. / Materials Science and En

Fo

−st

twbtdiid

The excellent toughness of the as-received steel is related

Fp

ig. 8. Load and impact energy versus displacement curves of one specimenbtained by instrumented impact test at −40 ◦C.

40 ◦C and can produce enough plastic deformation to release thetress in front of the main crack. Therefore, deep dimples withearing ridges were formed on fracture surface near the notch root.

Subsequently, the fracture crack propagates by the dimpled duc-ile mode. An example of this type of fracture is shown in Fig. 9(c),hich shows the fracture surface in crack propagation area. It can

e clearly seen that the fracture morphology is characterized byhe dimple depression in the surface. As a matter of fact, uneven

imples were frequently observed in crack propagation area. This

s most likely associated with the unstable crack propagation dur-ng impacting. Furthermore, the size and depth of dimples mainlyepend on the ductility of microstructure, the inclusion dimension

ig. 9. SEM fractographs of the CVN specimen tested at −40 ◦C: (a) near the notch root,

ropagation area.

gineering A 553 (2012) 96– 104 101

and the stress condition in front of crack [26]. The existence of dim-ples surrounded by tearing ridges implies that plastic deformationoccurs even at −40 ◦C during the main crack propagation. Thus,the resistance to crack propagation is very high in this steel andthe propagation of fracture crack requires large energy, which isconsistent with the analysis of impact curves in Fig. 8.

Fig. 10 shows typical examples of microvoids and cracks under-neath the fracture surface. The microvoids were nucleated atinterfaces between tempered martensite and ferrite, as shown witharrows in Fig. 10(a). This is consistent with previous works on frac-tographic investigations of dual phase steels, which indicate thatmicrovoids are nucleated preferentially either at ferrite–martensiteinterfaces or at precipitates in ferrite [7,25,27,28].

The load–displacement curve in Fig. 8 has revealed that a certainamount of plastic deformation is required for stress concentrationstrong enough for microvoids nucleation. During plasticity spread-ing in the soft phase (ferrite), interfaces between soft phases andhard phases (tempered martensite) are usually the locations forthe stress concentration, which can easily cause microvoids. How-ever, the precipitates in ferrite are finely dispersed in this steel(as in Fig. 4(c) and (d)), and the stress concentration at the pre-cipitates is not strong enough to tear apart the microstructure.Consequently, microvoids are nucleated predominantly at inter-faces between tempered martensite and ferrite during impacting,as frequently observed in Fig. 10(a).

partly to the presence of retained austenite. The film-like retainedaustenite is considered to weaken the stress concentration at thecrack tip through its plastic deformation [29,30], for which theimpact energy of this steel can be increased. Another point worth

(b) higher magnification of the area indicated by the rectangle in (a), and (c) crack

102 J. Kang et al. / Materials Science and Engineering A 553 (2012) 96– 104

F ractura in (a).

ccicio

tia

Fo

ig. 10. SEM micrographs showing the microvoids and cracks underneath the main fnd (b), (c) and (d) various paths of crack propagation. Arrows indicate microvoids

onsidering is that the energy of crack propagation makes the majorontribution to the impact energy (Fig. 8). In order to clarify thempact fracture mechanism, the crack propagation paths in theomposite microstructure have been observed by SEM, as shownn Fig. 10(b)–(d). In the schematic in Fig. 11 the characteristic pathsf crack propagation in this steel are summarized.

Cracks nucleating at ferrite–martensite interfaces were seeno propagate in two models. One path model can be character-zed by propagating along the specific direction of ferrite bands,s shown in Figs. 10(b) and 11(a). In this case, it is demonstrated

ig. 11. Schematic diagrams showing crack propagation paths in the microstructure comf ferrite bands and (b) across grains and ferrite–martensite interfaces.

e surface of the CVN specimen tested at −40 ◦C: (a) locations of microvoid nucleation

that the cracks grow in soft ferrite phases and propagate prefer-entially along the long axis direction of ferrite bands once cracksare nucleated, and then the further crack propagation might bearrested or deflected. This path model can be attributed mainly tothe stress localization conditions in ferrite. When the stress is notquite strong in front of the crack, it would propagate only along the

specific direction in the soft banded ferrite due to the restriction ofhard tempered martensite. Impact deformation could induce stresslocalization on the planes of weakness, i.e., cleavage planes, whichmake the crack cleavage propagation easy to occur. However, finely

posed of banded ferrite and tempered martensite: (a) along the specific direction

J. Kang et al. / Materials Science and En

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3

jeoTGas[hft

faanctoa[ima

ig. 12. TEM micrograph showing the interface between tempered martensite anderrite.

ispersed tempered martensite and interlocking arrangement oferrite bands play as the main barriers to the cleavage crack propa-ation and, as a matter of fact, contribute to the toughening in thisteel.

Fig. 11(b) shows the other path model for crack propagation.t can be found that cracks form at ferrite–martensite interfacesnd propagate across the grains and corresponding interfaces. Theirection of crack propagation seems to be changed by the inter-aces (Fig. 10(c)), while the relatively straight crack path has alsoeen detected (Fig. 10(d)). This can be attributed to the fact thatnce the crack-driving force is high enough cracks may propagaten the direction of growth and across the ferrite–martensite inter-aces. Moreover, the bunch of ferrite bands is very likely to possesshe same or similar orientation, and cracks probably propagatelong their common cleavage plane.

.6. Interface characteristics and cracking

The microstructure configuration of the experimental steel sub-ected to IHT process is the essential factor that determines impactnergy and crack propagation behavior. Among that, the interfacer grain boundary (GB) properties play quite an important role.he GB properties can be characterized by configuration, density,B energy, etc. GB energy �GB is a significant parameter for GBs,nd some investigations involving GB energetics and its relation-hip with cracking have been done in following cited literatures31–41]. The results of TEM and EBSD analysis in this investigationave displayed the importance of interface characteristics in impact

racture mechanism. So the effect of GBs on cracking behavior needso be discussed further here.

Fig. 12 shows the interface between tempered martensite anderrite. From Fig. 6 it is known that GBs of this kind are high-anglednd account for the greater proportion. The high misorientationngle is a very important characteristic from the view of high tough-ess [42–45]. GBs can be classified into three categories using theoincident site lattice (CSL) model: low-angle boundary (LAB) upo 15◦ misorientation, non-special CSL high-angle boundary (HAB)r general GBs, and special CSL GBs [31]. The GB energy of LABs

nd special GBs is small, while the energy of non-CSL HABs is high32–34]. The ferrite and martensite grains on different sides of thenterface in Fig. 12 have experienced quite different phase transfor-

ation histories. The interface has no obvious special configurationnd is considered to be non-special CSL high-angle boundary. Hence

gineering A 553 (2012) 96– 104 103

the GB energy of such GBs is higher than others and comes intoeffect during the fracture process.

It is shown that GBs with higher energy also have highermobility [35,36]. Non-CSL HABs are more sensitive to intergran-ular microcrack nucleation than LABs and special GBs [31,37],and cracks preferentially nucleate in non-CSL HABs [38,39]. In thepresent study, ferrite–martensite interfaces possess higher energyand are preferred positions for microcracks nucleation. During thepropagation of the main crack, strain concentration in front ofthe crack tip will induce microvoids at ferrite–martensite inter-faces. Microvoids are easier to connect with each other in highenergy GBs [40]. The main crack will reach to the microvoids andpropagate along the high-angle boundaries. Certainly, microvoidsare also generated at interfaces of inclusions and matrix. As fortransgranular cracking, propagation rate is also slowed by highboundary energy and the propagation rate is inversely proportionalto boundary energy [41]. Thereby, the crack cannot keep proceed-ing straightly forward in front of high-angle GBs and has to sufferobstacles and turn its way (Fig. 10(a) and (c)). As a result, the pre-vailing high-angle boundaries in this steel make the propagationroute winding and lengthened. They force the crack deflections andleave ligaments evidently, becoming effective barriers to the crackpropagation and slowing down the propagation rate in fact.

The composite structure of ferrite phase separated by fineirregular shape martensite is considered to be advantageous forbalancing plastic deformation between hard phase (temperedmartensite) and soft phase (ferrite). Furthermore, the banded fer-rite and dispersed martensite in this steel processed by quenching,IHT and tempering lead to more crack deflection, higher fracturesurface roughness and slower crack propagation.

4. Conclusions

(1) Dual phase steel with composite microstructure of ferrite phaseseparated by globular, rod and irregular shape martensite wasproduced by adding the IHT to the conventional heat treatmentof quenching and tempering.

(2) The film-like retained austenite was found along the lathboundaries. The as-received steel had a high ratio of high-angle(15◦ or more) boundaries versus low-angle ones.

(3) The average Charpy impact energy of specimens reached 212 Jat −40 ◦C. And the crack propagation energy made the majorcontribution to the total absorbed energy during impacting.

(4) Cracks nucleating at ferrite–martensite interfaces were seen topropagate in two models. One path model can be characterizedby propagating along the long axis direction of banded ferrite.The other one can be characterized by propagating across thegrains and corresponding interfaces.

(5) The improvement of impact toughness for the steel treated byintercritical heat treatment is attributed mainly to the retainedaustenite, the interlocking arrangement of banded ferrite andthe ferrite–martensite interfaces with high-angle misorienta-tion of 15◦ or more.

Acknowledgments

This work is supported by National Basic Research Program ofChina (no. 2010CB630800) and Fundamental Research Funds forthe Central Universities (no. N090607001).

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