jmpt 2013 secondary deformation of hot stamping specimens

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Journal of Materials Processing Technology 213 (2013) 818 825

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Journal of Materials Processing Technology

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Effect of thermo-mechanical process on the microstructure andsecondary-deformation behavior of 22MnB5 steels

Junying Mina, Jianping Lina,, Yongan Minb

a School of Mechanical Engineering, Tongji University, Shanghai 201804, Chinab School of Material Science and Engineering, Shanghai University, Shanghai 200072, China

a r t i c l e i n f o

Article history:Received 30 July 2012Received in revised form26 November 2012Accepted 29 December 2012Available online xxx

Keywords:Secondary-deformation22MnB5Digital image correlation (DIC)Ferrite transformationBainite transformation

a b s t r a c t

22MnB5 steel specimens were deformed at 923 K and 693 K to three strain levels to study the effectof applied strain level on the microstructure and secondary-deformation behavior. As the steel wasdeformed at 923 K, deformation induced ferrite transformation (DIFT) occurred even when a small strainof 0.044 was applied, and the volume fraction of deformation induced ferrite (DIF) increases with increas-ing applied strain level. When deformed at 693 K, deformation induced bainite transformation (DIBT) wasobserved when the applied strain was larger than 0.109. The incubation period for DIFT is shorter thanthat for DIBT, but the DIBT proceeds much faster than DIFT. Sub-size tensile specimens were cut fromthe hot deformed 22MnB5 steel specimens, and digital image correlation technique was employed toinvestigate the secondary-deformation behavior of the sub-size tensile specimens at room temperature.It is found that the appearance of DIF or DIB (deformation induced bainite) decreases the yield strengthand ultimate tensile strength (UTS) but increases the elongation and strengthductility product of thehot deformed 22MnB5 steel specimens compared with the as-quenched 22MnB5 steel specimen withfull martensite.

2013 Elsevier B.V. All rights reserved.

1. Introduction

The requirements on automotive light-weighting and increas-ing crashworthiness stimulate the application of hot stampingtechnology of 22MnB5 steel. According to Hein and Wilsius(2008), hot stamping technology has been developed to partial hotstamping, hot stamping of tailor-rolled blank, etc. to obtain hotstamped components with tailored mechanical properties, whilethe conventional hot stamping process is aimed to manufacturefull-martensite components with ultimate tensile strength (UTS) ashigh as 1500 MPa as introduced by Karbasian and Tekkaya (2010).For a B-pillar, the design of tailored microstructure with gradientmechanical properties, namely, a tailored structure of harder upperB-pillar and softer lower B-pillar with larger elongation, is bene-cial to increase both the resistance of deformation invasion andenergy absorption during side crash.

The microstructure of hot stamped components is closelyrelated to the thermo-mechanical process of the 22MnB5 steel.As stated by Bardelcik et al. (2010), a cooling rate of 25 K/s leadsto a constituent of 95% martensite and 5% bainite in the 22MnB5steel without hot deformation. At room temperature, the UTS of

Corresponding author. Tel.: +86 13901719457.E-mail address: [email protected] (J. Lin).

quenched 22MnB5 steel with full martensite shows little strainrate sensitivity when the strain rate increases from quasi-staticstate (0.003/s) to the high strain rate state (960/s). Nikravesh et al.(2012) investigated the effect of hot deformation on the martens-ite and bainite start temperature (Ms and Bs), and it was concludedthat the hot deformation decreases the Ms but increases the Bs from853 K to 873 K. Abbasi et al. (2012) reported isothermal deformationlowered the Ms and Mf. Eman et al. (2009) described the neckingprocess of press hardened boron steel by employing digital speck-led correlation method. Recently, Shi et al. (2012) reported thephase transformation in non-isothermally deformed boron steel,and it is concluded that the non-isothermal deformation promotesthe diffusional transformation, such as ferrite and bainite, andlowers the Ms. The microstructure of the hot formed boron steeldetermines its mechanical properties which is important to thecrash performance of hot stamped components. Especially, com-ponents with a mixture of ferrite and martensite or bainite andmartensite have been manufactured due to excellent combinedperformances of high strength and elongation. Therefore, it is ofsignicance to understand the secondary-deformation behavior ofboron steel after hot deformation, so as to determine the appropri-ate hot stamping process to manufacture hot stamped componentswith better mechanical performances.

In this paper, 22MnB5 steel specimens were austenized at1173 K and pulled at 923 K or 693 K to three strain levels and then

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J. Min et al. / Journal of Materials Processing Technology 213 (2013) 818 825 819

Table 1Chemical composition of 22MnB5 steel.

Alloy elements C Mn P S Si Al Ti B Cr

Content (wt./%) 0.221 1.211 0.019 0.003 0.258 0.036 0.039 0.0037 0.190

quenched to room temperature with a cooling rate of 30 K/s, and themicrostructures of the specimens were examined. Sub-size tensilespecimens were cut from the hot deformed specimens and pulledat room temperature with the aid of DIC technique to investigatethe secondary-deformation behavior of 22MnB5 steels.

2. Experiments

The experiments in this paper include uniaxial tensile tests of22MnB5 steel specimens at elevated temperatures, metallographicobservations and uniaxial tensile tests of hot deformed 22MnB5steel specimens at room temperature with the aid of DIC technique.

2.1. Uniaxial tensile tests at elevated temperatures

The chemical composition of the 22MnB5 steel is listed inTable 1. The dimensions of 22MnB5 steel specimens shown in Fig. 1follow the standard GB/T 4338-2006, and round grids with a diam-eter of 2.5 mm (d0) and a center-to-center distance of 3 mm werelaser etched on one surface of the specimens in order to measurethe strains after hot deformation. The depth of the grids is 30 mto make sure the grids keep clear after hot deformation and heattreatment.

All tensile tests were carried out on a thermo-mechanical simu-lation system Gleeble 3800. The specimens were heated to 1173 Kat a rate of 15 K/s for full austenitization and the soaking time was5 min. The specimens were quenched to the 923 K or 693 K at acooling rate of 30 K/s. Then isothermal deformation was performedon the specimens with a crosshead speed of 5 mm/s. The crossheaddisplacements were set as 2 mm, 4 mm and 10 mm, giving dwellingtime of 0.4 s, 0.8 s and 2 s at each temperature, respectively. Threespecimens were repeated for each experimental condition, whereone specimen was for metallographic microstructure observationand the other two were for secondary tensile tests at room tem-perature. Hence, there are 2 3 3 = 18 specimens in total. Afterdeformation, the hot deformed specimens were quenched to roomtemperature at a cooling rate of 30 K/s immediately, and the roundgrids on the specimens become ellipses with a length of d1 in themajor axis, namely, in the tensile direction. Then the major strain(1) can be calculated as ln(d1/d0).

Since this work focuses on the effect of hot deformation on themicrostructure and secondary deformation behavior of 22MnB5steels, it is worth to mention here whether ferrite or bainite trans-formation in 22MnB5 steel specimens with no hot deformationwhen the specimens statically dwell at 923 K or 693 K and theother heat treatment conditions are the same as those for hotdeformed specimens. Fig. 2 shows the dilatations of three 22MnB5steel specimens (10 mm 10 mm) with no hot deformation

Fig. 1. Illustration for the dimensions of specimens for hot tensile tests (unit in mm).The diameter and the center-to-center distance of the round grids laser etched onthe specimens are 2.5 mm and 3 mm, respectively.

during cooling, where the heating rate, soaking time and cool-ing rate are the same as described above. One specimen has nodwelling period at elevated temperature, namely, it was directlyquenched to room temperature after soaking. The other two speci-mens statically dwelled at 923 K and 693 K for 2 s, respectively. Thearrows in Fig. 2 indicate the Ms (669 K) in the three non-deformedspecimens. The static dwelling at 923 K or 693 K for 2 s has littleeffect on the dilatation vs. temperature curve and the Ms, and obvi-ous evidence indicating there is ferrite transformation or bainitetransformation is not observed in Fig. 2. Fig. 3 shows the metallo-graphic microstructures of the three non-deformed specimens. Forthe specimens with direct quenching and static dwelling at 693 Kfor 2 s, the microstructures (Fig. 3(a) and (c), respectively) are fullmartensite. Therefore, the incubation period for isothermal bainitetransformation at 693 K is larger than 2 s, and consequently, themicrostructure will be full martensite for the 22MnB5 steel speci-mens with static dwelling at 693 K for even shorter period (e.g. 0.4 sand 0.8 s). As regard to the specimen with static dwelling at 923 Kfor 2 s, the microstructure is composed of a very limited amountof ferrite (1% in volume fraction) and 99% (in volume fraction)martensite. Hence, the isothermal ferrite transformation needs anincubation period 2 s at 923 K, and the volume fraction of fer-rite is much less than 1% in the 22MnB5 steel specimens with staticdwelling at 923 K for a shorter period (e.g. 0.4 s and 0.8 s). The aboveresults will be referred in Section 3 to clarify the effect of staticdwelling time at 923 K or 493 K on the ferrite transformation orbainite transformation.

2.2. Metallographic experiments

Metallography specimens with dimensions of 4 mm 5 mmwere cut from the middle of the specimens by wire-electrodecutting and mounted. The 4 mm direction is the major strain direc-tion. The mounted specimens were ground by abrasive paper andpolished by woolen cloth with diamond paste. Then the Vickershardness of each specimen was measured on a hardness testerHXD-1000TC. Each specimen was tested three times and the meanvalue was calculated. The metallographic microstructure observa-tions of all specimens, which were chemically etched by 4% nitricacid, were performed on a scanning electron microscope.

Fig. 2. Dilatation vs. temperature during cooling of three non-deformed 22MnB5steel specimens. The arrows indicate the Ms of the three specimens.


820 J. Min et al. / Journal of Materials Processing Technology 213 (2013) 818 825

Fig. 3. Metallographic microstructures of non-deformed 22MnB5 steel specimenswith (a) direct quenching, (b) static dwelling at 923 K for 2 s and (c) static dwellingat 693 K for 2 s.

2.3. Secondary-tensile tests at room temperature with the aid ofDIC technique

Since there was a uniform temperature zone, which is con-sequently a uniform deformation zone, with a length of about30 mm positioned on the middle of each hot deformed speci-men, a sub-size tensile specimen was wire-electrode cut from eachhot deformed specimen for secondary-tensile test at room tem-perature as illustrated in Fig. 4(a). For comparison, two sub-sizetensile specimens were also cut from the as-quenched 22MnB5steel sheet. The cooling rate of the as-quenched 22MnB5 steel sheetis 30 K/s to obtain full martensite microstructure. The length of par-allel section on the sub-size tensile specimen is 18 mm to ensurethe material between two shoulders are homogeneous and to

Fig. 4. (a) Illustration for dimensions of sub-size tensile specimen (solid line, unit inmm) and (b) an image for a hot deformed specimen (upper) and a ground sub-sizetensile specimen (lower).

avoid heterogeneity in microstructure resulted from non-uniformthermo-mechanical process.

Due to slight oxidization on surfaces of specimens during hottensile tests, the sub-size tensile specimens were ground with 150#and 240# sand paper successively to remove the thin oxidized lay-ers on two surfaces and two sides, as shown in Fig. 4(b). Care wastaken during grinding to ensure the deviation in thickness is lessthan 0.5%. Then the thickness and width of the ground sub-sizetensile specimens were measured for stress calculation.

Prior to secondary-tensile testing, the surfaces of sub-size ten-sile specimens were cleaned with chloroform. One surface of eachspecimen (including two shoulder areas) was decorated with athin layer of white spray paint. Black spray paint droplets werethen applied but care was taken to ensure no single dropletexceeded the size of the chosen square pixel subset region dur-ing the DIC post-processing. The tensile tests were conducted ona ZWICK/ROELL Z050 universal testing machine with a crossheadspeed of 1.08 mm/min giving a nominal strain rate of 103/s. Allsub-size specimens were pulled till fracture. Stereo digital imagecorrelation detailed by Sutton et al. (2009) was employed to mea-sure full-eld strains during tensile test of each sub-size tensilespecimen. Digital images were recorded by two 5-mega-pixel cam-eras. The time step was set as 0.3 s corresponding to a frame rate of3.33 frame/s. Each image was tagged with a load and crosshead dis-placement from the analog output of the ZWICK/ROELL machine.The strain elds are computed from digital grids superimposed oneach image during post-processing of the images. The strain elddata should be converted to a uniaxial eld to compute true stress.In the present study, the uniaxial strain measurement was accom-plished with two methods, which have been successfully appliedin previous studies of metal deformation by Tong et al. (2005) andZavattieri et al. (2009). The rst corresponds to the conventionalmethod by use of an extensometer in tensile testing.

11 =1

M N

Nj=1

Mi=1

1(i, j) (1)

Here, M is grid point number along the tensile axis with a gaugelength (to be determined in a following section) positioned abovetensile specimen, and N is the grid point number transverseto the tensile axis. Therefore, Eq. (1) calculates the axial strain



Table 2The major strains (1) measured on hot deformed 22MnB5 steel specimens.

Deformation amount 2 mm 4 mm 10 mm

923 K 0.044 0.081 0.200693 K 0.041 0.109 0.206

integrated along both the length and the width of the specimen.The 2nd method corresponds to a localized strain measurement:

31 =1N

Nj=1

1(M, j) (2)

Eq. (2) is the axial strain integrated across a line running across thewidth of the specimen at any location along the specimen length.From volume constancy of plastic deformation and force equilib-rium along the tensile direction, the true stress can be computedby Eq. (3).

1 =F exp 1

w0 t0(3)

where F is the crosshead load, w0 and t0 are the initial width andthickness of the parallel section on the sub-size specimens.

3. Metallographic results

As reported in the authors previous study (Min et al., 2012), acooling rate of 30 K/s after full austenitization can inhibit ferriteand bainite transformations during cooling to room temperature,and full martensite (Fig. 3(a)) can be obtained in the 22MnB5steel specimen. In addition, according to Barcellona and Palmeri(2009), stored energy will be introduced in the austenite due to hotdeformation, which shortens the incubation period and promotesthe nucleation in ferrite transformation and bainite transforma-tion. Therefore, the austenite deformation at elevated temperaturewill induce ferrite transformation and bainite transformation andeven promote these transformations during following cooling afterdeformation. The strain levels of hot deformed 22MnB5 specimensare shown in Table 2. The metallographic microstructures of thespecimens deformed at 923 K and 693 K are shown in Figs. 5 and 6,respectively, and the corresponding Vickers hardness of the spec-imens is shown in Fig. 7. As detailed by Min et al. (2012) (923) Kis higher than the bainite transformation start temperature (Bs). Itcan be seen in Fig. 5 that ferrite transformation is induced by hotdeformation even when a small strain of 0.044 is applied. Hence,the incubation period for deformation induced ferrite transforma-tion (DIFT), which should be less than 0.4 s (calculated from thedeformation level and crosshead speed), is much shorter comparedwith conventional isothermal ferrite transformation (2 s). Fur-thermore, the volume fraction of deformation induced ferrite (DIF),which includes the ferrite transformed during deformation and fol-lowing cooling phase, increases with increasing the applied strainlevel. Consequently, the Vickers hardness decreases from 395HV5to 302HV5 as shown in Fig. 7. Qi et al. (2005) has demonstrated thatthe DIF prefers to nucleate on the grain boundaries since there isalways higher distortion energy at these locations. Therefore, thene ferrite distributes by a network in the case of a small strain(0.044) applied at 923 K, which can be seen in Fig. 5(a) and the mar-tensite still dominates the microstructure, and the volume fractionof DIF is 7.3%. As increasing deformation level at 923 K, the storedenergy in austenite grains increases, and the effect of deforma-tion inducing ferrite transformation is enhanced. When the ferritenucleation sites on the grain boundaries of the deformed austen-ite are exhausted, large amount of deformation bands in originalaustenite grains and the fresh ferrite front can provide plenty ofnucleation sites as stated by Qi et al. (2005). Hence, the nucleation

Fig. 5. Metallographic (SEM) images of the 22MnB5 steel specimens deformed at923 K with strain levels of (a) 0.044, (b) 0.081 and (c) 0.200. F indicates deformationinduced ferrite, and M indicates martensite.

sites of ferrite turn into austenite grain inside from grain boundariesof deformed austenite as the deformation level increases, whichcan be seen in Fig. 5(b) and (c). As a result, the volume fraction ofDIF increases to 58% and the martensite becomes a minor phasewhen the applied strain level increases to 0.2. Here it is worth toclarify that the effect of dwelling time (2 s) at 923 K on the ferritetransformation of 22MnB5 steel specimens can be negligible, sincethe volume fraction of ferrite in the non-deformed specimen withstatic dwelling at 923 K is 1% as stated in the Section 2.1, whichis a very small proportion compared with the ferrite amount in the22MnB5 steel specimens deformed at 923 K.

When the 22MnB5 steel was deformed at 693 K that is 24 Khigher than the Ms, almost full martensite is observed in thespecimen with an applied strain level of 0.041. Nevertheless, defor-mation induced bainite transformation (DIBT) can be found (8.7%in volume fraction) from the metallographic microstructures whenthe applied strain is 0.109. The deformation induced bainite (DIB)is acicular ferrite (AF) according to Zhang and Boyd (2010), which



Fig. 6. Metallographic (SEM) images of the 22MnB5 steel specimens deformed at693 K with strain levels of (a) 0.041, (b) 0.109 and (c) 0.206. B indicates deformationinduced bainite, and M indicates martensite.

is different from the conventional bainite (CB) generated from con-tinuous cooling transformation or isothermal transformation ofnon-deformed austenite, and the AF is formed as randomly dis-tributed ferrite laths, which are much ner than the CB and canbe seen in Fig. 6(b) and (c). When the applied strain increasesto 0.206, the volume fraction of DIB induced by hot deforma-tion increases signicantly to 86% but the hardness decreases to339HV5. Consequently, after a short incubation period less than0.8 s (a deformation level of 4 mm), the DIBT process proceeds muchfast, and it is even faster than the DIFT process. Here it is also worthto clarify that the effect of dwelling time (2 s) at 693 K on bainitetransformation can be ruled out since the incubation period thatisothermal bainite transformation at 693 K needs is larger than 2 sas stated in the Section 2.1.

Comparing the martensite transformed from non-deformedaustenite (Fig. 3(a)) with the martensite in Fig. 5(a) and 6(a), themartensite lath is rened by introducing deformation in austen-ite. Since the dislocation cells in deformed austenite will act as

Fig. 7. Vickers hardness of the hot deformed 22MnB5 steel specimens. The dashedline indicates the hardness of the as-quenched 22MnB5 steel.

barriers for the growth of martensite, the size of martensite lathis reduced and the lath renement is enhanced by increasing thestress in austenite. Therefore, the martensite lath in the specimensdeformed at 693 K is much ner than that deformed at 923 K due tohigher ow stress at a lower temperature. Certainly, the formationof ferrite or bainite in austenite grains and the recrystallization ofaustenite grains during deformation will also decrease the size ofmartensite lath.

4. Secondary-deformation behavior of hot deformed22MnB5 steels

4.1. Flow curves

The evolution of true strain eld on the as-quenched 22MnB5steel specimen is shown in Fig. 8(a). Each panel in Fig. 8(a) indi-cates the strain eld at the instant shown below and it coversthe parallel section of 18 mm on the sub-size tensile specimenas shown in Fig. 4(a). For comparison, each panel was set to bethe same size to eliminate the elongation effect after deformation.Figs. 8(b), 9(a) and (b) show the strain distribution along the 18 mm-section at different instants of the as-quenched and hot deformed22MnB5 steel specimens with a deformation level of 10 mm at923 K and 693 K, respectively. As described above, the micro-structures of the specimens corresponding to Figs. 8, 9(a) and (b)are full martensite, 58% ferrite with 42% martensite and 86% bainitewith 14% martensite, respectively.

As shown in Fig. 8(a), the necking band nucleated at t = 246 sfor the as-quenched 22MnB5 specimen. Before t = 246 s, the defor-mation in the 18 mm-section was nearly uniform, after that, thedeformation tended to be localized with a center at x = 7.2 mm. Thenthe strain at x = 7.2 mm increased faster than any other locationuntil fracture.

To obtain elongation of sub-size tensile specimens, the gagelength should be determined rstly. Since the dimensions of sub-size tensile specimens are non-standard, the length of parallelsection is 18 mm, smaller than the gage length (25 mm or 50 mm) ofthe commonly used extensometers. Hence, the 25 mm- or 50 mm-extensometers cannot be employed to measure the elongation ofthe sub-size tensile specimens in this paper. The necking bandwidth (wnb) is dened as: beyond the two edges of necking band,the strain at any location does not increase as deformation pro-ceeds, and the distance between two edges is the width of neckingband (wnb), as shown in Fig. 8(b), where the wnb is 8.3 mm. Hence,



Fig. 8. (a) True strain eld evolution and (b) true strain distribution along the 18 mmsection of the as-quenched 22MnB5 steel specimen.

the half necking band width (whnb) is the distance from the centerof a necking band to one of its edge. However, when the center ofa necking band is close to a gripper end (x = 0 mm or x = 18 mm),one edge of the necking band may locate out of the 18 mm parallelsection, namely, the 18 mm-section does not cover the full neckingband, which is similar as the case that a specimen fractures outsideof an extensometer in a conventional tensile test. Therefore, thegage length cannot be set as 18 mm which covers the whole paral-lel section. To avoid the situation described above and minimize theratio of necking band width to the gage length, the gage length isset as 9 mm and one end of the gage section is the center of neckingband, and another end locates between the center of necking bandand the further gripper end, as illustrated in Figs. 8(b), 9(a) and (b).The ratio of half necking band width (whnb) to gage length (9 mm)is between 0.45 and 0.5 according to Figs. 8(b), 9(a) and (b).

Once the gage length is determined, the average true strain(11) based on Eq. (1) can be calculated. Then the correspondingengineering strain (e11) and engineering stress (S1) are respectivelycomputed by

e11 = exp 11 1 (4)

S1 =F

w0 t0(5)

The S1 vs. e11 and 1 vs.

31 curves of hot formed and as quenched

22MnB5 steel specimens are shown in Fig. 10(a) and (b), respec-tively. Since the ratio of the whnb to the gage length is much higherthan the tensile specimens with standard dimensions (Zavattieriet al., 2009), the post necking part (beyond the maximum engi-neering stress) on S1 vs. e

11 curves is more obvious on the S1 vs. e

11

curves. The elongation to failure (Ef), yield strength (y) and ulti-mate tensile strength (b) of the specimens are determined based

Fig. 9. True strain distribution along the 18 mm section of hot deformed 22MnB5specimens with a deformation level of 10 mm at (a) 923 K and (b) 693 K.

on Fig. 10(a). The Ef and b are important indicators to evaluatingthe capability of energy absorption. The work hardening exponent(n-value) is tted from Fig. 10(b), in the strain range from 0.01 tothe maximum strain on each curve.

4.2. Mechanical properties

Fig. 11 shows the effect of volume fraction of DIF or DIB onthe y and b of the hot deformed 22MnB5 steel specimens, andthe dashed lines show the y and b of the as-quenched 22MnB5steel specimens. It can be seen in Fig. 11 that both the y and bare signicantly dependent on the volume fraction of DIF or DIB.When 22MnB5 steel was deformed at 693 K to a small strain of0.041, there is no soft phase like bainite transformed from austeniteand the microstructure is nearly full martensite. As aforemen-tioned, the introduction of dislocation into austenite due to hotdeformation renes the martensite lath. Hence, its hardness andstrength are higher than the as-quenched 22MnB5 steel as shownin Figs. 7 and 11. With increasing applied strain level at 693 K or923 K, the volume fraction of DIB or DIF increases. The appearanceof soft phase like DIF or DIB decreases the y and b signicantly. Itis worth noting that the specimen composed of 42% martensite and58% DIF shows comparable b with but signicantly lower y thanthe one composed of 14% martensite and 86% DIB. The former wasdeformed at 923 K to a strain of 0.2 and the latter was deformed at693 K to the same strain level. Therefore, it can be concluded thatthe DIF in the hot formed 22MnB5 is softer than the DIB. Since theapplied strain level shows similar effect on the hardness (H) as that



Fig. 10. (a) Engineering stress (S1) vs. engineering strain (e11) curve and (b) true

stress (1) vs. true strain (31) curve of the hot formed and as-quenched 22MnB5

steel specimens.

Fig. 11. Effect of volume fraction of DIF or DIB on the yield strength (y) and ultimatetensile strength (b) of hot deformed 22MnB5 steel. The dashed lines indicate they and b of as-quenched 22MnB5 steel.

Fig. 12. Relationships between hardness (H) and yield strength (y) and ultimatetensile strength (b).

on the y and b, the relationships y H and b H are plotted inFig. 12, where the lled symbols indicate the y and b of the as-quenched 22MnB5 steel specimens. For all hot deformed specimenswith different combinations of microstructures, the b shows goodlinear correlation with the hardness. When the hardness is lowerthan 400HV5, the linear correlation between y H is not good,since the DIB shows lower hardness but high yield strength.

Fig. 13 shows the effect of volume fraction of DIF or DIB onthe elongation to failure (Ef) and the work hardening exponent(n-value) of hot deformed 22MnB5 steel specimens. The dashedlines show the Ef and n-value of the as-quenched 22MnB5 steelspecimens. The Ef and n-value of all hot deformed specimens arelarger than those of the as-quenched specimens. In addition, theEf and n-value increase with increasing the volume fraction of DIFor DIB in specimens. For the specimens containing DIF which aredeformed at 923 K, they exhibit higher n-value and larger Ef, hencebetter secondary-formability, than the specimens containing thesame amount of DIB which are deformed at 693 K. It is also indi-cated that the secondary-formability of DIF is better than that ofDIB.

The strengthductility product () reects the energy absorp-tion directly and it is computed as = b Ef. The dependence ofthe on the volume fraction of DIF or DIB is shown in Fig. 14. Withincreasing the volume fraction of DIF, though the b decreases, the

Fig. 13. Effect of applied true strain on the elongation to failure (Ef) and work hard-ening exponent (n-value) of hot deformed 22MnB5 steel specimens. The dashedlines indicate the Ef and n-value of the as-quenched 22MnB5 steel specimens.



Fig. 14. Dependent of strengthductility product () of hot deformed 22MnB5 steelspecimens on the volume fraction of DIF or DIB. The dashed lines indicate the ofthe as-quenched 22MnB5 steel specimens.

increases to 15.4 GPa% since the Ef increases signicantly. However,the situation is not the same for specimens composed of DIB andmartensite, which were deformed at 693 K. As mentioned before,when the applied true strain is 0.109, 8.7% DIB is generated, whichdecreases the b but increases the Ef compared with the specimenswith full martensite. Consequently, the still increases to 14.9 GPa%from 13.5 GPa%. When the volume fraction of DIB increases to 86%,the increase of Ef cannot compensate the decrease of b, and thenthe decreases to 13.6 GPa%.

To this point, the appearance of DIF decreases the b signi-cantly, and the b decreases with increasing the volume fraction ofDIF, but the Ef and increases, which indicates better secondary-formability and energy absorptive capability. A mixture of 58%DIF and 42% martensite exhibits a b of 1000 MPa class and rela-tive larger elongation and strengthductility product. The DIB inmicrostructure decreases the b and increases the Ef. The b islarger but the Ef is smaller than the specimens containing the samevolume fraction of DIF. Hot deformed 22MnB5 steel specimen witha small volume fraction (8.7%) of DIB shows a b of 1382 MPa anda of 14.9 GPa%, hence, a better combination of b and .

5. Conclusions

22MnB5 steel specimens were deformed at 923 K and 693 Kto three strain levels, and microstructures of the specimenswere examined. Sub-size tensile specimens were cut from thehot deformed 22MnB5 steel specimens, and their secondary-deformation behaviors were studied by using of DIC technique.Following conclusions are yielded:

(1) When the 22MnB5 steel is deformed at 923 K, the incubationperiod for ferrite transformation is shortened signicantly toless than 0.4 s, and ferrite transformation is induced even whena small strain of 0.044 is applied. The volume fraction of DIFincreases with increasing the applied strain level.

(2) As deformed at 693 K, DIB is observed when the applied strainlevel is 0.109 but not in the 22MnB5 steel specimen with a strainof 0.041. Therefore, the incubation period for DIBT is short-ened to less than 0.8 s, but it is longer than that of DIFT. Withincreasing the applied strain to 0.206, the volume fraction ofDIB increases to 86%, which is much larger than the volumefraction of DIF when the specimen is deformed at 923 K to the

comparable strain level. Consequently, the DIBT proceeds fasterthan the DIFT once it starts.

(3) The appearance of DIF decreases the b of hot deformed22MnB5 steel specimens signicantly, and a larger volumefraction of DIF leads to a lower b but larger Ef and . As aresult, increasing the volume fraction of DIF indicates a bettersecondary-formability and energy absorptive capability. Fromthis aspect, the mixture of martensite and DIF is acceptable tothe lower B-pillar mentioned in the introduction, e.g. a mixtureof 58% DIF and 42% martensite exhibits a b of 1000 MPa, a Efof 15.5% and a of 15.4 GPa%.

(4) For the hot deformed 22MnB5 steel specimens composed ofmartensite and DIB, the b decrease and Ef increases withincreasing the volume fraction of DIB. However, the b is largerand the Ef is smaller than that in the specimens containing thesame volume fraction of DIF. A smaller volume fraction of DIBleads to a better combination of b and , e.g. the specimenwith 8.7% DIB shows a b of 1382 MPa and a of 14.9 GPa%.Full martensite or a mixture of martensite with a small amountof DIB is appropriate to the upper B-pillar as mentioned in theintroduction.

(5) The ultimate tensile strength of all hot deformed and as-quenched 22MnB5 specimens shows a good linear correlationwith the hardness. Due to the DIB exhibiting higher yieldstrength but low hardness, the yield strength correlatesthe hardness well only when the hardness is higher than400HV5.

Acknowledgment

The authors would like to thank the nancial support from theproject under grant no. 51075307 of the National Natural ScienceFoundation of China.

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jmpt 2013 secondary deformation of hot stamping specimens

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dib deformation

applied strain

c t22mnb5 steel specimens

subsize tensile specimens

volume fraction of deformation

hot stamping technology

shanghai university

strain levels