Ct

Wa

b

a

ARRA

KHSTT

1

tesotosX

mdambrp(acprb

0d

Materials Science and Engineering A 517 (2009) 369–374

Contents lists available at ScienceDirect

Materials Science and Engineering A

journa l homepage: www.e lsev ier .com/ locate /msea

hange of tensile behavior of a high-strength low-alloy steel withempering temperature

ei Yana, Lin Zhua, Wei Shab, Yi-yin Shana,∗, Ke Yanga

Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, ChinaMetals Research Group, School of Planning, Architecture and Civil Engineering, Queen’s University of Belfast, Belfast BT7 1NN, UK

r t i c l e i n f o

rticle history:eceived 18 February 2009

a b s t r a c t

The tensile behavior of a high-strength low-alloy (HSLA) steel after tempering at different temperaturesfrom 200 to 700 ◦C was investigated. The steel showed similar tensile behavior with almost no change in

eceived in revised form 30 March 2009ccepted 31 March 2009

eywords:igh-strength low-alloy steeltrain-hardening exponentensile property

strength for tempering below 400 ◦C. However, when the tempering temperature was increased from 500to 650 ◦C, the steel displayed not only a decrease in strength, but also gradually the upper yield points andlower strain-hardening ability. When the tempering temperature was increased up to 700 ◦C, the steelexhibited a “round roof” shaped tensile curve and a high strain-hardening exponent. These interestingphenomena of tensile behavior are well explained in view of the interactions of mobile dislocations anddissolved C and N atoms and their effects on the strain-hardening exponent.

empering

. Introduction

In order to reduce the materials cost and improve the transporta-ion efficiency, high-strength low-alloy (HSLA) steels are widelymployed in modern car manufacturing due to their excellenttrength-toughness combination and weldability [1–5]. Simultane-usly, thermomechanical control process (TMCP) substituting forhe traditional rolling process has effectively promoted the devel-pment of HSLA steels. Therefore, in recent decades the line-pipeteels have been developed from grade X60 to the current X80 and100 grades [6–10].

To achieve a combination of high strength and toughness, theicrostructure of lower bainite or ferrite plus martensite has been

esigned for HSLA steels [7–10]. Hence, in spite of additions oflloying elements such as Mo and B to enhance lower bainite andartensite in these steels, rapid cooling after finishing rolling has

een introduced. Generally, there are three cooling treatments afterolling [11,12]. The first one is direct quenching followed by tem-ering (DQT). The second one is accelerated continuous coolingACC), i.e., the as-rolled steel is cooled down to room temperaturet a given cooling rate. The third one is interrupted accelerated

ooling (IAC), i.e., the steel is subject to water-cooling in thehase transformation temperature region, and then air-cooled tooom temperature. Self-tempering could happen during air-coolingecause of the slower cooling rate in the component interior. A criti-

∗ Corresponding author.E-mail address: yyshan@imr.ac.cn (Y.-y. Shan).

921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2009.03.085

© 2009 Elsevier B.V. All rights reserved.

cal point in these three cooling treatments is that the rapidly cooledsteel should be tempered, which is necessary to achieve a goodstrength/toughness combination. However, as the lower bainitemicrostructure is only recently introduced in HSLA steels, studieson the mechanical properties have been rarely related to the tensilebehavior after tempering.

Tensile tests carried out on structural steels may provide valu-able information related to the microstructure. In a typical tensilecurve of an annealed low carbon steel, the upper and lower yieldpoints are well related to the interactions between dislocations andcarbon as well as nitrogen atoms. This theory may not be able toexplain the yield behavior in other metals with fcc or hcp latticestructures. However, a convincing explanation [13] should involvetwo aspects: (1) the density of mobile dislocations, (2) the rate ofdislocation glide. The strain rate of metals is related to the Burgervector, ‘�b’, the mobile dislocations density, ‘�’, and the rate of dislo-cation glide, v̄, as given in Eq. (1).

ε̇ = |�b|�v̄ (1)

where the rate of dislocation glide v̄ depends on the applied stress,as shown in Eq. (2).

v̄ = kv̄0

(�

�0

)m

(2)

where � is the shear stress in the sliding plane; �0 is the shearstress for dislocation glide of a unit speed; m is the stress expo-nent for dislocation glide, where the rate is thermally activated. Eq.(2) illustrates that a higher stress will produce a faster dislocationglide rate.

370 W. Yan et al. / Materials Science and Engineering A 517 (2009) 369–374

Table 1Chemical composition of the investigated steel, wt.%.

C Ni

0 0.7

drrtiswabsed

rTd

s

s

e

wieiddatp

st

2

sahatrdtrt2

TP

R

R

Mn Nb V Ti Mo Cr Cu

.046 1.79 0.049 0.03 0.023 0.31 0.31 0.20

In the as-recrystallized state and prior to tensile tests, the mobileislocations density may be relatively low, so a high dislocationate is necessary to meet the demand of plastic deformation. Cor-espondingly, a stress peak will occur at the upper yield point onhe tensile curve. Once moving, the mobile dislocations densityncreases quickly. Hence, a lower dislocation rate may be pos-ible to meet the demand of plastic deformation and the stressill correspondingly decrease. Consequently, the lower yield point

ppears on the tensile curve. When the moving dislocations will belocked (or impeded) or re-pinned and the mobile dislocations den-ity decreases, the same cycle described above happens again. Thisxplanation is in principle reasonable and suitable for the plasticeformation of most metals.

Another important aspect revealed in tensile tests (and theecorded stress–strain curves) is the strain-hardening exponent.he true stress and true strain relations of homogeneous plasticeformation can be described by Eqs. (3)–(5) [14].

= ken (3)

= �(1 + ε) (4)

= ln(1 + ε) (5)

here s is the true stress that can be calculated by Eq. (4); es the true strain and can be calculated by Eq. (5); � is thengineering stress; ε is the engineering strain; k is the harden-ng coefficient; and n is the strain-hardening exponent, whichemonstrates the amount of work hardening at an incrementaleformation strain. If n = 1, this shows that the material exhibitslinear work hardening characteristic. If n = 0, this indicates that

he material has no strain-hardening ability and behaves ideallylastic.

The present work is to investigate the tensile behavior of a HSLAteel with lower bainite microstructure after tempering at differentemperatures.

. Experimental

The steel under investigation with the composition as repre-ented in Table 1 was molten in a vacuum induction furnace. Thes-forged blocks with the size of 80 mm × 70 mm × 60 mm wereot-rolled to 7 mm thick plates in the process as shown in Table 2fter soaking at 1150 ◦C for 2 h. The samples for the tensile testo be carried out were cut from the plate perpendicular to theolling direction which is believed to be the mechanically weak

irection of as-rolled steels, and tempered at different tempera-ures for 30 min. Due to the anisotropy caused by hot-rolling, theolling direction usually has better mechanical properties than theransverse direction. The tempering temperatures were chosen as00, 300, 400, 500, 550, 575, 600, 625, 650, 700 ◦C. Temperature

able 2rocedure of the thermomechanical control process (TMCP).

Primary hot-rolling

eduction in thickness (mm) 70 → 60 60 → 42 42 → 27 27 → 22 22 → 1

olling temperature (◦C) 1100 1050 900

B Si S P Al O N

7 0.0014 0.10 0.002 0.0061 0.1 0.0004 0.003

around 600 ◦C is usually believed to have the strongest influenceon the mechanical properties of HSLA steels. Therefore, the tem-pering temperatures were frequently chosen to be around 600 ◦C.Standard tensile specimens with a gauge diameter of 3 mm and50 mm length (15 mm gauge length) were machined from theas-tempered samples. The tensile test was conducted at roomtemperature using a Zwick Z050 tensile machine equipped withan extensometer. The microstructures after tempering treatmentswere evaluated by using an optical microscope of Zeiss Axiovert200 MAT.

The double logarithmic form of Eq. (3) is shown in Eq. (6):

log s = log k + n log e (6)

According to Eqs. (4)–(6) and the recorded tensile test data of theuniform plastic deformation zone, the strain-hardening exponents‘n’ of the as-tempered samples of all investigated temperatureswere calculated.

3. Results

3.1. Microstructure

Fig. 1 shows that lower bainite microstructures with pancakeshaped grains were obtained for the steel with the employed TMCP.Each pancake grain is around 30 �m thick, 100 �m wide and severalhundred microns long. The bainite ferrite laths are characterized bythe carbides distributed along the lath boundaries. It is worth notic-ing that the laths in each grain mainly show only one direction.With increase of tempering temperature, the appearance of ferritelath boundaries became less clear. When the tempering tempera-ture was high up to 700 ◦C, small recrystallized grains appeared, asshown in Fig. 1e.

3.2. Tensile properties

Some of the recorded tensile curves are shown in Fig. 2. It can beseen that the stress–strain curves of the steel treated below 400 ◦Cdo not show obvious yield phenomenon. However, when the tem-pering temperature was 500 ◦C and higher, the upper yield pointappeared gradually. This phenomenon became obvious when thetempering temperature increased from 600 to 650 ◦C. Another pointworth considering is that the tensile curves of the tempered speci-mens at 500–650 ◦C exhibited plateaus after their yield point. Whenthe tempering temperature was as high as 700 ◦C, the upper yieldpoint disappeared and the tensile curve is characterized by a “roundroof” shape, indicating good formability.

In order to present the change in strength, both the yieldstrengths (YS) and ultimate tensile strengths (UTS) of the steel tem-pered at different temperatures are depicted in Fig. 3. This revealsthat the yield strength does not decrease with increasing temperingtemperatures below 650 ◦C, even possessing a small peak at 600 ◦C.

Final rolling and subsequentcooling procedure

Water-coolingrate (◦C/s)

Finishingtemperature

7 17 → 11 11 → 7 50 –

850 800 250

W. Yan et al. / Materials Science and Engineering A 517 (2009) 369–374 371

F . (a) at

N1wtspstii

tad4npTt

ig. 1. Microstructures of the as-processed steel in the as-rolled and tempered stateempered at 700 ◦C.

evertheless, the tensile strength remarkably decreases to about000 MPa when tempered at 500 ◦C and is close to the yield strengthhen the tempering temperature was up to 650 ◦C. At the highest

empering temperature of 700 ◦C, the yield strength decreases moreteeply than the tensile strength. Therefore, with increasing tem-erature the differences between the tensile strengths and the yieldtrengths become smaller and smaller, and they almost vanished inhe temperature range between 500 and 650 ◦C. When the temper-ng temperature reaches 700 ◦C, the difference in the stress valuesncreased.

The strain-hardening exponents calculated for differently heat-reated samples according to Eqs. (4)–(6) are presented in Fig. 4nd Table 3. The results show that the strain-hardening exponentecreased slightly when the tempering temperature was below

00 ◦C, but was subjected to an abrupt decrease to as low as= 0.021, when the tempering temperature was 650 ◦C. At the tem-ering temperature of 700 ◦C, the n value increased to n = 0.258.he strain-hardening exponent also displayed a small peak at theempering temperature of 600 ◦C.

s-rolled; (b) tempered at 200 ◦C; (c) tempered at 400 ◦C; (d) tempered at 600 ◦C; (e)

4. Discussion

4.1. The upper yield point

For tempering at 500 ◦C, dislocations in the steel should haveenough thermal activation energy to move and interact with eachother. Thus, many dislocations with opposite Burgers vectors, i.e.,the positive and negative dislocations, would interact and willannihilate. Consequently, the dislocations density will considerablydecrease. The ferrite lath boundaries would also begin to disappeardue to the movement of dislocations, as shown in Fig. 1c. Simulta-neously, iron carbide precipitates should occur at this temperaturetreatment, since the dissolved C atoms have been activated, result-ing in quite fast diffusion. Then, the precipitates act as strong

pinning obstacles for dislocations glide. Therefore, the mobile dis-locations density will be strongly reduced. On the other hand, thereduced dislocations density and disappearing lath boundaries inturn afford much space for moving dislocations. A higher stresswill be needed in order to drive the dislocations off the pinning to

372 W. Yan et al. / Materials Science and En

F

mbdlyt

Fd

ig. 2. Representative tensile curves of the steel tempered at different temperature.

ove. Thus, the upper yield strength occurs. When the dislocations

rake away from the pinning precipitates, the mobile dislocationsensity will increase again. Hence, the stress needed for gliding dis-

ocations decreases. Consequently, it is reasonable that the upperield point appears on the stress–strain curves of the steel whenhe tempering temperature will be above 500 ◦C, according to the

ig. 3. Ultimate tensile strength (UTS) and yield stress of the HSLA steel in depen-ence on the tempering temperature.

gineering A 517 (2009) 369–374

above-described mechanisms. With increasing tempering temper-ature, more ferrite laths disappeared and the dislocations densitywill be reduced, while the amount of precipitates increases. Fromthis, it is deduced that the upper yield point will become morepronounced with increasing of the tempering temperature. Thenucleation temperature for large amount of precipitates in HSLAsteels was believed to be about 600 ◦C [15], which is related withthe upper yield point and the yield strength peak. However, it isnoteworthy that there is only one peak in the tensile curves, not likethe typical oscillating yield points (Portevin–LeChatelier effect) onstress–strain curves of mild steels, because the mobile dislocationsdensity cannot be reduced effectively in this steel studied in thispaper.

However, tempering at 200 ◦C will lead to that the moving dis-locations in the steel are interacting with point defects such asvacancies and interstitial atoms. The mobile dislocations densitymay be much lower and will hardly increase. If the tempering tem-perature increases up to 300 and 400 ◦C, the dislocations will beable to move and interact with each other, but the density mightnot be strongly reduced, which indicates the small change in thestrength. Additionally, the ferrite lath boundaries are still stable,so the mean free path of gliding dislocations is restricted. Even ifthey will move, other immobile dislocations and the ferrite lathboundaries would immediately block them. Therefore, the amountof mobile dislocations will hardly increase. Hence, the upper yieldpoints did not appear in the steel tempered at 300 and 400 ◦C.

When the tempering temperature increases up to 700 ◦C, recrys-tallization occurs. Coarsening of precipitates in the steel takes placeand the effective pinning of dislocations diminishes. Therefore, theyield strength of the steel shows a strong decrease. The mobiledislocations density may be quite low and dislocation glide is inhib-ited by a large number of newly formed grain boundaries in therecrystallized grains. Thus, the density of mobile dislocations couldnot significantly increase. Therefore, the upper yield points do notappear on tensile curves.

4.2. Strain-hardening exponent

It should be noticed that the strain-hardening exponents werecalculated for the uniform plastic deformation range. A lowerstrain-hardening exponent indicates that the material possesses alower strain-hardening ability. Generally, the as-recrystallized met-als and alloys and the severely strain hardened material show thisfeature: a low n value. The primary reason in the case of the as-recrystallized metals and alloys is that the matrix has no enoughbarriers for the moving dislocations, and the reason in the case ofthe severely strain hardened material is that the dislocations can-not move any more in the matrix. In either case, the yield strengthis nearly equal to the tensile strength, for the described temperedsteel at 500–650 ◦C as shown in Fig. 3.

The phenomenon of low strain-hardening exponents of HSLA

steels was unexpected. The steel samples tempered at 500–650 ◦Cwere obviously not so strongly strain hardened, as described as thesecond reason above. No recrystallization did occur at these tem-peratures; see the primary reason as aforementioned. Therefore,

Table 3Strain-hardening exponent ‘n’ of the steel tempered at 500–650 ◦C.

Tempering temperature (◦C) Strain-hardening exponent ‘n’

500 0.036550 0.036575 0.029600 0.076625 0.029650 0.021

W. Yan et al. / Materials Science and Engineering A 517 (2009) 369–374 373

SLA s

tvottims

Fig. 4. The illustrated log true stress vs. true strain curves of the investigated H

here must be a third reason for the present steel. Lowering of nalue for a continuously annealed cold-rolled steel has also beenbserved because of boron additions [16,17]. It was reported that

he n value decreased significantly when the carbon content inhe steel was less than 0.0020%. The reason for this is explainedn terms of morphological changes of the carbide precipitates in

atrix and at grain boundaries. In another work [18], the relation-hip between n and the yield strength of a titanium containing IF

teel tempered at different temperatures revealing different stress exponents.

steel was analyzed. The results exhibited that the n value was con-trolled by dislocation-precipitate and dislocation-grain boundaryinteractions. It was found in the study as well that smaller amount of

dissolved C and N atoms reduced the efficiency of grain boundariesto block the dislocations movement and thus a decrease in the nvalue. Therefore, tempering at 500–650 ◦C can result in a depletionof the dissolved C and N atoms due to the precipitation of �-carbides(FeC2) or carbonitrides, which might be responsible for the low

374 W. Yan et al. / Materials Science and En

Ffp

sct

itasiaipbeitvFppctpeF

foi

[

[[[

[

ig. 5. Schematic representation of the strain-hardening exponent variation as aunction of the dissolved C and N atoms contents and the quantity and size of therecipitates.

train-hardening exponent. In addition, the decrease in the dislo-ation density as well as the disappearing of the ferrite lath duringempering would contribute to the occurrence of the low n value.

It should also be noticed that the precipitates could also pin glid-ng dislocations, and hence increase the strain-hardening rate. Fromhe above-described results, it is concluded that the dissolved Cnd N atoms and their interactions with dislocations have a muchtronger effect on the strain-hardening exponent than the precip-tates. The n value will increase gradually with an increase in themount of precipitates, but n decreases steeply with the decreasen the content of dissolved C and N atoms [17]. The formation of therecipitates would consume the dissolved C and N atoms, which cane demonstrated from the relation between the strain-hardeningxponent and the content of the dissolved C and N atoms as shownn Fig. 5. The increasing amount of precipitates is consistent withhe decreasing content of dissolved C and N atoms. Therefore, the nalue shows a considerable change as illustrated by the bold line inig. 5, i.e., the n value reaches a minimum at 575 ◦C. When the tem-ering temperature is above 600 ◦C and increases up to 650 ◦C, therecipitation kinetics was saturated and the growth rate of the pre-ipitates decreased. Hence, the n value decreases again. Therefore,he n value showed a small peak at 600 ◦C since the 600 ◦C tem-ering could produce the largest number of fine precipitates. Thisxplanation agrees quite well with the results shown in Table 3 and

ig. 3.

When the tempering was performed at 700 ◦C, the newlyormed, small, recrystallized grains not only increased the amountf grain boundaries, but also behaved as the second phase, resultingn high strain-hardening ability of this steel.

[[[

[

gineering A 517 (2009) 369–374

5. Conclusions

Lower bainite microstructure causing high strength of the HSLAsteel has been achieved by proper TMCP. The yield strength of thesteel shows a slight increase up to the tempering temperature of550 ◦C, then passing a maximum at 600 ◦C. Beyond 650 ◦C, the yieldstrength decreases steeply. The tensile strength decreases beyond200 ◦C and becomes close to the yield strength when the temperingtemperature ranges between 500 and 650 ◦C.

There are no pronounced yielding points appearing on the ten-sile curves of the steel samples tempered below 400 ◦C. However,the steel showed gradually developed upper yield points withrelatively low strain-hardening exponents when the temperingtemperature was increased from 500 to 650 ◦C. The low strain-hardening exponent reached a peak at 600 ◦C, as presented in theschematic drawing. This can be explained in view of the interac-tions of the mobile dislocations with the dissolved C and N atomsand their effect on the strain-hardening exponent.

The steel samples tempered at 700 ◦C show “round roof” shapedtensile curves possessing high strain-hardening exponents due tothe fine-grained microstructure in the as-recrystallized state. Thegoverning mechanism is the strong interactions of dislocations andinterstitials with grain boundaries and the finely-dispersed precip-itates of carbides and carbonitrides.

References

[1] W. Wang, Y.Y. Shan, K. Yang, Mater. Sci. Eng.: A 502 (2009) 38–44.[2] Y.M. Kim, S.K. Kim, Y.J. Lim, N.J. Kim, ISIJ Int. 42 (2002) 1571–1577.[3] Y.E. Smith, A.P. Coldren, R.L. Cryderman, Toward Improved Ductility and Tough-

ness, Climax Molybdenum Company (Japan) Ltd., Tokyo, 1972, pp. 119–142.[4] Y. Smith, A. Coldren, R. Cryderman, Met. Sci. Heat Treat. 18 (1976) 59–65.[5] M.C. Zhao, Y.Y. Shan, F.R. Xiao, K. Yang, Y.H. Li, Mater. Lett. 57 (2002) 141–145.[6] H.L. Li, Chin. Mech. Eng. 13 (2001) 349–352.[7] J.Y. Koo, M.J. Luton, N.V. Bangaru, R.A. Petkovic, Int. J. Offshore Polar Eng. 14

(2004) 2–10.[8] H. Asahi, T. Hara, M. Sugiyama, N. Maruyama, Y. Terada, H. Tamehiro, K. Koyama,

S. Ohkita, H. Morimoto, Int. J. Offshore Polar. Eng. 14 (2004) 11–17.[9] D.P. Fairchild, M.L. Macia, N.V. Bangaru, J.Y. Koo, Int. J. Offshore Polar Eng. 14

(2004) 18–28.10] H. Asahi, E. Tsuru, T. Hara, M. Sugiyama, Y. Terada, H. Shinada, S. Ohkita, H.

Morimoto, N. Doi, M. Murata, H. Miyazaki, E. Yamashita, T. Yoshida, N. Ayukawa,H. Akasaki, M.L. Macia, C.W. Petersen, J.Y. Koo, Int. J. Offshore Polar Eng. 14 (2004)36–41.

11] K.W. Huang, Steel Rolling 21 (2004) 55–58.12] C. Oouchi, T. Ookita, S. Yamomoto, Trans. Iron Steel Inst. Jpn. 22 (1982) 608.13] D.L. Shu, Metallic Mechanical Property, 2nd ed., China Machine Press, Beijing,

1997, pp. 12–14.14] J.H. Holloman, Trans. Met. Soc. AIME 162 (1945) 268–290.

15] Q.L. Yong, L. Zheng, Acta Metall. Sin. A 20 (1984) 9.16] Y. Funakawa, T. Inazumi, Y. Hosoya, ISIJ Int. 41 (2001) 900–907.17] Y. Hosoya, H. Kobayashi, T. Shimomura, K. Matsudo, K. Kurihara, Conf. Proc. On

Technology of Continuous Annealed Cold Rolled Sheet Steel, TMS-AIME, 1984,pp. 61–77.

18] P. Antoine, S. Vandeputte, J.B. Vogt, ISIJ Int. 45 (2005) 399–404.