Materials Science and Engineering A

Materials Science and Engineering A 517 (2009) 369 374 Contents lists available at ScienceDirect Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea Change of tensile behavior of a high-strength low-alloy steel with tempering temperature Wei Yan a, Lin Zhu a, Wei Sha b, Yi-yin Shan a,,keyang a a Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China b Metals Research Group, School of Planning, Architecture and Civil Engineering, Queen s University of Belfast, Belfast BT7 1NN, UK article info abstract Article history: Received 18 February 2009 Received in revised form 30 March 2009 Accepted 31 March 2009 Keywords: High-strength low-alloy steel Strain-hardening exponent Tensile property Tempering The tensile behavior of a high-strength low-alloy (HSLA) steel after tempering at different temperatures from 200 to 700 C was investigated. The steel showed similar tensile behavior with almost no change in strength for tempering below 400 C. However, when the tempering temperature was increased from 500 to 650 C, the steel displayed not only a decrease in strength, but also gradually the upper yield points and lower strain-hardening ability. When the tempering temperature was increased up to 700 C, the steel exhibited a round roof shaped tensile curve and a high strain-hardening exponent. These interesting phenomena of tensile behavior are well explained in view of the interactions of mobile dislocations and dissolved C and N atoms and their effects on the strain-hardening exponent. 2009 Elsevier B.V. All rights reserved. 1. Introduction In order to reduce the materials cost and improve the transportation efficiency, high-strength low-alloy (HSLA) steels are widely employed in modern car manufacturing due to their excellent strength-toughness combination and weldability [1 5]. Simultaneously, thermomechanical control process (TMCP) substituting for the traditional rolling process has effectively promoted the development of HSLA steels. Therefore, in recent decades the line-pipe steels have been developed from grade X60 to the current X80 and X100 grades [6 10]. To achieve a combination of high strength and toughness, the microstructure of lower bainite or ferrite plus martensite has been designed for HSLA steels [7 10]. Hence, in spite of additions of alloying elements such as Mo and B to enhance lower bainite and martensite in these steels, rapid cooling after finishing rolling has been introduced. Generally, there are three cooling treatments after rolling [11,12]. The first one is direct quenching followed by tempering (DQT). The second one is accelerated continuous cooling (ACC), i.e., the as-rolled steel is cooled down to room temperature at a given cooling rate. The third one is interrupted accelerated cooling (IAC), i.e., the steel is subject to water-cooling in the phase transformation temperature region, and then air-cooled to room temperature. Self-tempering could happen during air-cooling because of the slower cooling rate in the component interior. A criti- Corresponding author. E-mail address: yyshan@imr.ac.cn (Y.-y. Shan). cal point in these three cooling treatments is that the rapidly cooled steel should be tempered, which is necessary to achieve a good strength/toughness combination. However, as the lower bainite microstructure is only recently introduced in HSLA steels, studies on the mechanical properties have been rarely related to the tensile behavior after tempering. Tensile tests carried out on structural steels may provide valuable information related to the microstructure. In a typical tensile curve of an annealed low carbon steel, the upper and lower yield points are well related to the interactions between dislocations and carbon as well as nitrogen atoms. This theory may not be able to explain the yield behavior in other metals with fcc or hcp lattice structures. However, a convincing explanation [13] should involve two aspects: (1) the density of mobile dislocations, (2) the rate of dislocation glide. The strain rate of metals is related to the Burger vector, b, the mobile dislocations density,, and the rate of dislocation glide, v, as given in Eq. (1). ε = b v (1) where the rate of dislocation glide v depends on the applied stress, as shown in Eq. (2). v = k v 0 ( 0 ) m (2) where is the shear stress in the sliding plane; 0 is the shear stress for dislocation glide of a unit speed; m is the stress exponent for dislocation glide, where the rate is thermally activated. Eq. (2) illustrates that a higher stress will produce a faster dislocation glide rate. 0921-5093/$ see front matter 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.03.085

370 W. Yan et al. / Materials Science and Engineering A 517 (2009) 369 374 Table 1 Chemical composition of the investigated steel, wt.%. C Mn Nb V Ti Mo Cr Cu Ni B Si S P Al O N 0.046 1.79 0.049 0.03 0.023 0.31 0.31 0.20 0.77 0.0014 0.10 0.002 0.0061 0.1 0.0004 0.003 In the as-recrystallized state and prior to tensile tests, the mobile dislocations density may be relatively low, so a high dislocation rate is necessary to meet the demand of plastic deformation. Correspondingly, a stress peak will occur at the upper yield point on the tensile curve. Once moving, the mobile dislocations density increases quickly. Hence, a lower dislocation rate may be possible to meet the demand of plastic deformation and the stress will correspondingly decrease. Consequently, the lower yield point appears on the tensile curve. When the moving dislocations will be blocked (or impeded) or re-pinned and the mobile dislocations density decreases, the same cycle described above happens again. This explanation is in principle reasonable and suitable for the plastic deformation of most metals. Another important aspect revealed in tensile tests (and the recorded stress strain curves) is the strain-hardening exponent. The true stress and true strain relations of homogeneous plastic deformation can be described by Eqs. (3) (5) [14]. s = ke n (3) s = (1 + ε) (4) e = ln(1 + ε) (5) where s is the true stress that can be calculated by Eq. (4); e is the true strain and can be calculated by Eq. (5); is the engineering stress; ε is the engineering strain; k is the hardening coefficient; and n is the strain-hardening exponent, which demonstrates the amount of work hardening at an incremental deformation strain. If n = 1, this shows that the material exhibits a linear work hardening characteristic. If n = 0, this indicates that the material has no strain-hardening ability and behaves ideally plastic. The present work is to investigate the tensile behavior of a HSLA steel with lower bainite microstructure after tempering at different temperatures. 2. Experimental The steel under investigation with the composition as represented in Table 1 was molten in a vacuum induction furnace. The as-forged blocks with the size of 80 mm 70 mm 60 mm were hot-rolled to 7 mm thick plates in the process as shown in Table 2 after soaking at 1150 C for 2 h. The samples for the tensile test to be carried out were cut from the plate perpendicular to the rolling direction which is believed to be the mechanically weak direction of as-rolled steels, and tempered at different temperatures for 30 min. Due to the anisotropy caused by hot-rolling, the rolling direction usually has better mechanical properties than the transverse direction. The tempering temperatures were chosen as 200, 300, 400, 500, 550, 575, 600, 625, 650, 700 C. Temperature around 600 C is usually believed to have the strongest influence on the mechanical properties of HSLA steels. Therefore, the tempering temperatures were frequently chosen to be around 600 C. Standard tensile specimens with a gauge diameter of 3 mm and 50 mm length (15 mm gauge length) were machined from the as-tempered samples. The tensile test was conducted at room temperature using a Zwick Z050 tensile machine equipped with an extensometer. The microstructures after tempering treatments were evaluated by using an optical microscope of Zeiss Axiovert 200 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 the uniform plastic deformation zone, the strain-hardening exponents n of the as-tempered samples of all investigated temperatures were calculated. 3. Results 3.1. Microstructure Fig. 1 shows that lower bainite microstructures with pancake shaped grains were obtained for the steel with the employed TMCP. Each pancake grain is around 30 m thick, 100 m wide and several hundred microns long. The bainite ferrite laths are characterized by the carbides distributed along the lath boundaries. It is worth noticing that the laths in each grain mainly show only one direction. With increase of tempering temperature, the appearance of ferrite lath boundaries became less clear. When the tempering temperature was high up to 700 C, small recrystallized grains appeared, as shown in Fig. 1e. 3.2. Tensile properties Some of the recorded tensile curves are shown in Fig. 2.Itcanbe seen that the stress strain curves of the steel treated below 400 C do not show obvious yield phenomenon. However, when the tempering temperature was 500 C and higher, the upper yield point appeared gradually. This phenomenon became obvious when the tempering temperature increased from 600 to 650 C. Another point worth considering is that the tensile curves of the tempered specimens at 500 650 C exhibited plateaus after their yield point. When the tempering temperature was as high as 700 C, the upper yield point disappeared and the tensile curve is characterized by a round roof shape, indicating good formability. In order to present the change in strength, both the yield strengths (YS) and ultimate tensile strengths (UTS) of the steel tempered at different temperatures are depicted in Fig. 3. This reveals that the yield strength does not decrease with increasing tempering temperatures below 650 C, even possessing a small peak at 600 C. Table 2 Procedure of the thermomechanical control process (TMCP). Primary hot-rolling Final rolling and subsequent cooling procedure Water-cooling rate ( C/s) Finishing temperature Reduction in thickness (mm) 70 60 60 42 42 27 27 22 22 17 17 11 11 7 50 Rolling temperature ( C) 1100 1050 900 850 800 250

W. Yan et al. / Materials Science and Engineering A 517 (2009) 369 374 371 Fig. 1. Microstructures of the as-processed steel in the as-rolled and tempered state. (a) as-rolled; (b) tempered at 200 C; (c) tempered at 400 C; (d) tempered at 600 C; (e) tempered at 700 C. Nevertheless, the tensile strength remarkably decreases to about 1000 MPa when tempered at 500 C and is close to the yield strength when the tempering temperature was up to 650 C. At the highest tempering temperature of 700 C, the yield strength decreases more steeply than the tensile strength. Therefore, with increasing temperature the differences between the tensile strengths and the yield strengths become smaller and smaller, and they almost vanished in the temperature range between 500 and 650 C. When the tempering temperature reaches 700 C, the difference in the stress values increased. The strain-hardening exponents calculated for differently heattreated samples according to Eqs. (4) (6) are presented in Fig. 4 and Table 3. The results show that the strain-hardening exponent decreased slightly when the tempering temperature was below 400 C, but was subjected to an abrupt decrease to as low as n = 0.021, when the tempering temperature was 650 C. At the tempering temperature of 700 C, the n value increased to n = 0.258. The strain-hardening exponent also displayed a small peak at the tempering temperature of 600 C. 4. Discussion 4.1. The upper yield point For tempering at 500 C, dislocations in the steel should have enough thermal activation energy to move and interact with each other. Thus, many dislocations with opposite Burgers vectors, i.e., the positive and negative dislocations, would interact and will annihilate. Consequently, the dislocations density will considerably decrease. The ferrite lath boundaries would also begin to disappear due to the movement of dislocations, as shown in Fig. 1c. Simultaneously, iron carbide precipitates should occur at this temperature treatment, since the dissolved C atoms have been activated, resulting in quite fast diffusion. Then, the precipitates act as strong pinning obstacles for dislocations glide. Therefore, the mobile dislocations density will be strongly reduced. On the other hand, the reduced dislocations density and disappearing lath boundaries in turn afford much space for moving dislocations. A higher stress will be needed in order to drive the dislocations off the pinning to

372 W. Yan et al. / Materials Science and Engineering A 517 (2009) 369 374 above-described mechanisms. With increasing tempering temperature, more ferrite laths disappeared and the dislocations density will be reduced, while the amount of precipitates increases. From this, it is deduced that the upper yield point will become more pronounced with increasing of the tempering temperature. The nucleation temperature for large amount of precipitates in HSLA steels was believed to be about 600 C [15], which is related with the upper yield point and the yield strength peak. However, it is noteworthy that there is only one peak in the tensile curves, not like the typical oscillating yield points (Portevin LeChatelier effect) on stress strain curves of mild steels, because the mobile dislocations density cannot be reduced effectively in this steel studied in this paper. However, tempering at 200 C will lead to that the moving dislocations in the steel are interacting with point defects such as vacancies and interstitial atoms. The mobile dislocations density may be much lower and will hardly increase. If the tempering temperature increases up to 300 and 400 C, the dislocations will be able to move and interact with each other, but the density might not be strongly reduced, which indicates the small change in the strength. Additionally, the ferrite lath boundaries are still stable, so the mean free path of gliding dislocations is restricted. Even if they will move, other immobile dislocations and the ferrite lath boundaries would immediately block them. Therefore, the amount of mobile dislocations will hardly increase. Hence, the upper yield points did not appear in the steel tempered at 300 and 400 C. When the tempering temperature increases up to 700 C, recrystallization occurs. Coarsening of precipitates in the steel takes place and the effective pinning of dislocations diminishes. Therefore, the yield strength of the steel shows a strong decrease. The mobile dislocations density may be quite low and dislocation glide is inhibited by a large number of newly formed grain boundaries in the recrystallized grains. Thus, the density of mobile dislocations could not significantly increase. Therefore, the upper yield points do not appear on tensile curves. Fig. 2. Representative tensile curves of the steel tempered at different temperature. move. Thus, the upper yield strength occurs. When the dislocations brake away from the pinning precipitates, the mobile dislocations density will increase again. Hence, the stress needed for gliding dislocations decreases. Consequently, it is reasonable that the upper yield point appears on the stress strain curves of the steel when the tempering temperature will be above 500 C, according to the 4.2. Strain-hardening exponent It should be noticed that the strain-hardening exponents were calculated for the uniform plastic deformation range. A lower strain-hardening exponent indicates that the material possesses a lower strain-hardening ability. Generally, the as-recrystallized metals and alloys and the severely strain hardened material show this feature: a low n value. The primary reason in the case of the asrecrystallized metals and alloys is that the matrix has no enough barriers for the moving dislocations, and the reason in the case of the severely strain hardened material is that the dislocations cannot move any more in the matrix. In either case, the yield strength is nearly equal to the tensile strength, for the described tempered steel 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 C were obviously not so strongly strain hardened, as described as the second reason above. No recrystallization did occur at these temperatures; see the primary reason as aforementioned. Therefore, Table 3 Strain-hardening exponent n of the steel tempered at 500 650 C. Tempering temperature ( C) Strain-hardening exponent n Fig. 3. Ultimate tensile strength (UTS) and yield stress of the HSLA steel in dependence on the tempering temperature. 500 0.036 550 0.036 575 0.029 600 0.076 625 0.029 650 0.021

W. Yan et al. / Materials Science and Engineering A 517 (2009) 369 374 373 Fig. 4. The illustrated log true stress vs. true strain curves of the investigated HSLA steel tempered at different temperatures revealing different stress exponents. there must be a third reason for the present steel. Lowering of n value for a continuously annealed cold-rolled steel has also been observed because of boron additions [16,17]. It was reported that the n value decreased significantly when the carbon content in the steel was less than 0.0020%. The reason for this is explained in terms of morphological changes of the carbide precipitates in matrix and at grain boundaries. In another work [18], the relationship between n and the yield strength of a titanium containing IF steel was analyzed. The results exhibited that the n value was controlled by dislocation-precipitate and dislocation-grain boundary interactions. It was found in the study as well that smaller amount of dissolved C and N atoms reduced the efficiency of grain boundaries to block the dislocations movement and thus a decrease in the n value. Therefore, tempering at 500 650 C can result in a depletion of the dissolved C and N atoms due to the precipitation of -carbides (FeC 2 ) or carbonitrides, which might be responsible for the low

374 W. Yan et al. / Materials Science and Engineering A 517 (2009) 369 374 5. Conclusions Fig. 5. Schematic representation of the strain-hardening exponent variation as a function of the dissolved C and N atoms contents and the quantity and size of the precipitates. strain-hardening exponent. In addition, the decrease in the dislocation density as well as the disappearing of the ferrite lath during tempering would contribute to the occurrence of the low n value. It should also be noticed that the precipitates could also pin gliding dislocations, and hence increase the strain-hardening rate. From the above-described results, it is concluded that the dissolved C and N atoms and their interactions with dislocations have a much stronger effect on the strain-hardening exponent than the precipitates. The n value will increase gradually with an increase in the amount of precipitates, but n decreases steeply with the decrease in the content of dissolved C and N atoms [17]. The formation of the precipitates would consume the dissolved C and N atoms, which can be demonstrated from the relation between the strain-hardening exponent and the content of the dissolved C and N atoms as shown in Fig. 5. The increasing amount of precipitates is consistent with the decreasing content of dissolved C and N atoms. Therefore, the n value shows a considerable change as illustrated by the bold line in Fig. 5, i.e., the n value reaches a minimum at 575 C. When the tempering temperature is above 600 C and increases up to 650 C, the precipitation kinetics was saturated and the growth rate of the precipitates decreased. Hence, the n value decreases again. Therefore, the n value showed a small peak at 600 C since the 600 C tempering could produce the largest number of fine precipitates. This explanation agrees quite well with the results shown in Table 3 and Fig. 3. When the tempering was performed at 700 C, the newly formed, small, recrystallized grains not only increased the amount of grain boundaries, but also behaved as the second phase, resulting in high strain-hardening ability of this steel. Lower bainite microstructure causing high strength of the HSLA steel has been achieved by proper TMCP. The yield strength of the steel shows a slight increase up to the tempering temperature of 550 C, then passing a maximum at 600 C. Beyond 650 C, the yield strength decreases steeply. The tensile strength decreases beyond 200 C and becomes close to the yield strength when the tempering temperature ranges between 500 and 650 C. There are no pronounced yielding points appearing on the tensile curves of the steel samples tempered below 400 C. However, the steel showed gradually developed upper yield points with relatively low strain-hardening exponents when the tempering temperature was increased from 500 to 650 C. The low strainhardening exponent reached a peak at 600 C, as presented in the schematic drawing. This can be explained in view of the interactions of the mobile dislocations with the dissolved C and N atoms and their effect on the strain-hardening exponent. The steel samples tempered at 700 C show round roof shaped tensile curves possessing high strain-hardening exponents due to the fine-grained microstructure in the as-recrystallized state. The governing mechanism is the strong interactions of dislocations and interstitials with grain boundaries and the finely-dispersed precipitates 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 Toughness, Climax Molybdenum Company (Japan) Ltd., Tokyo, 1972, pp. 119 142. [4] Y. Smith, A. Coldren, R. Cryderman, Met. Sci. 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