Effects of microstructure and yield ratio on strain hardening and Bauschinger effect in two API X80 linepipe steels

Materials Science and Engineering A 551 (2012) 192–199

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Effects of microstructure and yield ratio on strain hardening and Bauschinger effect in two API X80 linepipe steels Seung Youb Han a , Seok Su Sohn a , Sang Yong Shin a , Jin-ho Bae b , Hyoung Seop Kim a , Sunghak Lee a,∗ a b

Center for Advanced Aerospace Materials, Pohang University of Science and Technology, Pohang 790-784, Republic of Korea Sheet Products & Process Research Group, Technical Research Laboratories, POSCO, Pohang 790-785, Republic of Korea

a r t i c l e

i n f o

Article history: Received 12 October 2011 Received in revised form 2 March 2012 Accepted 4 May 2012 Available online 15 May 2012 Keywords: API X80 linepipe steel Microstructure Bauschinger effect Strain hardening

a b s t r a c t In the present study, effects of microstructure and yield ratio on strain hardening and Bauschinger effect were investigated in two API X80 steel sheets fabricated by controlling the start cooling temperature. The steel whose start cooling temperature was lower had the higher fractions of granular bainite (GB) and martensite–austenite (MA) constituent and the lower fraction of acicular ferrite (AF), and showed the higher yield ratio. According to the results of the strain-reversal test composed of compressive and tensile tests at various compressive pre-strains, the reduction in yield strength of the steel having higher fractions of GB and MA was generally higher than that of the steel having lower fractions. This result could be explained by difference in density of mobile dislocations and by competing mechanisms between Bauschinger effect and strain hardening, which were susceptible to the minute change in pre-strain during the piping process. When the pre-strain was low, e.g., lower than 4%, the steels having low yield ratio and small Bauschinger effect were desirable to minimize the reduction in yield strength. © 2012 Elsevier B.V. All rights reserved.

1. Introduction Linepipe steels used for the long-range transportation of crude oil or natural gas generally require the high strength to endure the high-pressure. API X70 or X80 grade linepipe steels are widely used among linepipe steels, but recently developed high-strength linepipe steels such as API X120 steels start to be used [1–5]. Their hot-rolled coils are leveled (or uncoiled) into sheets, and then are shaped to be pipes. The American Gas Association (AGA) recommends the makers to ensure the strength of by pressure-resistant container tests, ring expansion tests, and tensile tests of sheets after flattened the pipe [6,7]. Most of linepipe makers evaluate the yield strength by the tensile test of sheets, which is an easy testing method. Since this test often gives the rough estimation of yield strength, it induces some difficulties in satisfying the exact yield strength standard required by the American Petroleum Institute (API). Unlike leveled or flattened sheets, pipes are subjected by tensile strains and compressive strains on exterior and interior wall areas, respectively. The increase in yield strength by strain hardening effect and the decrease in yield strength by Bauschinger effect work simultaneously in the same sheet, which gives differences in yield strength before and after the piping process [8–12].

∗ Corresponding author. Tel.: +82 54 279 2140; fax: +82 54 279 5887. E-mail address: [email protected] (S. Lee). 0921-5093/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.05.007

Zhonghua et al. [13] reported that the Bauschinger effect is the manifestation of non-isotropic strain hardening in an isotropic material such as a normalized steel. In other words, the Bauschinger effect is a result of the dislocation motion and structure generated during deformation, the result of which also causes the strain hardening during the initial deformation from the isotropic state. Strains applied on the sheet are also varied along the pipe thickness direction. Researches on yield strength variation with respect to compressive pre-strain applied on the sheet have been hardly conducted so far [14,15]. Particularly in the spiral piping process, in which the leveled sheet becomes twisted, the distribution of strains applied on the twisted sheet is complicated, and thus studies on difference in strengths before and after the spiral piping have been scarcely carried out. In order to achieve the stable construction and reliable maintenance of linepipes, variations in yield strength before and after the spiral piping should be precisely analyzed, and effects of microstructural factors, strain hardening, and Bauschinger effect on this yield strength variation should be essentially investigated. In the present study, two API X80 steels having different microstructures were fabricated by controlling cooling conditions, and their microstructures were analyzed. Tensile specimens taken from the leveled sheet were tested, and the yield strength and Bauschinger stress parameter were measured by the simulation test of piping, which was composed of compressive and tensile tests, with varying compressive pre-strain. From these results, variations in yield strength before and after the piping were analyzed,

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Table 1 Chemical compositions of the API X80 steels (wt.%). Steel

C

Si

Mn

Cr

Mo

Nb

Ti

API X80

0.05

0.3

1.8

0.3

0.2

0.1

0.02

and correlations between yield strength, Bauschinger parameter, and microstructural factor were investigated.

2. Experimental 2.1. API X80 linepipe steel sheet The steel used in this study was an API X80 grade linepipe steel having a minimum yield strength level of 552 MPa (80 ksi), and its chemical composition is shown in Table 1. An overall grain refinement effect was expected by rolling with a high rolling reduction ratio (over 80%) in the non-recrystallized region of austenite after austenitization at 1473 K (1200 ◦ C) [16,17]. This high rolling reduction ratio led to the increase in dislocation density and the subsequent grain refinement as dislocations acted as ferrite initiation sites during cooling [18–20]. The rolling of the steel was finished at the temperature of austenite region above Ar3 . After the finish rolling, the steel sheet was rapidly cooled from 1063 K (790 ◦ C) or 1113 K (840 ◦ C) down to 793 K (520 ◦ C) at a cooling rate of 20 K/s, and was coiled. The final width and thickness of the sheets were 1500 mm and 18.4 mm, respectively. For convenience, the steel sheets cooled from 1063 K (790 ◦ C) and 1113 K (840 ◦ C) are referred to as ‘A’ and ‘B’, respectively. A pipe (thickness: 18.4 mm; diameter: 960 mm) was fabricated by a spiral piping process after the hot-rolled coil was leveled by a leveler which is an uncoiling machine.

2.2. Microstructural analysis The steel sheet was polished and etched in a 2% nital solution, and microstructures of the longitudinal–transverse (L–T) plane (topside plane of the sheet) were observed by an optical microscope and a scanning electron microscope. Volume fractions of microstructures present in the steel sheet were measured by an image analyzer.

2.3. Tensile test and strain-reversal test for the piping Round specimens having a gage diameter of 6 mm and a gage length of 30 mm in the transverse direction were obtained from the 1/2 thickness location of the leveled sheet, and were tested at room temperature at a strain rate of 5 × 10−3 s−1 in accordance with the ASTM standard test method [21] by a universal testing machine of 100 kN capacity. The strain-reversal test composed of compressive and tensile tests for the piping was conducted on round specimens having a gage diameter of 6.35 mm and a gage length of 15 mm at room temperature at a strain rate of 5 × 10−3 s−1 by a universal testing machine [22]. The test specimen was prepared along the 30-deg-direction deviated from the rolling direction, in consideration of the direction of the spiral piping process. The tensile test or strain-reversal test was conducted three times for each datum point. The 0.2% off-set flow stress was determined to be the yield strength in the steels showing continuous yielding behavior, whereas the lower yield point was determined to be the yield strength in the steels showing discontinuous yielding behavior [8,16].

Fig. 1. SEM micrographs of the leveled sheet of the steels (a) A and (b) B, showing their L–T plane (topside plane of the sheet) microstructures (nital etched).

3. Results 3.1. Microstructure and hardness Fig. 1(a) and (b) are SEM microstructures of the leveled X80 steel sheets. Since both steel sheets were finish-rolled in the austenite region and water-cooled, they are composed mainly of acicular ferrite (AF) and granular bainite (GB), together with some secondary phases such as martensite–austenite (MA) constituents [1–5]. Average sizes of GBs and AFs are about 10 ␮m and 1–3 ␮m, respectively. Volume fractions of AF, GB, and MA were measured, and the results are shown in Table 2. The steel A whose start cooling temperature is lower than that of the steel B has the higher fractions of GB and MA and the lower fraction of AF than the steel B. The Vickers hardness was measured with intervals of 3 mm inward from the exterior surface of the sheet (before the piping) Table 2 Volume fractions of acicular ferrite (AF), granular bainite (GB), and secondary phases present in the API X80 steels (%). Steel

Acicular ferrite

Granular bainite

Secondary phasesa

A B

Bal. Bal.

23.8 15.6

3.6 0.7

a

Secondary phases mainly include martensite–austenite (MA) constituents.

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Sheet Exterior

Interior

Vickers Hardness (VHN)

280 270 260 250 240 230

A B

220 210

0

3

6

9

12

15

18

Distance from the Exterior Surface (mm)

Pipe Pipe Exterior

Interior

Vickers Hardness (VHN)

280 270 260 250 240 230

A B

220 210

0

3

6

9

12

15

18

Distance from the Exterior Surface (mm) Fig. 2. Vickers hardness vs distance from the exterior surface to interior surface of the (a) sheets and (b) pipes for the steels A and B.

3.2. Tensile properties of leveled sheets Room-temperature tensile stress–strain curves of the leveled sheets of the steels A and B are shown in Fig. 3, and tensile properties measured from these curves are shown in Table 3. Both the steels show the continuous yielding behavior, while the steel A shows the more continuous yielding behavior than the steel B. This subtle difference in the yielding behavior is related with the formation of mobile dislocations. In general, the yielding curves become smooth when hard phases are uniformly distributed to generate

mobile dislocations more easily [23,24]. That is, when hard MAs are finely distributed, mobile dislocations are actively generated along interfaces between MA and ferrite, which shows the smooth and continuous yielding behavior [23,24]. The yield strengths of the two steels are higher than 610 MPa, which meets the standard for the API X80 steel grade. The yield strength of the steel A is higher than that of the steel B, whereas the tensile strength is higher in the steel B than in the steel A. This is because the steel A has the higher fraction of GB having relatively low strain hardening capacity than the steel B [25–27]. Considering that the yield ratio (0.90) of the steel A is higher than that (0.81) of the steel B, the strain hardening capacity can be compared. The higher the yield ratio, the lower the strain hardening capacity becomes. Both the steels have the similar

800 700 600 Stress (MPa)

and the pipe under a 300 g load, and the results are shown in Fig. 2. The hardness values of the sheets of the two steels are almost same along the thickness direction within an error range of 10 VHN, which indicates that the steel sheets have the relatively homogeneous microstructures (Fig. 2(a)). On the other hand, the hardness of the pipe is higher than that of the sheet because of the strain hardening effect (Fig. 2(b)). The hardness is lowest at the center (9–12 mm distant from the exterior surface), and increases as the measurement location moves from the center to the surface because the steel sheet is subjected to the increased tensile or compressive strain. It is noticeable that the exterior wall shows the higher hardness by 10–20 VHN than the interior wall. This means that the exterior wall is subjected to the higher strain than the interior wall. In view of work hardening, the area whose hardness is lowest can be thought to be the weakly deformed location. The hardness of the steel A having higher volume fraction of hard GB and MA is higher by 10–20 VHN than that of the steel B. From the fact that the hardness is varied greatly with the thickness location, it can be predicted that the deformation amount is different at the sheet thickness location, which can affect mechanical properties of the pipe.

500 400

A

300

B

200 100 0

0

5

10

15

20

25

Strain (%) Fig. 3. Stress–strain curves obtained from the room-temperature tensile test of the steels (a) A and (b) B.

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Table 3 Room-temperature tensile properties of the API X80 steels. Steel

Yield strength (MPa)

Ultimate tensile strength (MPa)

Elongation (%)

Uniform elongation (%)

Yield ratio

A B

657 612

734 753

19 21

7.3 9.1

0.90 0.81

Table 4 Results of room-temperature strain-reversal tests composed of compressive and tensile tests for the piping of the API X80 steels. Steel

εpre (%)

 pre (MPa)

 y1 (MPa)

 y2 (MPa)

YS (MPa)

ˇ

A

0.5 1.1 2.4 3.8

637 660 715 746

620 627 630 612

477 438 452 477

143 189 178 135

0.25 0.34 0.37 0.36

B

0.5 1.1 2.4 3.8

596 637 692 754

578 593 584 602

485 472 452 467

93 121 132 135

0.19 0.26 0.35 0.38

* εpre ,  pre ,  y1 ,  y2 , YS, and ˇ are compressive pre-strain, compressive deformation stress, compressive yield stress, tensile yield stress, reduction in yield stress ( y1 −  y2 ), and Bauschinger stress parameter (( pre −  y2 )/ pre ), respectively.

elongation of 20%, and the uniform elongation is higher by 2% in the steel B than in the steel A. This difference in uniform elongation can be interpreted by the different strain hardening capacity caused by the higher GB fraction in the steel A than in the steel B.

3.3. Simulation test for the piping with varying compressive pre-strain Since the amount of deformation during the piping affects greatly on mechanical properties of the pipe, as aforementioned in Section 3.1, the strain-reversal test composed of compressive and tensile tests for the piping was conducted at various compressive pre-strains of 0.5, 1.1, 2.4, and 3.8%. Stress–strain curves of compressive and tensile tests of the steels A and B are shown in Fig. 4(a) through (h). The pre-strain portion of the stress vs strain data are plotted in the tensile domain as accumulated strains. In all the stress–strain curves, engineering stress and strain values are used, and the sum of pre-strain and tensile strain is determined to be 7% in order to clearly show the Bauschinger effect according to the pre-strain. In these curves, compressive pre-strain (εpre ), compressive deformation stress ( pre ), compressive yield stress ( y1 ), and tensile yield stress ( y2 ) are defined, as shown in Fig. 5. Bauschinger stress parameter (ˇ = ( pre −  y2 )/ pre ) [28] and reduction in yield stress (YS =  y1 −  y2 ) were calculated, and the results as well as εpre ,  pre ,  y1 , and  y2 are shown in Table 4. Though the steels A and B have different fractions of MA, they have basically similar microstructures. Thus, the difference in compressive and tensile yield stresses ( y1 and  y2 ) according to the difference in compressive pre-strain (εpre ) is similar at about 20 MPa in the two steels. As the pre-strain increases from 0.5% to 3.8%, the compressive deformation stress ( pre ) increases by about 110 MPa and 160 MPa in the steels A and B, respectively. This is because the steel A has the higher yield ratio as explained in Section 3.2, which implies the sharper strength increase in the steel B than in the steel A in a given strain range. Fig. 6 shows the reduction in yield stress (YS) data at various compressive pre-strains. In the steel A, the YS sharply increases from 143 MPa to 189 MPa, and then decreases to 135 MPa with increasing pre-strain, whereas it steadily increases from 93 MPa to 135 MPa in the steel B. When the pre-strain is low, the steel A has the larger YS than the steel B, but the YS is same in the steels A and B at the pre-strain of 3.8%.

4. Discussion 4.1. Correlation between Bauschinger effect and yield ratio The reduction in yield strength before and after the piping process is caused by the Bauschinger effect, which can be generally explained by the migration and pile-up of dislocations [29–35]. The Orowan model, in which the Bauschinger effect is explained on the basis of the dislocation theory [32], tells that the back stress increases as the density of mobile dislocations increases, the Bauschinger effect becomes evident, and the reduction in yield strength becomes large. Ormandy et al. [35] defined dislocations tangled by the dislocation pile-up as disordered dislocations, and reported that these dislocations hardly affected the Bauschinger effect when the reverse stress was applied. Thus, the Bauschinger effect is proportional to the density of mobile dislocations, but is not related with disordered dislocations. In order to compare the effect of deformation on the generation of mobile dislocations in the present steels, the expanded view near the yield point in tensile stress–strain curves in Fig. 3 is schematically shown in Fig. 7. The yield point phenomenon does not occur in the two steels because sufficient amounts of mobile dislocations are formed in the initial stage of deformation by the existence of MA. The elastic limit, where the elastic deformation ends and the yielding starts, is indicated as a point ‘1’, and the start point of plastic deformation, where the yielding ends and the plastic deformation starts, is indicated as a point ‘2’. The elastic limits (point 1) of the steels A and B are similar, whereas the difference between the elastic limit and start point of plastic deformation (point 2) of the steel A is twice larger than that of the steel B. This is because the mobile dislocation density in the steel A is more sharply increased than in the steel B in the early stage of deformation, and because disordered dislocations increase as the deformation is continuously proceeded, thereby resulting in the increased strength. The difference in generation of mobile dislocations between the steels A and B can also be compared in the continuous yielding behavior (Fig. 3). The curve of the steel A shows the more smooth and continuous yielding behavior than that of the steel B because the steel A has more mobile dislocations. Singh and Ramaswamy [12] studied that Bauschinger effect was related to yield ratio and uniform elongation, and that the material having larger Bauschinger effect showed the higher yield ratio and lower uniform elongation. The correlation between Bauschinger effect, yield ratio, and uniform elongation is also shown in the

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Fig. 4. Room-temperature compression/tension test curves of the leveled sheet of the steels (a) through (d) A and (e) through (h) B. The pre-strain portion of the stress vs strain data is plotted in the tensile domain as accumulated strains.

S.Y. Han et al. / Materials Science and Engineering A 551 (2012) 192–199

Fig. 5. Typical compression/tension test curve and definition of compressive prestrain (εpre ), compressive deformation stress ( pre ), compressive yield stress ( y1 ), and tensile yield stress ( y2 ).

197

present study, as shown in Tables 3 and 4. In the material having large Bauschinger effect, like the steel A, a number of mobile dislocations are generated and migrated in the early stage of deformation, which induces a great increase in strength. However, as the deformation is proceeded further, the dislocation pile-up increases, which inhibits the further increase in strength and results in high yield ratio and low uniform elongation. On the contrary, the material having small Bauschinger effect, like the steel B, has the low density of mobile dislocations generated in the early stage of deformation. The steel B generates few mobile dislocations, makes a seldom amount of dislocation pile-up, and does not have much increase in strength. As the deformation continues, the generation of mobile dislocations and their pile-up steadily take place so as to continuously increase the strength and uniform elongation, thereby resulting in low yield ratio. In Fig. 7, the curve of the steel B increases more sharply than that of the steel A after the start point of plastic deformation. This confirms the steady occurrence of the generation of mobile dislocations and their pile-up in the steel B. In order to decrease the reduction in yield strength after the piping, the Bauschinger effect needs to be reduced, and it is necessary to restrain the rapid generation of mobile dislocations during the early stage of deformation. This corresponds to inducing the steady strain hardening while the material is deformed, and the existing methodologies for lowering yield ratio which is proportionally related to the Bauschinger effect can be used. Swift [36] and Hollomon [37] expressed the yield ratio obtained from tensile tests as a function of proprietary material constants [38], and explained the microstructural requirements for low yield ratio as follows: N

Yield ratio = [ln (1 + b + ey )] ·

Fig. 6. Reduction in yield strength (YS) vs compressive pre-strain (εpre ).

exp (N − b) (1 + ey ) · N N

(1)

where b refers to material constant, ey to elongation at the time of yielding, and N to strain hardening exponent, among which b and N are controllable variables to decrease the yield ratio. Kim et al. [24,38] quantified the degree of strain hardening already present inside the material before plastic deformation, i.e., the material constant (b) which is a function of interior dislocation density, in terms of the following equation by using volume fractions of microstructures: b = (˛PF · XPF + ˛AF · XAF + ˛GB · XGB + ˛BF · XBF ) · exp(−k · XM ) + ˛M · XM

Fig. 7. Schematic diagram showing the expanded region near the yield point in tensile stress–strain curves of the steels A and B.

(2)

where ˛i refers to proprietary constant of microstructure, Xi to volume fraction of microstructure, k to constant related with size and distribution of secondary phases such as MA, ˛M to proprietary constant of MA, and XM to volume fraction of MA. k increases as MAs become finer and more homogeneously dispersed. Kim et al. [24,38] proposed the value of ˛ of polygonal ferrite (PF), acicular ferrite (AF), granular bainite (GB), bainitic ferrite (BF), and martensite (M) as 0.03, 0.015, 0.008, 0.003, and 0.0003, respectively. ˛ decreases in low-temperature transformed phases having high dislocation density, and the value of b tends to decrease as the fractions of microstructures having small ˛ increase. The effect of microstructure on strain hardening coefficient (N) is determined mainly by microstructural factors such as density of tangled dislocations, grain size, fraction of precipitates [24,38]. When put together Eqs. (1) and (2), the decrease in material constant (b) with increasing fraction of AF and the increase in strain hardening coefficient (N) by finely dispersing MAs are needed in order to decrease the yield ratio of the present steels composed of AF, GB, and MA.

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Fig. 8. Bauschinger stress parameter (ˇ = ( pre −  y2 )/ pre ) vs compressive prestrain (εpre ) on a logarithmic scale.

4.2. Correlation between Bauschinger effect and compressive pre-strain The difference in generation, migration, and pile-up of dislocations caused by the microstructural difference greatly affects the change in yield strength before and after the piping. Since this change in yield strength is susceptible to the strain applied on the sheet during the piping, it can be studied by varying the compressive pre-strain as shown in Fig. 6. The steel A shows the sharp reduction in yield strength after the low pre-strain of 1.1%, while the steel B shows the proportionally increasing behavior of yield strength with increasing pre-strain. In the steel A, when the prestrain is low as 0.5%, the number of mobile dislocations generated along boundaries between MA and ferrite explosively increases under the forward deformation, and the reverse deformation raises the back stress to greatly decrease the yield strength, thereby leading to the large YS. The steel B has the smaller YS than the steel A. This is because the steel A having higher fraction of MA than the steel B has more mobile dislocations generated by the early stage deformation, which results in the high back stress. When the prestrain is 1.1%, the YS of the steel A becomes largest, and that of the steel B slightly is larger than that of the pre-strain of 0.5%. At the pre-strain of 3.8%, the YS of both the steels A and B becomes same (135 MPa). The reason for the decreased YS in the steel A with increasing pre-strain is caused by the decrease in density of mobile dislocations, the increase in disordered dislocations, and the decrease in Bauschinger effect. In the steel B, as the pre-strain increases, the density of mobile dislocations steadily increases during the plastic deformation, and the YS continues to increase to become same to that of the steel A at the pre-strain of 3.8%. Fig. 8 shows the change in Bauschinger stress parameter according to compressive pre-strain on a logarithmic scale. When the pre-strain ranges 0.5–2.4%, the Bauschinger stress parameter of the steel A is higher than that of the Steel B, but becomes lower than that of the steel B at the pre-strain of 3.8%. This is associated with the fact that the abruptly decreased YS in the steel A at the pre-strain of 3.8% (Fig. 6). The increase in density of mobile dislocations is larger in the steel A than in the steel B at the pre-strain of 0.5–2.4%. At the increased pre-strain of 3.8%, the further generation of mobile dislocations is inhibited in the steel A, many dislocations are piled up to abruptly produce disordered dislocations, and the Bauschinger stress parameter decreases. On the contrary, the steel B has the lower density of mobile dislocations increased by deformation than the steel A, and has the lower Bauschinger stress parameter until

the pre-strain stays in the range of 0.5–2.4%. Even if the pre-strain increases to 3.8%, the increase of disordered dislocations is smaller in the steel B than in the steel A, and mobile dislocations are steadily generated, which makes the steel B have the higher Bauschinger stress parameter than the steel A. According to Lloyd et al. [39], the Bauschinger effect becomes saturated when the pre-strain is higher than 4–5%. The reverse of Bauschinger stress parameter shown in the steels A and B (Fig. 8) can be interpreted as a temporary phenomenon occurred in the steels of high yield ratio and high yield-point elongation. In the steels having high yield ratio, the Bauschinger effect becomes prominent in the early stage of deformation, and is saturated at low pre-strains. Thus, in the overall range of pre-strain, the steel having lower yield ratio has the lower Bauschinger effect, which minimizes the reduction in yield strength during the piping. In fact, since the deformation amount of API X80 steels during the spiral piping is lower than 4%, it is beneficial to use steels having low yield ratios at low pre-strains. The results of this study indicate that tensile and compressive stresses are applied on the exterior and interior walls of the pipe, respectively, and that the yield strength is determined by competing strain hardening and Bauschinger effect. The thickness center of the pipe has the lowest hardness, and the strain hardening increases as the deformation amount increases from the center to the surface. Since the pre-strain increases from the center to the surface, the Bauschinger effect becomes more prominent, and the yield strength is reduced. In order to minimize this reduction in yield strength after the piping, it is necessary to use the steels having small Bauschinger effect or the steels composed of microstructures of low yield ratios. In the high-strength steels such as API X80 linepipe steels composed of bainitic microstructures and secondary phases, the homogeneous distribution of fine MAs and the formation of a large amount of AF, which has the lower dislocation density than GB, are desirable for minimizing the Bauschinger effect.

5. Conclusions In the two API X80 steel sheets fabricated by controlling the start cooling temperature, effects of microstructure and yield ratio on strain hardening and Bauschinger effect were investigated.

(1) The API X80 steels consisted of acicular ferrite (AF), granular bainite (GB), and martensite–austenite (MA) constituent, and the fractions of these microstructures were varied with start cooling temperatures. The steel A having higher fraction of GB and MA showed the higher yield ratio than the steel B. (2) The hardness test results of the pipe indicated that mechanical properties could be affected by even a slight difference in strain. According to the strain-reversal test results at various compressive pre-strains, the reduction in yield strength of the steel A was higher than that of the steel B, but decreased with increasing pre-strain, and became equal to that of the steel B at the pre-strain of 3.8%. (3) The steel A having high yield ratio showed the high YS and Bauschinger effect due to the high density of mobile dislocations in the early stage of deformation. With further increase of pre-strain to 3.8%, the amount of immobile disordered dislocations increased, which led to the decrease in YS and Bauschinger stress parameter. The steel B having low yield ratio showed the similar YS with that of the steel A and the high Bauschinger stress parameter at the pre-strain of 3.8% due to the steady increase in mobile dislocations during the continuous deformation.

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(4) The Bauschinger effect and strain hardening effect during the piping process were susceptible to even the small change of pre-strain. When the pre-strain was low, e.g., lower than 4%, the steels having low yield ratios were desirable to minimize the reduction in yield strength by the Bauschinger effect. Acknowledgments This work was supported by POSCO under a contract No. 2010Y034 and the Ministry of Knowledge Economy under a contract number of M2007010007. The authors would like to thank for this support. References [1] C.W. Choi, H.J. Koh, S. Lee, Metall. Mater. Trans. 31A (2000) 2669–2674. [2] R. Deny, Pipeline Technology, Elsevier, Amsterdam, The Netherlands, 2000, pp. 1–116. [3] S.Y. Shin, K. Oh, S. Lee, N.J. Kim, Met. Mater. Int. 17 (2011) 29–40. [4] M.K. Gräf, H.G. Hillenbrand, C.J. Heckmann, K.A. Niederhoff, Proc. 13th Int. Offshore and Polar Engineering Conf., International Society of Offshore and Polar Engineers, Honolulu, 2003, pp. 97–104. [5] G. Mannucci, D. Harris, Fracture Properties of API X100 Gas Pipeline Steels, Final Report, European Commission, Brussels, Belgium, 2002, pp. 1–128. [6] Specification for High-Test Linepipe, Standard 5XL, 18th ed., American Petroleum Institute, New York, 1971. [7] Material Specification, Specification No. 2950-6-6, Canadian Arctic Gas Pipeline Limited, 1973. [8] G.E. Dieter, Mechanical Metallurgy, 3rd ed., McGraw-Hill, New York, 1986, pp. 236–237. [9] R. Sowerby, Y. Tomita, D.K. Uko, Mater. Sci. Eng. A41 (1979) 43–58. [10] G. Tither, M. Lavite, J. Metals 27 (1975) 15–23. [11] K.J. Pascoe, J. Strain Anal. 6 (1971) 167–180.

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