Study of high strength pipeline steels with different microstructures

Materials Science and Engineering A 502 (2009) 38–44

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Study of high strength pipeline steels with different microstructures Wei Wang a,b , Yiyin Shan a,∗ , Ke Yang a a b

Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, PR China Graduate School of Chinese Academy of Sciences, Beijing 100049, PR China

a r t i c l e

i n f o

Article history: Received 15 April 2008 Received in revised form 21 September 2008 Accepted 28 October 2008 Keywords: Pipeline steels Polygonal ferrite (PF) Acicular ferrite (AF) Mechanical properties Electron backscatter diffraction (EBSD) analysis

a b s t r a c t In the present study, comparison studies on mechanical properties were made on a commercial X70 grade polygonal ferrite (PF) dominated pipeline steel and a laboratory developed X90 grade acicular ferrite (AF) dominated pipeline steel obtained by optimum thermo-mechanical controlled processing (TMCP). Charpy impact test results indicated that the upper shelf energy (USE) of the AF pipeline steel was a little bit higher, but its energy transition temperature (ETT) was extremely low, about −162 ◦ C, much lower than that of the PF pipeline steel of about −121 ◦ C. It was analyzed that higher strength and better toughness of the AF pipeline steel came from its finer grain size and higher density of dislocations and subboundaries, which could be also further proved from the electron backscatter diffraction (EBSD) analysis for its finer effective grain size (EGS) and higher content of low angle grain boundaries (LAGBs), smaller cleavage fracture unit measured from the fracture surface around fracture origin fractured at −196 ◦ C, and its more bent crack propagation path in the fracture. © 2008 Elsevier B.V. All rights reserved.

1. Introduction In modern pipeline steel technology, in order to improve the transportation efficiency, the target for development is to obtain excellent comprehensive properties through innovations and optimizations on chemical composition design, microstructure design and related thermo-mechanical controlled processing (TMCP), etc. Since high strength low alloy pipeline steels with acicular ferrite (AF) microstructure were developed in the early 1970s [1,2], much work has been done to study this novel microstructure [3–5]. Regarding to the pipeline steels produced by TMCP, the ferrite microstructures can be mainly divided into polygonal ferrite (PF), quasi-polygonal ferrite (QF) or massive ferrite, granular bainitic ferrite (GF) and bainitic ferrite (BF), in the orders of decreasing transformation temperatures and increasing cooling rates [5–7]. The PF, transformed at the highest temperatures and the slowest cooling rates, is nucleated as grain boundary allotriomorphs and grows into equiaxed grains with smooth and continuous boundaries under the scale of the light microscopy [6,7]. The QF and parent austenite involved in a massive transformation ideally have the same composition, the transformation can be accomplished by short range diffusion across transformation interfaces [8]. However, interstitial or substitutional atom partitioning may occur at the migrating interfaces [9], causing the QF grains with irregu-

∗ Corresponding author. Tel.: +86 24 23971517; fax: +86 24 23971517. E-mail address: [email protected] (Y.Y. Shan). 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.10.042

lar and jagged boundaries, containing high density of dislocations, subboundaries, and even martensite/austenite (M/A) constituents [10,11]. Compared with the BF, the GF forms at the same transformation temperature range but at slower cooling rates [7,12], containing high density of dislocations, separated by low angle grain boundaries (LAGBs) [5], with granular and equiaxed retained austenite or M/A islands dispersed in the ferritic matrix. The BF consists of many elongated ferritic lath bundles with high density of dislocations, separated by large angle grain boundaries; and the lath bundles are composed of many parallel ferritic laths, separated by LAGBs [6]. It has been well accepted [1,4] that the AF microstructure comes from a mixed diffusion and shear transformation mode during continuous cooling beginning at a temperature slightly above the upper bainite temperature transformation region, presenting as an assemblage of interwoven non-parallel ferrite laths with high density of tangled dislocations. However, there are still controversies and uncertainties on the metallographical identification and classification of the phases. Sometimes, the AF microstructure is also considered as bainite [13] or QF [14]. However, from our previous work, the microstructure of AF is complex consisting of QF, GF and BF with dispersed islands of second phases in the matrix [12]. For a pipeline steel, if the AF microstructure can be achieved, it will be with better properties combination, such as high strength, excellent toughness, good H2 S resistance [15] and superior fatigue behavior [16], than the PF–pearlite (P) microstructure. The excellent combination of properties has led to the application of this type of steels in the manufactures of large dimension pipes for gas and oil transportations in the low temperature area [17–19].

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In the present study, a commercial PF dominated pipeline steel and a laboratory developed AF dominated pipeline steel through optimum TMCP were investigated in order to comparatively study the strength and toughness behaviors for the pipeline steels with different microstructures. In addition, the effective grain size (EGS) and distribution of grain boundary misorientation of steels were studied using the electron backscatter diffraction (EBSD) analysis, to better understand the relation of the AF microstructure with the high strength and excellent low temperature toughness for the steel.

2. Experimental procedure Steels used in this study were a commercial PF pipeline steel (A steel) and a laboratory developed AF pipeline steel (B steel), and chemical compositions of two steels are listed in Table 1. The commercial PF pipeline steel was reheated at 1180 ◦ C for 230 min, and then rolled in two stages, rough-rolling and finishrolling. The rough-rolling stage for the steel was started for rolling at 1130 ◦ C from thickness of 230–40 mm in seven steps; then the finish-rolling stage was started at 1020 ◦ C and finished rolling at 840 ◦ C from thickness of 40–7 mm in another seven steps, and after that the steel was cooled to 600 ◦ C at 20 ◦ C/s for coiling and finally air cooled to room temperature. The laboratory developed AF pipeline steel, designed by the authors, was smelted in a 100-kg vacuum induction furnace. According to our previous work, the continuous cooling transformation (CCT) diagrams of the experimental steel measured by Formastor-F and Gleeble-3500 [20,21], the hot rolling experiment was carried out on a pilot rolling mill with twin rolls of 370 mm in diameter. The AF pipeline steel was reheated at 1150 ◦ C for 50 min, started rolling at 1050 ◦ C and finished rolling at 750 ◦ C from thickness of 70–7 mm in seven steps, then cooled to 600 ◦ C at 20 ◦ C/s in order to obtain the AF microstructure, finally held at 600 ◦ C for 1 h and furnace cooled to room temperature to simulate the coiling process. Specimens taken from the transverse cross-section planes of the steels were mechanically polished and etched by a 3 pct nital solution, and then microstructures were observed by an S-3400N scanning electron microscope (SEM). Specimens for tensile tests were cut from the middle of the rolled plates in the longitudinal direction and tests were conducted at room temperature at a crosshead speed of 5 mm/min on a SCHENCK-100KN servo-hydraulic machine, according to the standard of ASTM E8M-04. Charpy impact tests were performed at temperatures of 0, −40, −100, −120, −150, −180, −196 and −269 ◦ C using the subsize Charpy V-notch (CVN) specimens with size of 5 mm × 10 mm × 55 mm, which were machined from the middle of the rolled plates in the transverse direction, in accordance with the standard of ASTM E23-02. In order to reduce errors in data interpretation, a regression analysis on CVN energy vs. test temperature was done by a hyperbolic tangent curve fitting method [22]. Based on the data from the regression analysis, the energy transition temperature (ETT), which corresponds to the average value of upper shelf energy (USE) and lower shelf energy, was determined. In order to examine the cleavage fracture unit and the crack propagation path, the fracture surface and the cross-sectional area beneath the fracture surface of the CVN specimens fractured at −196 ◦ C

Fig. 1. SEM micrographs of the (a) A and (b) B steels.

were observed by SEM after the fracture surface was coated by nickel. To investigate microstructural factors determining mechanical properties, the EBSD analysis was conducted on the transverse cross-section planes of specimens on a Carl–Zeiss LEO-1550 Schottky field-emission SEM, equipped with a Tex-SEM Laboratories orientation-imaging microscope (OIM) system with a step size of 0.15 ␮m for A steel and 0.10 ␮m for B steel, respectively. A more detailed microstructure examination was performed using transmission electron microscopy (TEM). Thin foils for TEM observation were taken from the transverse cross-section planes of the steel plates, mechanically thinned from 300 to 50 ␮m, electropolished by a twin-jet electropolisher in an electrolyte consisting of 10 vol pct perchloric acid and 90 vol pct ethanol, and then were examined on a JEM-2010 high resolution TEM operating at 200 kV. 3. Results 3.1. Microstructure Microstructures of Steels A and B are shown in Fig. 1(a) and (b), respectively. Misorientation maps of the two steels are shown in Fig. 2(a) and (b), respectively. It could be found that different ferrites

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

C

Si

Mn

Mo

NbVTi

A B

0.080 0.025

0.20 0.24

1.45 1.56

– 0.32

0.095 0.058

P (ppm) 120 20

S (ppm) 50 6

O (ppm)

N (ppm)

Fe

20 43

50 62

Bal. Bal.

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Fig. 2. Misorientation maps of the (a) A and (b) B steels.

of the PF, QF, GF and BF have been marked in Figs. 1 and 2, according to the different microstructural characteristics. The TEM micrographs of PF, QF, GF and BF are shown in Fig. 3(a)–(d), respectively. The PF grains have equiaxed, smooth and continuous boundaries, and the QF grains irregular and jagged boundaries, containing high

density of dislocations, subboundaries, and even M/A components. The GF grains include granular and equiaxed retained austenite or M/A islands dispersed in the ferritic matrix, also containing high density of dislocations, whereas the BF grains consist of many elongated ferritic lath bundles, with elongated and dispersed M/A

Fig. 3. TEM micrographs of (a) PF, (b) QF, (c) GF and (d) BF microstructures.

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Table 2 Room temperature tensile properties of the experimental steels. Steel

Yield strength (MPa)

Tensile strength (MPa)

Elongation (pct)

Y/T ratio (pct)

A B

530 626

600 716

26.0 23.1

88 87

Fig. 5. Optical fractographs of CVN specimens fractured at −196 ◦ C for the (a) A and (b) B steels. The size of each fracture surface is 8 mm × 5 mm.

strength/tensile strength) ratio and ductility of the two steels are almost the same despite of different microstructures. 3.3. CVN impact toughness Curves of CVN energy vs. test temperature of the two steels are shown in Fig. 4. Table 3 presents the CVN energy data in detail, from −269 to 0 ◦ C, in which the USE and ETT were obtained from Fig. 4. The two steels showed good toughness with high USE about 130 J and ETT lower than −120 ◦ C through the subsize CVN impact tests. B steel has a little bit higher USE but much lower ETT than those of A steel, showing excellent low temperature toughness. The difference in ETT could be confirmed from observations on the fracture surfaces, the cleavage fracture surfaces around fracture origin and the cleavage crack propagation paths fractured at −196 ◦ C, as shown in Figs. 5–7, respectively. From Fig. 6, the cleavage fracture unit of B steel with much low ETT was measured to be about 3 ␮m, much smaller than that of A steel, about 8 ␮m. The cleavage propagation path of B steel was much more bent than that of A steel, as shown in Fig. 7, indicating that B steel should have much smaller fracture unit and EGS than those of A steel.

Fig. 4. CVN energy vs. test temperature of the (a) A and (b) B steels.

components. Thus, it can be seen that A steel is mainly composed of PF, a few of QF and pearlite (P), while B steel mostly consists of the AF, i.e., QF, GF and BF, and a few of PF. 3.2. Tensile properties Tensile properties of the two steels are listed in Table 2. The yield strength of A steel is over 483 MPa (70 ksi), satisfying the strength requirement of API X70 grade pipeline steels. While the yield strength of B steel reached the strength level of ISO X90 grade pipeline steels, 625 MPa (90 ksi). However, the Y/T (yield Table 3 CVNa impact data of the experimental steels. Steel

◦

A B a

ETT (◦ C)

Absorbed energy (J) ◦

◦

◦

◦

◦

◦

◦

0 C

−40 C

−100 C

−120 C

−150 C

−180 C

−196 C

−269 C

USE

122.2 134.7

132.0 122.7

108.0 137.0

68.0 120.0

1.8 74.7

– 62.0

1.5 8.6

– 4.1

127 133

Specimen size is 5 mm × 10 mm × 55 mm.

−121 −162

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Fig. 6. SEM fractographs of CVN specimens around the cleavage fracture origin fractured at −196 ◦ C for the (a) A and (b) B steels.

Fig. 8. Distributions of grain boundary misorientation of the (a) A and (b) B steels.

4. Discussion

Fig. 7. SEM Micrographs of the cross-sectional area beneath the cleavage fracture surfaces of the CVN specimens fractured at −196 ◦ C for the (a) A and (b) B steels coated with Ni, showing the crack propagation path.

Higher strength of the AF pipeline steel, B steel, was due to its finer grain size and sub-grains, which could be confirmed from the EBSD analysis. The content of low angle grain boundaries (LAGBs) with misorientation less than 15◦ of B steel was 42.6 pct, higher than that of A steel, 21.0 pct, as calculated from Fig. 8, indicating that the AF pipeline steel contains higher density of dislocations and subboundaries, which is beneficial to the strength, than the PF pipeline steel. Generally, high angle grain boundaries (HAGBs) with misorientation of 15◦ or higher obtained from EBSD can be used as a crystallographic domain parameter showing the effective grain size (EGS) [23,24]. The image-quality maps and misorientation maps of the two steels are shown in Fig. 9, showing the grains having misorientation larger than 15◦ . From EBSD results, calculated by the number fraction with edge grains excluded, it could be noted that the EGS of the AF pipeline steel was 1.32 ␮m, much smaller than that of the PF pipeline steel, 3.29 ␮m. The higher content of LAGBs and the smaller EGS should contribute to the higher strength of the AF pipeline steel than that of the PF pipeline steel. The high strength of AF pipeline steel might be with no relation to the distribution of LAGBs. For the AF pipeline steel, high dislocation densities and subboundaries, which were LAGBs, could increase the strength because of dislocation and subboundary strengthening mechanism. The small EGS should also contribute to the high strength due to the grain boundary strengthening mechanism.

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Fig. 9. (a) and (c) Image-quality maps and (b) and (d) misorientation maps of the (a) and (b) A and (c) and (d) B steels, showing grains having misorientation larger than 15◦ .

From optical fractographs of CVN specimens fractured at −196 ◦ C for the two steels, as shown in Fig. 5, it can be obviously seen that the fracture surface of the PF pipeline steel was somewhat flat, whereas that of the AF pipeline steel was more curving. This indicated that the cleavages in the AF pipeline steel should experience more bent paths and absorb more energy at low temperatures. From the cross-sectional area beneath the cleavage fracture surfaces of the two steels, as marked by black arrow in Fig. 7, it should be noted that the cleavage crack propagation direction deflected effectively when it met the PF, QF, GF or BF grain boundaries, but with little change when met the BF laths in the BF grains. There were many fine grains in the AF pipeline steel, as shown in Fig. 7, causing the crack propagation path more bent than that in the PF pipeline steel. As listed in Table 3, CVN absorbed energies of the AF pipeline steel at low temperatures were higher than those of the PF pipeline steel. The USE of the AF dominated pipeline steel was a little bit higher, but its ETT was much lower than that of the PF pipeline steel, due to the much smaller EGS and fracture unit measured from the cleavage fracture surface around the fracture origin. The cleavage propagation path of the AF pipeline steel was much more bent than that of

the PF pipeline steel, which also implies that the fracture unit and EGS of the AF pipeline steel should be much smaller than those of the PF pipeline steel. 5. Conclusions From comparison studies on a commercial PF dominated pipeline steel and a laboratory developed AF dominated pipeline steel, the following conclusions could be made. (1) The strength of the AF pipeline steel obtained by optimum TMCP, reaching the level of ISO X90 grade pipeline steels, was higher than that of the PF pipeline steel which met the strength requirement of API X70 grade pipeline steels. However, elongation and Y/T ratio of the two steels were almost the same. (2) The USE of the AF pipeline steel was a little bit higher, but its ETT was extremely low, about −162 ◦ C, much lower than that of the PF pipeline steel which was about −121 ◦ C. (3) Higher strength and better toughness of the AF pipeline steel resulted from its finer grain size and higher density of disloca-

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tions and subboundaries, which could be also further proved from the EBSD analysis for its finer EGS of about 1.32 ␮m and higher content of LAGBs of about 42.6 pct, smaller cleavage fracture unit measured from the fracture surface around fracture origin fractured at −196 ◦ C, and its more bent crack propagation path in the fracture. Acknowledgments This work was financially supported by a fund from National Natural Sciences Foundation of China (No. 50471106) and a Key Technology R&D Program of Liaoning Province (No. 2005221002). The authors thank Dr. Mingchun Zhao for his help with some experiments. References [1] Y.E. Smith, A.P. Coldren, R.L. Cryderman, Toward Improved Ductility and Toughness, Climax Molybdenum Company (Japan) Ltd., Tokyo, 1972, pp. 119–142. [2] A.P. Coldren, Y.E. Smith, R.L. Cryderman, Processing and Properties of Low Carbon Steel, American Institute of Mining, Metallurgical and Petroleum Engineers, New York, 1973, pp. 163–189. [3] T. Tanaka, Int. Met. Rev. 26 (1981) 185–212. [4] J.L. Lee, M.H. Hon, G.H. Cheng, J. Mater. Sci. 22 (1987) 2767–2777. [5] S.W. Thompson, D.J. Colvin, G. Krauss, Metall. Trans. 21A (1990) 1493–1507. [6] H.L. Li, S.W. Guo, Y.R. Feng, C.Y. Huo, H.F. Chai, An Illustrative Collection of Microstructure Micrographs of High Strength Micro-alloyed Steels, Beijing, Petroleum Industry Press, 2001, p. 3. [7] G. Krauss, S.W. Thompson, ISIJ Int. 35 (1995) 937–945.

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