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Microstructure and high strength–toughness combination of a new 700 MPa Nb-microalloyed pipeline steel
Materials Science and Engineering A 478 (2008) 26–37
Microstructure and high strength–toughness combination of a new 700 MPa Nb-microalloyed pipeline steel S. Shanmugam a , N.K. Ramisetti a , R.D.K. Misra a,∗ , J. Hartmann b , S.G. Jansto c a
Center for Structural and Functional Materials and Department of Chemical Engineering, University of Louisiana at Lafayette, Lafayette, LA 70504-4130, USA b Arcelor Mittal, USA Research & Development Center, 3001 E. Columbus Drive, East Chicago, IN 46312, USA c Reference Metals, 1000 Old Pond Road, Bridgeville, PA 15017, USA Received 21 March 2007; received in revised form 14 May 2007; accepted 1 June 2007
Abstract A new ultrahigh strength niobium-microalloyed pipeline steel of yield strength ∼700 MPa has been processed. The Charpy impact toughness at 0 ◦ C was 27 J and tensile elongation was 16%. The ultrahigh strength is derived from the cumulative combination of fine grain size, solid solution strengthening with additional interstitial hardening, precipitation hardening from carbides, dislocation hardening, and mixed microstructure. The microstructure was characterized by polygonal ferrite, upper bainite, degenerated pearlite, and martensite–austenite (MA) constituents. The microstructure of weld and heat-affected zone (HAZ) was similar to the base metal such that the hardness is retained in the weld region implying insignificant softening in the weld zone. Niobium and titanium precipitates of different morphology and size range evolved during thermomechanical processing and include rectangular (∼500 nm), irregular (∼240–500 nm), cuboidal/spherical (∼125–300 nm), and very fine (<10 nm). They were generally MC type of carbides. An important aspect of the developed steel is significantly lean chemistry. © 2007 Elsevier B.V. All rights reserved. Keywords: Pipeline steel; Niobium-microalloyed; Microstructure
1. Introduction Currently there is a demand to transport crude oil and gas by pipeline at a higher operating pressure to increase the capacity. This requires the use of ultrahigh strength steels. Increasing the strength of the pipeline steel enables a significant reduction in wall thickness with consequent reduction in weight. Thus, a major goal within the steel industry is to develop ultrahigh strength microalloyed pipeline steels (∼700–800 MPa). It is, however, important that the increase in the yield strength is not accompanied by a decrease in fracture toughness and formability because a decrease in toughness will encourage stress-induced cracking, and reduced formability will cause difficulties in forming (e.g. pipe-bowing). Thus, high strength in association with high toughness and formability are important requirements of the pipeline industry [1–3]. High strength–toughness combination is essential for pipeline steels for transporting natural gas and crude oil over
∗
Corresponding author. E-mail address: [email protected] (R.D.K. Misra).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.06.003
a long distance under high pressure [2]. Other characteristics that are required include resistance to hydrogen-induced blister cracking in sour service environment [4,5], stress corrosion cracking resistance, especially in H2 S environment [6–8], and fatigue resistance [9,10]. It is important that the microstructure of the pipeline steel provides combination of aforementioned properties. The evolution of microstructure depends on alloy chemistry and thermomechanical processing. Alloying additions such as Mn, Nb, V, Ti, Mo, Ni, Cr and Cu are commonly employed in pipeline steels to obtain the desired microstructure and mechanical properties [5,6,11–13]. However, judicious selection of alloying elements is necessary to obtain beneficial effect on mechanical properties with reduced alloy cost. For instance, the numbers of alloying elements are reduced to achieve lower carbon equivalent (CE) to ensure good field weldability [2,3]. On the other hand alloy additions such as Cr, Cu and Ni are added to obtain strength in severe corrosive environment [2,5]. Controlled thermomechanical processing is considered as the primary route for the development of API grade pipeline steels because it provides the desirable and fine-grained microstructure. Furthermore, it allows high strength–toughness
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combination to be achieved with accelerated cooling [3]. The ultimate microstructure strongly depends on the processing parameters such as reheating temperature, percentage reduction, deformation temperature, cooling rate, and coiling temperature [14]. This is because fine austenite grains, substructure and dislocations in austenite effectively promote the transformation of fine ferrite. The primary grain refinement in controlled rolling is achieved through recrystallization of austenite during deformation and through the use of microalloying elements, such as Nb that precipitates as fine carbides and inhibits grain growth [14]. In thermomechanically processed pipeline steels [1,7–9,11,15–20], different combinations of microstructures are obtained. It is believed that acicular ferrite microstructure with uniform distribution of martensite/austenite (M/A) islands as a second phase, provides the desired mechanical properties [1,11]. Besides, acicular ferrite the preferred microstructure of pipeline steel is characterized by microstructural constituents that include bainite and quasi-polygonal ferrite or massive ferrite [14,21]. This study describes the microstructure and precipitation behavior with the aim to relate to the high strength–toughness combination of a new industrially processed
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700 MPa Nb-microalloyed pipeline steel with a lean chemistry. The precipitates were characterized in terms of morphology, size, chemistry, and crystallography. 2. Experimental The Nb-microalloyed pipeline steel described here was continuously cast and hot rolled to a minimum of ∼7.6 mm gage and subsequently fabricated to ∼150 mm diameter pipes. The alloy design and nominal chemical composition of the steel was based on our recently patented hot rolled steel [22] and consisted of Fe–(0.04–0.06)C–(0.035–0.05)Ti–(0.08–0.09)Nb–(0.3–0.4)Cr. The sulfur content was less than 0.005 wt.%. The steel was hot rolled using controlled rolling practice on Mittal Steel Indiana Harbor 84 in. hot strip mill. The thermomechanical processing details are not described because of proprietary reasons. Standard tensile tests were conducted at room temperature on longitudinal specimens machined according to ASTM E8 specification using computerized tensile testing system. The initial crosshead speed was 4.2 × 10−2 mm/s up to a strain of ∼5%,
Fig. 1. Low-magnification (a) and high-magnification (b)–(d) light micrographs showing general microstructure of base, HAZ, and weld metal zones.
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then the speed was increased to 4.2 × 10−1 mm/s. Impact toughness was measured using standard Charpy v-notch impact test (ASTM 23) at 0 ◦ C. Transmission electron microscopy was carried out on samples cut from the fabricated steel pipe and relate to regions corresponding base, HAZ, and weld metal zones. Thin foils were prepared by cutting thin wafers from the steel samples using precision saw (Buehler Isomet 1000) and grinding to ∼100 m in thickness. Three-millimeter discs were punched from the wafers and electropolished using a solution of 10% perchloric acid in acetic acid electrolyte. Carbon extraction replicas were also prepared for characterization of precipitates. The surface of the polished specimens was etched with 2% nital and carbon was evaporated onto the etched surface. Finally, the surface was scored to ∼3 mm squares and the sample etched first with 10% nital and then with 2% nital. Subsequently, the extracted replicas were rinsed with distilled water and placed on the copper grid and dried. Foils and carbon extraction replicas were examined with a Hitachi 7600 TEM operated at 100 kV.
Fig. 2. Knoop hardness distribution across the weld zones on a traverse perpendicular to weld zone.
Fig. 3. Bright field TEM micrographs illustrating the general microstructure of the base metal. (a) Non-equiaxed or quasi-polygonal ferrite, (b) quasi-polygonal ferrite region containing high density of dislocations and dislocation substructure, (c) quasi-polygonal ferrite region containing grain boundary carbides, and (d) degenerated pearlite or upper bainite.
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3. Results 3.1. Mechanical properties The yield and tensile strength of the hot end of the coil was ∼735 MPa (105 ksi) and ∼791 MPa (113 ksi), respectively. The Charpy v-notch impact toughness of Nb-microalloyed pipeline steel at 0 ◦ C was 27 J and the tensile elongation was 16%. It may be noted that the listed mechanical properties are the minimum values that were obtained on any coil. 3.2. Microstructure and hardness of base, HAZ, and weld metal zones Representative light micrographs of base, HAZ, and weld metal are presented in Fig. 1a–d. The microstructure was predominantly fine-grained ferrite in all the three zones. As expected, the base metal (as-hot rolled) contained some elongated ferrite grains (Fig. 1b). A hardness profile encompassing base, HAZ, and weld metal zones is presented in Fig. 2. The hardness–distance profile indicates that the weld zone and HAZ
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did not experience significant softening, consistent with the microstructure of the three zones, as described below. Representative low-magnification bright field TEM micrographs of the base metal are presented in Fig. 3a–d. In general, the microstructure was of mixed type consisting of ferrite, upper bainite, and degenerated pearlite. The ferrite grains were of nonequiaxed (Fig. 3a) and were characterized by sub-boundaries and dislocation substructures (Fig. 3b), and carbides at the ferrite grain boundaries (Fig. 3c). A region of the base metal characterized by upper bainite/degenerated pearlite is presented in Fig. 3d. Fig. 4a and b are high-magnification bright field TEM micrographs of ferrite grains that contained high density dislocations at the grain boundaries (Fig. 4a), dislocations in the body of the grain (Fig. 4b), and dislocation substructures (Fig. 4c). Representative bright field TEM micrographs of HAZ are presented in Fig. 5a–d. The ferrite microstructure was similar to the base metal and contained high density of dislocations and substructure (Fig. 5a and b). Small differences due to the welding process include coarse-grained quasi-polygonal ferrite with small M/A islands at grain boundaries (Fig. 5b). At high
Fig. 4. High-magnification bright field TEM micrographs of the base metal showing high density of (a) grain boundary dislocations, (b) matrix dislocations and (c) dislocation substructures.
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Fig. 5. Bright field transmission electron micrographs of HAZ showing (a) coarse-grained ferrite microstructure with dislocation substructures, (b) ferrite structure with small M/A islands at grain boundaries, (c) martensite and (d) upper bainite and the inset shows SAD pattern analysis.
magnification, the martensite contained microtwins of different orientations (Fig. 5c). The twinned martensite is an indication of high carbon content and the morphology is typical of that observed previously in pipeline steels [23]. Bainite was also observed. The SAD pattern analysis presented in the inset of Fig. 5d implied that the bainitic ferrite exhibits [0 0 1] orientation with the ferrite matrix. 3.3. Precipitation in the base metal Representative TEM micrographs of precipitates in the base metal are presented in Figs. 6–8. The precipitates were of different size and morphology and include cuboidal (∼150 nm), spherical (∼125 nm), irregular (∼240 nm), fine (∼10 nm) and very fine (<10 nm). The EDS analysis obtained for precipitates identified in Fig. 6a indicated that they were (Ti, Nb)C. Bright field TEM micrographs presenting illustrations of fine precipitates within the grain, at grain boundaries, and dark field micrograph of very fine scale precipitation are presented in Fig. 6b–d, respectively. The SAD pattern analysis (inset in Fig. 6d) obtained for fine scale precipitates confirmed that they
were MC type of cubic carbides, having [0 1 2]a //[0 1 1]MC orientation relationship with the ferrite matrix. Fig. 7a and b are bright field TEM micrographs showing regions of precipitation on dislocations and precipitate–dislocation interaction. Similar observations were made on carbon extraction replica and are briefly presented here. Examples of coarse and fine precipitates are presented in Fig. 8a and b. The coarse precipitates were identified as titanium and niobium rich carbides, while the SAD of fine precipitates confirmed that they were MC-type cubic carbides. A plot of precipitate size distribution corresponding to fine and very fine precipitates plot is presented in Fig. 8c, with average particle size of ∼6.5 nm. 3.4. Microstructure of the weld metal The microstructure of the weld metal predominantly consisted of coarse non-equiaxed ferrite grains with M/A islands at a number of grain boundaries. In addition, weld metal also contained small volume fraction of degenerated pearlite, bainite, and martensite. Representative TEM micrographs at low magni-
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Fig. 6. (a) Bright field TEM micrographs of the base metal obtained from thin foil showing coarse precipitates, (b) fine precipitates in the ferrite matrix, (c) at grain boundaries and (d) dark field TEM micrograph showing fine scale precipitation in ferrite matrix and the corresponding SAD pattern analysis in the inset.
Fig. 7. Bright field TEM micrographs obtained from thin foils of the base metal showing (a) fine scale precipitation on dislocations and (b) precipitate–dislocation interaction in the ferrite matrix.
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Fig. 8. (a) Bright field TEM micrographs obtained from carbon extraction replica of base metal showing coarse carbides, (b) dark field TEM micrograph showing MC type of carbide precipitates in ferrite with SAD pattern obtained for MC precipitate in the inset and (c) size distribution plot for the fine precipitates.
Fig. 9. Bright field TEM micrographs of the weld metal showing the general ferrite microstructure. (a and b) Non-equiaxed coarse-grained ferrite with M/A island distributed at grain boundaries, (c and d) non-equiaxed ferrite structure containing high density dislocations and substructures.
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Fig. 10. Bright field TEM micrographs of the weld metal showing (a) morphology of M/A islands in the ferrite matrix, (b) needle-type martensite at the ferrite grain boundary (c) degenerated pearlite and (d) upper bainite.
fication showing the general ferrite microstructure containing a number of M/A islands, sub-boundaries, and significant amount of dislocations and dislocation substructures are presented in Fig. 9a–d. In general, weld metal contained higher volume fraction of M/A constituent, bainite, and martensite as compared to the base metal. This may be a consequence of the different cooling rate experienced by the steel during the multi-pass welding process. Representative bright field TEM micrographs of various microstructural constituents present in the weld metal zone are presented in Fig. 10a–d. Fig. 10a and b describe typical characteristic morphology of martensite/austenite constituent and needle-like martensite, respectively, while Fig. 10c and d show typical morphology of degenerated pearlite and upper bainite, respectively. Since M/A constituent is predominantly present as a second phase in pipeline steel, a detailed TEM study was carried out to understand its microstructure. Fig. 11a and b show bright and dark field TEM micrographs of M/A constituent and the corresponding SAD pattern analysis is presented in Fig. 11c. The SAD pattern analysis revealed the presence of retained austenite together with martensite and the orientation
relationships were [1¯ 2 2]γ //[0 0 1]α and [1¯ 1 1]γ //[0 0 1]α . The M/A constituent formed as alternative layers of retained austenite and martensite. A schematic of M/A layers is depicted in Fig. 11d based on the morphological features observed from Fig. 11a and b. 3.5. Precipitation in the weld metal Fig. 12a presents representative bright field micrographs together with EDS analysis of coarse precipitates observed in the weld metal. The precipitates were larger in size as compared to base metal and exhibited rectangular (∼500 nm), irregular (∼500 nm), cuboidal (∼300 nm), spherical (∼200 nm) and fine (<10 nm) morphology. The EDS analysis revealed that these precipitates were rich in titanium, niobium, and contained small amount of chromium. Unlike the base metal, the weld metal typically contained precipitates of size range ∼18–28 nm in high density. Fig. 12b shows bright field TEM micrograph of these coarse precipitates. EDS analysis obtained on these precipitates indicated that they contained niobium, titanium, iron and manganese. The SAD pattern obtained on these precipitates revealed
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Fig. 11. (a and b) Bright and dark field TEM micrographs of the weld metal showing martensite–austenite (M/A) constituent, (c) SAD pattern analysis for the region shown in (a), and (d) schematic diagram showing the martensite–austenite (retained austenite) constituent.
a complex pattern as shown in the inset in Fig. 12b. The precipitate size distribution plot is presented in Fig. 12c and the average precipitate size was ∼14 nm. Fig. 13a and b are bright and dark field TEM micrographs of carbon extraction replica of weld metal illustrating the fine scale precipitation in weld metal. The SAD pattern analysis of fine precipitates presented in Fig. 13c indicated that these precipitates were MC-type cubic carbides. The precipitate size distribution plot is depicted in Fig. 13d. Majority of the fine precipitates were in the size range of 6–8 nm and the average precipitate size was ∼7 nm. 4. Discussion 4.1. Microstructures of Nb-microalloyed pipeline steel In summary, it can be said that the microstructure of base, HAZ, and weld metal was similar and characterized by a combination of non-equiaxed ferrite, degenerated pearlite/upper bainite, and martensite–austenite constituent (M/A) in the order of increasing cooling rate and decreasing transformation temperature [15–18,24,25]. The similarity in the microstructure of the three zones is consistent with the Knoop hardness (HK ) distribution across the weld zone on a traverse perpendicular to the weld line (Fig. 2). The small variation in the hardness is a conse-
quence of a number of factors that include recrystallization and grain refinement/growth, phase composition, and precipitates [26,27]. The non-equiaxed ferrite microstructure presented in Figs. 3, 5 and 9 is defined as quasi-polygonal ferrite or massive ferrite. Quasi-polygonal ferrite is the first high temperature (below Ae3, equilibrium temperature) ferrite phase to form during continuous cooling. It nucleates heterogeneously at the boundaries of the austenite grains. It is a reconstructive transformation involving diffusion of the atoms, such that the grains of ferrite grow freely across the austenite grain boundaries [27]. The transformation also occurs when the cooling is rapid enough such that the partitioning of solute atoms is minimized. The ferrite grows by diffusion and has a composition similar to that of the austenite. The grains formed under these conditions are likely to be coarse and are therefore referred as massive ferrite transformation [15,25]. The other microstructural constituents present in the investigated steel were degenerated pearlite/upperbainite (Figs. 3d, 5d, 10c and d), martensite (Figs. 5c and 10b), martensite/austenite constituent (M/A) (Figs. 5b, 9, 10a and 11) and they can be considered as microphases [27]. When major phases such as quasi-polygonal ferrite or massive ferrite, acicular or bainitic ferrite have formed the remaining fraction of
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Fig. 12. (a and b) Bright field TEM micrograph of weld metal obtained from carbon extraction replica showing coarse precipitates. The inset in (b) shows the SAD pattern obtained for the precipitates and (c) size distribution of precipitates in Fig. 12b.
untransformed microstructure is very small; hence they are designated as microphases. The excess carbon concentration of microphases is due to the partitioning of carbon because of the growth of the primary phase. Degenerated pearlite is formed by nucleation of cementite at ferrite/austenite interfaces followed by carbide-free ferrite layers enclosing the cementite particles. It nucleates in the transformation temperature range between normal pearlite and upper bainite [28]. Similar to lamellar pearlite, degenerated pearlite is also formed by diffusion process and considering its morphology, the difference is attributed to the insufficient diffusion of carbon to develop continuous lamellae [29]. Another possible situation when degenerated pearlite forms in the base and weld metal is during continuous cooling and involves transformation of M/A islands to degenerated pearlite along with massive ferrite transformation. Observations of degenerated pearlite formation have been reported in thermomechanically processed pipeline steels [1,15]. Similar to the degenerated pearlite, the formation of bainite in the weld microstructure can be attributed to the independent formation from austenite or the transformation of M/A constituent to bainite during continuous cooling [15]. Martensite–austenite constituent (M/A): In pipeline steel during continuous cooling austenite first transforms to ferrite and the remaining austenite becomes carbon rich. On further lowering of transformation temperature, the retained austenite transforms to non-equiaxed or lath ferrite or acic-
ular ferrite and the retained austenite is fully stabilized with highest carbon concentration. Thus, during cooling when the transformation temperature reaches martensite start temperature (Ms ), the high carbon austenite transforms into lenticular microtwinned martensite with different size and orientation (Fig. 5c) [1,10,15,23,24]. In Fig. 11 the typical M/A constituent consisting of alternate layers of martensite and retained austenite is illustrated. It is suggested that the martensite plate first formed will intersect and segment the austenite grain by blocking the plates that forms in the later period. Thus, the martensite formed at different stages results in different morphology and orientation [23]. The selected area diffraction obtained on these M/A plates indicated superimposed [1¯ 2 2]γ , [1¯ 1 1]γ and [0 0 1]α single crystal patterns. The dark field image (Fig. 11b) shows the retained austenite layer in between the martensite plates. Martensite–austenite constituent (M/A) is considered to play an important role in obtaining high strength–toughness combination in pipeline steels [10,15]. The fine and dispersed M/A islands are considered beneficial from the point of view of mechanical properties of pipeline steels [23]. However, in the investigated steel, it was predominantly present in the HAZ and weld zone. The results illustrated the formation of different microstructures in the order of increasing cooling rate and decreasing transformation temperature. Similar transformation kinetics (CCT diagram) was reported [15,24] for low carbon acicular ferrite pipe steels.
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Fig. 13. (a and b) Bright and dark field TEM micrographs of the weld metal obtained from carbon extraction replica showing fine scale precipitation, (c) SAD pattern analysis for the precipitates shown in (a), and (d) size-distribution of precipitates in (a) and (b).
Alloying elements used in the investigated steel such as Mn, Nb, Ti and Cr play an important role in the formation of resultant microstructure. For example, it is shown that the increase in Mn content (∼1.5%) of steel shifts the polygonal ferrite curve of CCT diagram to the right, promoting the formation of intermediate transformation products such as acicular ferrite and bainite [15,24,27]. Similarly, addition of Cr also shifts the polygonal ferrite transformation curve of CCT diagram to the right and causes decrease in diffusivity of carbon and Nb and delays the formation of carbides. Thus, the above results suggest that the complex microstructures observed in the investigated base, HAZ, and weld metal are a consequence of the cooling rates experienced by the steel during thermomechanical processing and subsequent continuous/accelerated cooling.
interchangeable in the precipitate lattice because of similarity in crystal structures (FCC NaCl type) and lattice parameter [34,35]. Fine strain-induced niobium carbides were observed in the ferrite phase of base steel (Fig. 6b and d), on dislocations (Fig. 7a and b), and in the weld metal (Fig. 13). Precipitate size distribution plots (Figs. 8c and 13d) indicated that the average particle size was in the range of ∼6–7 nm. This is an effective size range for precipitation hardening [30]. Thus, on the basis of size distribution and volume fraction of the precipitates, we can conclude that fine scale precipitation contributed to the strength. In summary, strength of the pipeline steel is a cumulative contribution of grain refinement, M/A constituent, bainite, precipitation hardening, and dislocation strengthening. 5. Conclusions
4.2. Precipitation in Nb-microalloyed pipeline steel The precipitates observed in the present pipeline steel are similar to those observed in conventionally hot rolled microalloyed steels and pipeline steels [21,30–38]. The cuboidal-type precipitates were earlier identified as titanium/niobium nitrides, [(Ti,Nb)N], spherical/irregular precipitates as titanium/niobium carbides [(Ti,Nb)C], and fine precipitates as carbides, [(Ti,Nb)C] [21,30,32,33]. Multi-microalloying design approach such as Nb–Ti, Ti–V, and Nb–V generally results in the formation of duplex carbonitrides [35]. Based on recent solubility product calculations carried out in the temperature range of 700–1300 ◦ C [32], it was shown that the microalloying elements, Ti, Nb, are
A new ultrahigh strength niobium-microalloyed steel of yield strength of ∼700 MPa consisting of polygonal ferrite together with upper bainite, degenerated pearlite, and martensite–austenite constituent has been developed. The microstructure of HAZ and weld zone was similar to that of the base metal, consistent with the hardness data. The precipitates in the niobium-microalloyed pipeline steel can be classified on the basis of size and morphology and include rectangular, irregular, cuboidal/spherical, and fine morphology with size range of ∼500, ∼240–500, 125–300 nm, and <10 nm, respectively. Strain-induced MC (NbC) type of carbides precipitated on dislocations, and at grain boundaries. The observed high
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strength–toughness combination is derived from the cumulative effect of fine-grained ferrite microstructure, M/A strengthening, dislocation strengthening, and precipitation strengthening. Acknowledgement The University of Louisiana at Lafayette gratefully acknowledges the financial support from CBMM, Brazil for the work presented here. References [1] Ming-Chun Zhao, Ke Yang, Y. Shan, Mater. Sci. Eng. A 335 (2002) 14– 20. [2] Ivani De S. Bott, Luis F.C.G. Teixeira, Paulo R. Rios, Metall. Trans. A 36A (2005) 443–454. [3] I.H.G. Hillenbrand, I.M. Graf, I.C. Kalwa, Proceedings of the Conference Niobium 2001, Orlando, FL, USA, 2001. [4] A. Takahashi, M. Iino, ISIJ Int. 36 (2) (1996) 235–240. [5] A. Takahashi, M. Iino, ISIJ Int. 36 (2) (1996) 241–245. [6] A. Takahashi, H. Ogawa, ISIJ Int. 36 (2) (1996) 334–340. [7] Ming-Chun Zhao, Ke Yang, Scripta. Mater. 52 (2005) 881–886. [8] A. Contreras, A. Albiter, M. Salazar, R. Perez, Mater. Sci. Eng. A 407 (2005) 45–52. [9] Y.M. Kim, S.K. Kim, Y.J. Lim, N.J. Kim, ISIJ Int. 42 (12) (2002) 1571–1577. [10] Y. Zhong, F. Xiao, J. Zhang, Y. Shan, W. Wang, Ke Yang, Acta Mater. 54 (2006) 435–443. [11] K. Junhua, Z. Lin, G. Bin, L. Pinghe, W. Aihua, X. Changsheng, Mater. Des. 25 (2004) 723–728. [12] W.B. Lee, S.G. Hong, C.G. Park, K.H. Kim, S.H. Park, Scripta Mater. 43 (2000) 319–324. [13] W. Sun, C. Lu, A.K. Tieu, Z. Jiang, X. Liu, G. Wang, Mater. Process. Technol. 125–126 (2002) 72–76. [14] F.R. Xiao, B. Liao, Y.Y. Shan, G.Y. Qiao, Y. Zhang, C. Zhang, K. Yang, Mater. Sci. Eng. A 431 (2006) 41–52. [15] Furen Xiao, Bo Liao, Deliang Ren, Yiying Shan, Ke Yang, Mater. Charact. 54 (2005) 305–314. [16] B. Hwang, S. Lee, Y.M. Kim, N.J. Kim, J.Y. Yoo, C.S. Woo, Mater. Sci. Eng. A 368 (2004) 18–27.
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[17] B. Hwang, Y.M. Kim, S. Lee, N.J. Kim, S.S. Ahn, Metall. Mater. Trans. 36A (2005) 371–387. [18] B. Hwang, Y.M. Kim, S. Lee, N.J. Kim, S.S. Ahn, Metall. Mater. Trans. 36A (2005) 725–739. [19] A.M. Elwazri, R. Varano, F. Siciliano, D. Bai, S. Yue, Metall. Trans. A 36A (2005) 2929–2936. [20] Ming-Chun Zhao, T. Hanamura, Hai Qui, Ke Yang, Mater. Sci. Eng. A 395 (2005) 327–332. [21] S. Shanmugam, R.D.K. Misra, J.E. Hartmann, S. Jansto, Mater. Sci. Eng. A 441 (2006) 215–229. [22] R.D.K. Misra, J.E. Hartman, A method of making a high strength low alloy steel, US patent no. 6488790, 2002. [23] C. Wang, X. Wu, J. Liu, N. Xu, Mater. Sci. Eng. A 438 (2006) 267–271. [24] Ming-Chun Zhao, Ke Yang, Yiying Shan, Mater. Sci. Engg. A 335 (2003) 126–136. [25] G. Krauss, S.W. Thompson, ISIJ Int. 35 (8) (1995) 937–945. [26] G.R. Stewart, A.M. Elwazri, R. Varano, N. Pokutylowicz, S. Yue, J.J. Jonas, Mater. Sci. Eng. A 420 (2006) 115–121. [27] H.K.D.H. Bhadeshia, L.E. Svensson, in: H. Cerjak, K.E. Easterling (Eds.), Mathematical Modelling of Weld Phenomena, Institute of Materials, London, 1993, pp. 109–182. [28] Y. Ohmori, Trans. ISIJ 11 (1971) 339–348. [29] Y. Ohmori, R.W.K. Honeycombe, Proc. ICSTIS (Suppl.) Trans. ISIJ 11 (1971) 1160–1165. [30] R.D.K. Misra, G.C. Weatherly, J.E. Hartmann, A.J. Boucek, Mater. Sci. Technol. 17 (2001) 1119–1129. [31] S. Shanmugam, M. Tanniru, R.D.K. Misra, D. Panda, S. Jansto, Mater. Sci. Technol. 21 (2005) 165–177. [32] S. Shanmugam, M. Tanniru, R.D.K. Misra, D. Panda, S. Jansto, Mater. Sci. Technol. 21 (2005) 883–892. [33] R.D.K. Misra, K.K. Tenneti, G.C. Weatherly, G. Tither, Metall Trans. 34A (2003) 2341–2351. [34] M. Charleux, W.J. Poole, M. Militizer, A. Deschamps, Metall Trans. 32A (2001) 1635–1646. [35] T.N. Baker, Y. Li, J.A. Wilson, A.J. Craven, D.N. Crowther, Mater. Sci. Technol. 20 (2004) 720–730. [36] M.J. Crooks, A.J. Garratt-Reed, J.B. Vander Sande, W.S. Owen, Metall Trans. 12A (1981) 1999–2013. [37] R.D.K. Misra, S.W. Thompson, T.A. Hylton, A.J. Boucek, Metall Trans. 32A (2001) 745–760. [38] C.P. Reip, S. Shanmugam, R.D.K. Misra, Mater. Sci. Eng. A 424 (2006) 307–317.
Microstructure and high strength–toughness combination of a new 700 MPa Nb-microalloyed pipeline steel S. Shanmugam a , N.K. Ramisetti a , R.D.K. Misra a,∗ , J. Hartmann b , S.G. Jansto c a
Center for Structural and Functional Materials and Department of Chemical Engineering, University of Louisiana at Lafayette, Lafayette, LA 70504-4130, USA b Arcelor Mittal, USA Research & Development Center, 3001 E. Columbus Drive, East Chicago, IN 46312, USA c Reference Metals, 1000 Old Pond Road, Bridgeville, PA 15017, USA Received 21 March 2007; received in revised form 14 May 2007; accepted 1 June 2007
Abstract A new ultrahigh strength niobium-microalloyed pipeline steel of yield strength ∼700 MPa has been processed. The Charpy impact toughness at 0 ◦ C was 27 J and tensile elongation was 16%. The ultrahigh strength is derived from the cumulative combination of fine grain size, solid solution strengthening with additional interstitial hardening, precipitation hardening from carbides, dislocation hardening, and mixed microstructure. The microstructure was characterized by polygonal ferrite, upper bainite, degenerated pearlite, and martensite–austenite (MA) constituents. The microstructure of weld and heat-affected zone (HAZ) was similar to the base metal such that the hardness is retained in the weld region implying insignificant softening in the weld zone. Niobium and titanium precipitates of different morphology and size range evolved during thermomechanical processing and include rectangular (∼500 nm), irregular (∼240–500 nm), cuboidal/spherical (∼125–300 nm), and very fine (<10 nm). They were generally MC type of carbides. An important aspect of the developed steel is significantly lean chemistry. © 2007 Elsevier B.V. All rights reserved. Keywords: Pipeline steel; Niobium-microalloyed; Microstructure
1. Introduction Currently there is a demand to transport crude oil and gas by pipeline at a higher operating pressure to increase the capacity. This requires the use of ultrahigh strength steels. Increasing the strength of the pipeline steel enables a significant reduction in wall thickness with consequent reduction in weight. Thus, a major goal within the steel industry is to develop ultrahigh strength microalloyed pipeline steels (∼700–800 MPa). It is, however, important that the increase in the yield strength is not accompanied by a decrease in fracture toughness and formability because a decrease in toughness will encourage stress-induced cracking, and reduced formability will cause difficulties in forming (e.g. pipe-bowing). Thus, high strength in association with high toughness and formability are important requirements of the pipeline industry [1–3]. High strength–toughness combination is essential for pipeline steels for transporting natural gas and crude oil over
∗
Corresponding author. E-mail address: [email protected] (R.D.K. Misra).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.06.003
a long distance under high pressure [2]. Other characteristics that are required include resistance to hydrogen-induced blister cracking in sour service environment [4,5], stress corrosion cracking resistance, especially in H2 S environment [6–8], and fatigue resistance [9,10]. It is important that the microstructure of the pipeline steel provides combination of aforementioned properties. The evolution of microstructure depends on alloy chemistry and thermomechanical processing. Alloying additions such as Mn, Nb, V, Ti, Mo, Ni, Cr and Cu are commonly employed in pipeline steels to obtain the desired microstructure and mechanical properties [5,6,11–13]. However, judicious selection of alloying elements is necessary to obtain beneficial effect on mechanical properties with reduced alloy cost. For instance, the numbers of alloying elements are reduced to achieve lower carbon equivalent (CE) to ensure good field weldability [2,3]. On the other hand alloy additions such as Cr, Cu and Ni are added to obtain strength in severe corrosive environment [2,5]. Controlled thermomechanical processing is considered as the primary route for the development of API grade pipeline steels because it provides the desirable and fine-grained microstructure. Furthermore, it allows high strength–toughness
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combination to be achieved with accelerated cooling [3]. The ultimate microstructure strongly depends on the processing parameters such as reheating temperature, percentage reduction, deformation temperature, cooling rate, and coiling temperature [14]. This is because fine austenite grains, substructure and dislocations in austenite effectively promote the transformation of fine ferrite. The primary grain refinement in controlled rolling is achieved through recrystallization of austenite during deformation and through the use of microalloying elements, such as Nb that precipitates as fine carbides and inhibits grain growth [14]. In thermomechanically processed pipeline steels [1,7–9,11,15–20], different combinations of microstructures are obtained. It is believed that acicular ferrite microstructure with uniform distribution of martensite/austenite (M/A) islands as a second phase, provides the desired mechanical properties [1,11]. Besides, acicular ferrite the preferred microstructure of pipeline steel is characterized by microstructural constituents that include bainite and quasi-polygonal ferrite or massive ferrite [14,21]. This study describes the microstructure and precipitation behavior with the aim to relate to the high strength–toughness combination of a new industrially processed
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700 MPa Nb-microalloyed pipeline steel with a lean chemistry. The precipitates were characterized in terms of morphology, size, chemistry, and crystallography. 2. Experimental The Nb-microalloyed pipeline steel described here was continuously cast and hot rolled to a minimum of ∼7.6 mm gage and subsequently fabricated to ∼150 mm diameter pipes. The alloy design and nominal chemical composition of the steel was based on our recently patented hot rolled steel [22] and consisted of Fe–(0.04–0.06)C–(0.035–0.05)Ti–(0.08–0.09)Nb–(0.3–0.4)Cr. The sulfur content was less than 0.005 wt.%. The steel was hot rolled using controlled rolling practice on Mittal Steel Indiana Harbor 84 in. hot strip mill. The thermomechanical processing details are not described because of proprietary reasons. Standard tensile tests were conducted at room temperature on longitudinal specimens machined according to ASTM E8 specification using computerized tensile testing system. The initial crosshead speed was 4.2 × 10−2 mm/s up to a strain of ∼5%,
Fig. 1. Low-magnification (a) and high-magnification (b)–(d) light micrographs showing general microstructure of base, HAZ, and weld metal zones.
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then the speed was increased to 4.2 × 10−1 mm/s. Impact toughness was measured using standard Charpy v-notch impact test (ASTM 23) at 0 ◦ C. Transmission electron microscopy was carried out on samples cut from the fabricated steel pipe and relate to regions corresponding base, HAZ, and weld metal zones. Thin foils were prepared by cutting thin wafers from the steel samples using precision saw (Buehler Isomet 1000) and grinding to ∼100 m in thickness. Three-millimeter discs were punched from the wafers and electropolished using a solution of 10% perchloric acid in acetic acid electrolyte. Carbon extraction replicas were also prepared for characterization of precipitates. The surface of the polished specimens was etched with 2% nital and carbon was evaporated onto the etched surface. Finally, the surface was scored to ∼3 mm squares and the sample etched first with 10% nital and then with 2% nital. Subsequently, the extracted replicas were rinsed with distilled water and placed on the copper grid and dried. Foils and carbon extraction replicas were examined with a Hitachi 7600 TEM operated at 100 kV.
Fig. 2. Knoop hardness distribution across the weld zones on a traverse perpendicular to weld zone.
Fig. 3. Bright field TEM micrographs illustrating the general microstructure of the base metal. (a) Non-equiaxed or quasi-polygonal ferrite, (b) quasi-polygonal ferrite region containing high density of dislocations and dislocation substructure, (c) quasi-polygonal ferrite region containing grain boundary carbides, and (d) degenerated pearlite or upper bainite.
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3. Results 3.1. Mechanical properties The yield and tensile strength of the hot end of the coil was ∼735 MPa (105 ksi) and ∼791 MPa (113 ksi), respectively. The Charpy v-notch impact toughness of Nb-microalloyed pipeline steel at 0 ◦ C was 27 J and the tensile elongation was 16%. It may be noted that the listed mechanical properties are the minimum values that were obtained on any coil. 3.2. Microstructure and hardness of base, HAZ, and weld metal zones Representative light micrographs of base, HAZ, and weld metal are presented in Fig. 1a–d. The microstructure was predominantly fine-grained ferrite in all the three zones. As expected, the base metal (as-hot rolled) contained some elongated ferrite grains (Fig. 1b). A hardness profile encompassing base, HAZ, and weld metal zones is presented in Fig. 2. The hardness–distance profile indicates that the weld zone and HAZ
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did not experience significant softening, consistent with the microstructure of the three zones, as described below. Representative low-magnification bright field TEM micrographs of the base metal are presented in Fig. 3a–d. In general, the microstructure was of mixed type consisting of ferrite, upper bainite, and degenerated pearlite. The ferrite grains were of nonequiaxed (Fig. 3a) and were characterized by sub-boundaries and dislocation substructures (Fig. 3b), and carbides at the ferrite grain boundaries (Fig. 3c). A region of the base metal characterized by upper bainite/degenerated pearlite is presented in Fig. 3d. Fig. 4a and b are high-magnification bright field TEM micrographs of ferrite grains that contained high density dislocations at the grain boundaries (Fig. 4a), dislocations in the body of the grain (Fig. 4b), and dislocation substructures (Fig. 4c). Representative bright field TEM micrographs of HAZ are presented in Fig. 5a–d. The ferrite microstructure was similar to the base metal and contained high density of dislocations and substructure (Fig. 5a and b). Small differences due to the welding process include coarse-grained quasi-polygonal ferrite with small M/A islands at grain boundaries (Fig. 5b). At high
Fig. 4. High-magnification bright field TEM micrographs of the base metal showing high density of (a) grain boundary dislocations, (b) matrix dislocations and (c) dislocation substructures.
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Fig. 5. Bright field transmission electron micrographs of HAZ showing (a) coarse-grained ferrite microstructure with dislocation substructures, (b) ferrite structure with small M/A islands at grain boundaries, (c) martensite and (d) upper bainite and the inset shows SAD pattern analysis.
magnification, the martensite contained microtwins of different orientations (Fig. 5c). The twinned martensite is an indication of high carbon content and the morphology is typical of that observed previously in pipeline steels [23]. Bainite was also observed. The SAD pattern analysis presented in the inset of Fig. 5d implied that the bainitic ferrite exhibits [0 0 1] orientation with the ferrite matrix. 3.3. Precipitation in the base metal Representative TEM micrographs of precipitates in the base metal are presented in Figs. 6–8. The precipitates were of different size and morphology and include cuboidal (∼150 nm), spherical (∼125 nm), irregular (∼240 nm), fine (∼10 nm) and very fine (<10 nm). The EDS analysis obtained for precipitates identified in Fig. 6a indicated that they were (Ti, Nb)C. Bright field TEM micrographs presenting illustrations of fine precipitates within the grain, at grain boundaries, and dark field micrograph of very fine scale precipitation are presented in Fig. 6b–d, respectively. The SAD pattern analysis (inset in Fig. 6d) obtained for fine scale precipitates confirmed that they
were MC type of cubic carbides, having [0 1 2]a //[0 1 1]MC orientation relationship with the ferrite matrix. Fig. 7a and b are bright field TEM micrographs showing regions of precipitation on dislocations and precipitate–dislocation interaction. Similar observations were made on carbon extraction replica and are briefly presented here. Examples of coarse and fine precipitates are presented in Fig. 8a and b. The coarse precipitates were identified as titanium and niobium rich carbides, while the SAD of fine precipitates confirmed that they were MC-type cubic carbides. A plot of precipitate size distribution corresponding to fine and very fine precipitates plot is presented in Fig. 8c, with average particle size of ∼6.5 nm. 3.4. Microstructure of the weld metal The microstructure of the weld metal predominantly consisted of coarse non-equiaxed ferrite grains with M/A islands at a number of grain boundaries. In addition, weld metal also contained small volume fraction of degenerated pearlite, bainite, and martensite. Representative TEM micrographs at low magni-
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Fig. 6. (a) Bright field TEM micrographs of the base metal obtained from thin foil showing coarse precipitates, (b) fine precipitates in the ferrite matrix, (c) at grain boundaries and (d) dark field TEM micrograph showing fine scale precipitation in ferrite matrix and the corresponding SAD pattern analysis in the inset.
Fig. 7. Bright field TEM micrographs obtained from thin foils of the base metal showing (a) fine scale precipitation on dislocations and (b) precipitate–dislocation interaction in the ferrite matrix.
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Fig. 8. (a) Bright field TEM micrographs obtained from carbon extraction replica of base metal showing coarse carbides, (b) dark field TEM micrograph showing MC type of carbide precipitates in ferrite with SAD pattern obtained for MC precipitate in the inset and (c) size distribution plot for the fine precipitates.
Fig. 9. Bright field TEM micrographs of the weld metal showing the general ferrite microstructure. (a and b) Non-equiaxed coarse-grained ferrite with M/A island distributed at grain boundaries, (c and d) non-equiaxed ferrite structure containing high density dislocations and substructures.
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Fig. 10. Bright field TEM micrographs of the weld metal showing (a) morphology of M/A islands in the ferrite matrix, (b) needle-type martensite at the ferrite grain boundary (c) degenerated pearlite and (d) upper bainite.
fication showing the general ferrite microstructure containing a number of M/A islands, sub-boundaries, and significant amount of dislocations and dislocation substructures are presented in Fig. 9a–d. In general, weld metal contained higher volume fraction of M/A constituent, bainite, and martensite as compared to the base metal. This may be a consequence of the different cooling rate experienced by the steel during the multi-pass welding process. Representative bright field TEM micrographs of various microstructural constituents present in the weld metal zone are presented in Fig. 10a–d. Fig. 10a and b describe typical characteristic morphology of martensite/austenite constituent and needle-like martensite, respectively, while Fig. 10c and d show typical morphology of degenerated pearlite and upper bainite, respectively. Since M/A constituent is predominantly present as a second phase in pipeline steel, a detailed TEM study was carried out to understand its microstructure. Fig. 11a and b show bright and dark field TEM micrographs of M/A constituent and the corresponding SAD pattern analysis is presented in Fig. 11c. The SAD pattern analysis revealed the presence of retained austenite together with martensite and the orientation
relationships were [1¯ 2 2]γ //[0 0 1]α and [1¯ 1 1]γ //[0 0 1]α . The M/A constituent formed as alternative layers of retained austenite and martensite. A schematic of M/A layers is depicted in Fig. 11d based on the morphological features observed from Fig. 11a and b. 3.5. Precipitation in the weld metal Fig. 12a presents representative bright field micrographs together with EDS analysis of coarse precipitates observed in the weld metal. The precipitates were larger in size as compared to base metal and exhibited rectangular (∼500 nm), irregular (∼500 nm), cuboidal (∼300 nm), spherical (∼200 nm) and fine (<10 nm) morphology. The EDS analysis revealed that these precipitates were rich in titanium, niobium, and contained small amount of chromium. Unlike the base metal, the weld metal typically contained precipitates of size range ∼18–28 nm in high density. Fig. 12b shows bright field TEM micrograph of these coarse precipitates. EDS analysis obtained on these precipitates indicated that they contained niobium, titanium, iron and manganese. The SAD pattern obtained on these precipitates revealed
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Fig. 11. (a and b) Bright and dark field TEM micrographs of the weld metal showing martensite–austenite (M/A) constituent, (c) SAD pattern analysis for the region shown in (a), and (d) schematic diagram showing the martensite–austenite (retained austenite) constituent.
a complex pattern as shown in the inset in Fig. 12b. The precipitate size distribution plot is presented in Fig. 12c and the average precipitate size was ∼14 nm. Fig. 13a and b are bright and dark field TEM micrographs of carbon extraction replica of weld metal illustrating the fine scale precipitation in weld metal. The SAD pattern analysis of fine precipitates presented in Fig. 13c indicated that these precipitates were MC-type cubic carbides. The precipitate size distribution plot is depicted in Fig. 13d. Majority of the fine precipitates were in the size range of 6–8 nm and the average precipitate size was ∼7 nm. 4. Discussion 4.1. Microstructures of Nb-microalloyed pipeline steel In summary, it can be said that the microstructure of base, HAZ, and weld metal was similar and characterized by a combination of non-equiaxed ferrite, degenerated pearlite/upper bainite, and martensite–austenite constituent (M/A) in the order of increasing cooling rate and decreasing transformation temperature [15–18,24,25]. The similarity in the microstructure of the three zones is consistent with the Knoop hardness (HK ) distribution across the weld zone on a traverse perpendicular to the weld line (Fig. 2). The small variation in the hardness is a conse-
quence of a number of factors that include recrystallization and grain refinement/growth, phase composition, and precipitates [26,27]. The non-equiaxed ferrite microstructure presented in Figs. 3, 5 and 9 is defined as quasi-polygonal ferrite or massive ferrite. Quasi-polygonal ferrite is the first high temperature (below Ae3, equilibrium temperature) ferrite phase to form during continuous cooling. It nucleates heterogeneously at the boundaries of the austenite grains. It is a reconstructive transformation involving diffusion of the atoms, such that the grains of ferrite grow freely across the austenite grain boundaries [27]. The transformation also occurs when the cooling is rapid enough such that the partitioning of solute atoms is minimized. The ferrite grows by diffusion and has a composition similar to that of the austenite. The grains formed under these conditions are likely to be coarse and are therefore referred as massive ferrite transformation [15,25]. The other microstructural constituents present in the investigated steel were degenerated pearlite/upperbainite (Figs. 3d, 5d, 10c and d), martensite (Figs. 5c and 10b), martensite/austenite constituent (M/A) (Figs. 5b, 9, 10a and 11) and they can be considered as microphases [27]. When major phases such as quasi-polygonal ferrite or massive ferrite, acicular or bainitic ferrite have formed the remaining fraction of
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Fig. 12. (a and b) Bright field TEM micrograph of weld metal obtained from carbon extraction replica showing coarse precipitates. The inset in (b) shows the SAD pattern obtained for the precipitates and (c) size distribution of precipitates in Fig. 12b.
untransformed microstructure is very small; hence they are designated as microphases. The excess carbon concentration of microphases is due to the partitioning of carbon because of the growth of the primary phase. Degenerated pearlite is formed by nucleation of cementite at ferrite/austenite interfaces followed by carbide-free ferrite layers enclosing the cementite particles. It nucleates in the transformation temperature range between normal pearlite and upper bainite [28]. Similar to lamellar pearlite, degenerated pearlite is also formed by diffusion process and considering its morphology, the difference is attributed to the insufficient diffusion of carbon to develop continuous lamellae [29]. Another possible situation when degenerated pearlite forms in the base and weld metal is during continuous cooling and involves transformation of M/A islands to degenerated pearlite along with massive ferrite transformation. Observations of degenerated pearlite formation have been reported in thermomechanically processed pipeline steels [1,15]. Similar to the degenerated pearlite, the formation of bainite in the weld microstructure can be attributed to the independent formation from austenite or the transformation of M/A constituent to bainite during continuous cooling [15]. Martensite–austenite constituent (M/A): In pipeline steel during continuous cooling austenite first transforms to ferrite and the remaining austenite becomes carbon rich. On further lowering of transformation temperature, the retained austenite transforms to non-equiaxed or lath ferrite or acic-
ular ferrite and the retained austenite is fully stabilized with highest carbon concentration. Thus, during cooling when the transformation temperature reaches martensite start temperature (Ms ), the high carbon austenite transforms into lenticular microtwinned martensite with different size and orientation (Fig. 5c) [1,10,15,23,24]. In Fig. 11 the typical M/A constituent consisting of alternate layers of martensite and retained austenite is illustrated. It is suggested that the martensite plate first formed will intersect and segment the austenite grain by blocking the plates that forms in the later period. Thus, the martensite formed at different stages results in different morphology and orientation [23]. The selected area diffraction obtained on these M/A plates indicated superimposed [1¯ 2 2]γ , [1¯ 1 1]γ and [0 0 1]α single crystal patterns. The dark field image (Fig. 11b) shows the retained austenite layer in between the martensite plates. Martensite–austenite constituent (M/A) is considered to play an important role in obtaining high strength–toughness combination in pipeline steels [10,15]. The fine and dispersed M/A islands are considered beneficial from the point of view of mechanical properties of pipeline steels [23]. However, in the investigated steel, it was predominantly present in the HAZ and weld zone. The results illustrated the formation of different microstructures in the order of increasing cooling rate and decreasing transformation temperature. Similar transformation kinetics (CCT diagram) was reported [15,24] for low carbon acicular ferrite pipe steels.
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Fig. 13. (a and b) Bright and dark field TEM micrographs of the weld metal obtained from carbon extraction replica showing fine scale precipitation, (c) SAD pattern analysis for the precipitates shown in (a), and (d) size-distribution of precipitates in (a) and (b).
Alloying elements used in the investigated steel such as Mn, Nb, Ti and Cr play an important role in the formation of resultant microstructure. For example, it is shown that the increase in Mn content (∼1.5%) of steel shifts the polygonal ferrite curve of CCT diagram to the right, promoting the formation of intermediate transformation products such as acicular ferrite and bainite [15,24,27]. Similarly, addition of Cr also shifts the polygonal ferrite transformation curve of CCT diagram to the right and causes decrease in diffusivity of carbon and Nb and delays the formation of carbides. Thus, the above results suggest that the complex microstructures observed in the investigated base, HAZ, and weld metal are a consequence of the cooling rates experienced by the steel during thermomechanical processing and subsequent continuous/accelerated cooling.
interchangeable in the precipitate lattice because of similarity in crystal structures (FCC NaCl type) and lattice parameter [34,35]. Fine strain-induced niobium carbides were observed in the ferrite phase of base steel (Fig. 6b and d), on dislocations (Fig. 7a and b), and in the weld metal (Fig. 13). Precipitate size distribution plots (Figs. 8c and 13d) indicated that the average particle size was in the range of ∼6–7 nm. This is an effective size range for precipitation hardening [30]. Thus, on the basis of size distribution and volume fraction of the precipitates, we can conclude that fine scale precipitation contributed to the strength. In summary, strength of the pipeline steel is a cumulative contribution of grain refinement, M/A constituent, bainite, precipitation hardening, and dislocation strengthening. 5. Conclusions
4.2. Precipitation in Nb-microalloyed pipeline steel The precipitates observed in the present pipeline steel are similar to those observed in conventionally hot rolled microalloyed steels and pipeline steels [21,30–38]. The cuboidal-type precipitates were earlier identified as titanium/niobium nitrides, [(Ti,Nb)N], spherical/irregular precipitates as titanium/niobium carbides [(Ti,Nb)C], and fine precipitates as carbides, [(Ti,Nb)C] [21,30,32,33]. Multi-microalloying design approach such as Nb–Ti, Ti–V, and Nb–V generally results in the formation of duplex carbonitrides [35]. Based on recent solubility product calculations carried out in the temperature range of 700–1300 ◦ C [32], it was shown that the microalloying elements, Ti, Nb, are
A new ultrahigh strength niobium-microalloyed steel of yield strength of ∼700 MPa consisting of polygonal ferrite together with upper bainite, degenerated pearlite, and martensite–austenite constituent has been developed. The microstructure of HAZ and weld zone was similar to that of the base metal, consistent with the hardness data. The precipitates in the niobium-microalloyed pipeline steel can be classified on the basis of size and morphology and include rectangular, irregular, cuboidal/spherical, and fine morphology with size range of ∼500, ∼240–500, 125–300 nm, and <10 nm, respectively. Strain-induced MC (NbC) type of carbides precipitated on dislocations, and at grain boundaries. The observed high
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