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Microstructure of high strength niobium-containing pipeline steel
Materials Science and Engineering A 441 (2006) 215–229
Microstructure of high strength niobium-containing pipeline steel S. Shanmugam a , R.D.K. Misra a,∗ , J. Hartmann b , S.G. Jansto c a
Center for Structural and Functional Materials, Department of Chemical Engineering, University of Louisiana at Lafayette, Lafayette, LA 70504-4130, USA b Mittal Steel, Indiana Harbor Works, 3001 Dickey Road, East Chicago, IN 46312, USA c Reference Metals, 1000 Old Pond Road, Bridgeville, PA 15017, USA Received 26 July 2006; accepted 4 August 2006
Abstract The paper describes the microstructural constituents in a industrially processed Nb-microalloyed pipeline steel having yield strength of ∼620 MPa. The microstructure of base, heat affected zone (HAZ), and weld metal of the fabricated steel pipe was examined by optical and transmission electron microscopy. The microstructure of thermomechanically processed pipeline steel primarily consisted of non-equiaxed ferrite of mixed morphologies with small fraction of degenerated pearlite. The microstructure contained high dislocation density, sub-boundaries and dislocation substructures. The HAZ was characterized by a combination of fine and coarse grained polygonal ferrite structure with high density of dislocations and fine cementite particles. In the weld metal, the constituents of complex ferrite were low temperature transformation products formed during continuous cooling such as quasi-polygonal or massive ferrite, acicular ferrite, bainitic ferrite and dispersion of coarse and fine cementite particles in the ferrite matrix. The precipitates in the investigated pipeline steel were of duplex type containing either Nb and Ti or Ti and Mo, even though the steel contained low concentration of titanium. Precipitates of different morphology and size range were observed and include rectangular (∼100–130 nm), cuboidal/spherical (∼20–100 nm), fine (∼10–20 nm) and very fine (<10 nm). They were MC type of carbides. © 2006 Elsevier B.V. All rights reserved. Keywords: Pipeline steels; Microstructure; Niobium
1. Introduction Modern pipeline steels are currently being produced to the strength level required for API grades using accelerated cooling and controlled rolling technology. The primary interest is to obtain the best possible combination of strength and toughness [1]. An increase in the strength of these steels is necessary for deep water applications and for remote areas where pipes experience transmission pressures exceeding 2500 psi. Weldability, formability, fracture toughness and transition temperature [2], resistance to hydrogen induced blister cracking in sour service environment [3,4], stress corrosion cracking resistance for underground service, especially with H2 S environment [5–7], and fatigue resistance are the additional requirements for transmission of oil and gas through pipelines [8,9]. Thus, it is imperative that the steel is characterized by a microstructure that provides the necessary strength–toughness combination. The microstructure depends on the alloy chemistry, thermo∗
Corresponding author. E-mail address: [email protected] (R.D.K. Misra).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.08.017
mechanical processing and degree of cooling. Alloy additions such as Mn, Nb, V, Ti, Mo, Ni and Cu are commonly added in pipeline steels in order to achieve the desired microstructure and mechanical properties [4,5,10–12]. However, judicious selection of alloying elements is necessary to obtain their beneficial effect on the evolution of microstructure and consequent on mechanical properties. For instance, Nb- and Mo-containing steels are commonly used in pipeline applications because it is observed that the HSLA steels containing Nb and Mo exhibit superior strength and toughness combination as compared to the HSLA steels containing Nb and V [11,12]. Also, careful control of alloying element content is equally important, according to the specific application of pipelines. For example, Mn is an important alloying element for solid solution strengthening, however, reduced Mn content in steels decreases the centerline microstructural banding [13]. Also, Mn content of greater than 0.3% in pipeline steels causes hydrogen induced blister cracking on being subjected to sour service environment. Thus, pipeline steels containing low Mn content with additional strength obtained from Cu has been considered for severe service condition [4].
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Controlled thermomechanical processing is another aspect of high strength pipeline steel for modern API grade steels [2,10]. The process parameters such as soaking temperature, rolling temperature, finishing temperature, and cooling rate play an important role in the ultimate microstructure. Higher cooling rate and lower interrupted cooling temperature are the parameters that determine the formation of acicular ferrite-based microstructure, a preferred microstructure for pipeline steels [10]. Different combinations of microstructures have been obtained in thermomechanically processed pipeline steels in laboratory simulated experiments and include ferrite–pearlite [7], polygonal ferrite–acicular ferrite [1], acicular ferrite–bainite [10], acicular ferrite–martensite–austenite constituent (M/A) [14], a mixed microstructure [8,15–17], acicular ferrite [6,14,18,19], bainite [3]. It is reported that a predominant acicular ferrite microstructure with martensite/austenite (M/A) islands, as a second phase exhibits optimum mechanical properties [1,10]. The acicular ferrite dominated microstructure can be obtained by optimizing processing conditions and without the need for significant alloying concentration [1,2,10,20]. The objective of this study is to report initial results concerning microstructure and precipitation behavior of industrially processed 630 MPa Nb-containing microalloyed pipeline steel that contributed to good strength–toughness combination. The precipitates were characterized in terms of morphology, size, chemistry, and crystallography. Furthermore, the orientation relationship between precipitates and ferrite were determined. 2. Experimental The Nb-microalloyed pipeline steels discussed here were industrial heats that were continuously cast and hot-rolled to a minimum of ∼8 mm gage and subsequently fabricated to ∼450 mm diameter pipes. The chemical composition of the investigated steel is given in Table 1. The thermomechanical processing details are not described here due to proprietary reason. Standard tensile tests were conducted at room temperature on longitudinal specimens machined according to ASTM E8 specification (dimensions 225 mm × 12.5 mm, gage length, 50 mm) using computerized tensile testing system. The initial crosshead speed was 4.2 × 10−2 mm/s up to a strain of ∼5%, then the speed was increased to 4.2 × 10−1 mm/s. Impact toughness was meaTable 1 Chemical composition of Nb-containing pipeline steel in wt.% C Mn Si P S Al Ti Nb Ca Mo
0.04/0.06 1.4/1.6 0.18/0.25 0.020 0.003/0.005 0.02/0.06 0.005/0.015 0.05/0.085 0.001/0.004 0.10/0.30
Table 2 Mechanical properties of Nb-containing pipeline steel Properties
Longitudinal direction
Transverse direction
Yield strength (MPa) Tensile strength (MPa) % Elongation Charpy v-notch impact toughness at 0 ◦ C (J)
613–659 692–716 26–30 –
591–593 722–734 28–29 173–266
sured using standard Charpy v-notch impact test (ASTM 23) at 0 ◦ C. Transmission electron microscopy was carried out on thin foils prepared by cutting thin wafers from the steel samples of base, HAZ and weld metal, 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. Foils were examined by HITACHI 7600 TEM operated at 120 kV. 3. Results 3.1. Mechanical properties The yield strength, tensile strength and % elongation data obtained in longitudinal and transverse direction for Nbmicroalloyed pipeline steels are summarized in Table 2. The longitudinal and transverse yield strength was in the range of 613–659 MPa and 591–593 MPa, respectively. The tensile strength in the longitudinal and transverse orientation was 692–716 MPa and 722–734 MPa, respectively. The transverse Charpy v-notch impact toughness of Nb-microalloyed pipe steel at 0 ◦ C was high and in the range of 173–266 J. 3.2. Microstructure: base metal and HAZ Microstructural studies on base, HAZ and weld metal of the fabricated pipe are important to identify any differences in microstructural constituents. Representative light micrographs of base, HAZ and weld metal are presented in Fig. 1a–c. The microstructure was predominantly fine ferrite structure in all the three regions, except that the HAZ contained large fraction of coarse polygonal ferrite grain structure (Fig. 1b). Representative low magnification bright field TEM micrographs of the base metal are presented in Figs. 2 and 3. It may be seen that the ferrite microstructure is complex and consists of non-equiaxed ferrite grains (Fig. 2a), sub-boundaries and dislocation substructures (Fig. 2b). These micrographs were obtained at different tilt angles (foil tilt) in order to reveal the image details due to different contrast among the grains. The selected area diffraction patterns (SAD) were also taken at corresponding tilt angles. The SAD patterns obtained from the grains identified as G1, G2 and G3 in Fig. 2a with the objective to determine the orientation of individual grains are presented in Fig. 2c–e. The SAD pattern analysis revealed that all the four grains were ori-
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Fig. 1. Light micrographs showing general microstructure of (a) base metal, (b) heat affected zone (HAZ), and (c) weld metal.
Fig. 2. Bright field TEM micrographs showing general microstructure of the base metal. (a) Non-equiaxed ferrite structure, (b) ferrite region containing high density of dislocations and dislocation substructures, and (c–e) SAD patterns obtained from grains identified as G1, G2 and G3 in image (a).
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Fig. 3. (a and b) Bright field TEM micrographs exhibit ferrite grains separated by wavy or curved grain boundary and (c) SAD pattern analysis corresponding to the image presented in (a).
ented close to [2 1 0] direction, suggesting that these grains form small angle or low angle grain boundaries with their neighboring grains. Fig. 3a and b shows massive ferrite grains separated by sharp edged wavy or curved grain boundary and containing high dislocation density. The SAD pattern analysis presented in Fig. 3c indicated that the ferrite phase shown in Fig. 3a corresponds to [0 0 1] zone axis. High density of dislocations and dislocation substructures in coarse ferrite grains are characteristics of massive ferrite [21]. Fig. 4a and b are bright field TEM micrographs of ferrite grains with acicular morphology. Majority of the acicular ferrite grains were observed at grain boundaries, while some were observed extending into the polygonal ferrite grains. They are identified with arrows in Fig. 4. These observations lead us to suggest that acicular ferrite nucleates at austenite grain boundaries and presumably, on non-metallic inclusions within the austenite grain. The SAD patterns obtained from individual grains identified as G1, G2 and G3 in Fig. 4b are presented in Fig. 4c–e. Their analysis revealed that all the three grains were oriented in [3 1 1] direction suggesting that the acicular ferrite
grains also form low angle grain boundaries similar to quasipolygonal ferrite grains presented above. Fig. 5a shows another ferrite region where the ferrite grains exhibit lath morphology. Such elongated parallel ferrite crystals with high dislocation density is characterized as bainitic ferrite [21]. Fig. 5b is a bright field TEM micrograph showing the non-lamellar pearlite colony that was occasionally present in the ferrite matrix of the base metal. The colony of non-parallel and broken up cementite platelets in ferrite matrix is characterized as degenerated pearlite [22,23]. Representative bright field TEM micrographs of HAZ are presented in Fig. 6a–d. It can be seen that the ferrite microstructure is complex and contained high density of dislocation density and substructure (Fig. 6a and b) similar to the base metal. In general, HAZ microstructure contained higher fraction of coarse grained polygonal ferrite (Fig. 6c and d) as compared to base metal. The HAZ microstructure did not contain significant fraction of acicular or bainitic ferrite and degenerated pearlite. However, dispersion of small cementite particles/platelets was observed throughout the ferrite matrix.
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Fig. 4. Bright field transmission electron micrographs (a and b) showing ferrite grains with acicular morphology in the microstructure of base metal and (c–e) SAD patterns corresponding to the ferrite grains identified as G1, G2 and G3 in image (b).
3.3. Precipitation: base metal Figs. 7 and 8 are representative bright field micrographs of precipitates in the base metal. The precipitates were classified as having spherical, cuboidal (∼20–100 nm) and fine (<10 nm) morphology. The precipitates identified as (i)–(iv) in Fig. 7a and b were noted to contain Ti, Nb and some Mo. Fig. 8a and b are bright and dark field TEM micrographs illustrating
the precipitation in ferrite and the corresponding SAD pattern is presented as inset in Fig. 8b. The SAD pattern analysis revealed that the fine precipitates in Fig. 8a were MC type of cubic carbides, while those in Fig. 8c were characterized as fine cementite (Fig. 8d). The MC type of carbides exhibited a [0 0 1]␣ //[0 1 1]MC Baker–Nutting orientation relationship with the ferrite matrix and the cementite precipitates exhib¯ Fe C //[0 1 1]α orientation relationship with the ferrite. ited [1 0 1] 3
Fig. 5. Bright field TEM micrographs showing (a) ferrite region with lath morphology and (b) degenerated pearlite.
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Fig. 6. Bright field transmission electron micrographs of HAZ (a and b) showing ferrite microstructure with dislocation substructures and (c and d) coarse ferrite grains with high dislocation density.
The precipitates observed in the HAZ were similar to the base metal in terms of size, morphology and distribution. 3.4. Microstructure: weld metal The microstructure of the weld metal was similar to the base metal and predominantly consisted of non-equiaxed fine grained ferrite with small volume fraction of degenerated pearlite, bainite, and martensite. Representative TEM micrographs at low magnification showing the ferrite microstructure containing a number of sub-boundaries and significant amount of dislocations and dislocation substructures are presented in Fig. 9. In general, weld metal contained higher dislocation density as compared to the base metal. This may be a consequence of the deformation introduced by thermal stresses in the steel during the welding process. In a manner similar to the base metal characterization, to reveal image details due to different contrast among the grains, the micrographs presented in Fig. 9a and b were obtained at different tilt angles. Also, the SAD patterns were taken at corresponding tilt angles. The SAD patterns were obtained from the individual grains identified as G1, G2, G3 and G4 in Fig. 9a to determine the orientation of individual grains and
are presented in Fig. 9c–f. The pattern analysis indicated that all four individual grain orientations were close to [3 1 1] direction and they form low angle grain boundaries with their neighboring grains. Similar orientation of individual grain obtained by the SAD pattern for the above mentioned grains suggests that the mis-orientation angle between them is within few degrees. Therefore, the boundary between such grains is characterized as low angle grain boundary. Fig. 10a and b presents bright field TEM micrographs of acicular ferrite grains in the weld metal. It is evident that the acicular ferrite grains are surrounded by large clusters of dislocations as compared to the base metal microstructure. Another characteristic feature of weld metal microstructure was dispersion of coarse precipitates along the grain boundary as well as within the grains, as shown in Fig. 10c. The SAD pattern obtained for these precipitates is presented in Fig. 10d and the analysis indicated that these precipitates were cementite particles. The orientation relationship between cementite and ferrite was [1¯ 0 2]Fe3 C //[2¯ 0 1]α . Fig. 11a and b are bright field TEM micrographs of twinned ferrite and degenerated pearlite, respectively. Fig. 12a and b shows bright and dark field micrographs of bainite in weld metal. Fig. 13a and b are bright field
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coarse precipitates observed in the weld metal. The EDS analysis revealed that these precipitates were rich in niobium and titanium and contained some molybdenum. Fig. 15a and b shows bright and dark field TEM micrographs of fine-scale precipitation in ferrite. The precipitates were characterized by fine (∼15–20 nm) morphology. The SAD pattern of fine precipitates (Fig. 15a) is presented as inset in Fig. 15b which indicated that these fine precipitates were MC type of cubic carbides. The MC type of carbides exhibited [0 1 1]MC //[0 0 1]␣ Baker–Nutting orientation relationship with the ferrite matrix. A dark field image of discrete precipitates of carbides (>5 nm) on dislocation is shown in Fig. 15c. 4. Discussion 4.1. Microstructures of Nb-microalloyed pipeline steel
Fig. 7. (a and b) Bright field TEM micrographs of coarse precipitates of base metal identified as (i), (ii), (iii) and (iv) (arrowed).
TEM micrographs of martensite islands in the ferrite matrix of weld metal. The SAD pattern obtained for the twinned martensite (Fig. 13b) is presented in Fig. 13c and the pattern analysis indicated the presence of austenite phase in between the martensite needles. Kurdjumov–Sachs orientation relationship [1¯ 1 1]γ //[0 1 1]α was obeyed between austenite and martensite. 3.5. Precipitation: weld metal The coarse precipitates similar to those observed in the base metal did not seem to have influenced the welding process. Fig. 14a and b presents representative bright field micrographs of
The microstructure of base, HAZ and weld metal of the investigated Nb-microalloyed pipeline steel is complex and consisted of quasi-polygonal or massive ferrite, bainitic ferrite, acicular ferrite, degenerated pearlite, bainite, and martensite–austenite constituent (M/A) in the order of decreasing transformation temperature [13–17,21]. In general, the microstructure of base, and weld metal predominantly consisted of ferrite of mixed morphologies together with small volume fraction of degenerated pearlite. The microstructure of HAZ was predominantly a combination of fine and coarse grained polygonal ferrite structure with dispersion of small cementite particles in it. On the other hand, the weld metal contained small volume fraction of bainite and M/A constituent, high dislocation density, sub-boundaries and dislocation substructures. The difference among the base, HAZ and the weld microstructures is attributed to the effect of heat flow and the cooling rate experienced by the steel during the welding process. In practice, the gap between the components to be joined is filled with layers of weld deposits, and during deposition of each layer, involves heating of the underlying microstructure. The temperature is high enough to cause reformation of austenite, which transforms to different microstructure during cooling part of the thermal cycle [24]. Also, the welding process introduces high residual stresses, which may influence the evolution of microstructure [24]. It is possible that the high residual and thermal stresses (which may exceed the yield stress) result in the overall microstructure characterized by high dislocation density in the weld zone as compared to the base metal. The different ferrite microstructures observed are individually discussed in the following paragraphs. The conventional acicular ferrite that nucleates on non-metallic inclusions is characterized by needle like grains [25]. The pipeline steel researchers have accepted the AF microstructure in hot-rolled plates as an independent microstructure different from the conventional one [14,19]. (i) Quasi-polygonal ferrite or massive ferrite The non-equiaxed ferrite microstructure presented in Figs. 2 and 9 is characterized as quasi-polygonal ferrite or massive ferrite. Quasi-polygonal ferrite is the first high temperature (below Ae3, equilibrium temperature) ferrite
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Fig. 8. (a and b) Bright and dark field TEM micrographs showing fine precipitates in ferrite, the inset in (b) shows SAD pattern obtained for precipitates shown in (a), (c) bright field TEM micrograph showing high density dislocation network containing fine precipitates, and (d) SAD pattern analysis for fine precipitates shown in (c).
phase to form during continuous cooling. It nucleates heterogeneously at the boundaries of the austenite grains. It is reconstructive transformation involving diffusion of the atoms, such that the grains of ferrite grow freely across the austenite grain boundaries [24]. Massive ferrite (Fig. 3) also forms at the transformation temperature range similar to quasi-polygonal ferrite during continuous cooling. The transformation occurs when the cooling is rapid enough such that the partitioning of solute atoms is minimized. The ferrite grows by diffusion and has the same composition as the austenite. The grains formed under these conditions are coarse and are therefore referred as massive ferrite transformation [14,21]. In contrast to the polygonal ferrite, the massive ferrite is characterized by the presence of high dislocation density [21]. (ii) Acicular or bainitic ferrite Acicular ferrite grains were observed in the base (Fig. 4) and weld (Fig. 10) metal. Acicular ferrite forms in the intermediate transformation temperature range during continuous cooling. It nucleates heterogeneously on austenite grain boundaries and non-metallic inclusions. The growth of acicular ferrite is diffusionless, with carbon partitioning in the
austenite after the transformation. It is a displacive transformation involving coordinated movement of atoms. The growth of acicular ferrite is accompanied by an invariantplane strain shape deformation. Since the shape change accommodates plastic deformation, the transformation is accompanied by the formation of high dislocation density in acicular ferrite and the residual austenite [21,24]. Bainitic ferrite (Fig. 5a) also forms in the intermediate transformation temperature range similar to acicular ferrite. The nucleation and growth mechanism of bainitic ferrite is same as that of the acicular ferrite formation involving displacive transformation. The morphology differs from the acicular ferrite because bainitic ferrite initially nucleates only at austenite grain boundaries and continues to grow by repeated formation of subunits, to generate sheaf/lath morphology, where as acicular ferrite nucleates mostly intragranularly at non-metallic inclusions [24]. (iii) Microphases The small amount of degenerated pearlite (Figs. 5b and 11b), bainite (Fig. 12), and martensite–austenite constituent (M/A) (Fig. 13) observed in the investigated steels can be considered as ‘microphases’ [24]. The fraction of
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Fig. 9. Bright field TEM micrographs of the weld metal showing general microstructure. (a) Non-equiaxed ferrite grain structure, (b) ferrite region containing high density of dislocations and dislocation substructures, and (c–f) SAD patterns corresponding to the ferrite grains identified as G1–G4 in micrograph (a).
untransformed microstructure after all the major phases (quasi-polygonal ferrite or massive ferrite, acicular or bainitic ferrite) have formed 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 primary phase. (a) Degenerated pearlite: It 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 [22]. 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 [23]. 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. Similar observations of degenerated pearlite formation have been reported in thermomechanically processed pipeline steels [1,14]. (b) Bainite: 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 [14]. (c) Martensite–austenite constituent (M/A): When the austenite is deformed in the non-crystallized region, large number of sub-boundaries and dislocations are introduced, which increase the nucleation rate of ferrite. As the ferrite nucleates and grows, partition-
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Fig. 10. (a and b) Bright field TEM micrographs of the weld metal showing ferrite grains with acicular morphology surrounded by high density of dislocations, (c) ferrite region containing cementite precipitates and high density of dislocations, and (d) SAD pattern corresponding to the region shown in (c).
Fig. 11. Bright field TEM micrographs of the weld metal zone. (a) Twinned ferrite and (b) degenerated pearlite.
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Fig. 12. (a) Bright and (b) dark field micrographs of bainite region of the weld metal.
ing of carbon takes place at austenite–ferrite interface. During accelerated cooling, part of the carbon enriched austenite would transform to martensite and the retained austenite would coexist, normally in
between the martensite laths or needles and it is referred as martensite–austenite constituent (M/A) [1,9,13,14]. In the present Nb-microalloyed steel, austenite phase of M/A constituent was identified through SAD pat-
Fig. 13. Transmission electron micrographs showing martensite–austenite constituent (M/A) in the weld metal. (a and b) Bright field micrographs of twinned martensite and (c) SAD pattern corresponding to the region shown in (b).
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or massive ferrite, acicular or bainitic ferrite, degenerated pearlite or martensite–austenite constituent in the order of increasing cooling rate and decreasing transformation temperature. Similar transformation kinetics (CCT diagram) was reported by Zhao et al. [13] for low carbon Mn–Nb–Mo acicular ferrite pipe steels. Alloying elements used in the investigated steel such as Mn, Nb, Ti and Mo play an important role in formation of 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 [13,14]. Similarly, addition of Mo 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, Mo strongly restrains the growth of polygonal ferrite and accelerates the formation of acicular ferrite [10,13,14,26]. Thus, the above results suggest that the complex microstructures observed in the investigated base, HAZ and weld metal are a consequence of different cooling rate experienced by the steel during thermomechanical processing and subsequent continuous/accelerated cooling.
Fig. 14. (a) Bright field TEM micrograph of coarse precipitates together with the EDX analysis of precipitates identified as (i) rectangular, (ii) cuboidal, (iii) spherical and (v) irregular morphologies. (b) Bright field TEM micrograph showing high density of coarse carbonitride precipitates in ferrite.
tern (Fig. 13c) obtained across the twinned martensite (Fig. 13b). Martensite–austenite constituent (M/A) is considered to play an important role in obtaining high strength–toughness combination in pipeline steels [9,14]. However, in the investigated steel, it was identified only in the welded zone. Xiao et al. [14] have reported the phase transformation kinetics diagram (CCT and TTT) and microstructures of low carbon Mn–Mo–Nb-microalloyed pipe steel. The results illustrated the formation of different microstructures such as quasi-polygonal
4.1.1. Low angle grain boundaries (LAGBs) The fine grained ferrite microstructure obtained after thermomechanical processing in Nb-microalloyed steel contained large number of low angle grain boundaries (LAGBs). This was confirmed by the analysis of electron diffraction patterns obtained on the group of individual grains which exhibited similar grain orientation. For ferrite grains shown in Figs. 2 and 4 of the base metal and in Fig. 9 of the weld metal, the orientations were [2 1 0], [3 1 1] and [3 1 1], respectively. The [3 1 1] grain orientation was predominant in the base and weld microstructure, among the orientations of various grains analyzed. This suggests the transformation of the deformation texture of austenite, {1 1 0}1 1 2, to {3 3 2}1 1 3 orientation in the transformed ferrite [27]. From the similar orientation of the grains, it can be stated that the steel exhibits a kind of local texture or mesotexture. Bhattacharjee et al. [28,29] have reported that the thermomechanically controlled rolled (TMCR) microalloyed steel can give rise to group of grains oriented close to one crystallographic direction and exhibit mesotexture. They have quantified mis-orientation angles of grain boundaries in TMCR steels and found that about 65% of grain boundaries had mis-orientation angle less than 10◦ . They also pointed out that a propagating crack front ignores the low angle grain boundary and propagates through it treating a group of grains as a single grain. Thus, the microstructure with LAGBs may be detrimental to toughness of pipeline steels. However, Zhong et al. [9] observed through in situ TEM studies that the LAGB with M/A film around it experienced similar effect on dislocation motion and crack propagation behavior as in high angle grain boundary (HAGB), in a ultra-fine grained microalloyed pipeline steel. This observation suggests that M/A film at the grain boundary can be an effective barrier to dislocation motion and crack propagation.
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Fig. 15. (a and b) Bright and dark field TEM micrographs showing MC type of carbide precipitates in ferrite. The inset in (b) shows the SAD pattern obtained for MC precipitate in (a). (c) Dark field TEM micrograph showing discrete carbides on dislocation network.
4.1.2. Curved grain boundaries Massive ferrite microstructure constituent was characterized by curved grain boundaries, high dislocation density and dislocation substructures (Fig. 3) in the base metal. The SAD pattern (Fig. 3c) obtained across these curved boundaries did not show any characteristic orientation relationship among the grain pairs. Partitioning of interstitials or substitutional atoms at the migrating austenite–ferrite interface may cause irregular growth and jogged grain boundaries [14,21], Massive transformation can result in the rapid growth of some grains and their interfaces can be facet-free with smooth curvatures [30,31]. In polycrystalline metals, flat and curved grain boundaries are characterized as singular and rough boundaries, respectively. The curved boundary is atomically rough and the singular boundary is atomically smooth [32,33]. When a flat surface undergoes roughening transition, it develops curves, this transition normally occurs at high temperatures [32,33]. A curved boundary may have the shape of ‘hill and valley’, where the shape distortion of grain boundary occurs due to impinging of precipitates that are coherent with one of the grain pairs [33]. In present study, it is possible that the curved boundaries (hill and valley type) are formed in the massive ferrite as a consequence of impingement of fine carbides (NbC) that are formed during TMP along the grain boundaries that subsequently cause migration of interfaces at different velocities.
4.2. Precipitation in Nb-microalloyed pipeline steel The size and morphology of the precipitates observed in the present pipeline steel are similar to those observed in conventionally hot-rolled microalloyed steels and pipeline steels processed through CSP thin-slab technology [34–42]. 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] [34,36,37]. However, EDS analysis indicated that some of the coarse precipitates in the investigated steel were rich in Ti and had small concentrations of Mo (Figs. 7a and 14a). There are no reports in the literature that Mo can dissolve in (Ti, Nb) carbides. The presence of Mo could come from a small Mo carbide particles, heterogeneously nucleated on the Nb carbides or it may be thermodynamically possible for Mo to substitute for Nb in the particles at low temperatures in this steel. A similar observation was made by Misra et al. in ultra-high strength hotrolled steel [34]. It is possible that titanium precipitates formed at high temperature provided nucleation site for molybdenum precipitates. Multi-microalloying design approach such as Nb–Ti, Ti–V, and Nb–V generally results in the formation of duplex carbonitrides [39]. Based on recent solubility product calculations carried out in the temperature range of 700–1300 ◦ C [36], it was shown that the microalloying elements, Ti, Nb,
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are interchangeable in the precipitate lattice because of similarity in crystal structures (fcc NaCl type) and lattice parameter [38,39]. Funakawa et al. [43] reported that (Ti, Mo)C precipitates contained equal amount of titanium and molybdenum in hot-rolled low carbon steel. Lee et al. [11] also observed that MC type of carbides contained Nb and Mo in a Nb–Mo hot-rolled HSLA steel. The present results suggest that similar to Ti–Nb, alloying elements, Ti–Mo and Nb–Mo may also be interchangeable in the precipitate lattice when both elements are present in MC (Ti, Mo)C type of precipitates. Fine carbides observed in the ferrite phase of base steel (Fig. 8a) and welded zone (Fig. 15a) exhibited Baker–Nutting (B–N) orientation relationship. In general, strain induced carbonitride precipitates of Nb in austenite exhibit cube–cube lattice orientation relationship and the austenite transforms to ferrite with Kurdjumove–Sachs relationship; subsequently precipitate formed in the austenite would exhibit K–S relationship with ferrite. The precipitates that nucleate in ferrite would exhibit B–N orientation relationship [11,44], The present results are consistent with the results reported in the literature. Our recent work on conventionally processed V–Nb–Ti and V–steels and on microalloyed pipeline steels processed through CSP thin-slab technology [42] exhibited similar kind of precipitation behavior and a detailed discussion on crystallography and ordering of the multi-microalloying precipitates has been previously discussed [35,36].
5. Conclusions The microstructure of thermomechanically processed Nbmicroalloyed pipeline steel primarily consisted of non-equiaxed ferrite of mixed morphologies together with small volume fraction of degenerated pearlite. The microstructure contained significant amount of dislocation density, sub-boundaries and dislocation substructures. The microstructure of HAZ contained fine and coarse grained polygonal ferrite structure with small cementite particles dispersed in it. In addition, the weld microstructure contained small volume fraction of bainite and martensite–austenite constituent (M/A) islands. The microstructural constituents include acicular ferrite, massive ferrite or quasi-polygonal ferrite, bainitic ferrite and dispersion of coarse and fine cementite particles in the ferrite matrix. The complex microstructures observed in the Nb-microalloyed steel is attributed to the alloy chemistry and varied cooling rate experienced by the steel. Electron diffraction pattern analysis indicated that the ferrite microstructure contained large number of low angle grain boundaries (LAGBs). The precipitates in the Nb-microalloyed pipeline steel were of duplex type containing Ti and Nb, or Ti and Mo even at low concentration of titanium present in the steel. The precipitates were rectangular, cuboidal/spherical, fine and very fine with size range of ∼100–130 nm, 20–100 nm, 10–20 nm and <10 nm, respectively. The alloying elements (Ti, Nb, and Mo) form MC type of carbides in the ferrite matrix. The observed high strength–toughness combination is a consequence of fine grained ferrite microstructure of mixed morphologies and precipitation strengthening.
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Microstructure of high strength niobium-containing pipeline steel S. Shanmugam a , R.D.K. Misra a,∗ , J. Hartmann b , S.G. Jansto c a
Center for Structural and Functional Materials, Department of Chemical Engineering, University of Louisiana at Lafayette, Lafayette, LA 70504-4130, USA b Mittal Steel, Indiana Harbor Works, 3001 Dickey Road, East Chicago, IN 46312, USA c Reference Metals, 1000 Old Pond Road, Bridgeville, PA 15017, USA Received 26 July 2006; accepted 4 August 2006
Abstract The paper describes the microstructural constituents in a industrially processed Nb-microalloyed pipeline steel having yield strength of ∼620 MPa. The microstructure of base, heat affected zone (HAZ), and weld metal of the fabricated steel pipe was examined by optical and transmission electron microscopy. The microstructure of thermomechanically processed pipeline steel primarily consisted of non-equiaxed ferrite of mixed morphologies with small fraction of degenerated pearlite. The microstructure contained high dislocation density, sub-boundaries and dislocation substructures. The HAZ was characterized by a combination of fine and coarse grained polygonal ferrite structure with high density of dislocations and fine cementite particles. In the weld metal, the constituents of complex ferrite were low temperature transformation products formed during continuous cooling such as quasi-polygonal or massive ferrite, acicular ferrite, bainitic ferrite and dispersion of coarse and fine cementite particles in the ferrite matrix. The precipitates in the investigated pipeline steel were of duplex type containing either Nb and Ti or Ti and Mo, even though the steel contained low concentration of titanium. Precipitates of different morphology and size range were observed and include rectangular (∼100–130 nm), cuboidal/spherical (∼20–100 nm), fine (∼10–20 nm) and very fine (<10 nm). They were MC type of carbides. © 2006 Elsevier B.V. All rights reserved. Keywords: Pipeline steels; Microstructure; Niobium
1. Introduction Modern pipeline steels are currently being produced to the strength level required for API grades using accelerated cooling and controlled rolling technology. The primary interest is to obtain the best possible combination of strength and toughness [1]. An increase in the strength of these steels is necessary for deep water applications and for remote areas where pipes experience transmission pressures exceeding 2500 psi. Weldability, formability, fracture toughness and transition temperature [2], resistance to hydrogen induced blister cracking in sour service environment [3,4], stress corrosion cracking resistance for underground service, especially with H2 S environment [5–7], and fatigue resistance are the additional requirements for transmission of oil and gas through pipelines [8,9]. Thus, it is imperative that the steel is characterized by a microstructure that provides the necessary strength–toughness combination. The microstructure depends on the alloy chemistry, thermo∗
Corresponding author. E-mail address: [email protected] (R.D.K. Misra).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.08.017
mechanical processing and degree of cooling. Alloy additions such as Mn, Nb, V, Ti, Mo, Ni and Cu are commonly added in pipeline steels in order to achieve the desired microstructure and mechanical properties [4,5,10–12]. However, judicious selection of alloying elements is necessary to obtain their beneficial effect on the evolution of microstructure and consequent on mechanical properties. For instance, Nb- and Mo-containing steels are commonly used in pipeline applications because it is observed that the HSLA steels containing Nb and Mo exhibit superior strength and toughness combination as compared to the HSLA steels containing Nb and V [11,12]. Also, careful control of alloying element content is equally important, according to the specific application of pipelines. For example, Mn is an important alloying element for solid solution strengthening, however, reduced Mn content in steels decreases the centerline microstructural banding [13]. Also, Mn content of greater than 0.3% in pipeline steels causes hydrogen induced blister cracking on being subjected to sour service environment. Thus, pipeline steels containing low Mn content with additional strength obtained from Cu has been considered for severe service condition [4].
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Controlled thermomechanical processing is another aspect of high strength pipeline steel for modern API grade steels [2,10]. The process parameters such as soaking temperature, rolling temperature, finishing temperature, and cooling rate play an important role in the ultimate microstructure. Higher cooling rate and lower interrupted cooling temperature are the parameters that determine the formation of acicular ferrite-based microstructure, a preferred microstructure for pipeline steels [10]. Different combinations of microstructures have been obtained in thermomechanically processed pipeline steels in laboratory simulated experiments and include ferrite–pearlite [7], polygonal ferrite–acicular ferrite [1], acicular ferrite–bainite [10], acicular ferrite–martensite–austenite constituent (M/A) [14], a mixed microstructure [8,15–17], acicular ferrite [6,14,18,19], bainite [3]. It is reported that a predominant acicular ferrite microstructure with martensite/austenite (M/A) islands, as a second phase exhibits optimum mechanical properties [1,10]. The acicular ferrite dominated microstructure can be obtained by optimizing processing conditions and without the need for significant alloying concentration [1,2,10,20]. The objective of this study is to report initial results concerning microstructure and precipitation behavior of industrially processed 630 MPa Nb-containing microalloyed pipeline steel that contributed to good strength–toughness combination. The precipitates were characterized in terms of morphology, size, chemistry, and crystallography. Furthermore, the orientation relationship between precipitates and ferrite were determined. 2. Experimental The Nb-microalloyed pipeline steels discussed here were industrial heats that were continuously cast and hot-rolled to a minimum of ∼8 mm gage and subsequently fabricated to ∼450 mm diameter pipes. The chemical composition of the investigated steel is given in Table 1. The thermomechanical processing details are not described here due to proprietary reason. Standard tensile tests were conducted at room temperature on longitudinal specimens machined according to ASTM E8 specification (dimensions 225 mm × 12.5 mm, gage length, 50 mm) using computerized tensile testing system. The initial crosshead speed was 4.2 × 10−2 mm/s up to a strain of ∼5%, then the speed was increased to 4.2 × 10−1 mm/s. Impact toughness was meaTable 1 Chemical composition of Nb-containing pipeline steel in wt.% C Mn Si P S Al Ti Nb Ca Mo
0.04/0.06 1.4/1.6 0.18/0.25 0.020 0.003/0.005 0.02/0.06 0.005/0.015 0.05/0.085 0.001/0.004 0.10/0.30
Table 2 Mechanical properties of Nb-containing pipeline steel Properties
Longitudinal direction
Transverse direction
Yield strength (MPa) Tensile strength (MPa) % Elongation Charpy v-notch impact toughness at 0 ◦ C (J)
613–659 692–716 26–30 –
591–593 722–734 28–29 173–266
sured using standard Charpy v-notch impact test (ASTM 23) at 0 ◦ C. Transmission electron microscopy was carried out on thin foils prepared by cutting thin wafers from the steel samples of base, HAZ and weld metal, 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. Foils were examined by HITACHI 7600 TEM operated at 120 kV. 3. Results 3.1. Mechanical properties The yield strength, tensile strength and % elongation data obtained in longitudinal and transverse direction for Nbmicroalloyed pipeline steels are summarized in Table 2. The longitudinal and transverse yield strength was in the range of 613–659 MPa and 591–593 MPa, respectively. The tensile strength in the longitudinal and transverse orientation was 692–716 MPa and 722–734 MPa, respectively. The transverse Charpy v-notch impact toughness of Nb-microalloyed pipe steel at 0 ◦ C was high and in the range of 173–266 J. 3.2. Microstructure: base metal and HAZ Microstructural studies on base, HAZ and weld metal of the fabricated pipe are important to identify any differences in microstructural constituents. Representative light micrographs of base, HAZ and weld metal are presented in Fig. 1a–c. The microstructure was predominantly fine ferrite structure in all the three regions, except that the HAZ contained large fraction of coarse polygonal ferrite grain structure (Fig. 1b). Representative low magnification bright field TEM micrographs of the base metal are presented in Figs. 2 and 3. It may be seen that the ferrite microstructure is complex and consists of non-equiaxed ferrite grains (Fig. 2a), sub-boundaries and dislocation substructures (Fig. 2b). These micrographs were obtained at different tilt angles (foil tilt) in order to reveal the image details due to different contrast among the grains. The selected area diffraction patterns (SAD) were also taken at corresponding tilt angles. The SAD patterns obtained from the grains identified as G1, G2 and G3 in Fig. 2a with the objective to determine the orientation of individual grains are presented in Fig. 2c–e. The SAD pattern analysis revealed that all the four grains were ori-
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Fig. 1. Light micrographs showing general microstructure of (a) base metal, (b) heat affected zone (HAZ), and (c) weld metal.
Fig. 2. Bright field TEM micrographs showing general microstructure of the base metal. (a) Non-equiaxed ferrite structure, (b) ferrite region containing high density of dislocations and dislocation substructures, and (c–e) SAD patterns obtained from grains identified as G1, G2 and G3 in image (a).
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Fig. 3. (a and b) Bright field TEM micrographs exhibit ferrite grains separated by wavy or curved grain boundary and (c) SAD pattern analysis corresponding to the image presented in (a).
ented close to [2 1 0] direction, suggesting that these grains form small angle or low angle grain boundaries with their neighboring grains. Fig. 3a and b shows massive ferrite grains separated by sharp edged wavy or curved grain boundary and containing high dislocation density. The SAD pattern analysis presented in Fig. 3c indicated that the ferrite phase shown in Fig. 3a corresponds to [0 0 1] zone axis. High density of dislocations and dislocation substructures in coarse ferrite grains are characteristics of massive ferrite [21]. Fig. 4a and b are bright field TEM micrographs of ferrite grains with acicular morphology. Majority of the acicular ferrite grains were observed at grain boundaries, while some were observed extending into the polygonal ferrite grains. They are identified with arrows in Fig. 4. These observations lead us to suggest that acicular ferrite nucleates at austenite grain boundaries and presumably, on non-metallic inclusions within the austenite grain. The SAD patterns obtained from individual grains identified as G1, G2 and G3 in Fig. 4b are presented in Fig. 4c–e. Their analysis revealed that all the three grains were oriented in [3 1 1] direction suggesting that the acicular ferrite
grains also form low angle grain boundaries similar to quasipolygonal ferrite grains presented above. Fig. 5a shows another ferrite region where the ferrite grains exhibit lath morphology. Such elongated parallel ferrite crystals with high dislocation density is characterized as bainitic ferrite [21]. Fig. 5b is a bright field TEM micrograph showing the non-lamellar pearlite colony that was occasionally present in the ferrite matrix of the base metal. The colony of non-parallel and broken up cementite platelets in ferrite matrix is characterized as degenerated pearlite [22,23]. Representative bright field TEM micrographs of HAZ are presented in Fig. 6a–d. It can be seen that the ferrite microstructure is complex and contained high density of dislocation density and substructure (Fig. 6a and b) similar to the base metal. In general, HAZ microstructure contained higher fraction of coarse grained polygonal ferrite (Fig. 6c and d) as compared to base metal. The HAZ microstructure did not contain significant fraction of acicular or bainitic ferrite and degenerated pearlite. However, dispersion of small cementite particles/platelets was observed throughout the ferrite matrix.
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Fig. 4. Bright field transmission electron micrographs (a and b) showing ferrite grains with acicular morphology in the microstructure of base metal and (c–e) SAD patterns corresponding to the ferrite grains identified as G1, G2 and G3 in image (b).
3.3. Precipitation: base metal Figs. 7 and 8 are representative bright field micrographs of precipitates in the base metal. The precipitates were classified as having spherical, cuboidal (∼20–100 nm) and fine (<10 nm) morphology. The precipitates identified as (i)–(iv) in Fig. 7a and b were noted to contain Ti, Nb and some Mo. Fig. 8a and b are bright and dark field TEM micrographs illustrating
the precipitation in ferrite and the corresponding SAD pattern is presented as inset in Fig. 8b. The SAD pattern analysis revealed that the fine precipitates in Fig. 8a were MC type of cubic carbides, while those in Fig. 8c were characterized as fine cementite (Fig. 8d). The MC type of carbides exhibited a [0 0 1]␣ //[0 1 1]MC Baker–Nutting orientation relationship with the ferrite matrix and the cementite precipitates exhib¯ Fe C //[0 1 1]α orientation relationship with the ferrite. ited [1 0 1] 3
Fig. 5. Bright field TEM micrographs showing (a) ferrite region with lath morphology and (b) degenerated pearlite.
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Fig. 6. Bright field transmission electron micrographs of HAZ (a and b) showing ferrite microstructure with dislocation substructures and (c and d) coarse ferrite grains with high dislocation density.
The precipitates observed in the HAZ were similar to the base metal in terms of size, morphology and distribution. 3.4. Microstructure: weld metal The microstructure of the weld metal was similar to the base metal and predominantly consisted of non-equiaxed fine grained ferrite with small volume fraction of degenerated pearlite, bainite, and martensite. Representative TEM micrographs at low magnification showing the ferrite microstructure containing a number of sub-boundaries and significant amount of dislocations and dislocation substructures are presented in Fig. 9. In general, weld metal contained higher dislocation density as compared to the base metal. This may be a consequence of the deformation introduced by thermal stresses in the steel during the welding process. In a manner similar to the base metal characterization, to reveal image details due to different contrast among the grains, the micrographs presented in Fig. 9a and b were obtained at different tilt angles. Also, the SAD patterns were taken at corresponding tilt angles. The SAD patterns were obtained from the individual grains identified as G1, G2, G3 and G4 in Fig. 9a to determine the orientation of individual grains and
are presented in Fig. 9c–f. The pattern analysis indicated that all four individual grain orientations were close to [3 1 1] direction and they form low angle grain boundaries with their neighboring grains. Similar orientation of individual grain obtained by the SAD pattern for the above mentioned grains suggests that the mis-orientation angle between them is within few degrees. Therefore, the boundary between such grains is characterized as low angle grain boundary. Fig. 10a and b presents bright field TEM micrographs of acicular ferrite grains in the weld metal. It is evident that the acicular ferrite grains are surrounded by large clusters of dislocations as compared to the base metal microstructure. Another characteristic feature of weld metal microstructure was dispersion of coarse precipitates along the grain boundary as well as within the grains, as shown in Fig. 10c. The SAD pattern obtained for these precipitates is presented in Fig. 10d and the analysis indicated that these precipitates were cementite particles. The orientation relationship between cementite and ferrite was [1¯ 0 2]Fe3 C //[2¯ 0 1]α . Fig. 11a and b are bright field TEM micrographs of twinned ferrite and degenerated pearlite, respectively. Fig. 12a and b shows bright and dark field micrographs of bainite in weld metal. Fig. 13a and b are bright field
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coarse precipitates observed in the weld metal. The EDS analysis revealed that these precipitates were rich in niobium and titanium and contained some molybdenum. Fig. 15a and b shows bright and dark field TEM micrographs of fine-scale precipitation in ferrite. The precipitates were characterized by fine (∼15–20 nm) morphology. The SAD pattern of fine precipitates (Fig. 15a) is presented as inset in Fig. 15b which indicated that these fine precipitates were MC type of cubic carbides. The MC type of carbides exhibited [0 1 1]MC //[0 0 1]␣ Baker–Nutting orientation relationship with the ferrite matrix. A dark field image of discrete precipitates of carbides (>5 nm) on dislocation is shown in Fig. 15c. 4. Discussion 4.1. Microstructures of Nb-microalloyed pipeline steel
Fig. 7. (a and b) Bright field TEM micrographs of coarse precipitates of base metal identified as (i), (ii), (iii) and (iv) (arrowed).
TEM micrographs of martensite islands in the ferrite matrix of weld metal. The SAD pattern obtained for the twinned martensite (Fig. 13b) is presented in Fig. 13c and the pattern analysis indicated the presence of austenite phase in between the martensite needles. Kurdjumov–Sachs orientation relationship [1¯ 1 1]γ //[0 1 1]α was obeyed between austenite and martensite. 3.5. Precipitation: weld metal The coarse precipitates similar to those observed in the base metal did not seem to have influenced the welding process. Fig. 14a and b presents representative bright field micrographs of
The microstructure of base, HAZ and weld metal of the investigated Nb-microalloyed pipeline steel is complex and consisted of quasi-polygonal or massive ferrite, bainitic ferrite, acicular ferrite, degenerated pearlite, bainite, and martensite–austenite constituent (M/A) in the order of decreasing transformation temperature [13–17,21]. In general, the microstructure of base, and weld metal predominantly consisted of ferrite of mixed morphologies together with small volume fraction of degenerated pearlite. The microstructure of HAZ was predominantly a combination of fine and coarse grained polygonal ferrite structure with dispersion of small cementite particles in it. On the other hand, the weld metal contained small volume fraction of bainite and M/A constituent, high dislocation density, sub-boundaries and dislocation substructures. The difference among the base, HAZ and the weld microstructures is attributed to the effect of heat flow and the cooling rate experienced by the steel during the welding process. In practice, the gap between the components to be joined is filled with layers of weld deposits, and during deposition of each layer, involves heating of the underlying microstructure. The temperature is high enough to cause reformation of austenite, which transforms to different microstructure during cooling part of the thermal cycle [24]. Also, the welding process introduces high residual stresses, which may influence the evolution of microstructure [24]. It is possible that the high residual and thermal stresses (which may exceed the yield stress) result in the overall microstructure characterized by high dislocation density in the weld zone as compared to the base metal. The different ferrite microstructures observed are individually discussed in the following paragraphs. The conventional acicular ferrite that nucleates on non-metallic inclusions is characterized by needle like grains [25]. The pipeline steel researchers have accepted the AF microstructure in hot-rolled plates as an independent microstructure different from the conventional one [14,19]. (i) Quasi-polygonal ferrite or massive ferrite The non-equiaxed ferrite microstructure presented in Figs. 2 and 9 is characterized as quasi-polygonal ferrite or massive ferrite. Quasi-polygonal ferrite is the first high temperature (below Ae3, equilibrium temperature) ferrite
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Fig. 8. (a and b) Bright and dark field TEM micrographs showing fine precipitates in ferrite, the inset in (b) shows SAD pattern obtained for precipitates shown in (a), (c) bright field TEM micrograph showing high density dislocation network containing fine precipitates, and (d) SAD pattern analysis for fine precipitates shown in (c).
phase to form during continuous cooling. It nucleates heterogeneously at the boundaries of the austenite grains. It is reconstructive transformation involving diffusion of the atoms, such that the grains of ferrite grow freely across the austenite grain boundaries [24]. Massive ferrite (Fig. 3) also forms at the transformation temperature range similar to quasi-polygonal ferrite during continuous cooling. The transformation occurs when the cooling is rapid enough such that the partitioning of solute atoms is minimized. The ferrite grows by diffusion and has the same composition as the austenite. The grains formed under these conditions are coarse and are therefore referred as massive ferrite transformation [14,21]. In contrast to the polygonal ferrite, the massive ferrite is characterized by the presence of high dislocation density [21]. (ii) Acicular or bainitic ferrite Acicular ferrite grains were observed in the base (Fig. 4) and weld (Fig. 10) metal. Acicular ferrite forms in the intermediate transformation temperature range during continuous cooling. It nucleates heterogeneously on austenite grain boundaries and non-metallic inclusions. The growth of acicular ferrite is diffusionless, with carbon partitioning in the
austenite after the transformation. It is a displacive transformation involving coordinated movement of atoms. The growth of acicular ferrite is accompanied by an invariantplane strain shape deformation. Since the shape change accommodates plastic deformation, the transformation is accompanied by the formation of high dislocation density in acicular ferrite and the residual austenite [21,24]. Bainitic ferrite (Fig. 5a) also forms in the intermediate transformation temperature range similar to acicular ferrite. The nucleation and growth mechanism of bainitic ferrite is same as that of the acicular ferrite formation involving displacive transformation. The morphology differs from the acicular ferrite because bainitic ferrite initially nucleates only at austenite grain boundaries and continues to grow by repeated formation of subunits, to generate sheaf/lath morphology, where as acicular ferrite nucleates mostly intragranularly at non-metallic inclusions [24]. (iii) Microphases The small amount of degenerated pearlite (Figs. 5b and 11b), bainite (Fig. 12), and martensite–austenite constituent (M/A) (Fig. 13) observed in the investigated steels can be considered as ‘microphases’ [24]. The fraction of
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Fig. 9. Bright field TEM micrographs of the weld metal showing general microstructure. (a) Non-equiaxed ferrite grain structure, (b) ferrite region containing high density of dislocations and dislocation substructures, and (c–f) SAD patterns corresponding to the ferrite grains identified as G1–G4 in micrograph (a).
untransformed microstructure after all the major phases (quasi-polygonal ferrite or massive ferrite, acicular or bainitic ferrite) have formed 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 primary phase. (a) Degenerated pearlite: It 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 [22]. 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 [23]. 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. Similar observations of degenerated pearlite formation have been reported in thermomechanically processed pipeline steels [1,14]. (b) Bainite: 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 [14]. (c) Martensite–austenite constituent (M/A): When the austenite is deformed in the non-crystallized region, large number of sub-boundaries and dislocations are introduced, which increase the nucleation rate of ferrite. As the ferrite nucleates and grows, partition-
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Fig. 10. (a and b) Bright field TEM micrographs of the weld metal showing ferrite grains with acicular morphology surrounded by high density of dislocations, (c) ferrite region containing cementite precipitates and high density of dislocations, and (d) SAD pattern corresponding to the region shown in (c).
Fig. 11. Bright field TEM micrographs of the weld metal zone. (a) Twinned ferrite and (b) degenerated pearlite.
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Fig. 12. (a) Bright and (b) dark field micrographs of bainite region of the weld metal.
ing of carbon takes place at austenite–ferrite interface. During accelerated cooling, part of the carbon enriched austenite would transform to martensite and the retained austenite would coexist, normally in
between the martensite laths or needles and it is referred as martensite–austenite constituent (M/A) [1,9,13,14]. In the present Nb-microalloyed steel, austenite phase of M/A constituent was identified through SAD pat-
Fig. 13. Transmission electron micrographs showing martensite–austenite constituent (M/A) in the weld metal. (a and b) Bright field micrographs of twinned martensite and (c) SAD pattern corresponding to the region shown in (b).
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or massive ferrite, acicular or bainitic ferrite, degenerated pearlite or martensite–austenite constituent in the order of increasing cooling rate and decreasing transformation temperature. Similar transformation kinetics (CCT diagram) was reported by Zhao et al. [13] for low carbon Mn–Nb–Mo acicular ferrite pipe steels. Alloying elements used in the investigated steel such as Mn, Nb, Ti and Mo play an important role in formation of 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 [13,14]. Similarly, addition of Mo 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, Mo strongly restrains the growth of polygonal ferrite and accelerates the formation of acicular ferrite [10,13,14,26]. Thus, the above results suggest that the complex microstructures observed in the investigated base, HAZ and weld metal are a consequence of different cooling rate experienced by the steel during thermomechanical processing and subsequent continuous/accelerated cooling.
Fig. 14. (a) Bright field TEM micrograph of coarse precipitates together with the EDX analysis of precipitates identified as (i) rectangular, (ii) cuboidal, (iii) spherical and (v) irregular morphologies. (b) Bright field TEM micrograph showing high density of coarse carbonitride precipitates in ferrite.
tern (Fig. 13c) obtained across the twinned martensite (Fig. 13b). Martensite–austenite constituent (M/A) is considered to play an important role in obtaining high strength–toughness combination in pipeline steels [9,14]. However, in the investigated steel, it was identified only in the welded zone. Xiao et al. [14] have reported the phase transformation kinetics diagram (CCT and TTT) and microstructures of low carbon Mn–Mo–Nb-microalloyed pipe steel. The results illustrated the formation of different microstructures such as quasi-polygonal
4.1.1. Low angle grain boundaries (LAGBs) The fine grained ferrite microstructure obtained after thermomechanical processing in Nb-microalloyed steel contained large number of low angle grain boundaries (LAGBs). This was confirmed by the analysis of electron diffraction patterns obtained on the group of individual grains which exhibited similar grain orientation. For ferrite grains shown in Figs. 2 and 4 of the base metal and in Fig. 9 of the weld metal, the orientations were [2 1 0], [3 1 1] and [3 1 1], respectively. The [3 1 1] grain orientation was predominant in the base and weld microstructure, among the orientations of various grains analyzed. This suggests the transformation of the deformation texture of austenite, {1 1 0}1 1 2, to {3 3 2}1 1 3 orientation in the transformed ferrite [27]. From the similar orientation of the grains, it can be stated that the steel exhibits a kind of local texture or mesotexture. Bhattacharjee et al. [28,29] have reported that the thermomechanically controlled rolled (TMCR) microalloyed steel can give rise to group of grains oriented close to one crystallographic direction and exhibit mesotexture. They have quantified mis-orientation angles of grain boundaries in TMCR steels and found that about 65% of grain boundaries had mis-orientation angle less than 10◦ . They also pointed out that a propagating crack front ignores the low angle grain boundary and propagates through it treating a group of grains as a single grain. Thus, the microstructure with LAGBs may be detrimental to toughness of pipeline steels. However, Zhong et al. [9] observed through in situ TEM studies that the LAGB with M/A film around it experienced similar effect on dislocation motion and crack propagation behavior as in high angle grain boundary (HAGB), in a ultra-fine grained microalloyed pipeline steel. This observation suggests that M/A film at the grain boundary can be an effective barrier to dislocation motion and crack propagation.
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Fig. 15. (a and b) Bright and dark field TEM micrographs showing MC type of carbide precipitates in ferrite. The inset in (b) shows the SAD pattern obtained for MC precipitate in (a). (c) Dark field TEM micrograph showing discrete carbides on dislocation network.
4.1.2. Curved grain boundaries Massive ferrite microstructure constituent was characterized by curved grain boundaries, high dislocation density and dislocation substructures (Fig. 3) in the base metal. The SAD pattern (Fig. 3c) obtained across these curved boundaries did not show any characteristic orientation relationship among the grain pairs. Partitioning of interstitials or substitutional atoms at the migrating austenite–ferrite interface may cause irregular growth and jogged grain boundaries [14,21], Massive transformation can result in the rapid growth of some grains and their interfaces can be facet-free with smooth curvatures [30,31]. In polycrystalline metals, flat and curved grain boundaries are characterized as singular and rough boundaries, respectively. The curved boundary is atomically rough and the singular boundary is atomically smooth [32,33]. When a flat surface undergoes roughening transition, it develops curves, this transition normally occurs at high temperatures [32,33]. A curved boundary may have the shape of ‘hill and valley’, where the shape distortion of grain boundary occurs due to impinging of precipitates that are coherent with one of the grain pairs [33]. In present study, it is possible that the curved boundaries (hill and valley type) are formed in the massive ferrite as a consequence of impingement of fine carbides (NbC) that are formed during TMP along the grain boundaries that subsequently cause migration of interfaces at different velocities.
4.2. Precipitation in Nb-microalloyed pipeline steel The size and morphology of the precipitates observed in the present pipeline steel are similar to those observed in conventionally hot-rolled microalloyed steels and pipeline steels processed through CSP thin-slab technology [34–42]. 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] [34,36,37]. However, EDS analysis indicated that some of the coarse precipitates in the investigated steel were rich in Ti and had small concentrations of Mo (Figs. 7a and 14a). There are no reports in the literature that Mo can dissolve in (Ti, Nb) carbides. The presence of Mo could come from a small Mo carbide particles, heterogeneously nucleated on the Nb carbides or it may be thermodynamically possible for Mo to substitute for Nb in the particles at low temperatures in this steel. A similar observation was made by Misra et al. in ultra-high strength hotrolled steel [34]. It is possible that titanium precipitates formed at high temperature provided nucleation site for molybdenum precipitates. Multi-microalloying design approach such as Nb–Ti, Ti–V, and Nb–V generally results in the formation of duplex carbonitrides [39]. Based on recent solubility product calculations carried out in the temperature range of 700–1300 ◦ C [36], it was shown that the microalloying elements, Ti, Nb,
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are interchangeable in the precipitate lattice because of similarity in crystal structures (fcc NaCl type) and lattice parameter [38,39]. Funakawa et al. [43] reported that (Ti, Mo)C precipitates contained equal amount of titanium and molybdenum in hot-rolled low carbon steel. Lee et al. [11] also observed that MC type of carbides contained Nb and Mo in a Nb–Mo hot-rolled HSLA steel. The present results suggest that similar to Ti–Nb, alloying elements, Ti–Mo and Nb–Mo may also be interchangeable in the precipitate lattice when both elements are present in MC (Ti, Mo)C type of precipitates. Fine carbides observed in the ferrite phase of base steel (Fig. 8a) and welded zone (Fig. 15a) exhibited Baker–Nutting (B–N) orientation relationship. In general, strain induced carbonitride precipitates of Nb in austenite exhibit cube–cube lattice orientation relationship and the austenite transforms to ferrite with Kurdjumove–Sachs relationship; subsequently precipitate formed in the austenite would exhibit K–S relationship with ferrite. The precipitates that nucleate in ferrite would exhibit B–N orientation relationship [11,44], The present results are consistent with the results reported in the literature. Our recent work on conventionally processed V–Nb–Ti and V–steels and on microalloyed pipeline steels processed through CSP thin-slab technology [42] exhibited similar kind of precipitation behavior and a detailed discussion on crystallography and ordering of the multi-microalloying precipitates has been previously discussed [35,36].
5. Conclusions The microstructure of thermomechanically processed Nbmicroalloyed pipeline steel primarily consisted of non-equiaxed ferrite of mixed morphologies together with small volume fraction of degenerated pearlite. The microstructure contained significant amount of dislocation density, sub-boundaries and dislocation substructures. The microstructure of HAZ contained fine and coarse grained polygonal ferrite structure with small cementite particles dispersed in it. In addition, the weld microstructure contained small volume fraction of bainite and martensite–austenite constituent (M/A) islands. The microstructural constituents include acicular ferrite, massive ferrite or quasi-polygonal ferrite, bainitic ferrite and dispersion of coarse and fine cementite particles in the ferrite matrix. The complex microstructures observed in the Nb-microalloyed steel is attributed to the alloy chemistry and varied cooling rate experienced by the steel. Electron diffraction pattern analysis indicated that the ferrite microstructure contained large number of low angle grain boundaries (LAGBs). The precipitates in the Nb-microalloyed pipeline steel were of duplex type containing Ti and Nb, or Ti and Mo even at low concentration of titanium present in the steel. The precipitates were rectangular, cuboidal/spherical, fine and very fine with size range of ∼100–130 nm, 20–100 nm, 10–20 nm and <10 nm, respectively. The alloying elements (Ti, Nb, and Mo) form MC type of carbides in the ferrite matrix. The observed high strength–toughness combination is a consequence of fine grained ferrite microstructure of mixed morphologies and precipitation strengthening.
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