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Keywords:

  • pipeline steel;
  • ultrafast cooling;
  • precipitation;
  • strength

Abstract

  1. Top of page
  2. Abstract
  3. 1 Introduction
  4. 2 Experimental Procedure
  5. 3 Results
  6. 4 Discussion
  7. 5 Conclusions
  8. Acknowledgments
  9. References

X70 pipe line steel with high mechanical strength and high ductility can be achieved through hot rolling and ultrafast cooling process at a low cost. A large amount of nano-sized precipitates and dislocations are identified near the surface layer of the steel applying ultrafast cooling, which contribute immeasurably to the improvement of mechanical properties. In addition, high cooling rate decreases the mean distance of pearlite laminae, and massive ferrite grains characterized by curved grain boundaries and high dislocation density are also transformed as one of the typical microstructures in the samples applying ultrafast cooling process.

1 Introduction

  1. Top of page
  2. Abstract
  3. 1 Introduction
  4. 2 Experimental Procedure
  5. 3 Results
  6. 4 Discussion
  7. 5 Conclusions
  8. Acknowledgments
  9. References

Modern pipeline steels with the best possible combination of strength and toughness are currently being developed to the strength level required for API grades using thermo-mechanical controlled process (TMCP).[1-4] Meanwhile, saving resources is greatly advocated for sustainable development. Generally, alloy elements such as Nb, Ti, Mo are added in X70 pipeline steels in order to achieve appropriate volume fraction of acicular ferrite and bainite ferrite through conventional TMCP.[5, 6] However, great attentions have been paid on the improvement of mechanical properties through New Generation Thermomechanical Control Process (NG-TMCP) which was characterized by ultrafast cooling in recent years.[7-9] As a result, diversified transformation behaviors and thus improved properties can be simply achieved by controlling cooling routes instead of adding more expensive alloys. The present study is to report initial results about microstructure and precipitation behavior of Nb-bearing pipeline steel applying ultrafast cooling which contributes to good strength-toughness combination.

2 Experimental Procedure

  1. Top of page
  2. Abstract
  3. 1 Introduction
  4. 2 Experimental Procedure
  5. 3 Results
  6. 4 Discussion
  7. 5 Conclusions
  8. Acknowledgments
  9. References

The chemical composition of the investigated steel is given in Table 1. The ingot used was cut from the industrial continuously cast with an initial thickness of 150 mm. After being austenitized at 1473 K for 1 h, the ingot was in turn subjected to two-step rolling at 1373 and 1223 K. The rolling reduction was about 60 and 80%, respectively. After finish rolling at 1173 K, two different cooling strategies were selected for executing of the microstructure control. The first one was laminar flow at an average cooling rate of 15 K s−1 with a finish cooling temperature 923 K (steel A), and the other was ultrafast cooling to 1053 K at a cooling rate of 50 K s−1 (steel B). Then, both steels were air-cooled to room temperature.

Table 1. Chemical composition of Nb-bearing pipeline steel in wt%
CSiMnPSNb
0.05–0.070.14–0.161.40–1.50≤0.01≤0.010.04–0.05

Micrographic observations were performed by scanning electron microscope (SEM, FEI QUANTA 600) operated at 20 kV and transmission electron microscope (TEM, FEI TECNAI G2 20) operated at 200 kV. SEM samples were prepared by mechanical grinding, polishing, and etching with 4% Nital. Thin foils for TEM observations were prepared by cutting into ∼0.4 mm thick slices using wire electrode discharge machine and mechanical grinding down to ∼30 µm thickness, and then were thinned to perforation by a Twin-jet Electro Polisher in a mixed solution of 5% perchloric and 95% acetic acid at 295 K. All the samples were sectioned along the transverse direction of the rolled plate and the surface micrographic observations were positioned at about 3 mm below the plate surface.

Standard tensile tests were conducted at room temperature on both longitudinal and transverse directions according to ASTM E8 specification (dimensions 225 mm × 12.5 mm, gage length, 50 mm) using computerized tensile testing system. The Charpy impact tests were performed at 253 K using the V-notched specimens measuring 10 × 10 × 55 mm3 according to ASTM E23.

3 Results

  1. Top of page
  2. Abstract
  3. 1 Introduction
  4. 2 Experimental Procedure
  5. 3 Results
  6. 4 Discussion
  7. 5 Conclusions
  8. Acknowledgments
  9. References

3.1 Mechanical Properties

The mechanical properties in both the longitudinal and transverse directions of Nb-bearing pipeline steels are summarized in Table 2. Most of the properties for steel B were more or less improved comparing to steel A. The transverse yield strength for steel A and steel B varied in the range of 455–470 MPa and 530–535 MPa, respectively. The transverse tensile strength for both steels was 540–555 MPa and 625–630 MPa, respectively. The Charpy v-notch impact toughness in both directions was almost higher than 300 J even at 253 K.

Table 2. Mechanical properties of Nb-bearing pipeline steels
 Steel ASteel B
Longitudinal directionTransverse directionLongitudinal directionTransverse direction
Yield strength [MPa]510–525455–470545–555530–535
Tensile strength [Mpa]590–600540–555605–615625–630
Total elongation [%]19.0–22.030.5–31.524.5–28.024.0–26.0
Charpy v-notch impact toughness at −20°C [J]307–332261–333309–343313–324

3.2 Microstructure: Surface and Half the Thickness

3.2.1 Ferrite and Pearlite

Figure Figure 1a through Figure Figure 1d show the optical micrographs of steel A and steel B. Major fraction of the microstructures of both steels is non-equiaxed ferrite with some amount of pearlite. Comparing with Figure Figure 1a and b, there is no obvious difference between the microstructures at the surface and half the thickness in steel A. In the case of steel B (Figure Figure 1c and d), both the ferrite grain size and the average block size of pearlite at half the thickness are larger than those near the surface. Meanwhile, the mean grain size in steel A is much smaller than that in steel B, especially that at half the thickness. In addition, both acicular ferrite and non-equiaxed ferrite are observed in steel A. The acicular ferrite grains in Figure 2 exhibit low angle boundaries with the adjacent grains, and high dislocation density inner the grains as presented in Figure 2b. However, amplified TEM micrographs in Figure Figure 3a and b present that dislocation density in steel B is much higher than that in steel A, especially at positions near the surface. This also manifests that ultrafast cooling process is great helpful to strengthen the materials through proliferation of dislocations in matrix without converting other process parameters. Figure Figure 3c and d show TEM micrographs of dislocations at half the thickness of steel A and steel B, respectively. The dislocation morphology seems no significant change at the surface and half the thickness of steel A, whereas the dislocation density becomes much lower at half the thickness of steel B. This means great temperature gradient through-thickness direction existing during the ultrafast cooling process. Moreover, Figure 4a and b show the mean distances of pearlite laminae near the surface of steel A and steel B measured to be about 120 and 60 nm, respectively. The finer the pearlite laminae, the higher the mechanical property. However, the magnified TEM micrographs at half the thickness of steel A and B showing the pearlite morphologies (Figure Figure 4c and d) indicate no significant difference between the pearlite inter-lamellar spacing of two steels. Besides, it is found that one pearlite block at half the thickness of steel B is composed of several pearlite lamellar packets with different orientations.

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Figure 1. Optical morphology at different positions of steel A a) surface, b) half the thickness, and steel B c) surface, d) half the thickness.

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Figure 2. Microstructures of acicular ferrite and dislocation near the surface of steel A.

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Figure 3. Dislocation morphology at different positions in steel A and steel B, a) and b) near the surface, c) and d) at half the thickness.

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Figure 4. Microstructures of pearlite with different lamellar spacing in steel A and steel B.

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Figure 5a shows acicular ferrite grains with high density of dislocations identified as G1 and G2, indicating that small angle or low angle grain boundaries were formed between each other. In addition, the polygonal ferrite with high dislocation density can also be observed. Figure 5b shows massive ferrite grains including high dislocation density in steel B, which were characterized by sharp-edged curved grain boundary. The massive ferrite grains were also reported by Shanmugam et al.[2]

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Figure 5. Typical TEM microstructures at half the thickness of steel A.

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3.2.2 Precipitation

It is interesting that a great number of precipitates were observed in the matrix ferrite grains of steel B as shown in Figure Figure 6a and b, while no obvious precipitates were detected in steel A. The precipitates can be classified into two types according to the distributional pattern. The first one is random precipitates in the matrix ferrite as observed in Figure 6a, and the other is interphase precipitates within pro-eutectoid ferrite (Figure 6b) caused by formation at γ/α transformation interfaces,[10, 11] which probably precipitated during the subsequent re-reddening and slow cooling basing on the great undercooling. The size of the precipitates ranged from 5 to 10 nm. The EDX analysis manifests that most of precipitates were probably cementites consisting of manganese and niobium. And for more detail, further measurement such as electron probe microanalysis (EPMA) should be implemented in order to avoid interference signals from the surrounding matrix.

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Figure 6. Two types of precipitates at surface of steel B and the corresponding EDX analysis.

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There are no visible precipitates in half the thickness of steel A and B. Moreover, remarkable difference in the surface and half the thickness of the samples definitely indicates that grain refinement strengthening and precipitation strengthening should be simultaneously achieved by ultrafast cooling, although more detailed work is required to reveal the relationship between chemical composition, control rolling, and ultrafast cooling process, as well as the related metallurgical principles.

4 Discussion

  1. Top of page
  2. Abstract
  3. 1 Introduction
  4. 2 Experimental Procedure
  5. 3 Results
  6. 4 Discussion
  7. 5 Conclusions
  8. Acknowledgments
  9. References

As is well known, the non-equiaxed ferrite mainly covers quasi-polygonal ferrite and massive ferrite. Since the quasi-polygonal ferrite is formed just below the Ae3 temperature during continuous cooling and the nucleation occurs heterogeneously at the austenite grain boundaries, and it is reconstructive transformation involving diffusion of the atoms,[12] the non-equiaxed ferrite in steel A was recognized as massive ferrite since the massive ferrite is usually characterized by the presence of high dislocation density in contrast to the polygonal ferrite,[13] which is just right in conformity with the microstructures in Figure 5. In addition, the massive ferrite transformation usually occurs when the cooling rate is much higher resulting in the slow-down of solute atoms distribution, and the grains formed under these conditions are coarse.[14, 15]

For steel B, the non-equiaxed ferrite was also recognized as the massive ferrite on account of much higher dislocation density in Figure 3b. However, it can be deduced that the subsequent cooling rate following ultrafast cooling process for steel B was lower than that for steel A, which was the main reason for the appearance of ferrite and pearlite in steel B.

Another typical morphology is acicular ferrite in steel A. The conventional acicular ferrite, which is characterized by needlelike grains,[15] nucleates on non-metallic inclusions and forms in the intermediate transformation temperature range during continuous cooling. It is a displacive transformation involving coordinated movement of atoms, which is similar to banite ferritic transformation. The most important difference between acicular ferritic and bainite ferritic transformations is that the accompanied carbide morphology with the former is much more irregular than that with the latter. The cementite in acicular ferrite is difficult to be identified for very small numbers.

A part of massive ferrite in half the thickness of steel B was characterized by curved grain boundaries and high dislocation density. A curved boundary usually has the shape of “hill and valley,” and it is usually supposed to be a result of transition from flat boundary. Additionally, flat and curved grain boundaries are characterized as singular and rough boundaries in polycrystalline metals, respectively. When a flat surface undergoes roughening transition, it develops curves, and this transition normally occurs at high temperatures.[16, 17] In the current study, the shape distortion of grain boundary occurs in the massive ferrite probably due to the impinging of precipitates (MC) that are coherent with one of the grain pairs.[2]

As stated previously, the most possible CCT curves for steel A and B can be summarized as Figure 7. The first kind of precipitation in steel B shown as Figure 6a was generated by the oversaturation of carbon due to ultrafast cooling and subsequently precipitated on the high density defects, which was also similarly detected by Pereloma et al.[18] The interphase precipitation is always obtained in vanadium alloyed carbon steels with isothermal transformation at about 873 K.[19] However, there are still few investigations about the interphase precipitation of Fe3C in Nb-bearing steels. The NANO-HITEN steel proposed by Funakawa et al.[20] with tensile strength of up to 780 MPa was also prepared through ultrafast cooling and isothermal coiling at different temperatures, and the strengthening mechanism was attributed to the interphase nanometer-sized carbides too. Hence, the interphase nanometer-sized cementite in the surface of steel B was probably a result of ultrafast cooling combining with subsequent air cooling. In addition, the microstructure evolving behavior is also dependent on the finishing temperature and re-reddening temperature of ultrafast cooling. In the present study, the finishing temperature of ultrafast cooling was definitely lower than A1 temperature which represents the decomposition of pearlite or cementite. However, further studies are still required to reveal the crystallograph relationship between the two types of carbides and ferrite matrix, also the effects of fraction, size, and distribution of carbides on other mechanical properties should be further discussed.

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Figure 7. Inferential CCT curve and cooling route for steel A and steel B.

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5 Conclusions

  1. Top of page
  2. Abstract
  3. 1 Introduction
  4. 2 Experimental Procedure
  5. 3 Results
  6. 4 Discussion
  7. 5 Conclusions
  8. Acknowledgments
  9. References

High mechanical strength and high ductility of Nb-bearing pipeline steel was simultaneously obtained with nanometer-sized cementite produced through hot rolling and ultrafasting cooling processes. The mechanical property can be improved to X70 grade (API Spec 5L-2004) by precipitation and dislocation strengthening even if the chemical composition was only alloyed with niobium and the matrix morphology was conventional ferrite and pearlite.

Acknowledgments

  1. Top of page
  2. Abstract
  3. 1 Introduction
  4. 2 Experimental Procedure
  5. 3 Results
  6. 4 Discussion
  7. 5 Conclusions
  8. Acknowledgments
  9. References

Project (HIT.NSRIF.2011102) supported by Natural Scientific Research Innovation Foundation in Harbin Institute of Technology, and Project (2012DXGJ09) supported by the Development of Science and Technology Plan of Weihai.

References

  1. Top of page
  2. Abstract
  3. 1 Introduction
  4. 2 Experimental Procedure
  5. 3 Results
  6. 4 Discussion
  7. 5 Conclusions
  8. Acknowledgments
  9. References