Microstructure and mechanical properties of Nb-API X70 low carbon steel

1 Microstructure and Mechanical Properties on plates of low-carbon Nb API 1 X70 Pipeline Steels 2 F. F. Conde 1, F. J. Pina¹, J. M. Giarola¹, G. S. Pereira¹, J.C. Francisco2 *J.A. Avila3, W. W. 3 Bose Filho¹ 4 1. USP – University of São Paulo, Av. Trabalhador São Carlense 400, São Carlos, SP, 13566- 5 590, Brazil. 6 2. Mechanical Engineering Department, UTFPR, Cornélio Procópio, Brazil 7 3. UNESP – São Paulo State University (UNESP), Campus of São João da Boa Vista, São João 8 da Boa Vista, SP, Brazil 9 *Corresponding author: * Prof. Dr. Julian A. Avila D. [email protected], +55 19 36382432, 10 Av. Profª Isette Corrêa Fontão, 505, Jardim das Flores, 13876-750 - São João da Boa Vista, SP, 11 Brazil. 12 13 14 Abstract 15 The oil and gas production has grown steadily, and the care with its transportation grid 16 needs to be maintained and expanded. The pipeline used in the transport of gas and oil 17 by-product require materials with high mechanical strength, toughness, and fatigue 18 resistance. In the production of low-carbon steels, the alloying elements and 19 thermomechanical processing aim to produce tough microstructures, usually finish 20 rolling at a temperature where the austenite or the austenite + ferrite phases are present, 21 to finally, during cooling, produce a fine ferrite and bainite. This study compares the 22 fracture toughness of the API 5L X70 steel between two plates of steel with different Nb 23 content, 0.06 and 0.09 Nb weight %. The mechanical response was correlated to the 24 grain size and distribution. The crack tip opening displacement (CTOD) parameter was 25 used to assess the fracture toughness at 25°C and -60°C. The microstructures and 26 microconstituents of the two plates of steel were similar. The matrices were composed 27 of ferrite, with bands of secondary products, including degenerated pearlite. Results 28 showed a prone to formation of a bimodal microstructure, coarse and fine grains, with 29 the addition of Nb. High-Nb steel presented better fracture toughness than normal-Nb 30 steel at low-temperature. 31 Keywords: fracture toughness; mechanical properties; low-carbon Nb steel; API X70 32 steel 33 34 1. Introduction 35 Pipelines can be located on land or submerged in seas, oceans, rivers, lakes, and 36 they are used for the transportation of petroleum and its derivatives, such as crude oil, 37 2 natural gas, liquid, and gaseous products. Pipeline grids are the most economical 38 transportation alternative [1–3]. Therefore, the development of steels for pipelines during 39 the last decades has been driven by the need for steels with improved combinations of 40 high strength, toughness and weldability at affordable prices [4,5]. To achieve these 41 goals, typical pipeline steels of API (American Petroleum Institute) grade, such as API 42 X70, up to API X120, rely upon the chemical alloying design and Thermo-Mechanical 43 Controlled Processing (TMCP) to provide high strength, toughness and ductility [6–8]. 44 The TMCP and alloying design help to improve mechanical properties by grain 45 refinement. Microalloying elements, such as Nb, V and Ti, cause significant increase of 46 the mechanical strength of steels by precipitation in the matrix, anchoring the austenite 47 grain boundaries in the austenite domain, resulting in a fine-grain microstructure after 48 cooling down[7,8]. The use of Nb has increased the hydrogen-induced cracking 49 resistance [9], improve mechanical strength, toughness, crack propagation resistance 50 [10], maintaining the ductility [11–13] and increasing hardenability, i.e., increasing 51 martensite content raising the martensite transformation starting (Ms) temperature [10]. 52 Nb in solid solution retards the pearlite and bainite reaction, stabilizing austenite phase, 53 and NbC precipitates act as nucleation sites for martensite and bainite [14]. Thus, 54 modern steel allows producing pipelines with thin walls maintaining the same working 55 pressure than pipes fabricated with older technologies, which results in a decrease of 56 weight and a reduction of the amount of deposited weld on each joint, lowering the overall 57 cost [5]. 58 Fracture toughness and crack propagation in steels are a function of their 59 chemical composition [7,15,16], microstructures [7,16–19], grain morphology [7,16], 60 inclusions, precipitates [7,15,16] and residual stresses induced after the TMCP process 61 [20]. Despite many studies regarding X70 and Nb influence on steels, most of them only 62 present tensile tests as mechanical assessment, and a few presents Charpy impact tests 63 impact as toughness information [20]. Crack Tip Opening Displacement (CTOD) tests 64 are well established as a highly reliable test for fracture toughness assessment [21]. A 65 pipeline under operative conditions is subjected to quasi-static stress and forces rather 66 than instantaneous impact loads [21]. CTOD tests are designed with samples geometry 67 and size able to guarantee a constant crack-tip triaxiality [22], with a crack propagation 68 on opening mode I and in quasi-static conditions. The use of CTOD to evaluate the effect 69 of Nb content in X70 pipelines, has not been documented in the literature to the best of 70 our knowledge. 71 This study aimed to conduct fracture mechanism assessment using the CTOD 72 parameter in an API 5L X70 steel with different Nb contents. Hence, we point towards to 73 determine whether Nb affects fracture toughness. Samples with 0.06 and 0.09 Nb weight 74 3 % were tested at 25°C and -60°C temperatures. Toughness results were discussed 75 alongside with the microstructural characterization, comparing the grain size and its 76 distribution. The microstructural analysis was conducted using the optical microscopy 77 (OM), scanning electron microscopy (SEM) and electron backscatter diffraction (EBSD). 78 79 80 2. Experimental Procedure 81 82 Plates of microalloyed steels API 5L X70, manufactured using hot strip mill 83 process, were used in the present study. The first material is called normal-Nb (0.06 wt% 84 Nb) and was obtained from 500 mm long, 500 mm wide, and 16 mm-thick sheets. The 85 second material with high Nb (0.09 wt% Nb), considered in the text as high-Nb, was 86 obtained from a 14 mm-thick helical welded pipe (SAW) with a diameter of 823 mm and 87 a length of 1500 mm, with the subsequent cutting process to obtain the plates. The 88 chemical composition of the studied steels was obtained through an atomic absorption 89 spectrophotometer analysis and is shown in table 1 and meets the requirement of API 90 5L X70 steel according to the standard [8]. 91 92 Table 1. Chemical composition (Wt %) steel the API 5L X70. * Unless otherwise agreed, 93 the sum of the niobium, vanadium, and titanium concentrations shall be < 0,15 %. 94 Elements C Nb Ti Mn Al Ni Si P S Cr Mo Fe Normal-Nb 0.07 0.06 0.021 1.77 0.033 0.01 0.22 0.019 0.004 0.15 0.06 Rel. High-Nb 0.04 0.09 0.017 1.65 0.033 0.26 0.23 0.017 0.004 0.24 0.01 Rel. API 5L X70 %maximum 0.18 * * 1.80 - - 0.45 0.025 0.015 - - Rel. 95 The TMCP scheme is illustrated in Figure 1. Processing involved heating over 96 1250 °C for full homogenization above the recrystallization temperature, dissolving all 97 nitrides and carbides, maintaining Nb in solid solution, followed by rough milling. The 98 finishing rolling temperature for the normal-Nb was at 750 °C, and for high-Nb was at 99 900 °C, near full austenitization. To produce a pipe of the high-Nb steel was necessary, 100 the plate free of deformation, which is the microstructure without strain-working as the 101 following step of the UOE process for pipe production. Both plates of steel were cooled 102 in the laminar section down to ~570 °C at approximately 10 °C/s, followed by air cooling. 103 104 105 4 106 Figure 1. a) Schematic thermomechanical processing schedule for normal and high Nb 107 steels. Both conditions were hot-rolled at 1250 °C and then normal-Nb with finish rolling 108 pass at 750 °C and high-Nb at 900 °C. In b) ThermoCalc simulation of a Fe-0.12C- 109 1.52Mn-0.058Nb-0.019Ti adapted from [22]. 110 111 The metallographic analysis was conducted on normal and high-Nb steel to 112 identify the rolling direction, and three pieces were cut in the transverse, normal, and 113 rolling planes. Samples were ground from 80 to 1200-grit SiC emery paper, polished with 114 CrO3 particles, according to the ASTM E3 standard [23]. The etching was conducted with 115 2 % HNO3 mixed with ethanol. Grain size measurements were performed using three 116 procedures: first, a microstructure analysis directly from the images obtained by OM 117 coupled to a computerized image analysis system; then, SEM images were taken using 118 an image analyzer. 119 Grain size analysis proved to be difficult to be evaluated by optical microscopy. 120 Therefore, EBSD maps were used instead. The Inverse Pole Figure (IPF) generated 121 from the EBSD technique represents a specific crystallographic orientation by a single 122 color. This single-colored region can be used as a crystallographic domain serving 123 effective grain size [24]. The detection of the bimodal microstructure was conducted 124 considering contours with orientation differences equal to or greater than 15°. A 125 comparative graphic was built. Low-angle grain-boundaries were found higher than 2° 126 and lower than 15° and computed from the EBSD data set. 127 Hardness assessment was conducted using five measurements for each sample 128 in the rolling and transverse planes, which were represented by the plane orthogonal 129 vector RD and TD. Mechanical properties were assessed with tensile tests in rectangular 130 dog-bone samples with a gauge section of 32 mm in length, and cross-section of 6x3 131 mm2, according to the ASTM E8 standard [25]. Samples were tested with longitudinal 132 and transverse samples orientation according to the rolling directions. 133 5 CTOD tests were conducted according to ASTM E1820 [26] in Single-edge- 134 notched bend - SE(B) - samples loaded in three-point bending, with notches oriented in 135 the L-T and T-L direction [27]. The samples were ground and polished to a mirror-like 136 condition to observe crack the growth during fatigue pre-cracking before CTOD tests. 137 The CTOD test was conducted with an increasing load, cross-head rate of 1 kN min-1 at 138 25°C and -60 °C. Samples were subjected to a soaking time at low temperatures to 139 ensure full specimen homogenization, and a calibrated Type K thermocouple was 140 welded to specimens to verify the temperature during CTOD tests. 141 142 143 3. Results 144 3.1 Microstructure analysis 145 146 Figures 2 show the microstructure of the studied steel. A 3D representation 147 shows the elongated grains due to the rolling process, which is better observed in the 148 detailed micrographs. Figure 2c and f reveals the formation of massive ferrite, polygonal 149 ferrite, and pancaked grains for normal-Nb and high-Nb 150 The normal-Nb steel presented pancaked ferrite and elongated grains mixed with 151 some fine and coarse polygonal ferrites. Some banding in the mid-plate region is 152 observed in Figure 2a-b, composed of secondary phases and microconstituents as 153 pearlite, degenerated pearlite, and martensite-austenite (M-A). A considerable fraction 154 of lath type morphology is also depicted by the formation of martensite and lath bainite, 155 also forming some package, conceiving package bainite, and granular bainite. As 156 normal-Nb was finished rolled at the intercritical field, some pancaked ferrite was 157 produced, and the remainder austenite enhanced the ferrite grain refinement during 158 cooling down to room temperature. 159 The High-Nb steel presented a microstructure mostly composed of polygonal 160 ferrite with some regions of secondary phases and some microconstituents, such as 161 pearlite, degenerated pearlite, and M-A. Hot rolling enabled recrystallization, resulting in 162 some equiaxed grains in the final microstructure, as depicted in Figure 2d-e. This steel 163 did not present banding in the mid-plate but rather only dispersed secondary phases and 164 microconstituents. High-Nb steel was finishing rolled at a higher temperature, above Ar3, 165 so only the austenite phase would be present. This heating led to the coarser grain in 166 comparison to normal-Nb. As the finishing rolling temperature was in the full 167 austenitization field, the subsequently formed ferrite free of deformation was developed 168 with a predominant polygonal morphology. 169 170 6 171 172 Figure 2. The microstructure of the normal-Nb, a), b) and c), and high-Nb, d), e) and f), 173 steels. a) and d) isometric view of the microstructure of the steel plate, b) and e) rolling 174 plane detail, c) and f) zoom-in images of the microstructure with high-angle grain 175 boundary, showing yellow arrows indicating massive ferrite, blue arrows indicating 176 polygonal ferrite and red arrows indicating pancake ferrite. The reference system is 177 based on RD, TD and ND, where the related planes are rolling, transversal and normal 178 planes that lie perpendicular to the respective directions. 179 180 Figure 3 shows the morphology of the grains in the all-Euler orientation maps with 181 the presence of LAGBs and HAGBs, with misorientation angles over 2° and 15°, 182 respectively. Figure 4 shows the grain size results measured by EBSD and optical 183 analysis on micrographs of the steels. Normal-Nb steel presented an amount of over 50 184 % of LAGBs in small grains (up to 4 µm size) and few LAGBs in large grain size. In 185 comparison to high-Nb steel, small grains (up to 4 µm size) accumulated over 80 % of 186 the LAGBs presented in the microstructure. The nearly 30 % of the area being 187 represented by grains larger than 12 µm in high-Nb steel resulted in a bimodal 188 microstructure, with fine grains mixed with coarse grains. This effect was also observed 189 in normal-Nb steel but a smaller proportion. 190 Regarding grain size, considering HAGBs, the high-Nb steel presented a higher 191 mean value of grain size than the normal-Nb steel. A similar distribution of grain size up 192 to 10 µm and a formation of grains larger than 12 µm was found in both plates of steel. 193 The optical measurements only detected the LAGBs due to the contour contrast caused 194 by etching and were performed using SEM images. The results showed a similar trend 195 7 of EBSD, revealing a higher mean value of grain size for high-Nb steel than normal-Nb. 196 The bimodal aspect is more intense on high-Nb steel, presenting a microstructure 197 composed of 41% of its area with grains larger than 10 µm, while normal-Nb shows 38% 198 of its microstructure composed of grains larger than 10 µm. By the optical analytical 199 method, it was observed grains up to 20 µm for high-Nb steel, while normal-Nb presented 200 grains up to 15 µm. 201 202 203 Figure 3: EBSD all Euler angle maps for the high and normal Nb steels. Normal-Nb with 204 a) LAGBs and HAGBs, and b) only contours of the HAGBs. High-Nb steel with c) LAGBs 205 and HAGBs, and d) only contours of the HAGBs. Coarser grains visualized in high-Nb 206 steel in d), revealing bimodal grain size aspect. Low misorientation grain boundaries 207 (LAGBs) are considered > 2°and high misorientation grain boundaries (HAGBs) > 15°. 208 209 210 Figure 4. EBSD and OM analyses for grain size and area distribution, a) normal-Nb steel, 211 and b) high-Nb steel. 212 213 8 214 3.2 Mechanical properties 215 216 Figure 5 shows the tensile test results. The normal-Nb presented a yield and 217 ultimate tensile stress of 556±3 MPa and 662±3 MPa with 25% of elongation, while high- 218 Nb presented yield and ultimate strength of 520±5 MPa and 619±2 MPa respectively, 219 with 30% of elongation, fulfilling the requirement for API 5L X70, PSL2 level grade. The 220 yield stress/ultimate tensile stress (Y/T) ratio shown in Figure 5 is high, above 0.8. Y/T 221 values below 0.7 are related to providing resistance to plastic collapse failures; however, 222 values up to 0.85 are satisfactory for structural steels [28]. 223 224 225 Figure 5. Tensile test results in the longitudinal (DL) and transverse (DT) directions. 226 227 3.3 Simulated Effect of Finishing Rolling Temperature on Elongation and Strength 228 229 The finishing rolling temperature is an essential aspect of the development of the 230 final microstructure, thus, influencing the mechanical properties of the steel. Low- 231 temperature interruption, i.e., temperature before air cooling, promotes the formation of 232 more bainite, which in turn increases the tensile strength of the steel [29], which was the 233 case of the normal-Nb steel in the present study. The previous study reported simulated 234 equations to indicate strength and elongation from applied TMCP temperatures as finish 235 rolling temperature [30]. 6 shows the divergences in mechanical properties between the 236 experimental and simulated results. A regression equation was used to find the yield 237 strength (Equation 1) and elongation (Equation 2) [30]. 238 The main difference between tested and simulated elongation is related to the 239 microstructure. The proposed Equation 2, as follows, shows an inverse relationship with 240 the start and finishing rolling temperature (Ts; Tf), as well as with the final cooling 241 9 temperature (Tc), while the only direct relationship is with the cooling rate (Vc). A ratio of 242 0.98 [30], where s.ΔL is the simulated elongation. 243 244 𝑠. 𝜎𝑌= 0.508𝑇𝑠− 0.231𝑇 𝑓− 0.334𝑇𝑐+ 1.905𝑉 𝑐+323.6 Equation 1 245 𝑠. ∆𝐿 = −0.002𝑇𝑠− 0.064𝑇 𝑓− 0.086 𝑇𝑐+ 0.325𝑉 𝑐+121.8 Equation 2 246 247 Equation 2 considers an increase of elongation towards a residual deformed 248 microstructure with fast cooling rates. Beyond some value, the formation of brittle phases 249 such as martensite is reached. This behavior is contradictory to the obtained results, in 250 which high-Nb presented an excellent ductility. Despite the posterior deformation for pipe 251 production, the prior microstructure was a low deformed polygonal ferrite, and afterward, 252 most of LAGBs were concentrated in grains smaller than 5 µm (Figure 4). This type of 253 microstructure led to a ductile matrix, capable of absorbed higher deformation than 254 normal-Nb steel. Three results showed that normal-Nb had accumulated more 255 deformation and defects up to its final condition than high-Nb. Higher distribution of 256 LAGBs on a wide range of grain size, higher yield strength, and consequently lower 257 elongation compared to high-Nb steel. 258 259 260 Figure 6. Comparison of results of yield, ultimate strength and elongation obtained from 261 the tensile test, and simulated data from literature [30]. Following symbols σy, σu, s.σy, 262 ΔL and s.ΔL represent measured yield strength, measured ultimate strength, simulated 263 yield strength, elongation, and simulated elongation. The prefix s symbolized the 264 simulated results. 265 266 267 16 Figure 2a-c. The residual plastic deformation guaranteed a higher constraint at the crack 457 tip by a saturated microstructure more densely full of dislocations. Thus, at room 458 temperature, good toughness values were obtained for normal-Nb steel, but slightly 459 lower than high-Nb. CTOD curves show more clearly this aspect in Figure 9b, where 460 high-Nb steel achieves higher forces for the same COD variation than normal-Nb. 461 The fracture surface of CTOD tests at low temperatures still presented a diffuse 462 and rough aspect in Figure 8 at a stable crack growth region, contrasting to a plane and 463 polished surface after a brittle failure. This result suggests that the energy necessary for 464 deformation at low temperatures was still lower than to cleavage the microstructure, not 465 triggering brittle crack propagation failure by cleavage at -60 °C, presenting ductile 466 deformation and failure with dimple formation. Bakshi et al. reported X70 steel subjected 467 to Charpy impact tests at 0 °C, -20 °C and -40 °C temperature, and lowering temperature 468 caused no significant change or increase of toughness results for RD sample [18], 469 maintaining ductile fracture aspect. 470 Analyzing the tensile values, the differences between steels of strain-hardening 471 aspect becomes inconclusive as high-Nb presented a Y/T ratio of 84% and normal-Nb 472 steel, 85%. Despite similar chemical composition of the studied steels, a variety of 473 phases and microconstituents occurred, mainly of ferrite and bainite, as suggested [50]. 474 Ligang et al. found that the increase of bainite content causes the increase of mechanical 475 strength, i.e., higher yield and ultimate strength while decreasing the impact energy [50]. 476 The reported results meet fine the presented steels, where the high-Nb had lower carbon 477 content, mainly polygonal morphology suggesting most of the microstructure be 478 composed of ferrite. In comparison, normal-Nb had higher carbon content, dominant 479 elongated, and lath morphology, suggesting pancaked ferrite and bainite. High-Nb 480 presented lower mechanical strength than normal-Nb, but higher fracture toughness, 481 which is related to the impact of energy absorption. 482 483 4.5 Modeling K to predict Bainite content 484 485 Ligang et al. [50] proposed a model to quantify bainite in the microstructure using 486 the mechanical strength of bainite and ferrite, as well as using CVN and KIC results. An 487 initial K value is termed KQ before validating the KIC according to the ASTM standard [51]. 488 Using conversion equations does not imply a KIC validation. Therefore, the converted 489 parameter in the present work is only termed K. Following the values of ferrite and bainite 490 yield strength of 420 and 726 MPa, respectively, attributed in [50] and finding the amount 491 of bainite for high and normal-Nb steels to match the yield strength. The proportional 492 respective of bainite in high and normal-Nb values were 33 and 46 % according to the 493 17 found data; however, this is a simulated data, and many other factors impact the 494 mechanical strength besides bainite formation. Thus, it is believed that the bainite 495 fraction in the presented steels differs from the obtained ones. Nonetheless, this result 496 indicates the formation of higher fractions of bainite and hard phases in normal-Nb steel, 497 which is consistent with the study and indeed presented a higher mechanical strength. 498 The following equations were used to find a simulated value of CVN and K. 499 500 𝐶𝑉𝑁 = [13.64 + 0.068(−28.12𝑓𝐵 + 6.64𝑅𝑎 + 4.48𝑁𝑓 )1.079 ]2.083 Equation 3 501 𝐾 = 58.75𝐶𝑉𝑁0.244 Equation 4 502 503 Where Ra is the particle aspect ratio given by b/a, where b and a are averaged 504 nodule sizes of the hard phase in the rolling and transverse direction, respectively, the 505 neighboring factor is defined as Nf=λL/λT, where λL and λT are average center-to-center 506 spacings of the hard phase in the rolling and the transverse direction, respectively [50]. 507 Substituting Equation 3 and 4, the following Equation 5 is generated. 508 509 𝐾 = 58.75 ∙ [3.773 + 0.255(−28.12𝑓𝐵+ 6.64𝑅𝑎+ 4.48𝑁𝑓)0.5484] Equation 5 510 511 From the obtained equation, values of Ra and Nf were found fixing 0.33 and 0.46 512 of fB, i.e., bainite fraction, to match the collected converted CTOD data to K values 513 according to ASTM1820 standard [51]. The simulated data matched the real data when 514 Ra and Nf were set 10 for high-Nb and 8 for normal-Nb. This behavior would mean that 515 normal-Nb steel presents a less spaced hard phase thought out the microstructure, 516 which would raise mechanical strength as seen. Therefore, this value is somewhat 517 consistent with the obtained results, and the simulated values are plotted in Figure 11. 518 519 520 18 521 Figure 11. Simulated data, according to [50]. In a), a proportional amount for yield 522 strength between ferrite and bainite phases was obtained to match the obtained tensile 523 results for both plates of steel. In b), the use of simulated files for KIC and bainite content 524 to find Ra and Nf. Simulated data shows similar results to the obtained ones. 525 526 527 5. Conclusions 528 The effect of different TMCP and change in the composition of two API X70 plates 529 of steel was studied. The following main conclusions were drawn: 530 • The high-Nb steel was subjected to a finishing rolling temperature at 900 °C 531 which caused the formation of a microstructure composed of ferrite with less 532 internal defects and deformation with bimodal microstructure, i.e., a mixture of 533 fine and coarse grains along with the bulk material with coarse grains larger than 534 12 μm composing over 30 % of the microstructure. A subsequent process of pipe 535 conformation was applied. In the UOE process, ferrite was deformed, and 536 produced of dislocation and defects occurred, and small fine grains (<5 μm) 537 concentrated up to 80 % of the total LAGBs. 538 • The normal-Nb steel was subjected to a finishing rolling temperature at 750 °C, 539 the intercritical field, generating a dominant pancaked and deformed ferrite. 540 LAGBs were distributed along with all sizes of grain in the microstructure. This 541 type of microstructures led to a more limited elongation and lower toughness at 542 room temperature. 543 • Both plates of steel, with slight differences in the microstructure, presented similar 544 mechanical properties. The increased mechanical strength of normal-Nb over 545 high-Nb is likely to be attributed to a lower finishing rolling temperature, more 546 deformed microstructure, and increased content of carbon, leading to the 547 19 formation of pearlite, degenerated pearlite colonies, bainite, and M-A 548 microconstituents. 549 • The CTOD tests showed good values for fracture toughness results for both 550 plates of steel, being that both showed higher toughness in the L-T configuration. 551 This behavior is due to the fact of a residual rolled microstructure, with more 552 HAGBs and LAGBs parallel to the rolling plane, and the crack propagation in 553 transverse direction must overcome more obstacles as grain boundaries. 554 • The change of Nb in the composition did not play a significant role in controlling 555 the fracture toughness value. 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