热轧管线用钢板中的针状铁素体组织素体(中英文翻译).doc

热轧管线用钢板中的针状铁素体组织素体通过连续冷却曲线(CCT)和热加工模拟试验，大批量的管线钢中的过冷奥氏体转变得到进一步研究。基于研究的结果，人们提出一种能够生产出针状铁素体控制的混合微观结构的热变形控制过程(TMCP)。结果显示：在现有实验条件下冷却速度的增加可以改善最后微观组织中的针状铁素体的百分比率。因此，针状铁素体所控制的微观结构可以通过两个阶段来完成，即奥氏体再结晶合肥再结晶区的控制轧制阶段及以30k/s冷却速度进行的控制冷却阶段。介绍针状铁素体（AF）具有相当高的强度、硬度，因此它主要用于管线用钢生产。自20世纪70年代以来，人们对高强度、低合金管线用钢做了大量的关于针状铁素体的工作。现在已被证实针状铁素体来源于一个混合组织和剪切变形模型，开始的温度比变换区域的临界稍高，它包含一个共梧或半共梧的奥氏体铁素体界面，这个界面遵循NWKS 规律。迄今为止，主要对TiO针状铁素体钢和大量焊接金属做了大量关于针状铁素体的工作。日本钢铁市场声称TiO是特色组织，当一段时间以后针状铁素体核在非金属包含物中的奥氏体细化物中产生。热轧板带钢中的针状铁素体不同于前面所提到的这些，因为它不需要一个特殊的化学成分。Tanaka 指出对于含有0.07C2.0Mn0.6Nb0.5Mo的针状铁素体钢支配微观组织，这种组织优化了机械性能，它包含在不平衡奥氏体和马氏体的稳定区域中的细非等轴铁素体。此外，针状铁素体的微观特性表现在一个不平衡奥氏体区域中的交汇处，而这个区域是由一些过量的高密度、杂乱的碳化物组成。最近，针状铁素体支配微观组织已经形成。正如一个传统的微观组织，比如在热轧管线用钢板中的一个铁素体珠光体微观组织。在以前的学术中已经忽略了在热轧管线钢板中针状铁素体的形成。为了增加工业产品和有针状铁素体支配微观结构的管线钢的应用。通过CCT 曲线、热模拟试验和理想的热变形控制过程（TMCP），针状铁素体在现有的商业管线钢中的形成已经得到研究。实验步骤在当前的研究中，这种材料用于石油、天然气输送管道用钢。它的化学成分是Fe0.076C0.24Si1.33Mn0.014P0.0032S0.03Al0.0048N0.04Nb0.02Ti0.04V。过冷奥氏体的形成是由（FF）向支配的。试样用于淬火长10mm，直径3mm的圆管是由实验厚钢板连铸中切割出来的。以.0.1、0.4、1.、3、4、10、50和100k/s的成线性冷却速度被用于构造CCT曲线。这个热模拟试验由1500mm热循环模拟装置操作，分别在奥氏体再结晶区、在奥氏体非再结晶区、混合奥氏体再结晶合非再结晶区中形成。热模拟试验试件是高20mm，直径8mm，它是由30mm30mm的连续锻造出的试样钢，并且加热3分钟到1150。热轧试验是由两个直径为37mm 轧辊的轧机操作的。规格6080140的热轧试样是由连续的试样钢切割出来并且最终轧制成8mm的钢板。表1 列出了热轧试验中的压下率分配。微观组织的由光学显微镜观测的，显微镜型号为S360SEM 和JEM2000FXTEM。对于光学显微镜和SEM，试样被机械地磨光，并在3%的溶液中蚀刻。对于TEM 观察，金属片是由300um厚盘提供的。它第一次被机械切割达到厚50um，然后在10%和90%的酸溶液中，用两个喷气的电动磨光器磨光。结论和分析 过冷奥氏体的转变在钢铁产业中，CCT 曲线对于获得微观结构和获得设计中最适应的TMCP 是有很大用途的。图1显示：冷却速率分别为0.1、1、10和50k/s时所对应的微观结构照片。图2显示当试件首先在1150奥氏体化10分钟，然后在以复杂的冷却速度冷却，图2也表明了在这种情况下运用线性冷却速度膨胀方法和金相学构造CCT 曲线。在冷却速度为50k/s时，试样微观组织的转变主要在之上，它占有平行铁素体组织区域和碳化物相混合的边界区域。另外，先前的粗状奥氏体边界网状组织可以被清除的看到。冷却速度低于10k/s时，试样的微观组织被针状铁素体所占据。与相比较，针状铁素体表现出一种特别的不规则的结构。这种结构含有大量微小的尺寸，以混乱的方式，错综复杂的方向分配着。同时，先前的粗状奥氏体边界网状组织已经消失，而碳化物被分配在铁素体区域，当冷却速度减少到1k/s，试样的微观组织包含着多边奥氏体及许多来自单独生长在奥氏体稳定区域和铁素铁区域中的伪珠光体。当冷却速度减少到1k/s 以下时，试样的微观组织包含着多边奥氏体和许多珠光体。图2中的CCT 曲线显示非等轴奥氏体针状铁素体转换在实验中发生的在冷却速度为110k/s。此外，一个占有针状铁素体的微观组织会在冷却速度为10k/s中获得。冷却速度低于0.4k/s会促使主要产生一种针状珠光体和珠光体的形成物。当冷却速度高于0.4k/s珠光体消失时，多边形奥氏体也消失。当试样在以一个较宽的冷却速度范围0.41k/s 连续冷却时，获得的微观组织不可能消除为伪珠光体，这个微观组织不是一个理想的区域，因为机械性能对它有决定性影响。根据前面所讨论的，对于试验钢在CCT 中的过冷奥氏体的冷却速度在110k/s 时，一定量的针状铁素体可以被获得。热模拟试验众所周知，对于同种类钢，由于应变导致变形，所以与静态转换相比，动态转换会缓慢的向左上角移动。根据Manohar和Chandra的研究，在针状铁素体的转变区域为了得到动态转换和静态转换间最适当的微观结构，前面的冷却速度应高于后面的冷却速度1020k/s。前面所提到的静态转换中，针状铁素体会在冷却速度为110k/s间形成，对于试验钢针状铁素体希望在动态转换中以冷却速度为1030k/s区域中形成。因此，在热变形试验中，所有的试件为了产生一个在1030k/s阶段的冷却速度，它们都会在变形后被水喷淋或压缩空气。图3显示出多种微观组织中的变形温度。在奥氏体的再结晶区域，细化晶粒是不显著的，因为晶粒存在着高温变形。在很大程度上奥氏体非再结晶区和非结晶区的变形不能减小晶粒尺寸，因为存在着高温变形。在混合奥氏体再结晶区和非结晶区的变形产生一种针状铁素体和针状珠光体的混合微观组织。此外，这是大量的变形是类似的，在此区域的变形产生有利的微观组织。在此变形过程中最有可能的解释是再结晶精化，产生大量的变形带和基地，从而增加和位置。因此，在试验钢中一定数量好的针状铁素体会在奥氏体结晶区和非结晶区域，冷却速度在1030k/s阶段中通过变形而得到。这对设计最优化TMCP 很有帮助。热轧实验根据前面的研究，为了得到一定量好的针状铁素体，所设计的热轧规程含有两个阶段，控制轧制变形和以1030k/s的冷却速度控制冷却。轧制第一阶段控制奥氏体结晶区域，第二阶段控制奥氏体非再结晶区域。以上的规程在表4用图解法显示。热轧试件的微观组织是用SEM 观察的。结果显示所有的微观组织主要都是由针状珠光体和针状铁素体组成的。多边铁素体的外形有些扁平。针状铁素体组织构成一大批不规则的铁素体，组织显示出奇特和不规则结构，大量的晶粒以混乱的方向分布着。针状铁素体的细节特征由图6显示出来。针状铁素体的密度高度混乱，且是由一系列的非平行铁素体条形板的相交网格组成。其中，在条形板中还有微粒或孤立区域。经离子探测器探测，孤立区域的碳含量比铸膜周围的碳含量要高。这也就意味着孤立区是碳聚集的M/A 孤立区，下面介绍其组成。在奥氏体非再结晶区轧制中，控轧后通过快冷的手段，经历的得到细化，大量的针状珠光体开始减少，针状铁素体群在混合的微观结构中增加。试件取10k/s 的冷却速度，针状珠光体能够清楚地被观察到。而针状珠光体少量的试孤立的。试件取30k/s 冷却速度下，大量的针状珠光体明显减少，而且针状铁素体的聚集组织急剧增加导致针状铁素体占据微观组织的大部分。试验结果说明对于试验钢，含有大部分针状铁素体的微观组织通过TMCP 可以获得。例如，在奥氏体结晶区和奥氏体非再结晶区并且冷却速度控制在30k/s下，多道次变形会在两个阶段中控制轧制，这种轧制已被CCT 曲线和热模拟试验中优化。在工业产品中一个铁素体珠光体的混合微观组织可以在试验中形成。这表明在现有的TMCP 程序下最大程度上可得到一定量的针状铁素体，而且将来还会应用在工业产品中。结论对于管线钢，包含针状铁素体的混合微观组织可以通过最佳的TMCP 得到，TMCP 通过连续的冷却变形图和热变形试验构造，在现有的试验条件下，在较高的冷却速度下，更多的针状铁素体会在一个混合的微观组织中存在。此外，占有大多数针状铁素体的微观组织在最优TMCP 中得到是合适的因为多道次变形在两个阶段中控制，在奥氏体结晶区和非结晶区域，并且冷却速度控制在30k/s。Acicular ferrite formation during hot plate rolling for pipeline steelsZhao M.-C.;Shan Y.-Y.;Xiao F.-R.;Yang K. Abstract:The transformation of super cooled austenite in a commercial pipeline steel was investigated by means of continuous cooling transformation (CCT) and hot simulation experiments. Based on the obtained results, an improved thermo mechanical control process (TMCP) was proposed, which could produce a mixed microstructure dominated by acicular ferrite. Results indicated that an increase in the cooling rate could improve the percentage of acicular ferrite in the final microstructure under the present experimental conditions. Furthermore, the acicular ferrite dominated microstructure could be obtained by a two stage controlled rolling in the austenite recrystallisation region plus the non-recrystallisation region and controlled cooling at a cooling rate of 30 K s-1. MST/5230 The authors are in the Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China (). Manuscript received 3 September 2001; accepted 16 April 2002.# 2003 IoM Communications Ltd. Published by Money for the Institute of Materials, Minerals and Mining.Introduction Acicular ferrite (AF) is of considerable commercial importance for line pipes because of its relatively high strength and toughness. Since AF high strength low alloy pipeline steels were developed in the early 1970s,1,2 much work has been doneonthisstructure.3 5Itisaccepted1,5 that AF results from a mixed diffusion and shear transformation mode beginning at a temperature slightly above the upper bainite temperature transformation range, and the growth of AF generally involves a coherent or semi coherent austenite/ ferrite (c /a ) interface that always obeys the Nishiyama Wasserman or Kurdjumov Sachsrelationship. So far, most work concerning AF has dealt with either the TiO AF steels or weld metals.6 9 The Japanese steelmakers claimed that TiO is an essential feature of the former, while for the latter, nucleation of AF takes place inside the austenite grains at non-metallic inclusions. However, AF in hot plate rolling is very different from these mentioned above, which does not demand for a special chemistry composition.Tanaka10 showed that for a0.07C 2.0Mn 0.6Nb 0.5MoAF steel, an AF dominated microstructure, which has optimum mechanical properties, consists of . Ne non-equiaxed ferrite dispersed with cementite and marten site/austenite (M/A) islands. Moreover, the micro structural characteristic of AF presents an assemblage of interwoven non-parallel ferrite laths with high density tangled dislocations pinned by some ultra. Ne cabinetries. Recently, an AF dominated microstructure was developed as an alter native to the conventional microstructure such as a ferrite paralytic microstructure in the hot plate rolling of pipeline steels.4 The formation of AF during hot plate rolling of pipeline steels has been overlooked in previous studies. In order to increase industrial production and application of pipeline steels possessing an AF dominated microstructure, AF formation was studied by the continuous cooling transformation (CCT) diagram, the hot simulation experiment and the optimized thermo mechanical control process (TMCP) for a commercial pipe line steel in the present work. Experimental procedure The material used in the present investigation is a commercial pipeline steel applied in the oil andgasin dustry. Its chemical composition was Fe 0.076C 0.24Si 1.33Mn 0.014P 0.0032S 0.03Al 0.0048N 0.04Nb 0.02Ti 0.04V (wt-%).The transformation of super cooled austenite was conducted by a Form Astor F dilatometer. Specimens used for quench dilatometry, 10 mm long cylindrical tubes with adiameterof3 mm, were cut from the continuous cast slabs of the tested steel. The linear cooling rates of 0.1, 0.4, - 11, 3, 4, 10, 50, and 100 K swere employed to construct CCT diagram. The hot simulation experiments were performed on a Gleeble 1500 hot simulator, and deformed with mul-pass in the austenite recrystallisation region, in austenite non-recrystallisation region, and in the mixed austenite recrystallisation plus on-recrystallisation region, respectively. Hot simulation specimens with a height of 20 mm and the diameter of 8 mm, which were machined from the 30630 mm forged slabs of the continuous cast slabs of the tested steel, were reheated at 1150 C for 3 min. The hot rolling experiments were carried out on a pilot rolling mill with 370 mm diameter twin rolls. The hot rolling specimens of 606806140 mm were cut from the continuous cast slabs of the tested steel and. Nelly rolled to 8 mm thick plates. Table 1 lists the interpose reduction distribution for the hot rolling experiments. Microstructures were examined by optical microscope, a CambridgeS360SEMandaJEM2000FX IITEM. For optical microscope and SEM, specimens were mechanically polished and etched in 3 vol.-% nital solution. For TEM observation, thin foils were prepared from 300 mm thick discs, which were rst mechanically thinned to 50 mm and then electro polished by a twinjet electro polisher in a solution of 10 vol.-% perchloric acid and 90 vol.-% acetic acid. Results and discussion TRANSFORMATIONOF SUPERCOOLED AUSTENITE ACCT diagram is believed to be useful in gaining an insight into the microstructures and designing an optimum TMCP Table1 Interpass reduction distribution in hot rolling experiments Technological parameter High temperature region Pass 1 Pass 2 Low temperature region Pass 3 Pass 4 Pass 5 Thickness, mm Reduction, % Strain Strain rate, s-1 44 26.7 0.36 2.56 32 27.3 0.37 3.03 22 31.3 0.43 3.86 12 45.5 0.70 6.29 8 33.3 0.47 6.68 DOI10.1179/026708303225010641 Materials Scienceand Technology March2003 Vol.19 355 356 Zhao et al. Acicular ferriteformation during hotplate rolling for pipeline steels a 50K s-1, upper bainite; b 10K s-1, acicular ferrite; c 1K s-1, polygonal ferrite + pseudoopearlite; d 0.1K s-1, polygonal ferrite + pearlite 1 Microstructures of super cooled austenite obtained at different cooling rates during industrial processing of steels. Figure 1 show the corresponding microstructures, which were obtained under - 1the representative cooling rates of 0.1, 1, 10 and 50 K s. Figure 2 shows CCT diagram constructed using the linear cooling rate dynamometric method and metallographic when the specimens were . rst austenitised for 10 min at 1150 C, then cooled at the various cooling rates. At the cooling rate of 50 K s- 1, the transformed microstructure(Fig. 1a)of the specimen was mainly upper bainite (BC2), which possessed the parallel ferrite with lath structure and carbide precipitates along the lath boundaries. In addition, the prior austenite grain boundary network could be seen clearly. With the decrease of the cooling rate down 1 to 10 K s , the microstructure (Fig. 1b) of the specimen was dominated by AF. In comparison with B2C, AF showed a unique and irregular con. guration, which had various grain sizes distributed in a chaotic manner with random orientations. At the same time, the prior austenite grain boundary network was eliminated, and carbide precipitates were distributed into the ferrite laths. When the cooling rate 1 was reduced to 1 K s , the microstructure (Fig. 1c) of the specimen consisted of polygonal ferrite (PF) with some pseudo operalite (B3P) that came from the independent growth of cementite and ferrite from austenite. 11 Decreasing the 1 cooling rate to 0.1Ks, the microstructure(Fig.1d) of the specimen was composed of PF plus some pearlite (P). The CCT diagram presented in Fig.2 shows that nonesuch rmalaustenite AF transformation in the experiment - 1 tal steel occurred at a cooling rate range from 1 to 10 K s . Furthermore, an AF dominated microstructure was obtained at the cooling rate of 10K s- 1. Cooling rates lower than 1 0.4 K should promote the predominant formation of a mixture of PF and pearlite. Polygonal ferrite was avoided at - 1 a cooling rate of more than 4 K awhile pearlite was 1 suppressed at a cooling rate of more than 0.4 K s. When the specimens were continuously cooled at a slightly wider - 1cooling raterangefrom0.4to1Ks,theobtained microstructure failed to avoid pseudo pearlite, which was not a desired phase owing to its detrimental effect in mechanical properties. Based on the discussion mentioned above, a certain amount of AF can be obtained at a cooling rate range from. HOT SIMULATIONEXPERIMENT It is known that for the same steel, CCT diagram of dynamic transformation will move slightly toward the top left corner compared with that of static transformation because of strain induced transformation. Based on the work of Manohar and Chandra,12 in order to achieve the approximately corresponding microstructure between dynamic transformation and static transformation at the transformation range of AF, the cooling rate of the former should be 10 20K s- 1 higher than that of the latter. As a result AF, which forms at a cooling rate range from 1 to 1 10 K sduring static transformation mentioned above, is expected to form in dynamic transformation at a cooling- 1 rate range from 10 to 30 K sfor the experimental steel. Therefore, in the hot simulation experiments, all specimens were cooled by water spraying or compressed air after deforming in order to produce a cooling rate in the range- 1from10to30Ks. Figure 3 shows the variation of the microstructure with the deformation temperature. During deformation in the austenite recrystallation region, grain re. nement was unremarkable because of the high growth of the recrystallised 2 Continuous cooled transformation diagram of experi1 to10K s- 1 duringthe CCT of supercooling austenite for mental steel: BC2 upper bainite; BP3pseudopearlite; AF the experimental steel. acicular ferrite; PF polygonal ferrite Materials ScienceandTechnology March2003 Vol.19 Zhao et al. Acicular ferrite formation during hot plate rolling for pipeline steels 357 a upper bainite, 1150C (80%); b upper bainite +a small amount of polygonal ferrite, 850C (70%); c acicular ferrite +polygonal ferrite, 1150C(40%) 850C(50%); d acicular ferrite +a small amount of polygonal ferrite, 1150C (25%) 1100C (20%) 950C (25%) 900C (33%) 3 Microstructures of specimens after hot simulation testing at different temperatures and reductions grains, and B2C was formed after cooling (Fig. 3a). The deformation in the austenite non-recrystallisation region could not decrease grain size to a large extent owing to the absence of high temperature deformation re. nement, and (BC2+PF) was obtained after cooling(Fig.3b). The deformation in the mixed austenite recrystallation plus non recrystallation region resulted in a mixed microstructure of AF plus PF(Fig. 3c).Moreover, when the amount of total deformation was similar, the mul-pass deformations in this region produced bene. cial microstructure re. nement (Fig. 3d). The likely explanation is in that this deformation process could retain the effect of recrystallisation re. ning and produce large numbers of deformation bands and substructures to increase the nucleation sites. Thus, a certain amount of . ne AF in the experimental steel could be achieved by themul-pass deformations in the mixed austenite recrystallisation plus non-recrystallisation region and in controlled cooling at a cooling rate range 1 from 10 to 30 K s . This will be signi. cantly helpful in designing the optimized TMCP. HOT ROLLING EXPERIMENT Based on the above work, in order to achieve a certain amount of ne AF, the hot rolling schedule was designed as a two stage controlled rolling with themul-pass deformation and controlled cooling at the cooling rates of 10 30K s- 1.The . rst stage rolling was controlled in the austenite recrystallisation region and the second stage in the austenite non-recrystallisation region. The above schedule is presented schematically in Fig. 4. Microstructures of the hot rolled specimens were observed using SEM. Result shows microstructures of all the specimens are mainly composed of PF and/or AF colonies. Polygonal ferrite is somewhat polygonal or like a plate in shape (Fig. 5a). In contrast, AF colonies constitute an irregular ferrite mass, showing unique and irregular on. gurations, and various grain sizes distributed in a chaotic manner with random orientations (Fig. 5b). The detailed TEM features of AF are revealed in Fig. 6. Acicular ferrite is composed of an assemblage of interwoven non-parallel ferrite laths with fairly high dislocation densities and there are ultra. ne particles or islands inside and/or among the laths (Fig. 6a).The carbon content in the islands was found to be higher than that in the surrounding matrix using ion probe. This means that the islands are carbon enriched marten site/austenite (M/A) islands, whose formation may be explained as follows. During the rolling process in the austenite non-recrystallation region, high dislocation densities are formed in the austenite, which impedes the growth of the coherent and/or semi coherent c /a interface and accelerates diffusion of carbon to the c /a interface. All these result in the partitioning of carbon and produce carbon enriched austenite. During the following cooling process, part of the carbon enriched austenite may transform to martensite and the retained austenite will coexist with the martensite. As a result, M/A island was observed in hot rolling specimen, as shown in Fig. 6a. In addition, some dispersed carbon it ride precipitates are also observed in the matrix of AF (Fig. 6b). All these micro structural characteristics are highly consistent with the de. nition of AF of pipeline steels given by Smith.1 4 Schematic diagram of experimental program for the rmomechanical control process Materials Science and Technology March2003 Vol.19 358 Zhao et al. Acicular ferrite formation during hot plate rolling for pipelines teels a b a polygonal ferrite +small amount of pseudo opearlite, 10K s-1; b acicular ferrite, 30K s-1 5 Micrographs of partial specimens after the rmomechanical control process after different cooling rates (SEM) With the increase of the cooling rate after hot rolling, the grain size becomes smaller,