中国科学院金属研究所 沈阳材料科学国家(联合)实验室 沈阳 110016
文献标识码: 分类号 TG161 文章编号 1005-3093(2016)08-0561-07
收稿日期: 2015-11-11
网络出版日期: 2016-09-28
版权声明: 2016 《材料研究学报》编辑部 《材料研究学报》编辑部
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摘要
应用扫描电子显微镜、透射电子显微镜和能谱技术等手段研究了G18CrMo2-6钢正火组织中第二相在680℃下保温一系列时间段的转变。结果表明, 在正火后贝氏体中析出相主要为马氏体/奥氏体(M/A)组元和合金渗碳体(M3C); 在回火保温初期M/A分解为铁素体(α)与M3C组织, 随着保温时间延长M3C逐渐球化并溶解, M23C6在晶界析出并长大, 同时基体上有细小弥散的MC相析出。即回火保温过程中组织随时间的延长发生M/A→α+M3C, M3C→M23C6+MC的变化。
关键词：
Abstract
The evolution of the secondary phase in G18CrMo2-6 heat-resistant steel induced by tempering at 680℃ for a series of durations was investigated by optical microscope (OM), scanning electron microscope (SEM), transmission electron microscope (TEM) and energy dispersive X-ray detector (EDX). It demonstrates that martensite/austenite (M/A) particles and M3C precipitate on bainite matrix after normalizing. During the tempering, decomposition of M/A particles into the M3C carbides in bainite matrix can be found. The increasing of tempering time results in the precipitation of MC, the spheroidization and refinement of M3C carbide as well as the precipitation and coarsening of M23C6 carbide at the grain boundaries.
Keywords：
低合金CrMo耐热钢的含碳量低, 合金成分以Cr、Mo为主。该类钢具有较高的热强性、优异的抗氧化性和抗氢脆性能、优良的加工工艺性及较高的性价比, 广泛应用于核动力、石油和化工等领域的高温作业压力容器中[1, 2]。G18CrMo2-6钢是一种典型的低合金CrMo钢, 碳含量较低, 主要用于核电汽轮机缸体铸件的制造。汽轮机缸体的工作条件十分严苛, 既要承受本身和装在其内部的零部件的重量及内外压差产生的作用力, 还要承受沿缸体轴向和径向温度分布不均匀而产生的热应力, 因此要求材料在服役过程中具有较高的强度、良好的塑性、韧性及加工性能。
在低合金耐热钢高温回火以及服役过程中, 耐热钢中的碳化物的形态、成分、结构不断变化, 使碳化物向动力学上更加稳定的类型、分布状态转变[3-7], 碳化物的演变对此类耐热钢的力学性能有重要影响。因此, 低合金CrMo耐热钢在回火过程中的析出相演化、表征研究一直受到关注[8-10]。在一般情况下, 合金元素、温度均影响碳化物的演化过程。Fujita等[11]研究了2.25Cr1Mo钢和3Cr1.5Mo钢在600℃时效过程中碳化物的演化过程, 发现在2.25Cr1Mo钢的高温回火过程中, 随着回火保温时间的延长碳化物相结构演化的大致次序为: M3C→M3C+M2C→M3C+M2C+M7C3→M3C+M2C+
M7C3+M23C6。而在3Cr1.5Mo钢的保温过程中, 随着保温时间的延长碳化物相结构演化的次序大致为: M3C→M3C+M7C3→M3C+M2C+M7C3+M23C6→M7C3+M2C+M23C6。陶鹏等[12, 13]研究了2.25Cr1Mo钢以及2.25Cr1MoV钢焊缝在700℃高温回火时碳化物的转变过程, 发现在高温回火保温过程中碳化物演化的次序为: M3C→M3C+M7C3→M3C+M7C3+M23C6→M7C3+M23C6。由此可见, 回火条件不同则不同成分的低合金CrMo耐热钢具有不同的第二相演化次序。同样, G18CrMo2-6钢也因回火工艺的不同而表现不同的第二相析出行为, 并且第二相析出是影响其组织稳定性的重要因素。基于此, 本文对G18CrMo2-6钢在680℃进行了一系列时间间隔的回火实验, 研究其中第二相的变化规律。
实验用G18CrMo2-6钢用真空感应炉中熔炼, 测定其成分列于表1所示。将铸件在1100℃均匀化退火20 h, 然后将其切割成尺寸为10 mm×10 mm×10 mm的热处理试块, 为了防止试样在热处理过程中氧化, 热处理试样用石英管封起来。将G18CrMo2-6钢在940℃保温2 h后以正火的方式冷却到室温, 随后进行回火处理一系列时间保温(0, 2, 10, 20, 100, 200, 500和1000 h)后空冷至室温进行组织分析。
表1 标准中G18CrMo2-6钢及实验用G18CrMo2-6钢的化学成分
Table 1 Chemical compositions of the steel G18CrMo2-6 in standard and this work (%, mass fraction)
| Material | C% | Mn% | Si% | Cr% | Mo% | Ni% |
|---|---|---|---|---|---|---|
| Standard | 0.15-0.2 | 0.5-0.9 | 0.2-0.6 | 0.4-0.65 | 0.45-0.7 | 0.3-0.5 |
| G18CrMo2-6 | 0.16 | 0.75 | 0.45 | 0.61 | 0.61 | 0.46 |
对正火试样磨抛处理后, 用4%硝酸酒精腐蚀, 经扫描电镜观察后用Lepera试剂(1％偏重亚硫酸钠水溶液+4％苦味酸酒精溶液, 按体积1∶1混合)对试样进行着色腐蚀。在着色腐蚀条件下M/A组元为白色, 铁素体基体为土黄色。将着色试样在激光共聚焦显微镜(Laser Scanning Confocal Microscope LSCM)下观察。其他金相试样经研磨、抛光和4%硝酸酒精溶液浸蚀后, 用AXIOVERT 200MAT光学显微镜(OM)和扫描电子显微镜(SEM)观察试样组织。用Tenupol-5电解双喷减薄仪制备透射电镜(TEM)样品, 用Tecnai G2 20及F20透射电镜观察样品中碳化物的形貌和分布, 并对组织结构进行分析, 同时用EDS分析碳化物的化学成分。
图1给出了试样正火后微观组织, 正火后材料组织为67%铁素体+33%贝氏体, 其中贝氏体中有大块状相及少量条状、颗粒状析出相出现。对正火样品腐蚀后采用显微硬度仪进行标记, 用扫描电镜观察(图2a)。贝氏体基体上出现大块状相(图2a中的圆圈中部分)。样品进而进行着色腐蚀后通过激光共聚焦显微观察(图2b), 灰黑色基体上分布着白色岛状M/A组元, 与扫描电镜下块状第二相分布基本一致(图2b中圆圈中部分), 确定正火后大块状岛状组织为M/A组元, 其中M/A组元由马氏体和残余奥氏体组成[14]。
图1 G18CrMo2-6钢940℃×2 h正火后的微观组织
Fig.1 OM (a) and SEM (b, c) images of 940℃×2 h normalized G18CrMo2-6 steel
图2 G18CrMo2-6 钢正火后的微观组织
Fig.2 SEM (a) and LSCM (b) images of air-cooling normalized G18CrMo2-6 steel
低合金CrMo耐热钢回火保温时间对碳化物的演化有着重要影响。图3给出了回火热处理后金相组织。随着回火时间的延长金相显微镜下铁素体与贝氏体对比度降低。随着回火保温时间的延长白色部分所占比例逐渐增多, 似乎铁素体比例增多, 贝氏体比例减少。这是在回火保温过程中贝氏体基体上的M/A组元分解、M3C溶解以及其他碳化物析出而造成金相下的假象。样品正火后贝氏体区局部区域有较多的大块状M/A组元分布(图4a), 在680℃回火保温1 h后可见M/A组元已经分解, 但分解相仍在局部聚集(图4 b)。随着保温时间的延长, 当保温时间达到2 h后分解析出的第二相均匀分布在贝氏体区基体上(图4 c)。此外, 对不同时间回火热处理后的试样中贝氏体区域进行扫描电镜观察, 结果如图5所示。由图5可见, 在保温时间较短时, 除M/A组元发生溶解外贝氏体中析出相在形态上未发生明显变化, 第二相主要呈长条状和块状分布, 且第二相颗粒差异较大, 不均匀分布在贝氏体铁素体的基体上, 第二相主要由大块状M/A组元和长条状的第二相组织组成。短时间保温后, M/A组元基本全部完成分解, 随着回火时间的延长合金渗碳体发生球化, 长条状第二相数量逐渐减少, 细小颗粒状第二相数量增加, 弥散分布于铁素体贝氏体基体上(图5d-f)。随着回火时间的继续延长, 1000 h回火后的颗粒状第二相数量显著减少(图5f)。不同时间回火热处理后试样基体和晶界区域的扫描电镜观察结果如图6所示。由图6可见, 在回火保温刚开始时晶界上几乎没有第二相的出现, 铁素体基体上同样没有析出相出现(图6a)。随着回火时间的延长晶界上有析出相析出(图6b, c), 且逐渐长大粗化(图6d, e, f), 同时铁素体基体上有细小弥散的析出相析出(图6d, e, f)。
图3 在680℃保温不同时间的金相组织
Fig.3 The OM images of G18CrMo2-6 steel tempered at 680℃ for 0 h (a), 100 h (b), 500 h (c) and 1000 h (d)
图4 M/A组元随着保温时间的变化
Fig.4 SEM images of M/A island after normalized (a) and tempered at 680℃ for 1 h (b), 2 h (c)
图5 G18CrMo2-6 钢经过不同回火时间后贝氏体区的微观组织
Fig.5 SEM images of bainite region after tempered at 680℃ for 0 h (a), 20 h(b), 100 h (c), 200 h (d), 500 h (e) and 1000 h (f)
图6 G18CrMo2-6钢经过不同回火时间后晶界的微观组织
Fig.6 SEM images of grain boundary after normalized and tempered at 680℃ for 0 h (a), 20 h (b), 100 h (c), 200 h (d), 500 h (e) and 1000 h (f)
图7a和b给出了大块状第二相的TEM像。低温回火后组织中大块状组织通过衍射斑点分析晶格类型与基体铁素体一样, 呈bcc结构, 其原始组织确认为M/A组元。图8a和b给出了680℃回火保温2 h、100 h回火后长条状、颗粒状第二相的TEM像, 由图可见, 回火保温时间延长后长条状析出相长度变短, 发生球化。根据衍射斑点(图 8c), 长条状及颗粒状第二相大多为合金渗碳体M3C, EDS(图8d)。分析结果表明, 碳化物的合金元素主要为Fe, 有少量的Cr和Mn。由此可以看出, 随着回火保温时间的延长合金渗碳体的形貌随回火温度升高发生球化溶解, 由长条状变为细小的颗粒状。根据对晶界析出相进行透射电镜表征结果(图9), 可确定晶界析出相为M23C6。根据对SEM、EDS、TEM的分析, 正火后贝氏体区析出的M3C属于亚稳相, 在高温回火保温过程中, 长条状合金渗碳体的平直面以及弯曲面的化学势不一致, 这一化学势差为C原子的扩散提供了驱动力, 扩散会导致平直面渗碳体(M3C)的剥落与球化[2]。同时, 由于晶界处于高能态, 在高温下Cr、Mo、Mn等合金元素扩散能力加强, 并开始向晶界扩散, 与碳结合后形成M23C6在晶界析出, 从而使晶界能量降低。
图7 G18CrMo2-6钢的大块状第二相TEM像及衍射斑点分析
Fig.7 TEM micrographs (a) and SAED pattern analysis of carbides (b) for G18CrMo2-6 steel normalized
图8 G18CrMo2-6 钢回火组织的TEM像及碳化物衍射斑点分析
Fig.8 TEM micrographs of rod-like particle with small size within bainite (tempering with 2 h (a), 100 h (b)), SAED pattern analysis (c) and EDS result
图9 G18CrMo2-6 钢回火后晶界析出相的 TEM像及衍射斑点分析
Fig.9 TEM micrographs of M23C6carbide at grain boundary (a) and SAED pattern analysis (b)
在回火保温过程中, 随着保温时间的延长基体上有大量细小的颗粒状或针状的析出相析出。如图6(e, f)所示, 根据对析出相进行透射电镜微衍射分析, 确定其为MC相[14]。钢中以固态存在的MoC的析出相有利于阻止杂质元素P和S在晶界偏聚, 这对改善或抑制钢的回火脆性是有利的[15]。
根据一系列的形貌观察及表征分析, 在正火后贝氏体中析出相主要为M/A组元和合金渗碳体(M3C), 在回火加热过程中M/A组元分解, M/A→α+M3C。在回火保温过程中M/A组元分解为α和M3C, 随着回火时间的延长M3C发生球化溶解, 弥散细小的MC在铁素体基体上析出, M23C6在晶界析出并迅速长大, 即在回火过程中组织发生M/A→α+M3C, M3C→M23C6+MC的演化。
1. 将正火后的G18CrMo2-6钢在680℃保温不同时间, 材料的组织为铁素体与粒状贝氏体的混合组织, 其中粒状贝氏体由M/A组元及合金渗碳体(M3C)组成。
2. 在材料的回火过程中随着保温时间的延长大块状M/A组元分解, 贝氏体基体上第二相逐步球化, 晶界上开始析出第二相并长大, 铁素体基体上逐步又第二相析出。在整个回火保温过程中第二相演变过程为M/A→α+M3C, M3C→M23C6+MC。
The authors have declared that no competing interests exist.
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Carbide reactions and phase equilibria of 2.4Cr–0.7Mo, 2.6Cr–0.7Mo–0.14V, 2.6Cr–0.7Mo–0.3V, and 2.6Cr–1.0Mo–0.3V (wt%) steels tempered at 773–933 K for a maximum of 1000 h were investigated. Carbide particles extracted in carbon replica were analysed by means of electron diffraction and energy-dispersive X-ray spectroscopy. Thermodynamic calculations were used to complement and make accurate the experimental results. The time–temperature diagrams of carbide phase stability were constructed for all investigated steels. Here the M 3 C, M 2 C, M 7 C 3 and MC carbides in various combinations are lined up in sequences, the carbides M 23 C 6 and M 6 C form separate areas. The shape of the M 6 C area is similar to that of the Laves phase in medium alloy steels. In the investigated steels three types of phase equilibria were found: ferrite+M 7 C 3 , ferrite+M 7 C 3 +MC and ferrite+M 7 C 3 +M 6 C+MC.
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| [11] |
A method has recently been developed to estimate the speed with which precipitation reactions occur in power plant steels. It is based on Avrami theory but with an adaptation that allows the treatment of simultaneous reactions. In the present work, a number of approximations and inconsistencies in the theory have been eliminated and this kinetic theory for simultaneous reactions has been modified with the treatments of both diffusion-controlled growth and capillarity effect in multicomponent systems. The modified model can predict not only volume fraction changes of each carbide but also particle sizes. New experimental results on alloy carbide in 3Cr1.5Mo and 214;Cr1Mo steels are reported and shown to be comparable to the modified theory.
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| [12] |
Evolution and coarsening of carbides in 2.25Cr-1Mo steel weld metal during high temperature tempering,
Transformation and coarsening of carbides in 2.25Cr-1Mo steel weld metal during tempering at 700 oC for different time intervals ranging from 1 to 150 h has been examined by transmission electron microscopy and scanning electron microscopy. M3C carbides were observed in the as-welded specimens and when tempered the precipitates were mainly composed of M3C, M7C3 and M23C6 carbides. A sequence for corresponding carbide transformation during tempering with initial precipitation of M3C and followed by M7C3 and M23C6 has been proposed. The precipitation of M7C3 with higher Cr content is the main factor contributing to the decrease of coarsening rate of precipitates after prolonged tempering. The decrease of hardness in the tempered specimens agreed well with the prediction of the weakening of precipitation strengthening due to the coarsening of carbides.
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| [13] |
Evolution of second phase in 225Cr-1Mo-0.25V steel weld metal during high temperature tempering, 2.25Cr-1Mo-0.25V钢焊缝中第二相高温回火转变,
<p>利用TEM和EDX技术, 研究了2.25Cr-1Mo-0.25V钢焊缝中第二相在700 ℃下保温一系列时间段的转变. 结果表明, 焊后空冷焊缝中的第二相为<em>M<sub>3</sub></em>C, 回火过程中第二相主要组成为<em>M<sub>3</sub></em>C+<em>M</em><sub>7</sub>C<sub>3</sub>+<em>M</em><sub>23</sub>C<sub>6</sub>. 长时间回火后<em>M<sub>3</sub></em>C发生球化并消失, 初期析出的低<em>w</em><sub>Cr</sub>/<u>w</u><sub>Fe</sub>(Cr与Fe的质量比)的<em>M</em><sub>7</sub>C<sub>3</sub>是亚稳相, <em>M</em><sub>23</sub>C<sub>6</sub>只能在较低的回火参数范围内存在, 后期形成了<em>w</em><sub>Cr</sub><em>w</em><sub>Fe</sub>更高的<em>M</em><sub>7</sub>C<sub>3</sub>稳定的第二相. <em>M<sub>3</sub></em>C多呈球状,<em>M</em><sub>7</sub>C<sub>3</sub>和<em>M</em><sub>23</sub>C<sub>6</sub>多为长条状和块状.</p>
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| [14] |
Effect of microstructure evolution on strength and impact toughness of G18CrMo2-6 heat-resistant steel during tempering,
G18CrMo2-6 steel is one kind of low alloy CrMo steel with ferrite/bainite phase constituent. It is widely adopted in the large casting parts of pressure vessels. This paper aims at investigating evolution of G18CrMo2-6 steel microstructure, tensile strength and impact toughness during the tempering (680聽掳C). The tempered microstructure characterization shows that the tempering time mainly affects the precipitation of three kinds of carbides, naming MC, M 3 C and M 23 C 6 . Both tensile strength and impact toughness do not obey the monotonic evolution with the tempering time but the complex changes. The shift of mechanical properties results from the interaction among matrix phases and the evolution of three different kinds of carbides. The softening of matrix due to the dislocation recovery and the decrease of the amout of carbon and other alloy elements in matrix lattice due to the precipitation of M 23 C 6 cause the decreasing tensile strength for the whole tempering except for that an increase during the short-term tempering controlled by the precipitation of MC and the spheroidization and refinement of M 3 C carbide. On the other hand, both the softening of matrix and the spheroidization and refinement of M 3 C carbide are responsible for the increasing impact toughness. However, the precipitation and coarsening of M 23 C 6 at ferrite grain boundaries during long-term tempering results in the sharp deterioration of impact toughness. It attributes to the larger grain boundary carbides result in the lower critical fracture stress of a carbide-ferrite interface.
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| [15] |
Phase transformation in Fe-Mo-C and Fe-W-C steels-I, the structural evolution during tempering at 700℃ ,
Mechanism and kinetics of carbide transformation during tempering at 700°C have been studied in Fe-Mo-C and Fe-W-C steels (with up to 2.5% W or Mo) by transmission electron microscopy (TEM) and X-ray diffraction. The sequence of carbide formation is FeC→MoC→(FeMoC, MC) in molybdenum steels and FeC→MC→MCin tungsten steels. Increasing the alloying element level increases the rate of carbide replacement reaction. In Fe-Mo-C steels the FeMoC carbides nucleate preferentially at the FeC-α interface and grow into cementite, whereas in Fe-W-C steels the MC carbides usually precipitate inside cementite giving rise to the transformation FeC→MC. The software DICTRA and THERMO-CALC were used to simulate cementite growth and to show the possibility of the transformation. The MCcarbide is first confined to prior austenite grain boundaries and penetrates into the grains with increasing tempering time. During growth the MCcarbide absorbs surrounding pre-existing carbides. As a result, after tempering for 500 h, patches of two-phase areas with (MC+ α) or (MC + α) are observed in tungsten steels, and patches of (MC+ α) or (FeMoC + α) in molybdenum steels. The alloying element partitioning between α and precipitated carbides was determined using TEM-EDS. It was established that the MCcarbide is stable at 700°C in both investigated steels. The FeMoC carbide is stable in the Fe-Mo-C system at this temperature. The MC carbide was not observed even after tempering for 3000 h. 08 .Mécanisme et cinétique de précipitation de carbures dans des aciers Fe-W-C et Fe-Mo-C (jusqu'à 2.5% de Mo ou W) après recuit à 700°C ont étéétudié par microscopie électronique de transmission (TEM) et par rayons X. La formation des carbures suit la séquence FeC→MoC→(FeMoC, MC) dans des alliages Fe-Mo-C et la séquence FeC→MC→MCdans des alliages Fe-W-C. La vitesse de remplacement augmente avec le contenu en Mo et W. Dans les alliages Fe-Mo-C la germination du carbure FeMoC se passe de préférence aux interfaces FeC-α, suivie d'une croissance vers l'intérieur de la cémentite. Dans les alliages Fe-W-C le carbure MC est germiné à l'intérieur de la cémentite, donnant lieu à la transformation Fe→MC. Les logiciels DICTRA et THERMO-CALC ontété utilisé pour simuler la croissance de cémentite et pour montrer la possibilité de la transformation . Le carbure MCest germiné aux anciens joints de grains d'austénite. Lors de la croissance vers l'intérieur des grains des carbures préformés sont absorbés. Après 500 h de recuit des plaques biphasés de type (MC+ α) et (MC + α) sont obtenues dans les alliages Fe-W-C, dans les alliages Fe-Mo-C des plaques biphasées (MC+ α) et (FeMoC + α). La distribution de W et Mo dans les carbures et dans la matrice ont été déterminé par TEM-EDS. Il est conclu que le carbure MCest stable à 700°C dans les deux systèmes, et que le carbure FeMoC est également stable dans le système Fe-Mo-C. Le carbure MC n'a puêtre révélé même après 3000 h de recuit.Mechanismus und Kinetik der Karbidausscheidung in Fe-W-C und Fe-Mo-C St01hlen (Massengehalte bis 2.5% W oder Mo) bei 700°C wurde mit Hilfe von Transmissionselektronenmikroskopie und R02ntgenbeugung analysiert. Die Ausscheidungssequenz istFeC→MoC→(FeMoC, MC) in Mo-St01hlen und FeC→MC→MCin W-St01hlen. Mit zunehmendem Legierungsgehalt steigt die Rate der Karbidumwandlungen. In Fe-Mo-C Legierungen erfolgt die Keimbildung von FeMoC bevorzugt an der FeC-α Phasengrenze mit anschlie08endem Wachstum in den Zementit, w01hrend die Keimbildung von MC in Fe-W-C Legierungen im Inneren des Zementits erfolgt mit anschlie08ender Reaktion Fe→MC. Mit der Software DICTRA und TERMO-CALC wurde das Wachstum und die M02glichkeit der Umwandlung analysiert. Das MCKarbid bildet sich zun01chst an den früheren Austenitkorngrenzen mit anschlie08endem Wachstum in das Korn, wobei umliegende Karbide absorbiert werden. Nach 500 h Auslagerung liegen dann in den Fe-W-C Legierungen zweiphasige Bereiche mit (MC+ α) und (MC + α), in den Fe-Mo-C Legierungen Bereiche mit (MC+ α) und (FeMoC + α) nebeneinander vor. Mit TEM-EDS wurde die Zusammensetzung der Gefügebestandteile bestimmt. Es konnte best01tigt werden, da08 das MCKarbid bei 700°C in beiden Legierungstypen stabil ist. Im System Fe-MoC ist das FeMoC Karbid ebenfalls stabil. Das MC Karbid dagegen war auch nach 3000 h Auslagerung nicht nachweisbar.
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