Research Progress on Initiation Mechanism of Local Corrosion Induced by Inclusions in Low Alloy Steel
LIU Chao, CHEN Tianqi, LI Xiaogang,
National Materials Corrosion and Protection Data Center, Institute of Advanced Materials & Technology, University of Science and Technology Beijing, Beijing 100083, China
Inclusions are inevitable metallurgical defects in steel that can significantly impact the corrosion resistance of materials by inducing local corrosion initiation. The mechanisms related with the initiation and development of inclusion-induced localized corrosion have been the subject of controversy in recent years. This paper provides a comprehensive review of the various mechanisms of inclusion-induced localized corrosion, including electrochemical corrosion, chemical dissolution, and electrochemical-chemical dissolution mechanisms. In addition, controlling the formation and behavior of inclusions is crucial for improving the corrosion resistance of steel, while the chemical composition, size and shape of the inclusions are the key influencing factors for inducing localized corrosion. Finally, the future research directions for the study of inclusions-induced local corrosion mechanism and the regulation of corrosion-resistant steel are discussed.
LIU Chao, CHEN Tianqi, LI Xiaogang. Research Progress on Initiation Mechanism of Local Corrosion Induced by Inclusions in Low Alloy Steel. Journal of Chinese Society for Corrosion and Protection[J], 2023, 43(4): 746-754 DOI:10.11902/1005.4537.2023.147
Fig.1
Linear scan results of LEIS for three metallurgical defects (a) [7], A, B and C being Si-rich inclusions, micropores and carbides, respectively, where the resistance at both A and B points is less than the substrate, whereas the resistance at C is greater than the substrate; mechanism diagram of pitting initiation and development induced by Y-S-O inclusions (b) [12], in which YS is preferentially dissolved as anodic phase and formed a corrosion microcouple with substrate and surface passivation film to induce local pitting initiation
第二种机理为化学溶解机制，该机制认为钢中绝缘非金属夹杂物 (如Al2O3[16]等) 由于自身导电性较差，无法与钢基体构成腐蚀电偶，因此腐蚀主要由化学溶解过程引发。根据夹杂物自身的化学稳定性可分为两种情形：若夹杂物自身化学稳定性差，其主要通过自身的化学溶解形成点蚀坑，诱发点蚀萌生和发展。若夹杂物自身稳定性好，则夹杂物自身不发生溶解，而通过夹杂物与基体之间的性能差异，如塑性和热膨胀系数等，造成钢基体局部高电化学活性，最终导致夹杂物周围钢基体的溶解。Liu等研究表明，Al2O3[16]，ZrO2-Ti2O3-Al2O3[17]和 (RE)AlO3-(RE)2O2S-(RE) x S y[18, 19]等均为绝缘体，图2a中Al2O3的CSAFM结果表明夹杂物处的电流值仅为0~2 pA，而基体部分的电流值为5 nA，远大于夹杂物处的电流值，证实Al2O3为绝缘性夹杂，无法和钢基体构成腐蚀电偶。图2b和c中ZrO2-Ti2O3-Al2O3夹杂和(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3夹杂物同样被证明为绝缘体，不能与基体构成腐蚀电偶。而图2d中可以看到，在西沙模拟液中浸泡30 min后，ZrO2-Ti2O3-Al2O3夹杂并未溶解，这可以归因于该夹杂物自身具有较高的化学稳定性。对比之下，(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3中(RE)2O2S-(RE) x S y 由于较差的化学稳定性而自身溶解。Liu的工作让研究者们认识到[16~19]，仅仅采用电偶腐蚀理论解释和理解夹杂物诱发局部腐蚀是不可接受的，绝缘非金属夹杂物诱发局部腐蚀可能是一种不同于电偶腐蚀的全新机理。根据Gutman的电化学理论[20]，图2e中绝缘体夹杂物ZrO2-Ti2O3-Al2O3附近的高晶格畸变区域增强了基体的反应活性，高活性钢基体与周围非高活性区域钢基体之间可能形成电偶腐蚀，这为非金属夹杂物诱发局部腐蚀的解释提供了新的思路。从上述文献的结果中可以看出，对于不导电或导电性差的非金属夹杂物，其诱导局部腐蚀萌生的机理不同于MnS等导电夹杂物，微缝隙或夹杂物周围高化学活性的区域可能成为诱发点蚀的主因。Hu等[21]通过SKPFM测试证实了 (Al, Mg, Ca, Mn)-氧硫复合夹杂物与基体之间的局部电化学不均匀性，结果表明该夹杂物的Volta电位高于基体，但电流密度仅为8 pA，因此传统的微电偶腐蚀机制并不适用。相比之下，EBSD测试结果表明该夹杂物周围存在高晶格畸变和位错密度，这使得夹杂物周围基体成为热力学不稳定区域，成为优先腐蚀的部分。
图2
Al2O3，ZrO2-Ti2O3-Al2O3和(RE)2O2S-(RE) x S y -(RE, Zr, Ti)O x 的SEM谱，CAAFM以及相应的高度/电流分布图；在pH=4.9的西沙模拟液中浸泡30 min后ZrO2-Ti2O3-Al2O3和(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3的腐蚀形貌及Al2O3和ZrO2-Ti2O3-Al2O3的KAM图，高KAM代表钢基体的局部塑性变形发生在不易变形的夹杂物周围，且机械形变会导致表面电化学异质性的重新分布[16~19]
Fig.2
SEM and CAAFM images of Al2O3(a), ZrO2-Ti2O3-Al2O3 (b) and (Re)2O2S-(Re) x S y -(Re, Zr, Ti)O x (c) and their corresponding height/current distribution diagrams, corrosion morphology of ZrO2-Ti2O3-Al2O3 and ((RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3 after soaking in simulated Xisha solution with pH=4.9 for 30 min (d), and KAM diagrams of Al2O3 and ZrO2-Ti2O3-Al2O3 (e). The high KAM indicates that the local plastic deformation of the steel matrix occurs around the inclusion which is difficult to deform, and the mechanical deformation will lead to the redistribution of the surface electrochemical heterogeneity [16~19]
Fig.3
Typical corrosion pits formed on Al-Ca-O-S inclusions after 24 h immersion in NS4 solution: (a) optical image, (b) 2D image, (c) three-dimensional profile, (d) height line, where the position of the cross-sectional profile is marked in Fig.3b as a red line [22], (e) schematic of dissolution kinetics of sulfide-oxide complex inclusions and the resulting local corrosion process: the original inclusions underwent the galvanic corrosion stage and the corrosion spreading stage, respectively [23]
Fig.4
Surface potential distribution of MnS (a), surface work function of different crystal faces of MnS (b), the results show that the work function of MnS is smaller than that of Fe matrix, and mechanism diagram of local corrosion induced by MnS (c), MnS is used as an anodic corrosion couple with steel substrate to form steady or metastable pitting pits under the synergistic effect of Cl-and S, and the corrosion product “S” shell covers the surface of pitting pits to form concentration difference cells, further promote the corrosion process [29]
Fig.6
SEM image and EDS results (a), TEM image and SAD result (b) of TiN inclusions, interface morphology between TiN inclusions and SiO2 (c), and XPS spectra of TiO2 2p3/2 and TiO2 2p1/2 peaks of passivation film formed on pure Ti sample (d) [25]
Fig.7
Composition distribution of active/inactive inclusions in the phase diagram of the Al2O3-MgO-CaO system (a) [45] and MgS (b), MgY2S4 (c), Y2O3 (d) and YS (e) energy band structure diagram [12]
Corrosion is a ubiquitous and costly problem for a variety of industries. Understanding and reducing the cost of corrosion remain primary interests for corrosion professionals and relevant asset owners. The present study summarises the findings that arose from the landmark “Study of Corrosion Status and Control Strategies in China”, a key consulting project of the Chinese Academy of Engineering in 2015, which sought to determine the national cost of corrosion and costs associated with representative industries in China. The study estimated that the cost of corrosion in China was approximately 2127.8 billion RMB (~ 310 billion USD), representing about 3.34% of the gross domestic product. The transportation and electronics industries were the two that generated the highest costs among all those surveyed. Based on the survey results, corrosion is a major and significant issue, with several key general strategies to reduce the cost of corrosion also outlined.
The early corrosion development of ultra-low carbon bainitic (ULCB) low alloy steel in NaCl solution was studied by ex situ imaging of corrosion morphology and in situ monitoring of microarea current density and potential, and the corrosion mechanism from initial localized corrosion to uniform corrosion was interpreted. The results indicate that the corrosion development of ULCB steel from initial localized corrosion around inclusions to the uniform corrosion on the whole steel surface is controlled by the galvanic couple effect between different phases resulting from their electrode potential difference in electrolyte solution. The early localized corrosion of steel matrix is initiated and accelerated by the galvanic couple effect between MnS inclusions and steel matrix to form the initial corrosion gaps and the circular corrosion spots around inclusions. The ohmic drop caused by solution resistance influences the acceleration effect of the galvanic couple. With the separation of inclusion from steel matrix, this galvanic couple effect becomes invalid, which results in the expansion from localized corrosion to uniform corrosion. The microgalvanic couple between martensite/residual austenite (M/A) islands and bainite ferrite also accelerates the anodic dissolution of bainite ferrite phase; however, its acceleration corrosion effect is much weaker than that caused by MnS inclusion.
In situ characterization by localized electrochemical impedance spectroscopy of the electrochemical activity of microscopic inclusions in an X100 steel
The surface properties of weathering steel (WS) is very important for its service performance and safety, and the localized corrosion induced by inclusions is closely related to the surface properties of WS and its application. In the current work, a common spherical (Al, Mg, Ca, Mn)-oxy-sulfide inclusion was selected to investigate the corrosion evolution of complex inclusion and its effect on localized corrosion on WS surface. The results indicate the inclusion in WS consists of (Ca, Mn) sulfides part and (Ca, Al, Mg) oxides part with complex core-shell structure. Locally preferential dissolution occurs in (Ca, Mn) sulfides part as well as metal matrix around the inclusions. Furthermore, both parts of the inclusions with poor conductivity and high-density dislocation at metal matrix around the inclusions was found, which suggests that traditional micro-galvanic corrosion cell may not be the cause of inclusion-induced localized corrosion on WS surface at initial stage of corrosion. The variation in maximum and average depth around the inclusion or selected region with immersion time indicates that localized corrosion induced by inclusions is overwhelmed by uniform corrosion of WS in the latter stage of immersion, then the rust formed on WS surface consists of two layers.
WangL W, XinJ C, ChengL J, et al.
Influence of inclusions on initiation of pitting corrosion and stress corrosion cracking of X70 steel in near-neutral pH environment
The present work systematically investigated the initiation mechanism of localized corrosion induced by Al2O3-MnS composite inclusion in E690 steel under a simulated marine environment. The results showed that a micro-gap exists between the Al2O3-MnS inclusion and the matrix, and electron backscattered diffraction (EBSD) analysis revealed significant lattice dislocation zones around the Al2O3-MnS composite inclusion. The presence of the micro-gap and the lattice dislocation both promoted the localized corrosion initiation. The Volta potential of Al2O3 detected by scanning Kelvin probe force microscopy (SKPFM) was approximately 149.33 mV higher than that of the steel matrix, and the Volta potential of MnS was 10 mV lower than that of the steel matrix. The current-sensing atomic force microscopy (CSAFM) results showed that the Al2O3 was not conductive, while the MnS had good conductive properties. Therefore, it was not possible for a galvanic couple to be formed between Al2O3 and the adjacent steel matrix. A galvanic couple effect between the MnS and the adjacent steel matrix was directly demonstrated for the first time. The MnS acted as the anode phase for preferential dissolution in the corrosion process. The in situ immersion experiments and the Pourbaix diagram results confirmed that the dissolution of MnS was an electrochemical reaction process and the dissolution of Al2O3 was a chemical reaction.
LiuC, RevillaR I, LiX, et al.
New insights into the mechanism of localised corrosion induced by TiN-containing inclusions in high strength low alloy steel
This work investigated the chemical and electrochemical mechanisms of localised corrosion triggered by CaS·xMgO·yAl2O3·TiN complex inclusions in high strength low alloy steel (HSLAS) under a simulated marine environment. Special focus was given to the role of the TiN portion of the inclusion on the initiation and growth of the corrosion pits. The thermodynamic process of pitting initiation was investigated by Gibbs free energy, Pourbaix diagram and first principle calculation. Localised corrosion is mainly induced by inclusions and triggered by dissolution of adjacent distorted matrix. Chemical dissolution of CaS portion in CaS·xMgO·yAl2O3·TiN complex inclusion creates an acidic aggressive environment that accelerates the further dissolution of inclusion and matrix. Galvanic coupling effect between TiN inclusion and matrix is directly verified. TiN covered with a TiO2 film acts as the cathodic phase in galvanic corrosion, although it has a lower Volta potential than the matrix. This is an unusual correlation with the scanning Kelvin probe force microscopy result, which has been explained for this special system.
ZhangX W, YangC F, ZhangL F.
Effects of cooling rate and isothermal holding on the characteristics of MnS particles in high-carbon heavy rail steels
The characteristics of MnS particles were intensively investigated at three different cooling rate of 80.4 K · s−1 (water cooling), 3.8 K · s−1 (air cooling) and 1.8 K · s−1 (furnace cooling) as well as the different isothermal holding temperature and time in laboratory experiments. The three-dimensional (3D) morphology of MnS particles was extracted from steel samples using non-aqueous solution electrolysis. The results showed that the 3D morphology of MnS changed from a nearly spherical into rod-like and the area fraction and average diameter of MnS increased with decreasing cooling rate. During isothermal holding process, the morphology of MnS changed little at 1473 K (1200 °C), but their shape profiles varied from a nearly spherical and spindle-like to irregular at higher holding temperature 1673 K (1400 °C) when the holding time exceeded 60 min. Moreover, the number density and area fraction of MnS decreased with increasing holding time at 1573 K (1300 °C) and 1673 K (1400 °C), respectively. Especially at 1573 K (1300 °C), the 1 ∼ 3 µm MnS inclusions were dissolved and lead to decreasing of number density, but that > 3 µm one occurs growth and resulted in increasing of average diameter. The calculation results show that the starting temperature of precipitation of MnS was about 1627 K (1354 °C) and effect of cooling rate on the segregation of Mn and S is insignificant. Considering the segregation of solutes, MnS formation and growth takes place in the solid/liquid interface of steel when the solid fraction is close to 0.9567 during solidification. It has been found that the increase of cooling rate gives rise to the decreased of MnS diameter because the growth time of MnS is short. Furthermore, thermodynamic calculations of MnS solid solubility product were carried out to reveal the high holding temperature and long holding time favors the dissolution of MnS particles. It is necessary to decrease the sulfur content by less than 16 ppm in order to assure that the larger MnS which formed during solidification redissolves in the steel matrix, rather than relying on increasing the heating temperature which is above 1649 K (1376 °C). Subsequent, the MnS will precipitate again in a finely dispersive state during rolling process, and it can hinder annealing grain growth and finally make for the improvement of the toughness property of the steel.
VuilleminB, PhilippeX, OltraR, et al.
SVET, AFM and AES study of pitting corrosion initiated on MnS inclusions by microinjection
A microelectrochemical system for in situ high-resolution optical microscopy: morphological characteristics of pitting at MnS inclusion in stainless steel
Transient inclusion evolution during modification of alumina inclusions by calcium in liquid steel: part I. Background, experimental techniques and analysis methods
This paper aims to investigate the atmospheric corrosion mechanism of structural materials to develop more advanced corrosion-control technologies and cost-reduction strategies. As a second phase in steels, the non-metallic oxide inclusions are considered to not only affect the mechanical properties of steel but also the corrosion resistance of steel. So, an important research goal in this paper is to investigate the indoor accelerated corrosion kinetics of Q450NQR1 weathering steel, analyzing the galvanic polarity of different inclusions in electrochemical corrosion microcell between the inclusion and steel matrix and then elucidating the influence mechanism of inclusions on corrosion resistance of weathering steel.
MaH C, LiuZ Y, DuC W, et al.
Stress corrosion cracking of E690 steel as a welded joint in a simulated marine atmosphere containing sulphur dioxide
In situ characterization by localized electrochemical impedance spectroscopy of the electrochemical activity of microscopic inclusions in an X100 steel
... [7]及Y-S-O夹杂物诱发点蚀萌生和发展的机理[12]Linear scan results of LEIS for three metallurgical defects (a) <sup>[<xref ref-type="bibr" rid="R7">7</xref>]</sup>, A, B and C being Si-rich inclusions, micropores and carbides, respectively, where the resistance at both A and B points is less than the substrate, whereas the resistance at C is greater than the substrate; mechanism diagram of pitting initiation and development induced by Y-S-O inclusions (b)<sup> [<xref ref-type="bibr" rid="R12">12</xref>]</sup>, in which YS is preferentially dissolved as anodic phase and formed a corrosion microcouple with substrate and surface passivation film to induce local pitting initiationFig.1
第二种机理为化学溶解机制，该机制认为钢中绝缘非金属夹杂物 (如Al2O3[16]等) 由于自身导电性较差，无法与钢基体构成腐蚀电偶，因此腐蚀主要由化学溶解过程引发.根据夹杂物自身的化学稳定性可分为两种情形：若夹杂物自身化学稳定性差，其主要通过自身的化学溶解形成点蚀坑，诱发点蚀萌生和发展.若夹杂物自身稳定性好，则夹杂物自身不发生溶解，而通过夹杂物与基体之间的性能差异，如塑性和热膨胀系数等，造成钢基体局部高电化学活性，最终导致夹杂物周围钢基体的溶解.Liu等研究表明，Al2O3[16]，ZrO2-Ti2O3-Al2O3[17]和 (RE)AlO3-(RE)2O2S-(RE) x S y[18, 19]等均为绝缘体，图2a中Al2O3的CSAFM结果表明夹杂物处的电流值仅为0~2 pA，而基体部分的电流值为5 nA，远大于夹杂物处的电流值，证实Al2O3为绝缘性夹杂，无法和钢基体构成腐蚀电偶.图2b和c中ZrO2-Ti2O3-Al2O3夹杂和(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3夹杂物同样被证明为绝缘体，不能与基体构成腐蚀电偶.而图2d中可以看到，在西沙模拟液中浸泡30 min后，ZrO2-Ti2O3-Al2O3夹杂并未溶解，这可以归因于该夹杂物自身具有较高的化学稳定性.对比之下，(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3中(RE)2O2S-(RE) x S y 由于较差的化学稳定性而自身溶解.Liu的工作让研究者们认识到[16~19]，仅仅采用电偶腐蚀理论解释和理解夹杂物诱发局部腐蚀是不可接受的，绝缘非金属夹杂物诱发局部腐蚀可能是一种不同于电偶腐蚀的全新机理.根据Gutman的电化学理论[20]，图2e中绝缘体夹杂物ZrO2-Ti2O3-Al2O3附近的高晶格畸变区域增强了基体的反应活性，高活性钢基体与周围非高活性区域钢基体之间可能形成电偶腐蚀，这为非金属夹杂物诱发局部腐蚀的解释提供了新的思路.从上述文献的结果中可以看出，对于不导电或导电性差的非金属夹杂物，其诱导局部腐蚀萌生的机理不同于MnS等导电夹杂物，微缝隙或夹杂物周围高化学活性的区域可能成为诱发点蚀的主因.Hu等[21]通过SKPFM测试证实了 (Al, Mg, Ca, Mn)-氧硫复合夹杂物与基体之间的局部电化学不均匀性，结果表明该夹杂物的Volta电位高于基体，但电流密度仅为8 pA，因此传统的微电偶腐蚀机制并不适用.相比之下，EBSD测试结果表明该夹杂物周围存在高晶格畸变和位错密度，这使得夹杂物周围基体成为热力学不稳定区域，成为优先腐蚀的部分. ...
... [7], A, B and C being Si-rich inclusions, micropores and carbides, respectively, where the resistance at both A and B points is less than the substrate, whereas the resistance at C is greater than the substrate; mechanism diagram of pitting initiation and development induced by Y-S-O inclusions (b) [12], in which YS is preferentially dissolved as anodic phase and formed a corrosion microcouple with substrate and surface passivation film to induce local pitting initiationFig.1
第二种机理为化学溶解机制，该机制认为钢中绝缘非金属夹杂物 (如Al2O3[16]等) 由于自身导电性较差，无法与钢基体构成腐蚀电偶，因此腐蚀主要由化学溶解过程引发.根据夹杂物自身的化学稳定性可分为两种情形：若夹杂物自身化学稳定性差，其主要通过自身的化学溶解形成点蚀坑，诱发点蚀萌生和发展.若夹杂物自身稳定性好，则夹杂物自身不发生溶解，而通过夹杂物与基体之间的性能差异，如塑性和热膨胀系数等，造成钢基体局部高电化学活性，最终导致夹杂物周围钢基体的溶解.Liu等研究表明，Al2O3[16]，ZrO2-Ti2O3-Al2O3[17]和 (RE)AlO3-(RE)2O2S-(RE) x S y[18, 19]等均为绝缘体，图2a中Al2O3的CSAFM结果表明夹杂物处的电流值仅为0~2 pA，而基体部分的电流值为5 nA，远大于夹杂物处的电流值，证实Al2O3为绝缘性夹杂，无法和钢基体构成腐蚀电偶.图2b和c中ZrO2-Ti2O3-Al2O3夹杂和(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3夹杂物同样被证明为绝缘体，不能与基体构成腐蚀电偶.而图2d中可以看到，在西沙模拟液中浸泡30 min后，ZrO2-Ti2O3-Al2O3夹杂并未溶解，这可以归因于该夹杂物自身具有较高的化学稳定性.对比之下，(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3中(RE)2O2S-(RE) x S y 由于较差的化学稳定性而自身溶解.Liu的工作让研究者们认识到[16~19]，仅仅采用电偶腐蚀理论解释和理解夹杂物诱发局部腐蚀是不可接受的，绝缘非金属夹杂物诱发局部腐蚀可能是一种不同于电偶腐蚀的全新机理.根据Gutman的电化学理论[20]，图2e中绝缘体夹杂物ZrO2-Ti2O3-Al2O3附近的高晶格畸变区域增强了基体的反应活性，高活性钢基体与周围非高活性区域钢基体之间可能形成电偶腐蚀，这为非金属夹杂物诱发局部腐蚀的解释提供了新的思路.从上述文献的结果中可以看出，对于不导电或导电性差的非金属夹杂物，其诱导局部腐蚀萌生的机理不同于MnS等导电夹杂物，微缝隙或夹杂物周围高化学活性的区域可能成为诱发点蚀的主因.Hu等[21]通过SKPFM测试证实了 (Al, Mg, Ca, Mn)-氧硫复合夹杂物与基体之间的局部电化学不均匀性，结果表明该夹杂物的Volta电位高于基体，但电流密度仅为8 pA，因此传统的微电偶腐蚀机制并不适用.相比之下，EBSD测试结果表明该夹杂物周围存在高晶格畸变和位错密度，这使得夹杂物周围基体成为热力学不稳定区域，成为优先腐蚀的部分. ...
Pitting corrosion behavior in advanced high strength steels
... [12]Linear scan results of LEIS for three metallurgical defects (a) <sup>[<xref ref-type="bibr" rid="R7">7</xref>]</sup>, A, B and C being Si-rich inclusions, micropores and carbides, respectively, where the resistance at both A and B points is less than the substrate, whereas the resistance at C is greater than the substrate; mechanism diagram of pitting initiation and development induced by Y-S-O inclusions (b)<sup> [<xref ref-type="bibr" rid="R12">12</xref>]</sup>, in which YS is preferentially dissolved as anodic phase and formed a corrosion microcouple with substrate and surface passivation film to induce local pitting initiationFig.1
第二种机理为化学溶解机制，该机制认为钢中绝缘非金属夹杂物 (如Al2O3[16]等) 由于自身导电性较差，无法与钢基体构成腐蚀电偶，因此腐蚀主要由化学溶解过程引发.根据夹杂物自身的化学稳定性可分为两种情形：若夹杂物自身化学稳定性差，其主要通过自身的化学溶解形成点蚀坑，诱发点蚀萌生和发展.若夹杂物自身稳定性好，则夹杂物自身不发生溶解，而通过夹杂物与基体之间的性能差异，如塑性和热膨胀系数等，造成钢基体局部高电化学活性，最终导致夹杂物周围钢基体的溶解.Liu等研究表明，Al2O3[16]，ZrO2-Ti2O3-Al2O3[17]和 (RE)AlO3-(RE)2O2S-(RE) x S y[18, 19]等均为绝缘体，图2a中Al2O3的CSAFM结果表明夹杂物处的电流值仅为0~2 pA，而基体部分的电流值为5 nA，远大于夹杂物处的电流值，证实Al2O3为绝缘性夹杂，无法和钢基体构成腐蚀电偶.图2b和c中ZrO2-Ti2O3-Al2O3夹杂和(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3夹杂物同样被证明为绝缘体，不能与基体构成腐蚀电偶.而图2d中可以看到，在西沙模拟液中浸泡30 min后，ZrO2-Ti2O3-Al2O3夹杂并未溶解，这可以归因于该夹杂物自身具有较高的化学稳定性.对比之下，(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3中(RE)2O2S-(RE) x S y 由于较差的化学稳定性而自身溶解.Liu的工作让研究者们认识到[16~19]，仅仅采用电偶腐蚀理论解释和理解夹杂物诱发局部腐蚀是不可接受的，绝缘非金属夹杂物诱发局部腐蚀可能是一种不同于电偶腐蚀的全新机理.根据Gutman的电化学理论[20]，图2e中绝缘体夹杂物ZrO2-Ti2O3-Al2O3附近的高晶格畸变区域增强了基体的反应活性，高活性钢基体与周围非高活性区域钢基体之间可能形成电偶腐蚀，这为非金属夹杂物诱发局部腐蚀的解释提供了新的思路.从上述文献的结果中可以看出，对于不导电或导电性差的非金属夹杂物，其诱导局部腐蚀萌生的机理不同于MnS等导电夹杂物，微缝隙或夹杂物周围高化学活性的区域可能成为诱发点蚀的主因.Hu等[21]通过SKPFM测试证实了 (Al, Mg, Ca, Mn)-氧硫复合夹杂物与基体之间的局部电化学不均匀性，结果表明该夹杂物的Volta电位高于基体，但电流密度仅为8 pA，因此传统的微电偶腐蚀机制并不适用.相比之下，EBSD测试结果表明该夹杂物周围存在高晶格畸变和位错密度，这使得夹杂物周围基体成为热力学不稳定区域，成为优先腐蚀的部分. ...
... [12], in which YS is preferentially dissolved as anodic phase and formed a corrosion microcouple with substrate and surface passivation film to induce local pitting initiationFig.1
第二种机理为化学溶解机制，该机制认为钢中绝缘非金属夹杂物 (如Al2O3[16]等) 由于自身导电性较差，无法与钢基体构成腐蚀电偶，因此腐蚀主要由化学溶解过程引发.根据夹杂物自身的化学稳定性可分为两种情形：若夹杂物自身化学稳定性差，其主要通过自身的化学溶解形成点蚀坑，诱发点蚀萌生和发展.若夹杂物自身稳定性好，则夹杂物自身不发生溶解，而通过夹杂物与基体之间的性能差异，如塑性和热膨胀系数等，造成钢基体局部高电化学活性，最终导致夹杂物周围钢基体的溶解.Liu等研究表明，Al2O3[16]，ZrO2-Ti2O3-Al2O3[17]和 (RE)AlO3-(RE)2O2S-(RE) x S y[18, 19]等均为绝缘体，图2a中Al2O3的CSAFM结果表明夹杂物处的电流值仅为0~2 pA，而基体部分的电流值为5 nA，远大于夹杂物处的电流值，证实Al2O3为绝缘性夹杂，无法和钢基体构成腐蚀电偶.图2b和c中ZrO2-Ti2O3-Al2O3夹杂和(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3夹杂物同样被证明为绝缘体，不能与基体构成腐蚀电偶.而图2d中可以看到，在西沙模拟液中浸泡30 min后，ZrO2-Ti2O3-Al2O3夹杂并未溶解，这可以归因于该夹杂物自身具有较高的化学稳定性.对比之下，(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3中(RE)2O2S-(RE) x S y 由于较差的化学稳定性而自身溶解.Liu的工作让研究者们认识到[16~19]，仅仅采用电偶腐蚀理论解释和理解夹杂物诱发局部腐蚀是不可接受的，绝缘非金属夹杂物诱发局部腐蚀可能是一种不同于电偶腐蚀的全新机理.根据Gutman的电化学理论[20]，图2e中绝缘体夹杂物ZrO2-Ti2O3-Al2O3附近的高晶格畸变区域增强了基体的反应活性，高活性钢基体与周围非高活性区域钢基体之间可能形成电偶腐蚀，这为非金属夹杂物诱发局部腐蚀的解释提供了新的思路.从上述文献的结果中可以看出，对于不导电或导电性差的非金属夹杂物，其诱导局部腐蚀萌生的机理不同于MnS等导电夹杂物，微缝隙或夹杂物周围高化学活性的区域可能成为诱发点蚀的主因.Hu等[21]通过SKPFM测试证实了 (Al, Mg, Ca, Mn)-氧硫复合夹杂物与基体之间的局部电化学不均匀性，结果表明该夹杂物的Volta电位高于基体，但电流密度仅为8 pA，因此传统的微电偶腐蚀机制并不适用.相比之下，EBSD测试结果表明该夹杂物周围存在高晶格畸变和位错密度，这使得夹杂物周围基体成为热力学不稳定区域，成为优先腐蚀的部分. ...
... 随着第三代氧化物冶金技术的广泛应用，Al脱氧Ca处理和Mg-Al脱氧Ca处理技术日趋成熟，钙处理对钢中夹杂物的改性作用对钢基体耐蚀性产生了重要影响.Ca处理后，钢中夹杂物的尺寸明显减小，对夹杂物的成分、形状和尺寸的控制作用明显[39].经Ca处理后，钢中高硬度、不规则的Al2O3夹杂物转变为xCaO-yAl2O3复合氧化物夹杂物，其硬度低、呈球状，并可能伴随着xCaS-yCaO-zAl2O3复合氧化物夹杂物的形成[40].在各种经Ca处理改性的铝酸钙和xCaO-yAl2O3化合物中，12CaO-7Al2O3(最低熔点1455 ℃)是一种理想的改性产物，可以降低原夹杂物与钢基体之间的应力[41, 42].Wang等[43]研究了EH36钢中(Ca, Mg, Al)-O x -S y 夹杂物和局部腐蚀的关联，图7a显示了Al2O3-MgO-CaO三元相图，CaO:Al2O3比例约为1.25是区分活性和非活性夹杂物的界限，且通过对比和计算夹杂物和基体之前的热膨胀系数和残余应力，验证了非活性夹杂物3CaO·Al2O3周围不存在明显的局部塑形变形，说明其并不会诱发点蚀.Ca处理不仅可以改性非金属氧化物夹杂物，也可以改性硫化物夹杂物.相关研究表明在X80钢管生产线中，Ca处理还可以将条状的MnS夹杂物转化为球形的CaS夹杂物和 (Ca, Mn)S复合夹杂物，降低夹杂物与钢基体之间的应力.同时，含Ca的夹杂物在腐蚀过程中可通过水解生成OH-，降低了近表面腐蚀介质的酸度，减缓了腐蚀速度[44].Zhu等[45]也得到了类似的结论，即CaS等夹杂物有助于降低近表面腐蚀介质的酸度，提高pH值，防止酸性腐蚀介质对钢基体的进一步腐蚀.经过Ca处理后，在Ca氧化物外层析出的带有硫化物和氮化物的微纳复合夹杂物可以有效地促进针状铁素体的成核，诱导针状铁素体的形成[34].针状铁素体的纵横交错的结构可以有效地抑制裂纹的增长，减少应力腐蚀开裂的风险.Ca处理后会产生大量的 (Ca、Mg、Mn)S、SiO2和CaS夹杂物，其中CaS夹杂物的耐腐蚀性最差，这被认为是诱发焊接区腐蚀萌生和扩展的主要因素[46,47].
SEM image and EDS results (a), TEM image and SAD result (b) of TiN inclusions, interface morphology between TiN inclusions and SiO<sub>2</sub> (c), and XPS spectra of TiO<sub>2</sub> 2p3/2 and TiO<sub>2</sub> 2p1/2 peaks of passivation film formed on pure Ti sample (d)<sup> [<xref ref-type="bibr" rid="R25">25</xref>]</sup>Fig.6
Composition distribution of active/inactive inclusions in the phase diagram of the Al<sub>2</sub>O<sub>3</sub>-MgO-CaO system (a)<sup> [<xref ref-type="bibr" rid="R45">45</xref>]</sup> and MgS (b), MgY<sub>2</sub>S<sub>4</sub> (c), Y<sub>2</sub>O<sub>3</sub> (d) and YS (e) energy band structure diagram <sup>[<xref ref-type="bibr" rid="R12">12</xref>]</sup>Fig.7
Role of Al2O3 inclusions on the localized corrosion of Q460NH weathering steel in marine environment
6
2018
... 第二种机理为化学溶解机制，该机制认为钢中绝缘非金属夹杂物 (如Al2O3[16]等) 由于自身导电性较差，无法与钢基体构成腐蚀电偶，因此腐蚀主要由化学溶解过程引发.根据夹杂物自身的化学稳定性可分为两种情形：若夹杂物自身化学稳定性差，其主要通过自身的化学溶解形成点蚀坑，诱发点蚀萌生和发展.若夹杂物自身稳定性好，则夹杂物自身不发生溶解，而通过夹杂物与基体之间的性能差异，如塑性和热膨胀系数等，造成钢基体局部高电化学活性，最终导致夹杂物周围钢基体的溶解.Liu等研究表明，Al2O3[16]，ZrO2-Ti2O3-Al2O3[17]和 (RE)AlO3-(RE)2O2S-(RE) x S y[18, 19]等均为绝缘体，图2a中Al2O3的CSAFM结果表明夹杂物处的电流值仅为0~2 pA，而基体部分的电流值为5 nA，远大于夹杂物处的电流值，证实Al2O3为绝缘性夹杂，无法和钢基体构成腐蚀电偶.图2b和c中ZrO2-Ti2O3-Al2O3夹杂和(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3夹杂物同样被证明为绝缘体，不能与基体构成腐蚀电偶.而图2d中可以看到，在西沙模拟液中浸泡30 min后，ZrO2-Ti2O3-Al2O3夹杂并未溶解，这可以归因于该夹杂物自身具有较高的化学稳定性.对比之下，(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3中(RE)2O2S-(RE) x S y 由于较差的化学稳定性而自身溶解.Liu的工作让研究者们认识到[16~19]，仅仅采用电偶腐蚀理论解释和理解夹杂物诱发局部腐蚀是不可接受的，绝缘非金属夹杂物诱发局部腐蚀可能是一种不同于电偶腐蚀的全新机理.根据Gutman的电化学理论[20]，图2e中绝缘体夹杂物ZrO2-Ti2O3-Al2O3附近的高晶格畸变区域增强了基体的反应活性，高活性钢基体与周围非高活性区域钢基体之间可能形成电偶腐蚀，这为非金属夹杂物诱发局部腐蚀的解释提供了新的思路.从上述文献的结果中可以看出，对于不导电或导电性差的非金属夹杂物，其诱导局部腐蚀萌生的机理不同于MnS等导电夹杂物，微缝隙或夹杂物周围高化学活性的区域可能成为诱发点蚀的主因.Hu等[21]通过SKPFM测试证实了 (Al, Mg, Ca, Mn)-氧硫复合夹杂物与基体之间的局部电化学不均匀性，结果表明该夹杂物的Volta电位高于基体，但电流密度仅为8 pA，因此传统的微电偶腐蚀机制并不适用.相比之下，EBSD测试结果表明该夹杂物周围存在高晶格畸变和位错密度，这使得夹杂物周围基体成为热力学不稳定区域，成为优先腐蚀的部分. ...
... [16]，ZrO2-Ti2O3-Al2O3[17]和 (RE)AlO3-(RE)2O2S-(RE) x S y[18, 19]等均为绝缘体，图2a中Al2O3的CSAFM结果表明夹杂物处的电流值仅为0~2 pA，而基体部分的电流值为5 nA，远大于夹杂物处的电流值，证实Al2O3为绝缘性夹杂，无法和钢基体构成腐蚀电偶.图2b和c中ZrO2-Ti2O3-Al2O3夹杂和(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3夹杂物同样被证明为绝缘体，不能与基体构成腐蚀电偶.而图2d中可以看到，在西沙模拟液中浸泡30 min后，ZrO2-Ti2O3-Al2O3夹杂并未溶解，这可以归因于该夹杂物自身具有较高的化学稳定性.对比之下，(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3中(RE)2O2S-(RE) x S y 由于较差的化学稳定性而自身溶解.Liu的工作让研究者们认识到[16~19]，仅仅采用电偶腐蚀理论解释和理解夹杂物诱发局部腐蚀是不可接受的，绝缘非金属夹杂物诱发局部腐蚀可能是一种不同于电偶腐蚀的全新机理.根据Gutman的电化学理论[20]，图2e中绝缘体夹杂物ZrO2-Ti2O3-Al2O3附近的高晶格畸变区域增强了基体的反应活性，高活性钢基体与周围非高活性区域钢基体之间可能形成电偶腐蚀，这为非金属夹杂物诱发局部腐蚀的解释提供了新的思路.从上述文献的结果中可以看出，对于不导电或导电性差的非金属夹杂物，其诱导局部腐蚀萌生的机理不同于MnS等导电夹杂物，微缝隙或夹杂物周围高化学活性的区域可能成为诱发点蚀的主因.Hu等[21]通过SKPFM测试证实了 (Al, Mg, Ca, Mn)-氧硫复合夹杂物与基体之间的局部电化学不均匀性，结果表明该夹杂物的Volta电位高于基体，但电流密度仅为8 pA，因此传统的微电偶腐蚀机制并不适用.相比之下，EBSD测试结果表明该夹杂物周围存在高晶格畸变和位错密度，这使得夹杂物周围基体成为热力学不稳定区域，成为优先腐蚀的部分. ...
... [16~19]，仅仅采用电偶腐蚀理论解释和理解夹杂物诱发局部腐蚀是不可接受的，绝缘非金属夹杂物诱发局部腐蚀可能是一种不同于电偶腐蚀的全新机理.根据Gutman的电化学理论[20]，图2e中绝缘体夹杂物ZrO2-Ti2O3-Al2O3附近的高晶格畸变区域增强了基体的反应活性，高活性钢基体与周围非高活性区域钢基体之间可能形成电偶腐蚀，这为非金属夹杂物诱发局部腐蚀的解释提供了新的思路.从上述文献的结果中可以看出，对于不导电或导电性差的非金属夹杂物，其诱导局部腐蚀萌生的机理不同于MnS等导电夹杂物，微缝隙或夹杂物周围高化学活性的区域可能成为诱发点蚀的主因.Hu等[21]通过SKPFM测试证实了 (Al, Mg, Ca, Mn)-氧硫复合夹杂物与基体之间的局部电化学不均匀性，结果表明该夹杂物的Volta电位高于基体，但电流密度仅为8 pA，因此传统的微电偶腐蚀机制并不适用.相比之下，EBSD测试结果表明该夹杂物周围存在高晶格畸变和位错密度，这使得夹杂物周围基体成为热力学不稳定区域，成为优先腐蚀的部分. ...
... [16~19]SEM and CAAFM images of Al<sub>2</sub>O<sub>3</sub>(a), ZrO<sub>2</sub>-Ti<sub>2</sub>O<sub>3</sub>-Al<sub>2</sub>O<sub>3</sub> (b) and (Re)<sub>2</sub>O<sub>2</sub>S-(Re) <i><sub>x</sub></i> S <i><sub>y</sub></i> -(Re, Zr, Ti)O <i><sub>x</sub></i> (c) and their corresponding height/current distribution diagrams, corrosion morphology of ZrO<sub>2</sub>-Ti<sub>2</sub>O<sub>3</sub>-Al<sub>2</sub>O<sub>3</sub> and ((RE)<sub>2</sub>O<sub>2</sub>S-(RE) <i><sub>x</sub></i> S <i><sub>y</sub></i> -(RE,Zr,Ti)O <i><sub>x</sub></i> -(RE)AlO<sub>3</sub> after soaking in simulated Xisha solution with pH=4.9 for 30 min (d), and KAM diagrams of Al<sub>2</sub>O<sub>3</sub> and ZrO<sub>2</sub>-Ti<sub>2</sub>O<sub>3</sub>-Al<sub>2</sub>O<sub>3</sub> (e). The high KAM indicates that the local plastic deformation of the steel matrix occurs around the inclusion which is difficult to deform, and the mechanical deformation will lead to the redistribution of the surface electrochemical heterogeneity <sup>[<xref ref-type="bibr" rid="R16">16</xref>~<xref ref-type="bibr" rid="R19">19</xref>]</sup>Fig.2
Influence of rare earth metals on mechanisms of localised corrosion induced by inclusions in Zr-Ti deoxidised low alloy steel
2
2020
... 第二种机理为化学溶解机制，该机制认为钢中绝缘非金属夹杂物 (如Al2O3[16]等) 由于自身导电性较差，无法与钢基体构成腐蚀电偶，因此腐蚀主要由化学溶解过程引发.根据夹杂物自身的化学稳定性可分为两种情形：若夹杂物自身化学稳定性差，其主要通过自身的化学溶解形成点蚀坑，诱发点蚀萌生和发展.若夹杂物自身稳定性好，则夹杂物自身不发生溶解，而通过夹杂物与基体之间的性能差异，如塑性和热膨胀系数等，造成钢基体局部高电化学活性，最终导致夹杂物周围钢基体的溶解.Liu等研究表明，Al2O3[16]，ZrO2-Ti2O3-Al2O3[17]和 (RE)AlO3-(RE)2O2S-(RE) x S y[18, 19]等均为绝缘体，图2a中Al2O3的CSAFM结果表明夹杂物处的电流值仅为0~2 pA，而基体部分的电流值为5 nA，远大于夹杂物处的电流值，证实Al2O3为绝缘性夹杂，无法和钢基体构成腐蚀电偶.图2b和c中ZrO2-Ti2O3-Al2O3夹杂和(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3夹杂物同样被证明为绝缘体，不能与基体构成腐蚀电偶.而图2d中可以看到，在西沙模拟液中浸泡30 min后，ZrO2-Ti2O3-Al2O3夹杂并未溶解，这可以归因于该夹杂物自身具有较高的化学稳定性.对比之下，(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3中(RE)2O2S-(RE) x S y 由于较差的化学稳定性而自身溶解.Liu的工作让研究者们认识到[16~19]，仅仅采用电偶腐蚀理论解释和理解夹杂物诱发局部腐蚀是不可接受的，绝缘非金属夹杂物诱发局部腐蚀可能是一种不同于电偶腐蚀的全新机理.根据Gutman的电化学理论[20]，图2e中绝缘体夹杂物ZrO2-Ti2O3-Al2O3附近的高晶格畸变区域增强了基体的反应活性，高活性钢基体与周围非高活性区域钢基体之间可能形成电偶腐蚀，这为非金属夹杂物诱发局部腐蚀的解释提供了新的思路.从上述文献的结果中可以看出，对于不导电或导电性差的非金属夹杂物，其诱导局部腐蚀萌生的机理不同于MnS等导电夹杂物，微缝隙或夹杂物周围高化学活性的区域可能成为诱发点蚀的主因.Hu等[21]通过SKPFM测试证实了 (Al, Mg, Ca, Mn)-氧硫复合夹杂物与基体之间的局部电化学不均匀性，结果表明该夹杂物的Volta电位高于基体，但电流密度仅为8 pA，因此传统的微电偶腐蚀机制并不适用.相比之下，EBSD测试结果表明该夹杂物周围存在高晶格畸变和位错密度，这使得夹杂物周围基体成为热力学不稳定区域，成为优先腐蚀的部分. ...
Towards a better understanding of localised corrosion induced by typical non-metallic inclusions in low-alloy steels
5
2021
... 第二种机理为化学溶解机制，该机制认为钢中绝缘非金属夹杂物 (如Al2O3[16]等) 由于自身导电性较差，无法与钢基体构成腐蚀电偶，因此腐蚀主要由化学溶解过程引发.根据夹杂物自身的化学稳定性可分为两种情形：若夹杂物自身化学稳定性差，其主要通过自身的化学溶解形成点蚀坑，诱发点蚀萌生和发展.若夹杂物自身稳定性好，则夹杂物自身不发生溶解，而通过夹杂物与基体之间的性能差异，如塑性和热膨胀系数等，造成钢基体局部高电化学活性，最终导致夹杂物周围钢基体的溶解.Liu等研究表明，Al2O3[16]，ZrO2-Ti2O3-Al2O3[17]和 (RE)AlO3-(RE)2O2S-(RE) x S y[18, 19]等均为绝缘体，图2a中Al2O3的CSAFM结果表明夹杂物处的电流值仅为0~2 pA，而基体部分的电流值为5 nA，远大于夹杂物处的电流值，证实Al2O3为绝缘性夹杂，无法和钢基体构成腐蚀电偶.图2b和c中ZrO2-Ti2O3-Al2O3夹杂和(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3夹杂物同样被证明为绝缘体，不能与基体构成腐蚀电偶.而图2d中可以看到，在西沙模拟液中浸泡30 min后，ZrO2-Ti2O3-Al2O3夹杂并未溶解，这可以归因于该夹杂物自身具有较高的化学稳定性.对比之下，(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3中(RE)2O2S-(RE) x S y 由于较差的化学稳定性而自身溶解.Liu的工作让研究者们认识到[16~19]，仅仅采用电偶腐蚀理论解释和理解夹杂物诱发局部腐蚀是不可接受的，绝缘非金属夹杂物诱发局部腐蚀可能是一种不同于电偶腐蚀的全新机理.根据Gutman的电化学理论[20]，图2e中绝缘体夹杂物ZrO2-Ti2O3-Al2O3附近的高晶格畸变区域增强了基体的反应活性，高活性钢基体与周围非高活性区域钢基体之间可能形成电偶腐蚀，这为非金属夹杂物诱发局部腐蚀的解释提供了新的思路.从上述文献的结果中可以看出，对于不导电或导电性差的非金属夹杂物，其诱导局部腐蚀萌生的机理不同于MnS等导电夹杂物，微缝隙或夹杂物周围高化学活性的区域可能成为诱发点蚀的主因.Hu等[21]通过SKPFM测试证实了 (Al, Mg, Ca, Mn)-氧硫复合夹杂物与基体之间的局部电化学不均匀性，结果表明该夹杂物的Volta电位高于基体，但电流密度仅为8 pA，因此传统的微电偶腐蚀机制并不适用.相比之下，EBSD测试结果表明该夹杂物周围存在高晶格畸变和位错密度，这使得夹杂物周围基体成为热力学不稳定区域，成为优先腐蚀的部分. ...
... [18]Surface morphology (a), cross-sectional morphology (b), EBSD test region (c) and corresponding KAM plots (d) of Al<sub>2</sub>O<sub>3</sub><sup> [<xref ref-type="bibr" rid="R18">18</xref>]</sup>Fig.5<strong>2.3</strong> 氮化物系夹杂
Effect of inclusions modified by rare earth elements (Ce, La) on localized marine corrosion in Q460NH weathering steel
4
2017
... 第二种机理为化学溶解机制，该机制认为钢中绝缘非金属夹杂物 (如Al2O3[16]等) 由于自身导电性较差，无法与钢基体构成腐蚀电偶，因此腐蚀主要由化学溶解过程引发.根据夹杂物自身的化学稳定性可分为两种情形：若夹杂物自身化学稳定性差，其主要通过自身的化学溶解形成点蚀坑，诱发点蚀萌生和发展.若夹杂物自身稳定性好，则夹杂物自身不发生溶解，而通过夹杂物与基体之间的性能差异，如塑性和热膨胀系数等，造成钢基体局部高电化学活性，最终导致夹杂物周围钢基体的溶解.Liu等研究表明，Al2O3[16]，ZrO2-Ti2O3-Al2O3[17]和 (RE)AlO3-(RE)2O2S-(RE) x S y[18, 19]等均为绝缘体，图2a中Al2O3的CSAFM结果表明夹杂物处的电流值仅为0~2 pA，而基体部分的电流值为5 nA，远大于夹杂物处的电流值，证实Al2O3为绝缘性夹杂，无法和钢基体构成腐蚀电偶.图2b和c中ZrO2-Ti2O3-Al2O3夹杂和(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3夹杂物同样被证明为绝缘体，不能与基体构成腐蚀电偶.而图2d中可以看到，在西沙模拟液中浸泡30 min后，ZrO2-Ti2O3-Al2O3夹杂并未溶解，这可以归因于该夹杂物自身具有较高的化学稳定性.对比之下，(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3中(RE)2O2S-(RE) x S y 由于较差的化学稳定性而自身溶解.Liu的工作让研究者们认识到[16~19]，仅仅采用电偶腐蚀理论解释和理解夹杂物诱发局部腐蚀是不可接受的，绝缘非金属夹杂物诱发局部腐蚀可能是一种不同于电偶腐蚀的全新机理.根据Gutman的电化学理论[20]，图2e中绝缘体夹杂物ZrO2-Ti2O3-Al2O3附近的高晶格畸变区域增强了基体的反应活性，高活性钢基体与周围非高活性区域钢基体之间可能形成电偶腐蚀，这为非金属夹杂物诱发局部腐蚀的解释提供了新的思路.从上述文献的结果中可以看出，对于不导电或导电性差的非金属夹杂物，其诱导局部腐蚀萌生的机理不同于MnS等导电夹杂物，微缝隙或夹杂物周围高化学活性的区域可能成为诱发点蚀的主因.Hu等[21]通过SKPFM测试证实了 (Al, Mg, Ca, Mn)-氧硫复合夹杂物与基体之间的局部电化学不均匀性，结果表明该夹杂物的Volta电位高于基体，但电流密度仅为8 pA，因此传统的微电偶腐蚀机制并不适用.相比之下，EBSD测试结果表明该夹杂物周围存在高晶格畸变和位错密度，这使得夹杂物周围基体成为热力学不稳定区域，成为优先腐蚀的部分. ...
... ~19]，仅仅采用电偶腐蚀理论解释和理解夹杂物诱发局部腐蚀是不可接受的，绝缘非金属夹杂物诱发局部腐蚀可能是一种不同于电偶腐蚀的全新机理.根据Gutman的电化学理论[20]，图2e中绝缘体夹杂物ZrO2-Ti2O3-Al2O3附近的高晶格畸变区域增强了基体的反应活性，高活性钢基体与周围非高活性区域钢基体之间可能形成电偶腐蚀，这为非金属夹杂物诱发局部腐蚀的解释提供了新的思路.从上述文献的结果中可以看出，对于不导电或导电性差的非金属夹杂物，其诱导局部腐蚀萌生的机理不同于MnS等导电夹杂物，微缝隙或夹杂物周围高化学活性的区域可能成为诱发点蚀的主因.Hu等[21]通过SKPFM测试证实了 (Al, Mg, Ca, Mn)-氧硫复合夹杂物与基体之间的局部电化学不均匀性，结果表明该夹杂物的Volta电位高于基体，但电流密度仅为8 pA，因此传统的微电偶腐蚀机制并不适用.相比之下，EBSD测试结果表明该夹杂物周围存在高晶格畸变和位错密度，这使得夹杂物周围基体成为热力学不稳定区域，成为优先腐蚀的部分. ...
... ~19]SEM and CAAFM images of Al<sub>2</sub>O<sub>3</sub>(a), ZrO<sub>2</sub>-Ti<sub>2</sub>O<sub>3</sub>-Al<sub>2</sub>O<sub>3</sub> (b) and (Re)<sub>2</sub>O<sub>2</sub>S-(Re) <i><sub>x</sub></i> S <i><sub>y</sub></i> -(Re, Zr, Ti)O <i><sub>x</sub></i> (c) and their corresponding height/current distribution diagrams, corrosion morphology of ZrO<sub>2</sub>-Ti<sub>2</sub>O<sub>3</sub>-Al<sub>2</sub>O<sub>3</sub> and ((RE)<sub>2</sub>O<sub>2</sub>S-(RE) <i><sub>x</sub></i> S <i><sub>y</sub></i> -(RE,Zr,Ti)O <i><sub>x</sub></i> -(RE)AlO<sub>3</sub> after soaking in simulated Xisha solution with pH=4.9 for 30 min (d), and KAM diagrams of Al<sub>2</sub>O<sub>3</sub> and ZrO<sub>2</sub>-Ti<sub>2</sub>O<sub>3</sub>-Al<sub>2</sub>O<sub>3</sub> (e). The high KAM indicates that the local plastic deformation of the steel matrix occurs around the inclusion which is difficult to deform, and the mechanical deformation will lead to the redistribution of the surface electrochemical heterogeneity <sup>[<xref ref-type="bibr" rid="R16">16</xref>~<xref ref-type="bibr" rid="R19">19</xref>]</sup>Fig.2
Effects of rare earth metals addition on the resistance to pitting corrosion of super duplex stainless steel-Part 1
1
2010
... 第二种机理为化学溶解机制，该机制认为钢中绝缘非金属夹杂物 (如Al2O3[16]等) 由于自身导电性较差，无法与钢基体构成腐蚀电偶，因此腐蚀主要由化学溶解过程引发.根据夹杂物自身的化学稳定性可分为两种情形：若夹杂物自身化学稳定性差，其主要通过自身的化学溶解形成点蚀坑，诱发点蚀萌生和发展.若夹杂物自身稳定性好，则夹杂物自身不发生溶解，而通过夹杂物与基体之间的性能差异，如塑性和热膨胀系数等，造成钢基体局部高电化学活性，最终导致夹杂物周围钢基体的溶解.Liu等研究表明，Al2O3[16]，ZrO2-Ti2O3-Al2O3[17]和 (RE)AlO3-(RE)2O2S-(RE) x S y[18, 19]等均为绝缘体，图2a中Al2O3的CSAFM结果表明夹杂物处的电流值仅为0~2 pA，而基体部分的电流值为5 nA，远大于夹杂物处的电流值，证实Al2O3为绝缘性夹杂，无法和钢基体构成腐蚀电偶.图2b和c中ZrO2-Ti2O3-Al2O3夹杂和(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3夹杂物同样被证明为绝缘体，不能与基体构成腐蚀电偶.而图2d中可以看到，在西沙模拟液中浸泡30 min后，ZrO2-Ti2O3-Al2O3夹杂并未溶解，这可以归因于该夹杂物自身具有较高的化学稳定性.对比之下，(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3中(RE)2O2S-(RE) x S y 由于较差的化学稳定性而自身溶解.Liu的工作让研究者们认识到[16~19]，仅仅采用电偶腐蚀理论解释和理解夹杂物诱发局部腐蚀是不可接受的，绝缘非金属夹杂物诱发局部腐蚀可能是一种不同于电偶腐蚀的全新机理.根据Gutman的电化学理论[20]，图2e中绝缘体夹杂物ZrO2-Ti2O3-Al2O3附近的高晶格畸变区域增强了基体的反应活性，高活性钢基体与周围非高活性区域钢基体之间可能形成电偶腐蚀，这为非金属夹杂物诱发局部腐蚀的解释提供了新的思路.从上述文献的结果中可以看出，对于不导电或导电性差的非金属夹杂物，其诱导局部腐蚀萌生的机理不同于MnS等导电夹杂物，微缝隙或夹杂物周围高化学活性的区域可能成为诱发点蚀的主因.Hu等[21]通过SKPFM测试证实了 (Al, Mg, Ca, Mn)-氧硫复合夹杂物与基体之间的局部电化学不均匀性，结果表明该夹杂物的Volta电位高于基体，但电流密度仅为8 pA，因此传统的微电偶腐蚀机制并不适用.相比之下，EBSD测试结果表明该夹杂物周围存在高晶格畸变和位错密度，这使得夹杂物周围基体成为热力学不稳定区域，成为优先腐蚀的部分. ...
Insight into the triggering effect of (Al, Mg, Ca, Mn)-oxy-sulfide inclusions on localized corrosion of weathering steel
1
2021
... 第二种机理为化学溶解机制，该机制认为钢中绝缘非金属夹杂物 (如Al2O3[16]等) 由于自身导电性较差，无法与钢基体构成腐蚀电偶，因此腐蚀主要由化学溶解过程引发.根据夹杂物自身的化学稳定性可分为两种情形：若夹杂物自身化学稳定性差，其主要通过自身的化学溶解形成点蚀坑，诱发点蚀萌生和发展.若夹杂物自身稳定性好，则夹杂物自身不发生溶解，而通过夹杂物与基体之间的性能差异，如塑性和热膨胀系数等，造成钢基体局部高电化学活性，最终导致夹杂物周围钢基体的溶解.Liu等研究表明，Al2O3[16]，ZrO2-Ti2O3-Al2O3[17]和 (RE)AlO3-(RE)2O2S-(RE) x S y[18, 19]等均为绝缘体，图2a中Al2O3的CSAFM结果表明夹杂物处的电流值仅为0~2 pA，而基体部分的电流值为5 nA，远大于夹杂物处的电流值，证实Al2O3为绝缘性夹杂，无法和钢基体构成腐蚀电偶.图2b和c中ZrO2-Ti2O3-Al2O3夹杂和(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3夹杂物同样被证明为绝缘体，不能与基体构成腐蚀电偶.而图2d中可以看到，在西沙模拟液中浸泡30 min后，ZrO2-Ti2O3-Al2O3夹杂并未溶解，这可以归因于该夹杂物自身具有较高的化学稳定性.对比之下，(RE)2O2S-(RE) x S y -(RE,Zr,Ti)O x -(RE)AlO3中(RE)2O2S-(RE) x S y 由于较差的化学稳定性而自身溶解.Liu的工作让研究者们认识到[16~19]，仅仅采用电偶腐蚀理论解释和理解夹杂物诱发局部腐蚀是不可接受的，绝缘非金属夹杂物诱发局部腐蚀可能是一种不同于电偶腐蚀的全新机理.根据Gutman的电化学理论[20]，图2e中绝缘体夹杂物ZrO2-Ti2O3-Al2O3附近的高晶格畸变区域增强了基体的反应活性，高活性钢基体与周围非高活性区域钢基体之间可能形成电偶腐蚀，这为非金属夹杂物诱发局部腐蚀的解释提供了新的思路.从上述文献的结果中可以看出，对于不导电或导电性差的非金属夹杂物，其诱导局部腐蚀萌生的机理不同于MnS等导电夹杂物，微缝隙或夹杂物周围高化学活性的区域可能成为诱发点蚀的主因.Hu等[21]通过SKPFM测试证实了 (Al, Mg, Ca, Mn)-氧硫复合夹杂物与基体之间的局部电化学不均匀性，结果表明该夹杂物的Volta电位高于基体，但电流密度仅为8 pA，因此传统的微电偶腐蚀机制并不适用.相比之下，EBSD测试结果表明该夹杂物周围存在高晶格畸变和位错密度，这使得夹杂物周围基体成为热力学不稳定区域，成为优先腐蚀的部分. ...
Influence of inclusions on initiation of pitting corrosion and stress corrosion cracking of X70 steel in near-neutral pH environment
... [22,23]Typical corrosion pits formed on Al-Ca-O-S inclusions after 24 h immersion in NS4 solution: (a) optical image, (b) 2D image, (c) three-dimensional profile, (d) height line, where the position of the cross-sectional profile is marked in Fig.3b as a red line <sup>[<xref ref-type="bibr" rid="R22">22</xref>]</sup>, (e) schematic of dissolution kinetics of sulfide-oxide complex inclusions and the resulting local corrosion process: the original inclusions underwent the galvanic corrosion stage and the corrosion spreading stage, respectively<sup> [<xref ref-type="bibr" rid="R23">23</xref>]</sup>Fig.3<strong>2</strong> 几种典型夹杂物对低合金钢局部腐蚀的影响规律<strong>2.1</strong> 硫化物系夹杂
... [22], (e) schematic of dissolution kinetics of sulfide-oxide complex inclusions and the resulting local corrosion process: the original inclusions underwent the galvanic corrosion stage and the corrosion spreading stage, respectively [23]Fig.3<strong>2</strong> 几种典型夹杂物对低合金钢局部腐蚀的影响规律<strong>2.1</strong> 硫化物系夹杂
... ,23]Typical corrosion pits formed on Al-Ca-O-S inclusions after 24 h immersion in NS4 solution: (a) optical image, (b) 2D image, (c) three-dimensional profile, (d) height line, where the position of the cross-sectional profile is marked in Fig.3b as a red line <sup>[<xref ref-type="bibr" rid="R22">22</xref>]</sup>, (e) schematic of dissolution kinetics of sulfide-oxide complex inclusions and the resulting local corrosion process: the original inclusions underwent the galvanic corrosion stage and the corrosion spreading stage, respectively<sup> [<xref ref-type="bibr" rid="R23">23</xref>]</sup>Fig.3<strong>2</strong> 几种典型夹杂物对低合金钢局部腐蚀的影响规律<strong>2.1</strong> 硫化物系夹杂
SEM image and EDS results (a), TEM image and SAD result (b) of TiN inclusions, interface morphology between TiN inclusions and SiO<sub>2</sub> (c), and XPS spectra of TiO<sub>2</sub> 2p3/2 and TiO<sub>2</sub> 2p1/2 peaks of passivation film formed on pure Ti sample (d)<sup> [<xref ref-type="bibr" rid="R25">25</xref>]</sup>Fig.6
Composition distribution of active/inactive inclusions in the phase diagram of the Al<sub>2</sub>O<sub>3</sub>-MgO-CaO system (a)<sup> [<xref ref-type="bibr" rid="R45">45</xref>]</sup> and MgS (b), MgY<sub>2</sub>S<sub>4</sub> (c), Y<sub>2</sub>O<sub>3</sub> (d) and YS (e) energy band structure diagram <sup>[<xref ref-type="bibr" rid="R12">12</xref>]</sup>Fig.7
Composition distribution of active/inactive inclusions in the phase diagram of the Al<sub>2</sub>O<sub>3</sub>-MgO-CaO system (a)<sup> [<xref ref-type="bibr" rid="R45">45</xref>]</sup> and MgS (b), MgY<sub>2</sub>S<sub>4</sub> (c), Y<sub>2</sub>O<sub>3</sub> (d) and YS (e) energy band structure diagram <sup>[<xref ref-type="bibr" rid="R12">12</xref>]</sup>Fig.7
... [29]Surface potential distribution of MnS (a), surface work function of different crystal faces of MnS (b), the results show that the work function of MnS is smaller than that of Fe matrix, and mechanism diagram of local corrosion induced by MnS (c), MnS is used as an anodic corrosion couple with steel substrate to form steady or metastable pitting pits under the synergistic effect of Cl<sup>-</sup>and S, and the corrosion product “S” shell covers the surface of pitting pits to form concentration difference cells, further promote the corrosion process<sup> [<xref ref-type="bibr" rid="R29">29</xref>]</sup>Fig.4
... [29]Surface potential distribution of MnS (a), surface work function of different crystal faces of MnS (b), the results show that the work function of MnS is smaller than that of Fe matrix, and mechanism diagram of local corrosion induced by MnS (c), MnS is used as an anodic corrosion couple with steel substrate to form steady or metastable pitting pits under the synergistic effect of Cl<sup>-</sup>and S, and the corrosion product “S” shell covers the surface of pitting pits to form concentration difference cells, further promote the corrosion process<sup> [<xref ref-type="bibr" rid="R29">29</xref>]</sup>Fig.4
A microelectrochemical system for in situ high-resolution optical microscopy: morphological characteristics of pitting at MnS inclusion in stainless steel
Transient inclusion evolution during modification of alumina inclusions by calcium in liquid steel: part I. Background, experimental techniques and analysis methods
1
2011
... 随着第三代氧化物冶金技术的广泛应用，Al脱氧Ca处理和Mg-Al脱氧Ca处理技术日趋成熟，钙处理对钢中夹杂物的改性作用对钢基体耐蚀性产生了重要影响.Ca处理后，钢中夹杂物的尺寸明显减小，对夹杂物的成分、形状和尺寸的控制作用明显[39].经Ca处理后，钢中高硬度、不规则的Al2O3夹杂物转变为xCaO-yAl2O3复合氧化物夹杂物，其硬度低、呈球状，并可能伴随着xCaS-yCaO-zAl2O3复合氧化物夹杂物的形成[40].在各种经Ca处理改性的铝酸钙和xCaO-yAl2O3化合物中，12CaO-7Al2O3(最低熔点1455 ℃)是一种理想的改性产物，可以降低原夹杂物与钢基体之间的应力[41, 42].Wang等[43]研究了EH36钢中(Ca, Mg, Al)-O x -S y 夹杂物和局部腐蚀的关联，图7a显示了Al2O3-MgO-CaO三元相图，CaO:Al2O3比例约为1.25是区分活性和非活性夹杂物的界限，且通过对比和计算夹杂物和基体之前的热膨胀系数和残余应力，验证了非活性夹杂物3CaO·Al2O3周围不存在明显的局部塑形变形，说明其并不会诱发点蚀.Ca处理不仅可以改性非金属氧化物夹杂物，也可以改性硫化物夹杂物.相关研究表明在X80钢管生产线中，Ca处理还可以将条状的MnS夹杂物转化为球形的CaS夹杂物和 (Ca, Mn)S复合夹杂物，降低夹杂物与钢基体之间的应力.同时，含Ca的夹杂物在腐蚀过程中可通过水解生成OH-，降低了近表面腐蚀介质的酸度，减缓了腐蚀速度[44].Zhu等[45]也得到了类似的结论，即CaS等夹杂物有助于降低近表面腐蚀介质的酸度，提高pH值，防止酸性腐蚀介质对钢基体的进一步腐蚀.经过Ca处理后，在Ca氧化物外层析出的带有硫化物和氮化物的微纳复合夹杂物可以有效地促进针状铁素体的成核，诱导针状铁素体的形成[34].针状铁素体的纵横交错的结构可以有效地抑制裂纹的增长，减少应力腐蚀开裂的风险.Ca处理后会产生大量的 (Ca、Mg、Mn)S、SiO2和CaS夹杂物，其中CaS夹杂物的耐腐蚀性最差，这被认为是诱发焊接区腐蚀萌生和扩展的主要因素[46,47]. ...
Effects of calcium modification on the electrochemical and corrosion properties of weathering steel
1
2002
... 随着第三代氧化物冶金技术的广泛应用，Al脱氧Ca处理和Mg-Al脱氧Ca处理技术日趋成熟，钙处理对钢中夹杂物的改性作用对钢基体耐蚀性产生了重要影响.Ca处理后，钢中夹杂物的尺寸明显减小，对夹杂物的成分、形状和尺寸的控制作用明显[39].经Ca处理后，钢中高硬度、不规则的Al2O3夹杂物转变为xCaO-yAl2O3复合氧化物夹杂物，其硬度低、呈球状，并可能伴随着xCaS-yCaO-zAl2O3复合氧化物夹杂物的形成[40].在各种经Ca处理改性的铝酸钙和xCaO-yAl2O3化合物中，12CaO-7Al2O3(最低熔点1455 ℃)是一种理想的改性产物，可以降低原夹杂物与钢基体之间的应力[41, 42].Wang等[43]研究了EH36钢中(Ca, Mg, Al)-O x -S y 夹杂物和局部腐蚀的关联，图7a显示了Al2O3-MgO-CaO三元相图，CaO:Al2O3比例约为1.25是区分活性和非活性夹杂物的界限，且通过对比和计算夹杂物和基体之前的热膨胀系数和残余应力，验证了非活性夹杂物3CaO·Al2O3周围不存在明显的局部塑形变形，说明其并不会诱发点蚀.Ca处理不仅可以改性非金属氧化物夹杂物，也可以改性硫化物夹杂物.相关研究表明在X80钢管生产线中，Ca处理还可以将条状的MnS夹杂物转化为球形的CaS夹杂物和 (Ca, Mn)S复合夹杂物，降低夹杂物与钢基体之间的应力.同时，含Ca的夹杂物在腐蚀过程中可通过水解生成OH-，降低了近表面腐蚀介质的酸度，减缓了腐蚀速度[44].Zhu等[45]也得到了类似的结论，即CaS等夹杂物有助于降低近表面腐蚀介质的酸度，提高pH值，防止酸性腐蚀介质对钢基体的进一步腐蚀.经过Ca处理后，在Ca氧化物外层析出的带有硫化物和氮化物的微纳复合夹杂物可以有效地促进针状铁素体的成核，诱导针状铁素体的形成[34].针状铁素体的纵横交错的结构可以有效地抑制裂纹的增长，减少应力腐蚀开裂的风险.Ca处理后会产生大量的 (Ca、Mg、Mn)S、SiO2和CaS夹杂物，其中CaS夹杂物的耐腐蚀性最差，这被认为是诱发焊接区腐蚀萌生和扩展的主要因素[46,47]. ...
Observation of calcium aluminate inclusions at interfaces between Ca-treated, Al-killed steels and slags
1
2003
... 随着第三代氧化物冶金技术的广泛应用，Al脱氧Ca处理和Mg-Al脱氧Ca处理技术日趋成熟，钙处理对钢中夹杂物的改性作用对钢基体耐蚀性产生了重要影响.Ca处理后，钢中夹杂物的尺寸明显减小，对夹杂物的成分、形状和尺寸的控制作用明显[39].经Ca处理后，钢中高硬度、不规则的Al2O3夹杂物转变为xCaO-yAl2O3复合氧化物夹杂物，其硬度低、呈球状，并可能伴随着xCaS-yCaO-zAl2O3复合氧化物夹杂物的形成[40].在各种经Ca处理改性的铝酸钙和xCaO-yAl2O3化合物中，12CaO-7Al2O3(最低熔点1455 ℃)是一种理想的改性产物，可以降低原夹杂物与钢基体之间的应力[41, 42].Wang等[43]研究了EH36钢中(Ca, Mg, Al)-O x -S y 夹杂物和局部腐蚀的关联，图7a显示了Al2O3-MgO-CaO三元相图，CaO:Al2O3比例约为1.25是区分活性和非活性夹杂物的界限，且通过对比和计算夹杂物和基体之前的热膨胀系数和残余应力，验证了非活性夹杂物3CaO·Al2O3周围不存在明显的局部塑形变形，说明其并不会诱发点蚀.Ca处理不仅可以改性非金属氧化物夹杂物，也可以改性硫化物夹杂物.相关研究表明在X80钢管生产线中，Ca处理还可以将条状的MnS夹杂物转化为球形的CaS夹杂物和 (Ca, Mn)S复合夹杂物，降低夹杂物与钢基体之间的应力.同时，含Ca的夹杂物在腐蚀过程中可通过水解生成OH-，降低了近表面腐蚀介质的酸度，减缓了腐蚀速度[44].Zhu等[45]也得到了类似的结论，即CaS等夹杂物有助于降低近表面腐蚀介质的酸度，提高pH值，防止酸性腐蚀介质对钢基体的进一步腐蚀.经过Ca处理后，在Ca氧化物外层析出的带有硫化物和氮化物的微纳复合夹杂物可以有效地促进针状铁素体的成核，诱导针状铁素体的形成[34].针状铁素体的纵横交错的结构可以有效地抑制裂纹的增长，减少应力腐蚀开裂的风险.Ca处理后会产生大量的 (Ca、Mg、Mn)S、SiO2和CaS夹杂物，其中CaS夹杂物的耐腐蚀性最差，这被认为是诱发焊接区腐蚀萌生和扩展的主要因素[46,47]. ...
Modification of MgO·Al2O3 spinel inclusions in Al-killed steel by Ca-treatment
1
2011
... 随着第三代氧化物冶金技术的广泛应用，Al脱氧Ca处理和Mg-Al脱氧Ca处理技术日趋成熟，钙处理对钢中夹杂物的改性作用对钢基体耐蚀性产生了重要影响.Ca处理后，钢中夹杂物的尺寸明显减小，对夹杂物的成分、形状和尺寸的控制作用明显[39].经Ca处理后，钢中高硬度、不规则的Al2O3夹杂物转变为xCaO-yAl2O3复合氧化物夹杂物，其硬度低、呈球状，并可能伴随着xCaS-yCaO-zAl2O3复合氧化物夹杂物的形成[40].在各种经Ca处理改性的铝酸钙和xCaO-yAl2O3化合物中，12CaO-7Al2O3(最低熔点1455 ℃)是一种理想的改性产物，可以降低原夹杂物与钢基体之间的应力[41, 42].Wang等[43]研究了EH36钢中(Ca, Mg, Al)-O x -S y 夹杂物和局部腐蚀的关联，图7a显示了Al2O3-MgO-CaO三元相图，CaO:Al2O3比例约为1.25是区分活性和非活性夹杂物的界限，且通过对比和计算夹杂物和基体之前的热膨胀系数和残余应力，验证了非活性夹杂物3CaO·Al2O3周围不存在明显的局部塑形变形，说明其并不会诱发点蚀.Ca处理不仅可以改性非金属氧化物夹杂物，也可以改性硫化物夹杂物.相关研究表明在X80钢管生产线中，Ca处理还可以将条状的MnS夹杂物转化为球形的CaS夹杂物和 (Ca, Mn)S复合夹杂物，降低夹杂物与钢基体之间的应力.同时，含Ca的夹杂物在腐蚀过程中可通过水解生成OH-，降低了近表面腐蚀介质的酸度，减缓了腐蚀速度[44].Zhu等[45]也得到了类似的结论，即CaS等夹杂物有助于降低近表面腐蚀介质的酸度，提高pH值，防止酸性腐蚀介质对钢基体的进一步腐蚀.经过Ca处理后，在Ca氧化物外层析出的带有硫化物和氮化物的微纳复合夹杂物可以有效地促进针状铁素体的成核，诱导针状铁素体的形成[34].针状铁素体的纵横交错的结构可以有效地抑制裂纹的增长，减少应力腐蚀开裂的风险.Ca处理后会产生大量的 (Ca、Mg、Mn)S、SiO2和CaS夹杂物，其中CaS夹杂物的耐腐蚀性最差，这被认为是诱发焊接区腐蚀萌生和扩展的主要因素[46,47]. ...
Correlation between active/inactive (Ca, Mg, Al)-O x -S y inclusions and localised marine corrosion of EH36 steels
1
2021
... 随着第三代氧化物冶金技术的广泛应用，Al脱氧Ca处理和Mg-Al脱氧Ca处理技术日趋成熟，钙处理对钢中夹杂物的改性作用对钢基体耐蚀性产生了重要影响.Ca处理后，钢中夹杂物的尺寸明显减小，对夹杂物的成分、形状和尺寸的控制作用明显[39].经Ca处理后，钢中高硬度、不规则的Al2O3夹杂物转变为xCaO-yAl2O3复合氧化物夹杂物，其硬度低、呈球状，并可能伴随着xCaS-yCaO-zAl2O3复合氧化物夹杂物的形成[40].在各种经Ca处理改性的铝酸钙和xCaO-yAl2O3化合物中，12CaO-7Al2O3(最低熔点1455 ℃)是一种理想的改性产物，可以降低原夹杂物与钢基体之间的应力[41, 42].Wang等[43]研究了EH36钢中(Ca, Mg, Al)-O x -S y 夹杂物和局部腐蚀的关联，图7a显示了Al2O3-MgO-CaO三元相图，CaO:Al2O3比例约为1.25是区分活性和非活性夹杂物的界限，且通过对比和计算夹杂物和基体之前的热膨胀系数和残余应力，验证了非活性夹杂物3CaO·Al2O3周围不存在明显的局部塑形变形，说明其并不会诱发点蚀.Ca处理不仅可以改性非金属氧化物夹杂物，也可以改性硫化物夹杂物.相关研究表明在X80钢管生产线中，Ca处理还可以将条状的MnS夹杂物转化为球形的CaS夹杂物和 (Ca, Mn)S复合夹杂物，降低夹杂物与钢基体之间的应力.同时，含Ca的夹杂物在腐蚀过程中可通过水解生成OH-，降低了近表面腐蚀介质的酸度，减缓了腐蚀速度[44].Zhu等[45]也得到了类似的结论，即CaS等夹杂物有助于降低近表面腐蚀介质的酸度，提高pH值，防止酸性腐蚀介质对钢基体的进一步腐蚀.经过Ca处理后，在Ca氧化物外层析出的带有硫化物和氮化物的微纳复合夹杂物可以有效地促进针状铁素体的成核，诱导针状铁素体的形成[34].针状铁素体的纵横交错的结构可以有效地抑制裂纹的增长，减少应力腐蚀开裂的风险.Ca处理后会产生大量的 (Ca、Mg、Mn)S、SiO2和CaS夹杂物，其中CaS夹杂物的耐腐蚀性最差，这被认为是诱发焊接区腐蚀萌生和扩展的主要因素[46,47]. ...
Characterization of inclusions of X80 pipeline steel and its correlation with hydrogen-induced cracking
1
2011
... 随着第三代氧化物冶金技术的广泛应用，Al脱氧Ca处理和Mg-Al脱氧Ca处理技术日趋成熟，钙处理对钢中夹杂物的改性作用对钢基体耐蚀性产生了重要影响.Ca处理后，钢中夹杂物的尺寸明显减小，对夹杂物的成分、形状和尺寸的控制作用明显[39].经Ca处理后，钢中高硬度、不规则的Al2O3夹杂物转变为xCaO-yAl2O3复合氧化物夹杂物，其硬度低、呈球状，并可能伴随着xCaS-yCaO-zAl2O3复合氧化物夹杂物的形成[40].在各种经Ca处理改性的铝酸钙和xCaO-yAl2O3化合物中，12CaO-7Al2O3(最低熔点1455 ℃)是一种理想的改性产物，可以降低原夹杂物与钢基体之间的应力[41, 42].Wang等[43]研究了EH36钢中(Ca, Mg, Al)-O x -S y 夹杂物和局部腐蚀的关联，图7a显示了Al2O3-MgO-CaO三元相图，CaO:Al2O3比例约为1.25是区分活性和非活性夹杂物的界限，且通过对比和计算夹杂物和基体之前的热膨胀系数和残余应力，验证了非活性夹杂物3CaO·Al2O3周围不存在明显的局部塑形变形，说明其并不会诱发点蚀.Ca处理不仅可以改性非金属氧化物夹杂物，也可以改性硫化物夹杂物.相关研究表明在X80钢管生产线中，Ca处理还可以将条状的MnS夹杂物转化为球形的CaS夹杂物和 (Ca, Mn)S复合夹杂物，降低夹杂物与钢基体之间的应力.同时，含Ca的夹杂物在腐蚀过程中可通过水解生成OH-，降低了近表面腐蚀介质的酸度，减缓了腐蚀速度[44].Zhu等[45]也得到了类似的结论，即CaS等夹杂物有助于降低近表面腐蚀介质的酸度，提高pH值，防止酸性腐蚀介质对钢基体的进一步腐蚀.经过Ca处理后，在Ca氧化物外层析出的带有硫化物和氮化物的微纳复合夹杂物可以有效地促进针状铁素体的成核，诱导针状铁素体的形成[34].针状铁素体的纵横交错的结构可以有效地抑制裂纹的增长，减少应力腐蚀开裂的风险.Ca处理后会产生大量的 (Ca、Mg、Mn)S、SiO2和CaS夹杂物，其中CaS夹杂物的耐腐蚀性最差，这被认为是诱发焊接区腐蚀萌生和扩展的主要因素[46,47]. ...
Effects of inclusion on corrosion resistance of weathering steel in simulated industrial atmosphere
3
2016
... 随着第三代氧化物冶金技术的广泛应用，Al脱氧Ca处理和Mg-Al脱氧Ca处理技术日趋成熟，钙处理对钢中夹杂物的改性作用对钢基体耐蚀性产生了重要影响.Ca处理后，钢中夹杂物的尺寸明显减小，对夹杂物的成分、形状和尺寸的控制作用明显[39].经Ca处理后，钢中高硬度、不规则的Al2O3夹杂物转变为xCaO-yAl2O3复合氧化物夹杂物，其硬度低、呈球状，并可能伴随着xCaS-yCaO-zAl2O3复合氧化物夹杂物的形成[40].在各种经Ca处理改性的铝酸钙和xCaO-yAl2O3化合物中，12CaO-7Al2O3(最低熔点1455 ℃)是一种理想的改性产物，可以降低原夹杂物与钢基体之间的应力[41, 42].Wang等[43]研究了EH36钢中(Ca, Mg, Al)-O x -S y 夹杂物和局部腐蚀的关联，图7a显示了Al2O3-MgO-CaO三元相图，CaO:Al2O3比例约为1.25是区分活性和非活性夹杂物的界限，且通过对比和计算夹杂物和基体之前的热膨胀系数和残余应力，验证了非活性夹杂物3CaO·Al2O3周围不存在明显的局部塑形变形，说明其并不会诱发点蚀.Ca处理不仅可以改性非金属氧化物夹杂物，也可以改性硫化物夹杂物.相关研究表明在X80钢管生产线中，Ca处理还可以将条状的MnS夹杂物转化为球形的CaS夹杂物和 (Ca, Mn)S复合夹杂物，降低夹杂物与钢基体之间的应力.同时，含Ca的夹杂物在腐蚀过程中可通过水解生成OH-，降低了近表面腐蚀介质的酸度，减缓了腐蚀速度[44].Zhu等[45]也得到了类似的结论，即CaS等夹杂物有助于降低近表面腐蚀介质的酸度，提高pH值，防止酸性腐蚀介质对钢基体的进一步腐蚀.经过Ca处理后，在Ca氧化物外层析出的带有硫化物和氮化物的微纳复合夹杂物可以有效地促进针状铁素体的成核，诱导针状铁素体的形成[34].针状铁素体的纵横交错的结构可以有效地抑制裂纹的增长，减少应力腐蚀开裂的风险.Ca处理后会产生大量的 (Ca、Mg、Mn)S、SiO2和CaS夹杂物，其中CaS夹杂物的耐腐蚀性最差，这被认为是诱发焊接区腐蚀萌生和扩展的主要因素[46,47]. ...
... [45]和能带结构图[12]Composition distribution of active/inactive inclusions in the phase diagram of the Al<sub>2</sub>O<sub>3</sub>-MgO-CaO system (a)<sup> [<xref ref-type="bibr" rid="R45">45</xref>]</sup> and MgS (b), MgY<sub>2</sub>S<sub>4</sub> (c), Y<sub>2</sub>O<sub>3</sub> (d) and YS (e) energy band structure diagram <sup>[<xref ref-type="bibr" rid="R12">12</xref>]</sup>Fig.7
