Microstructure and corrosion resistance of nanocrystalline

114Univ. Chem. 2023, 38 (12), 114–119收稿：2023-06-27；录用：2023-08-01；网络发表：2023-08-11*通讯作者，Email:*****************.cn基金资助：2021年基础学科拔尖学生培养计划2.0研究课题(20211014)；天津市首批虚拟教研室试点建设项目(化学类交叉人才培养课程建设虚拟教研室)•专题• doi: 10.3866/PKU.DXHX202306051 当表面活性剂遇到大环分子阮文娟，李悦，耿文超，郭东升*南开大学化学学院，天津 300071摘要：近年来，表面和胶体化学与大环化学的结合引起了科学家的普遍关注。

将多样的大环结构引入表面活性剂分子，不仅极大地丰富了表面活性剂分子的种类，还可以赋予其大环的主客体识别功能。

由此所开发出的大环两亲和超两亲分子已在生物成像和药物递送中表现出很高的应用潜力。

从传统表面活性剂到大环两亲和超两亲分子的发展、应用表明，不同领域的交叉融合对科学研究的发展是非常重要的。

关键词：表面活性剂；胶束；大环结构；大环两亲分子；超两亲分子中图分类号：G64；O6Encountering of Surfactants with Macrocyclic MoleculesWen-Juan Ruan, Yue Li, Wen-Chao Geng, Dong-Sheng Guo *College of Chemistry, Nankai University, Tianjin 300071, China.Abstract: In recent years, the combination of surface and colloid chemistry with macrocyclic chemistry has garnered widespread attention among scientists. The integration of diverse macrocyclic structures into surfactant molecules not only greatly enriches the diversity of surfactants, but also imparts them with the host-guest recognition functionality of macrocycles. Macrocyclic amphiphiles and supra-amphiphiles, developed from this approach, have demonstrated high potential in applications such as bioimaging and drug delivery. The evolution from traditional surfactants to macrocyclic amphiphiles and supra-amphiphiles underscores the importance of interdisciplinary integration in advancing scientific research.Key Words: Surfactants; Micelles; Macrocycles; Macrocyclic amphiphiles; Supra-amphiphiles表面活性剂及其所构筑的胶束是表面和胶体化学中所涉及的一类非常重要的体系。

第27卷第2期粉末冶金材料科学与工程2022年4月V ol.27 No.2 Materials Science and Engineering of Powder Metallurgy Apr. 2022DOI:10.19976/ki.43-1448/TF.2021090激光熔覆马氏体/铁素体涂层的组织与抗磨耐蚀性能张磊1, 2，陈小明1, 2，霍嘉翔1，张凯1, 2，曹文菁1, 2，程新闯3(1. 水利部产品质量标准研究所浙江省水利水电装备表面工程技术研究重点实验室，杭州 310012；2. 水利部杭州机械设计研究所水利机械及其再制造技术浙江省工程实验室，杭州 310012；3. 绍兴市曹娥江大闸管理局，绍兴 312000)摘要：为提高液压活塞杆的耐腐蚀和抗磨损性能，在45号钢表面采用激光熔覆技术在不同激光功率下制备具有马氏体/铁素体组织的Fe基合金熔覆层。

利用X射线衍射仪、扫描电镜、X射线能谱仪等手段表征涂层的物相组成、微观形貌和元素分布，采用维氏硬度计和干滑动摩擦试验机对涂层的显微硬度和抗磨损性能进行测试，并通过电化学工作站研究熔覆层的耐腐蚀性能。

结果表明：Fe基合金熔覆层的主要物相为α-Fe、Ni-Cr-Fe、γ-(Fe,C)和Fe9.7Mo0.3等，主要组织为马氏体、铁素体和少量残余奥氏体。

熔覆层的枝晶态组织均匀致密，无裂纹和孔隙缺陷，涂层与基体呈冶金结合。

涂层的硬度与耐磨性能随激光功率增大而提高，当功率为2.4 kW时，涂层的平均显微硬度(HV)为647.64，耐磨性能为45号钢的9.37倍，磨损机制为磨粒磨损。

随激光功率提高，Fe基合金熔覆层的耐腐蚀性能先升高后降低，当激光功率为2.0 kW时涂层具有最佳耐腐蚀性能，显著高于活塞杆常用碳钢、不锈钢以及电镀硬铬等材料，可在相关领域替代电镀铬。

关键词：激光熔覆；Fe基合金；组织；磨损；腐蚀；活塞杆中图分类号：TG174.44文献标志码：A 文章编号：1673-0224(2022)02-196-09All Rights Reserved.Microstructure and wear-corrosion resistance performance oflaser cladding martensite/ferrite coatingZHANG Lei1, 2, CHEN Xiaoming1, 2, HUO Jiaxiang1, ZHANG Kai1, 2, CAO Wenjing1, 2, CHENG Xinchuang3(1. Key Laboratory of Surface Engineering of Equipment for Hydraulic Engineering of Zhejiang Province, Standard &Quality Control Research Institute, Ministry of Water Resources, Hangzhou 310012, China;2. Water Machinery and Remanufacturing Technology Engineering Laboratory of Zhejiang Province, HangzhouMechanical Research Institute, Ministry of Water Resources, Hangzhou 310012, China;3. Shaoxing Municipal Cao’e River Floodgate Construction Administration Committee, Shaoxing 312000, China)Abstract: To improve the corrosion resistance and wear resistance of piston rod, Fe-based coatings with martensite andferrite structure were prepared on 45# steel by laser cladding. The phase compositions, microstructure and elementsdistribution of the coatings were characterized by X-ray diffractometer, scanning electron microscope and X-ray energydispersive spectrometer. The microhardness and wear resistance of the coatings were tested by Vickers hardness testerand dry sliding friction wear tester. Furthermore, the corrosion resistance of laser cladding Fe-based coatings was studiedby electrochemical workstation. The results show that the phase of laser cladding Fe-based alloy coating is mainlycomposed of α-Fe, Ni-Cr-Fe, γ-(Fe,C), Fe9.7Mo0.3. The main microstructure is martensite, ferrite and a small amount ofresidual austenite. The dendritic structure of coating is uniform, compact, without cracks or pores. The coating and thesubstrate are bonded metallurgically. The hardness and wear resistance of the coatings increase with increasing基金项目：浙江省“一带一路”国际科技合作项目(2019C04019)；浙江省公益性技术应用研究计划资助项目(GC22E017317，LGC19E090001，2018C37029)收稿日期：2021−11−02；修订日期：2021−12−23通信作者：张磊，工程师，硕士。

NORMEINTERNATIONALECEI IEC INTERNATIONALSTANDARD 61854Première éditionFirst edition1998-09Lignes aériennes –Exigences et essais applicables aux entretoisesOverhead lines –Requirements and tests for spacersCommission Electrotechnique InternationaleInternational Electrotechnical Commission Pour prix, voir catalogue en vigueurFor price, see current catalogue© IEC 1998 Droits de reproduction réservés Copyright - all rights reservedAucune partie de cette publication ne peut être reproduite niutilisée sous quelque forme que ce soit et par aucunprocédé, électronique ou mécanique, y compris la photo-copie et les microfilms, sans l'accord écrit de l'éditeur.No part of this publication may be reproduced or utilized in any form or by any means, electronic or mechanical,including photocopying and microfilm, without permission in writing from the publisher.International Electrotechnical Commission 3, rue de Varembé Geneva, SwitzerlandTelefax: +41 22 919 0300e-mail: inmail@iec.ch IEC web site http: //www.iec.chCODE PRIX PRICE CODE X– 2 –61854 © CEI:1998SOMMAIREPages AVANT-PROPOS (6)Articles1Domaine d'application (8)2Références normatives (8)3Définitions (12)4Exigences générales (12)4.1Conception (12)4.2Matériaux (14)4.2.1Généralités (14)4.2.2Matériaux non métalliques (14)4.3Masse, dimensions et tolérances (14)4.4Protection contre la corrosion (14)4.5Aspect et finition de fabrication (14)4.6Marquage (14)4.7Consignes d'installation (14)5Assurance de la qualité (16)6Classification des essais (16)6.1Essais de type (16)6.1.1Généralités (16)6.1.2Application (16)6.2Essais sur échantillon (16)6.2.1Généralités (16)6.2.2Application (16)6.2.3Echantillonnage et critères de réception (18)6.3Essais individuels de série (18)6.3.1Généralités (18)6.3.2Application et critères de réception (18)6.4Tableau des essais à effectuer (18)7Méthodes d'essai (22)7.1Contrôle visuel (22)7.2Vérification des dimensions, des matériaux et de la masse (22)7.3Essai de protection contre la corrosion (22)7.3.1Composants revêtus par galvanisation à chaud (autres queles fils d'acier galvanisés toronnés) (22)7.3.2Produits en fer protégés contre la corrosion par des méthodes autresque la galvanisation à chaud (24)7.3.3Fils d'acier galvanisé toronnés (24)7.3.4Corrosion causée par des composants non métalliques (24)7.4Essais non destructifs (24)61854 © IEC:1998– 3 –CONTENTSPage FOREWORD (7)Clause1Scope (9)2Normative references (9)3Definitions (13)4General requirements (13)4.1Design (13)4.2Materials (15)4.2.1General (15)4.2.2Non-metallic materials (15)4.3Mass, dimensions and tolerances (15)4.4Protection against corrosion (15)4.5Manufacturing appearance and finish (15)4.6Marking (15)4.7Installation instructions (15)5Quality assurance (17)6Classification of tests (17)6.1Type tests (17)6.1.1General (17)6.1.2Application (17)6.2Sample tests (17)6.2.1General (17)6.2.2Application (17)6.2.3Sampling and acceptance criteria (19)6.3Routine tests (19)6.3.1General (19)6.3.2Application and acceptance criteria (19)6.4Table of tests to be applied (19)7Test methods (23)7.1Visual examination (23)7.2Verification of dimensions, materials and mass (23)7.3Corrosion protection test (23)7.3.1Hot dip galvanized components (other than stranded galvanizedsteel wires) (23)7.3.2Ferrous components protected from corrosion by methods other thanhot dip galvanizing (25)7.3.3Stranded galvanized steel wires (25)7.3.4Corrosion caused by non-metallic components (25)7.4Non-destructive tests (25)– 4 –61854 © CEI:1998 Articles Pages7.5Essais mécaniques (26)7.5.1Essais de glissement des pinces (26)7.5.1.1Essai de glissement longitudinal (26)7.5.1.2Essai de glissement en torsion (28)7.5.2Essai de boulon fusible (28)7.5.3Essai de serrage des boulons de pince (30)7.5.4Essais de courant de court-circuit simulé et essais de compressionet de traction (30)7.5.4.1Essai de courant de court-circuit simulé (30)7.5.4.2Essai de compression et de traction (32)7.5.5Caractérisation des propriétés élastiques et d'amortissement (32)7.5.6Essais de flexibilité (38)7.5.7Essais de fatigue (38)7.5.7.1Généralités (38)7.5.7.2Oscillation de sous-portée (40)7.5.7.3Vibrations éoliennes (40)7.6Essais de caractérisation des élastomères (42)7.6.1Généralités (42)7.6.2Essais (42)7.6.3Essai de résistance à l'ozone (46)7.7Essais électriques (46)7.7.1Essais d'effet couronne et de tension de perturbations radioélectriques..467.7.2Essai de résistance électrique (46)7.8Vérification du comportement vibratoire du système faisceau/entretoise (48)Annexe A (normative) Informations techniques minimales à convenirentre acheteur et fournisseur (64)Annexe B (informative) Forces de compression dans l'essai de courantde court-circuit simulé (66)Annexe C (informative) Caractérisation des propriétés élastiques et d'amortissementMéthode de détermination de la rigidité et de l'amortissement (70)Annexe D (informative) Contrôle du comportement vibratoire du systèmefaisceau/entretoise (74)Bibliographie (80)Figures (50)Tableau 1 – Essais sur les entretoises (20)Tableau 2 – Essais sur les élastomères (44)61854 © IEC:1998– 5 –Clause Page7.5Mechanical tests (27)7.5.1Clamp slip tests (27)7.5.1.1Longitudinal slip test (27)7.5.1.2Torsional slip test (29)7.5.2Breakaway bolt test (29)7.5.3Clamp bolt tightening test (31)7.5.4Simulated short-circuit current test and compression and tension tests (31)7.5.4.1Simulated short-circuit current test (31)7.5.4.2Compression and tension test (33)7.5.5Characterisation of the elastic and damping properties (33)7.5.6Flexibility tests (39)7.5.7Fatigue tests (39)7.5.7.1General (39)7.5.7.2Subspan oscillation (41)7.5.7.3Aeolian vibration (41)7.6Tests to characterise elastomers (43)7.6.1General (43)7.6.2Tests (43)7.6.3Ozone resistance test (47)7.7Electrical tests (47)7.7.1Corona and radio interference voltage (RIV) tests (47)7.7.2Electrical resistance test (47)7.8Verification of vibration behaviour of the bundle-spacer system (49)Annex A (normative) Minimum technical details to be agreed betweenpurchaser and supplier (65)Annex B (informative) Compressive forces in the simulated short-circuit current test (67)Annex C (informative) Characterisation of the elastic and damping propertiesStiffness-Damping Method (71)Annex D (informative) Verification of vibration behaviour of the bundle/spacer system (75)Bibliography (81)Figures (51)Table 1 – Tests on spacers (21)Table 2 – Tests on elastomers (45)– 6 –61854 © CEI:1998 COMMISSION ÉLECTROTECHNIQUE INTERNATIONALE––––––––––LIGNES AÉRIENNES –EXIGENCES ET ESSAIS APPLICABLES AUX ENTRETOISESAVANT-PROPOS1)La CEI (Commission Electrotechnique Internationale) est une organisation mondiale de normalisation composéede l'ensemble des comités électrotechniques nationaux (Comités nationaux de la CEI). La CEI a pour objet de favoriser la coopération internationale pour toutes les questions de normalisation dans les domaines de l'électricité et de l'électronique. A cet effet, la CEI, entre autres activités, publie des Normes internationales.Leur élaboration est confiée à des comités d'études, aux travaux desquels tout Comité national intéressé par le sujet traité peut participer. Les organisations internationales, gouvernementales et non gouvernementales, en liaison avec la CEI, participent également aux travaux. La CEI collabore étroitement avec l'Organisation Internationale de Normalisation (ISO), selon des conditions fixées par accord entre les deux organisations.2)Les décisions ou accords officiels de la CEI concernant les questions techniques représentent, dans la mesuredu possible un accord international sur les sujets étudiés, étant donné que les Comités nationaux intéressés sont représentés dans chaque comité d’études.3)Les documents produits se présentent sous la forme de recommandations internationales. Ils sont publiéscomme normes, rapports techniques ou guides et agréés comme tels par les Comités nationaux.4)Dans le but d'encourager l'unification internationale, les Comités nationaux de la CEI s'engagent à appliquer defaçon transparente, dans toute la mesure possible, les Normes internationales de la CEI dans leurs normes nationales et régionales. Toute divergence entre la norme de la CEI et la norme nationale ou régionale correspondante doit être indiquée en termes clairs dans cette dernière.5)La CEI n’a fixé aucune procédure concernant le marquage comme indication d’approbation et sa responsabilitén’est pas engagée quand un matériel est déclaré conforme à l’une de ses normes.6) L’attention est attirée sur le fait que certains des éléments de la présente Norme internationale peuvent fairel’objet de droits de propriété intellectuelle ou de droits analogues. La CEI ne saurait être tenue pour responsable de ne pas avoir identifié de tels droits de propriété et de ne pas avoir signalé leur existence.La Norme internationale CEI 61854 a été établie par le comité d'études 11 de la CEI: Lignes aériennes.Le texte de cette norme est issu des documents suivants:FDIS Rapport de vote11/141/FDIS11/143/RVDLe rapport de vote indiqué dans le tableau ci-dessus donne toute information sur le vote ayant abouti à l'approbation de cette norme.L’annexe A fait partie intégrante de cette norme.Les annexes B, C et D sont données uniquement à titre d’information.61854 © IEC:1998– 7 –INTERNATIONAL ELECTROTECHNICAL COMMISSION––––––––––OVERHEAD LINES –REQUIREMENTS AND TESTS FOR SPACERSFOREWORD1)The IEC (International Electrotechnical Commission) is a worldwide organization for standardization comprisingall national electrotechnical committees (IEC National Committees). The object of the IEC is to promote international co-operation on all questions concerning standardization in the electrical and electronic fields. To this end and in addition to other activities, the IEC publishes International Standards. Their preparation is entrusted to technical committees; any IEC National Committee interested in the subject dealt with may participate in this preparatory work. International, governmental and non-governmental organizations liaising with the IEC also participate in this preparation. The IEC collaborates closely with the International Organization for Standardization (ISO) in accordance with conditions determined by agreement between the two organizations.2)The formal decisions or agreements of the IEC on technical matters express, as nearly as possible, aninternational consensus of opinion on the relevant subjects since each technical committee has representation from all interested National Committees.3)The documents produced have the form of recommendations for international use and are published in the formof standards, technical reports or guides and they are accepted by the National Committees in that sense.4)In order to promote international unification, IEC National Committees undertake to apply IEC InternationalStandards transparently to the maximum extent possible in their national and regional standards. Any divergence between the IEC Standard and the corresponding national or regional standard shall be clearly indicated in the latter.5)The IEC provides no marking procedure to indicate its approval and cannot be rendered responsible for anyequipment declared to be in conformity with one of its standards.6) Attention is drawn to the possibility that some of the elements of this International Standard may be the subjectof patent rights. The IEC shall not be held responsible for identifying any or all such patent rights. International Standard IEC 61854 has been prepared by IEC technical committee 11: Overhead lines.The text of this standard is based on the following documents:FDIS Report on voting11/141/FDIS11/143/RVDFull information on the voting for the approval of this standard can be found in the report on voting indicated in the above table.Annex A forms an integral part of this standard.Annexes B, C and D are for information only.– 8 –61854 © CEI:1998LIGNES AÉRIENNES –EXIGENCES ET ESSAIS APPLICABLES AUX ENTRETOISES1 Domaine d'applicationLa présente Norme internationale s'applique aux entretoises destinées aux faisceaux de conducteurs de lignes aériennes. Elle recouvre les entretoises rigides, les entretoises flexibles et les entretoises amortissantes.Elle ne s'applique pas aux espaceurs, aux écarteurs à anneaux et aux entretoises de mise à la terre.NOTE – La présente norme est applicable aux pratiques de conception de lignes et aux entretoises les plus couramment utilisées au moment de sa rédaction. Il peut exister d'autres entretoises auxquelles les essais spécifiques décrits dans la présente norme ne s'appliquent pas.Dans de nombreux cas, les procédures d'essai et les valeurs d'essai sont convenues entre l'acheteur et le fournisseur et sont énoncées dans le contrat d'approvisionnement. L'acheteur est le mieux à même d'évaluer les conditions de service prévues, qu'il convient d'utiliser comme base à la définition de la sévérité des essais.La liste des informations techniques minimales à convenir entre acheteur et fournisseur est fournie en annexe A.2 Références normativesLes documents normatifs suivants contiennent des dispositions qui, par suite de la référence qui y est faite, constituent des dispositions valables pour la présente Norme internationale. Au moment de la publication, les éditions indiquées étaient en vigueur. Tout document normatif est sujet à révision et les parties prenantes aux accords fondés sur la présente Norme internationale sont invitées à rechercher la possibilité d'appliquer les éditions les plus récentes des documents normatifs indiqués ci-après. Les membres de la CEI et de l'ISO possèdent le registre des Normes internationales en vigueur.CEI 60050(466):1990, Vocabulaire Electrotechnique International (VEI) – Chapitre 466: Lignes aériennesCEI 61284:1997, Lignes aériennes – Exigences et essais pour le matériel d'équipementCEI 60888:1987, Fils en acier zingué pour conducteurs câblésISO 34-1:1994, Caoutchouc vulcanisé ou thermoplastique – Détermination de la résistance au déchirement – Partie 1: Eprouvettes pantalon, angulaire et croissantISO 34-2:1996, Caoutchouc vulcanisé ou thermoplastique – Détermination de la résistance au déchirement – Partie 2: Petites éprouvettes (éprouvettes de Delft)ISO 37:1994, Caoutchouc vulcanisé ou thermoplastique – Détermination des caractéristiques de contrainte-déformation en traction61854 © IEC:1998– 9 –OVERHEAD LINES –REQUIREMENTS AND TESTS FOR SPACERS1 ScopeThis International Standard applies to spacers for conductor bundles of overhead lines. It covers rigid spacers, flexible spacers and spacer dampers.It does not apply to interphase spacers, hoop spacers and bonding spacers.NOTE – This standard is written to cover the line design practices and spacers most commonly used at the time of writing. There may be other spacers available for which the specific tests reported in this standard may not be applicable.In many cases, test procedures and test values are left to agreement between purchaser and supplier and are stated in the procurement contract. The purchaser is best able to evaluate the intended service conditions, which should be the basis for establishing the test severity.In annex A, the minimum technical details to be agreed between purchaser and supplier are listed.2 Normative referencesThe following normative documents contain provisions which, through reference in this text, constitute provisions of this International Standard. At the time of publication of this standard, the editions indicated were valid. All normative documents are subject to revision, and parties to agreements based on this International Standard are encouraged to investigate the possibility of applying the most recent editions of the normative documents indicated below. Members of IEC and ISO maintain registers of currently valid International Standards.IEC 60050(466):1990, International Electrotechnical vocabulary (IEV) – Chapter 466: Overhead linesIEC 61284:1997, Overhead lines – Requirements and tests for fittingsIEC 60888:1987, Zinc-coated steel wires for stranded conductorsISO 34-1:1994, Rubber, vulcanized or thermoplastic – Determination of tear strength – Part 1: Trouser, angle and crescent test piecesISO 34-2:1996, Rubber, vulcanized or thermoplastic – Determination of tear strength – Part 2: Small (Delft) test piecesISO 37:1994, Rubber, vulcanized or thermoplastic – Determination of tensile stress-strain properties– 10 –61854 © CEI:1998 ISO 188:1982, Caoutchouc vulcanisé – Essais de résistance au vieillissement accéléré ou à la chaleurISO 812:1991, Caoutchouc vulcanisé – Détermination de la fragilité à basse températureISO 815:1991, Caoutchouc vulcanisé ou thermoplastique – Détermination de la déformation rémanente après compression aux températures ambiantes, élevées ou bassesISO 868:1985, Plastiques et ébonite – Détermination de la dureté par pénétration au moyen d'un duromètre (dureté Shore)ISO 1183:1987, Plastiques – Méthodes pour déterminer la masse volumique et la densitérelative des plastiques non alvéolairesISO 1431-1:1989, Caoutchouc vulcanisé ou thermoplastique – Résistance au craquelage par l'ozone – Partie 1: Essai sous allongement statiqueISO 1461,— Revêtements de galvanisation à chaud sur produits finis ferreux – Spécifications1) ISO 1817:1985, Caoutchouc vulcanisé – Détermination de l'action des liquidesISO 2781:1988, Caoutchouc vulcanisé – Détermination de la masse volumiqueISO 2859-1:1989, Règles d'échantillonnage pour les contrôles par attributs – Partie 1: Plans d'échantillonnage pour les contrôles lot par lot, indexés d'après le niveau de qualité acceptable (NQA)ISO 2859-2:1985, Règles d'échantillonnage pour les contrôles par attributs – Partie 2: Plans d'échantillonnage pour les contrôles de lots isolés, indexés d'après la qualité limite (QL)ISO 2921:1982, Caoutchouc vulcanisé – Détermination des caractéristiques à basse température – Méthode température-retrait (essai TR)ISO 3417:1991, Caoutchouc – Détermination des caractéristiques de vulcanisation à l'aide du rhéomètre à disque oscillantISO 3951:1989, Règles et tables d'échantillonnage pour les contrôles par mesures des pourcentages de non conformesISO 4649:1985, Caoutchouc – Détermination de la résistance à l'abrasion à l'aide d'un dispositif à tambour tournantISO 4662:1986, Caoutchouc – Détermination de la résilience de rebondissement des vulcanisats––––––––––1) A publierThis is a preview - click here to buy the full publication61854 © IEC:1998– 11 –ISO 188:1982, Rubber, vulcanized – Accelerated ageing or heat-resistance testsISO 812:1991, Rubber, vulcanized – Determination of low temperature brittlenessISO 815:1991, Rubber, vulcanized or thermoplastic – Determination of compression set at ambient, elevated or low temperaturesISO 868:1985, Plastics and ebonite – Determination of indentation hardness by means of a durometer (Shore hardness)ISO 1183:1987, Plastics – Methods for determining the density and relative density of non-cellular plasticsISO 1431-1:1989, Rubber, vulcanized or thermoplastic – Resistance to ozone cracking –Part 1: static strain testISO 1461, — Hot dip galvanized coatings on fabricated ferrous products – Specifications1)ISO 1817:1985, Rubber, vulcanized – Determination of the effect of liquidsISO 2781:1988, Rubber, vulcanized – Determination of densityISO 2859-1:1989, Sampling procedures for inspection by attributes – Part 1: Sampling plans indexed by acceptable quality level (AQL) for lot-by-lot inspectionISO 2859-2:1985, Sampling procedures for inspection by attributes – Part 2: Sampling plans indexed by limiting quality level (LQ) for isolated lot inspectionISO 2921:1982, Rubber, vulcanized – Determination of low temperature characteristics –Temperature-retraction procedure (TR test)ISO 3417:1991, Rubber – Measurement of vulcanization characteristics with the oscillating disc curemeterISO 3951:1989, Sampling procedures and charts for inspection by variables for percent nonconformingISO 4649:1985, Rubber – Determination of abrasion resistance using a rotating cylindrical drum deviceISO 4662:1986, Rubber – Determination of rebound resilience of vulcanizates–––––––––1) To be published.。

个人科研情况简介胡捷，男，34岁，一直从事腐蚀电化学、混凝土耐久性和钢筋混凝土腐蚀与防护的研究工作。

目前，项目申请人主持一项国家自然科学基金青年项目、一项广东省自然科学基金项目和一项高性能土木工程材料国家重点实验室开放课题的研究工作；近几年来，作为骨干成员参与一项国家自然科学基金项目（编号：50471027/E011101）、一项美国自然科学基金项目（编号：CBET-0933246）和一项荷兰IOP（Innovation program）基金项目（编号：IOP SHM08743）的研究工作。

曾获第十三届全国青年腐蚀与防护科技论文讲评会特等奖，第七届中国腐蚀与防护协会优秀论文奖和第一届武汉市自然科学优秀论文三等奖，已发表论文20余篇，其中SCI收录论文16篇（第一作者9篇，二区4篇），EI收录论文2篇（第一作者1篇）。

近三年承担的科研项目：1.新型核壳型纳米缓蚀剂的制备与优化及其对钢筋混凝土腐蚀的自修复作用与机理，国家自然科学基金青年基金项目，25万元，2013年1月-2015年12月（项目负责人）2.胶囊型有机纳米缓蚀剂对钢筋混凝土腐蚀的自修复作用，广东省自然科学基金博士启动基金项目，3万元，2012年10月-2014年9月（项目负责人）3.基于核壳型有机纳米缓蚀剂的钢筋混凝土腐蚀自修复作用研究，高性能土木工程材料国家重点实验室开放课题，5万元，2013年7月-2015年7月（项目完成人）4.Nano-Materials with Tailored Properties for Self-healing of Corrosion Damages inReinforced Concrete，荷兰IOP（Innovation program）基金项目，30欧元，2009年8月-2011年8月（第二完成人）近三年发表的代表性论文：1. J. Hu, D.A. Koleva, Y. Ma, E. Schlangen, K. van Breugel, P. Petrov. The influence of admixed polymeric micelles on the microstructural properties and global performance of cement-based materials. Cement and Concrete Research, 42: 1122-1133, 2012. (IF=3.11, 二区)2. J. Hu*, D.A. Koleva, P. Petrov, K.van Breugel. Polymeric vesicles for corrosion control in reinforced mortar: Electrochemical behavior, steel surface analysis and bulk matrix properties. Corrosion Science, 65: 414–430, 2012. (IF=3.62, 二区)3. J. Hu*, D.A. Koleva, J.H.W. de Wit, H. Kolev, K. van Breugel, Corrosion performance of carbon steel in simulated pore solution in the presence of micelles. Journal of Electrochemical Society, 158 (3): C76-C87, 2011. (IF=2.59, 二区)4. J. Hu,Y. Ma, L. Zhang, F. Gan, Y.S. Ho. A historical review and bibliometric analysis of research on lead in drinking water field from 1991 to 2007. Science of the Total Environment, 408: 1738-1744, 2010. (IF=2.83, 二区)5. J. Hu*, D.A.Koleva, K.van Breugel. Corrosion performance of reinforced mortar in the presence of polymeric nano-aggregates: Electrochemical behavior, surface analysis and properties of the steel/cement paste interface. Journal of Materials Science, 47: 4568-4578, 2012. (IF=2.13, 三区)6. J. Hu, F. Gan, S. Triantafyllidou, C.K. Nguyen, M.A. Edwards. Copper-induced metal release from lead pipe into drinking water. Corrosion, 68 (11): 1037-1048, 2012. (IF=1.77, 三区)7. J. Hu, D.A. Koleva, Y. Ma, E. Schlangen, K. van Breugel. Early Age Hydration, Microstructure and Micromechanical Properties of Cement Paste Modified with Polymeric V esicles. Journal of Advanced Concrete Technology, 11: 291-300, 2013. (IF=0.55, 四区)8. J. Hu, H. Zhu, Y. Ma, T. Yi, X. Mao, A. Lin, F. Gan. Corrosion protection of stainless steel by separate polypyrrole electrode in acid solutions. Materials and Corrosion - Werkstoffe und Korrosion, 62 (1): 68-73, 2011. (IF=1.21, 四区)9. J. Hu*, D.A. Koleva, K. van Breugel. Effect of admixed micelles on the microstructure alterations of reinforced mortar subjected to chloride induced corrosion. Procedia Engineering, 14: 344-352, 2011. (EI)。

冷冻浇注成型多孔陶瓷的研究进展刘波涛1,2，骆兵2，许壮志2，刘贵山1(1 大连工业大学化工与材料学院，大连 1 1 6034；2 辽宁省轻工科学研究院，沈阳110036）【摘要】冷冻浇注成型是近年来发展起来的制备具有定向直孔结构多孔陶瓷的成型工艺，该工艺制备的多孔陶瓷具有高孔隙率、孔的方向和尺寸可控等特点。

本文简单介绍了多孔陶瓷材料的几种传统制备方法及其特点，重点对冷冻浇注成型工艺以及在国内外的最新研究进展进行了详细的综述，并对未来研究方向进行了展望。

【关键词】冷冻浇注成型，多孔陶瓷，孔径，定向贯通直孔道其多孔特性，制备的关键和难点是形成多孔结构。

不同的制备方法所得到制品的气孔形态、尺寸及其分布各不相同，且对其性能和功能有着重大影响。

因此，多孔陶瓷的制备方法成为人们研究的热点，其中应用比较成功、研究比较活跃的有，挤压成型法，添加造孔剂法，发泡法，溶胶凝胶发，有机泡沫浸渍法，表 1 比较了这几种工艺的特点[4-9]。

近年来，随着多孔陶瓷使用范围的不断扩大，人们对简化工艺、提高效率、降低成本的要求日益提高，传统的制备方法已不能满足人们的要求。

近年来新的制备技术不断被提出，都在努力朝着高孔隙率、结构均匀并可控制、力学性能优良的方向发展。

在这种形势下，冷冻浇注成型法由于具有孔隙率、孔的方向和尺寸可控、在垂直于孔方向上力学性能好等优点，成为了人们研究的热点[10-13]。

中图分类号：TQ174.6+2文献标示码：A1 引言多孔陶瓷是一种含有较多孔隙的新型陶瓷材料，在过滤膜、生物材料、热或声绝缘散装材料以及敏感材料等[1-3] 领域有着巨大的经济效益和社会效益。

我国从20世纪80 年代初开始研制多孔陶瓷。

多孔陶瓷首要特征是表 1多孔陶瓷传统制备工艺Table1The traditional f or min g techni ques of por ous cera mics2 冷冻浇注成型陶瓷材料的冷冻浇注成型工艺是在冷冻干燥技术基础上发展而来的。

定向生长碳化5堆焊复合耐磨板及其应用王微伟1，乔明亮2,赵春燕'，刘爱国2（1.江苏瑞米克金属技术有限公司，江苏常州213172；2.沈阳理工大学，辽宁沈阳110159）摘要:采用药芯焊丝明弧自保焊技术在Q235C低碳钢基体上进行了双层耐磨堆焊,采用强制冷却技术对M'C3碳化物的生长方向进行了控制,制备出了定向生长碳化辂堆焊复合耐磨板#采用光学显微镜、硬度仪、针盘磨损试验机对耐磨板堆焊层的的显微组织、硬度和耐磨性能进行了测试分析。

结果表明，第一堆焊层为亚共晶成分，显微组织为马氏体（+残余奥氏体）+共晶组织;第二堆焊层为过共晶成分，显微组织为初生碳化物M7C3+共晶组织;从熔合区到堆焊表面硬度逐渐升高，堆焊表面硬度达到60.2HRC#通过强制冷却技术控制,第二堆焊层中M7C3碳化物的生长方向垂直于堆焊表面，有利于提高堆焊层的耐磨性，针盘磨损试验中磨损失重仅为作为对比试样的Q235C基体的1/30#在磨煤机筒体防磨板上的实际使用结果显示，其使用寿命达到某市售堆焊耐磨板的3倍#关键词：耐磨板；堆焊；碳化辂;定向生长中图分类号：TG455Hardfacel wear resistant composite plate witt orientate'ctromium carbide and is applicationWang Weiwl1,Qiao Mingliang2,Zhao Chunyan1,Lin Aiguo2$1.Jiangsu RemakeMeeaeTethn/egyC/.，Led.，Changzh/u2131'2，Jiangsu，China;2.Shenyang Ligong University,Shenyang110159,Liaoning,China）Abstrace:Low corbon steel Q235C plates were hardfaced wita cored wires,as welO as self-shielded open arc.Growth oeentation of M7C3corbides was controlled with forced cooling.Doubie-layered overlays w W orientat­ed theomium taebideweeepeoduted on eheBubBeeaee.MiteoBeeuteueeB,haedneBand wea e eeBi ean te of ehe o A verlay were analyzed with opticoi microscope,hadness testee and pin-on-dWk weae testee.The results showed that composition of the first overlay was hypoeutectic and its microstructure consisted of martensite with residuai austenite and e utectic position of the second overlay was hypereutectic, and its microstructures consisted of primare corbides M7C3and eutectic microstructures.Hadness of tae ovee-lays incressed from the fsion boundare te tae surfaco of the second overlay.Hadness of the surfaco of tae o­verlay was60.2HRC.Controlled with forced cooling technology,growta oeentation of M7C3corbides in the second overlay was perpendiculae te the surfaco.Wee resistanco of tte overlay was incressed with oeentated-growing M7C3corbines,and wer loss in tae pin-on-disk test was only1/30of that of tte Q235C substrate.Applicotion in cool mili btrel showed0x0its life was3times of a cortain commercialiy aveilablehardfaced composite plate.Key words：w—resistant plate;hardfacing；chromium corbide;olentated growth62020年第5期0前言工业系统设备耐磨防护中常用的材料包括铸石、氧化铝陶瓷、超高分子量聚乙烯、合金钢耐磨板等，但这些材料存在脆性大、难以进行切割及焊接加工等一系列问题［1］#采用堆焊或熔敷技术制造的带有碳化物增强相的复合耐磨板具有耐磨性好、抗冲击、进行切割及焊接加工等优点，在耐磨防护中获得了广泛的应用［1"3］#在众多的碳化物中，碳化锯由于具有耐磨性能优异、价格低廉、可以在Fe-Cr-C合金体系中原位生成的特点，成为铁基耐磨材料堆焊、熔敷的首选增强材料。

Microstructure and corrosion resistance of the layers formed on the surface of precipitation hardenable plastic mold steel by plasma-nitridingDong Cherng Wen*Department of Mechanical Engineering,China University of Science and Technology,No.245,Section3,Academia Road,Nangang District,Taipei City11581,Taiwan,ROC1.IntroductionIn the injection molding of polymers,the mold is a critical element.Continuous growing activity of the engineering plastics led to necessity of developing low cost and high-performance mold steels.Precipitation hardenable tool steels are now used for this application with good results giving better quality,low costs,and increased performance;for instance,NAK80mold steel,a prehardened steel used for making plastic molds,developed by Daido Steel,Japan.This steel is usually supplied in the solution treated condition,with hardness ranging from30to32HRC.After aging at475–5258C for5h,the hardness reaches39–41HRC[1]. With its excellent machinability,this steel is intended to replace AISI P-20modified steel.Although reasonably high mechanical properties can be attained in this steel,it is possible to extend the lifetime of plastic injection molds by means of surface treatments such as nitriding.Nitriding is a surface treatment technique used to introduce nitrogen into metallic materials to improve their surface hardness,mechanical properties,as well as wear and corrosion resistance[2–4].Conventional gas-and liquid-nitriding processes are not suitable for precipitation hardenable steels because the high temperature used in these processes(more than5508C),which is above the aging temperature of precipitation hardenable steels, would result in the overaging of the core.However,plasma-nitriding can be carried out at a temperature lower than the aging temperature.In the plasma-nitriding process,by means of a glow-discharge using a mixture of N2and H2,with the steel at a temperature of the order of magnitude5008C,nitrogen can penetrate the surface and diffuse into the steel[5].Under such conditions,in the structure of the nitrided layers,the top nitrided layer,known as the compound layer,is composed mainly of e-Fe2À3N and g0-Fe4N as well as nitrides with alloying elements.The layer beneath the compound layer is known as the so-called diffusion layer,which consists mainly of interstitial atoms in solid solution andfine,coherent nitride precipitates when the solubility limit is reached.The thickness and compositions of these layers are functions of treatment temperature,and time and composition of the base material[6–9].The diffusion layer determines the strength of the nitrided layer,as well as its fatigue strength[10,11],while the compound layer determines the tribological characteristics andApplied Surface Science256(2009)797–804A R T I C L E I N F OArticle history:Received11May2009Received in revised form12August2009 Accepted17August2009Available online21August2009PACS:81.65.Lp82.45.Bb61.05.cpKeywords:Plasma nitridingPlastic mold steelNitridesCorrosion A B S T R A C TPlasma-nitriding is used to improve the wear resistance and corrosion resistance of plastic mold steels by modifying the surface layers of these steels.In this study,a precipitation hardenable plastic mold steel (NAK80)was plasma-nitrided at470,500,and5308C for4,8,and12h under25%N2+75%H2 atmosphere in an industrial nitriding facility.The microstructures of the base material and nitrided layers as well as the core hardness were examined,and various phases present were determined by X-ray diffraction.The corrosion behaviors were evaluated using anodic polarization tests and salt fog spray tests in3.5%NaCl solution.The results had shown that plasma-nitriding does not cause the core to soften by overaging.Nitriding and aging could be achieved simultaneously in the same treatment cycle.Plasma-nitriding of NAK80 mold steel produced a nitrided layer composed of an outer compound layer constituting a mixture of e-nitride and g0-nitride and an adjacent nitrogen diffusion layer on the steel surface.The amount of e-nitride and total nitrides increased with an increase in nitriding temperature and nitriding time. Corrosion study revealed that plasma-nitriding significantly improved the corrosion resistance in terms of corrosion potential,corrosion and pitting current density,and corrosion rate.This improvement was found to be directly related to the increase in the amount of e-nitride at the surface,indicating the amount of e-nitride controlling the corrosion resistance.ß2009Elsevier B.V.All rights reserved.*Tel.:+886227867048;fax:+886227867253. E-mail address:dcwen@.tw.Contents lists available at ScienceDirect Applied Surface Sciencej o u r n a l h o m e p a g e:w w w.e l s e v i e r.c o m/l o c a t e/a p s u s c0169-4332/$–see front matterß2009Elsevier B.V.All rights reserved.doi:10.1016/j.apsusc.2009.08.062corrosion resistance[12–14].The e-nitride shows better corrosion resistance when compared with g0-nitride because of its crystalline structure and higher nitrogen content[15].Another investigation reported that g0-nitride formed after iron nitriding in the surface layer is relatively stable in corrosive medium[16].The corrosion resistance of the nitrided layer,thus,depends on the type of nitride formed in the compound layer.However,for a deeper under-standing of precipitation hardenable tool steels the effects of plasma-nitriding on their microstructure,and corrosion behavior are essential,but very little work has been done on these aspects.In the present work,NAK80mold steel was plasma-nitrided at three different temperatures,that is,470,500,and5308C for4,8 and12h.The metallurgical structures and corrosion behavior of the surface nitrided layer are evaluated,analyzed,and discussed.2.Experimental detailsThe material used for this study was NAK80mold steel,which was obtained commercially in the form of a plate of15mm thickness.The specified chemical composition and the actual chemical analysis of the material obtained using a Leco GDS-750 glow-discharge optical emission spectroscope(GDOS)are shown in Table1.As can be seen,the chemical composition of the material falls within the composition range specified for NAK80steel.The material was received in the solution treated condition and had been aged at5008C for5h and had a hardness of40.7HRC.This received material was designed as unnitrided for the purpose of comparison.Test specimens of12.7mmÂ12.7mmÂ12.7mm were cut from the NAK80steel plate.These specimens were initially solubilized at9008C and then aged at470,500,and5308C for periods ranging from1to20h to determine the aging curves and guide the nitriding treatments.Before nitriding,the solubilized specimen surfaces were ground using#1200silicon carbide(SiC)paper and were then ultra-sonically cleaned.Plasma-nitriding was carried out in an industrial nitriding facility(DC-pulsed).The specimens were sputter cleaned in an atmosphere of80%Ar+20%H2at about2508C for1h,to remove the oxide layers formed on their surfaces.They were then plasma-nitrided in an atmosphere of25%N2+75%H2at temperatures of470,500,and5308C for4,8,and12h,with the chamber pressure maintained at600Pa.The plasma-nitriding temperature was controlled to an accuracy ofÆ58C,using a photoelectric thermometer.After the nitriding process was com-pleted the specimens were allowed to cool in the chamber to room temperature in a stream offlowing nitrogen to protect the heated surfaces from oxidation.The parameters used in experiment are summarized in Table2.The microstructures of the base material and the nitrided layers were examined using an Olympus BHM optical microscope(OM) and a JEOL JSM-5600scanning electron microscope(SEM);X-ray diffraction(XRD)investigations of the compound layer were carried out using a Shimadzu X-ray diffractometer(LabX XRD-6000)with Cu K a radiation.Quantitative analysis was carried out using the normalized relative(integrated)intensity ratio(RIR) method or the matrixflushing method[17].The core Rockwell‘‘C’’hardness was measured for solubilized,aged,and nitrided specimens.An average hardness value of each specimen was obtained based on six measurements.A potentiostat(EG&G362)outfitted with analysis software (CorrWare)was used to evaluate electrochemical behavior.Theflat cell is a three-electrode setup consisting of the specimen as the working electrode,a saturated calomel electrode(SCE)as the reference electrode,and a platinum sheet used as the counter electrode.The electrolyte used was3.5%NaCl solution.Potentio-dynamic polarization was swept at afixed rate of1.0mV/s.The salt fog spray test was performed in accordance with ASTM B117.The solution used for salt spray was the same as in the electrochemical test.The test temperature was maintained at35Æ28C,and the humidity was maintained at94Æ4%.After the electrochemical andTable1Chemical composition of NAK80mold steel(wt.%).Element Specified ObtainedC0.05–0.180.12Si0.15–1.00.26Mn 1.0–2.0 1.43Ni 2.5–3.5 3.18Al0.5–1.5 1.02Cu0.7–1.50.95Mo0.1–0.40.22Fe Balance Balance Table2Parameters for plasma-nitriding treatments.Parameter Sputter cleaning Plasma nitridingTemperature(˚C)250470;500;530 Voltage(V)400600Current density(mA cmÀ2)0.8 1.2Total pressure(Pa)200600Time(h)14;8;12 Treatment gas(vol)80%Ar+20%H225%N2+75%H2Fig.1.Microstructure of NAK80mold steel in the solubilized and aged state.Nital4%.Fig.2.Aging curves of the specimens aged at various temperatures.D.C.Wen/Applied Surface Science256(2009)797–804 798salt fog spray tests,the surface morphologies were inspected using SEM to reveal the severity of corrosion degradation in terms of pit morphology,distribution,and density.3.Results and discussion3.1.MicrostructureFig.1shows the microstructure of the steel in the solution treated condition and had been aged at5008C for5h.A bainite plus martensite microstructure was visible.The average hardness in the solubilized condition was31.5HRC.The aging curves of the specimens aged at various temperatures are shown in Fig.2.A high hardness level was reached in4–12h of aging,and the average values of this hardness were found to be40,40.5,and41.5HRC for aging at470,500,and5308C,respectively.In addition,no signs of overaging were observed before14h of treatment.Fig.3shows the XRD patterns obtained from the surface of unnitrided and plasma-nitrided specimens.The unnitrided speci-men showed only a-Fe,whereas the plasma-nitrided specimens exhibited additional peaks due to different nitrides.In the specimen treated at4708C(Fig.3(b)),some e-nitride and g0-nitride peaks accompanied the a-Fe peaks.In plasma-nitriding at 5008C(Fig.3(c)),the nitrided surface layer was dominantly e-nitride along with g0-nitride,and the amount of a-Fe phase was much lesser than that in the plasma-nitriding at4708C.TheFig.3.XRD diffraction patterns of(a)unnitrided specimen,(b)nitrided at4708C for12h,(c)nitrided at5008C for12h,(d)nitrided at5308C for12h,(e)nitrided at5308C for 8h,and(f)nitrided at5308C for4h.D.C.Wen/Applied Surface Science256(2009)797–804799amount of e -nitride and g 0-nitride in 5308C nitrided surface layer was further increased,and the amount of a -Fe phase was decreased apparently (Fig.3(d)).Therefore,the peaks of a -Fe phase almost disappeared.According to these patterns,it was obvious that the intensity of the a -Fe peaks progressivelydecreased,and the intensity of the nitride peaks progressively increased with an increase in nitriding temperature.Thus,the iron nitride formed in the surface hardened layer on plasma-nitriding was compound layer,in agreement with other previous observa-tions [18–20].Fig.3(d)–(f)shows that the relative intensity of Fe xNFig.4.(a)Variation of amount (volume percent)of nitrides with nitriding time at different nitriding temperatures.(b)The amount of total nitrides (outlined symbols)and the proportion of epsilon to gamma nitrides (solid symbols)with nitriding time at different nitridingtemperatures.Fig.5.Cross-sectional microstructures of the plasma-nitrided specimens:(a)nitrided at 5308C for 4h (OM)and (b)nitrided at 5308C for 12h(SEM).Fig.6.Polarization curves of the unnitrided and plasma-nitrided specimens in 3.5%NaClsolution.Fig.7.Variations of corrosion potential and corrosion rate with nitriding time at different nitriding temperatures.(The corrosion rate of the unnitrided specimen was 17Â10À3mg/h).D.C.Wen /Applied Surface Science 256(2009)797–804800peaks increased with an increase in nitriding time at a given temperature,which indicated an increase in the amount of the nitride phases at the expense of parent a -Fe matrix.Quantitative analysis obtained from XRD patterns revealed that the amount of e -and g 0-nitride increased with increasing nitriding temperature (Fig.4(a)).The amount of e -nitride was also found to increase with an increase in nitriding time,but the amount of g 0-nitride did not change significantly with nitriding time.However,the amount of e -nitride was much larger than that of g 0-nitride for all nitrided specimens.Fig.4(b)shows the variations of e -to g 0-nitride ratio and the amount of total nitrides as a function of nitriding time at different nitriding temperatures.The proportion of e -to g 0-nitride phases increased with increasing nitriding time at a given nitriding temperature,but the variation in this proportion was not directly related to the nitriding temperature.The amount of total nitride increased with increasing nitriding time and nitriding temperature.This increase was primarily ascribed to the increase of e -nitride.A maximum amount of total nitrides was formed when nitriding was carried out at 5308C for 12h (95.1%).A typical microstructure of plasma-nitrided specimen is shown in Fig.5(a).It consisted of an internal nucleus of bainite plus martensite and a nitrided layer on the external surface.The nitrided surface consisted of two layers:a compound layer at the top,being very thin,and a nitrogen diffusion layer beneath.Fig.5(b)shows the scanning electron microscopy cross-section images of the nitrided layer of plasma-nitrided specimen treated at 5308C for 12h.The formation of thick and dense compound layer was consistent with the fact that the peaks of a -Fe phase almost disappeared,as shown in Fig.3(d).The thickness of the compound layer and diffusion layer of all plasma-nitrided specimens is shown in Table 3.It is observed that both the thickness of the compound layer and the depth of the diffusion layer increased with an increase in nitriding temperature and nitriding time.The core hardness of all plasma-nitrided specimens is also presented in Table 3.It should be noted that the core hardness at all treatment conditions was in the same level as the hardness of the unnitrided specimen.This confirmed that the core does not soften by overaging during plasma-nitriding,that is,microstructure of the core remains unaltered afterplasma-nitriding.Fig.8.Post corrosion SEM morphologies of specimens after electrochemical behavior tests in 3.5%NaCl solution:(a)unnitrided,(b)nitrided at 4708C for 12h,(c)nitrided at 5008C for 12h,(d)nitrided at 5308C for 12h,(e)unnitrided specimens showing the general corrosion,and (f)specimen nitrided at 5308C for 12h showing the pitting corrosion.D.C.Wen /Applied Surface Science 256(2009)797–8048013.2.Corrosion behavior3.2.1.Electrochemical behavior testsFig.6shows the anodic polarization curves for unnitrided and plasma-nitrided specimens treated at different temperatures for 12h.It was seen that the unnitrided specimen did not have an evident passivation region and suffered progressive dissolution, and the corrosion potential(E corr)wasÀ0.50V(SCE).In contrast, all nitriding processes shifted the corrosion potential in the noble direction and promoted passivation resulting in very low corrosion current density(I corr)and very high corrosion potential.Between the corrosion potential and1.25V(SCE)there appeared an evident passivation region.It was observed that with an increase in nitriding temperature,the corrosion potential increased,whereas the corrosion current density decreased,which means nitriding could improve the corrosion resistance.Similar trend was also observed when nitriding was carried out for4and8h.The enhanced corrosion resistance was attributed to the presence of a dense nitride layer rich in e-nitride on the surface.Because nitride is a noble phase,formation of more nitride helps to protect the surface from corrosion attack,and for this reason,with an increase in nitriding time or temperature,corrosion resistance improves. This result agrees quite well with the plasma-nitriding study mentioned earlier[19],which showed that the nitrided layer with the e-phase has excellent corrosion resistance.A corrosion current of1mA is equal to a weight loss of 1.04mg/h according to Faraday’s law,assuming a density of7.8g/cm3and atom weight of iron of55.85for NAK80mold steel.Fig.7shows the variation of corrosion potential and corrosion rate with nitriding time at different nitriding temperatures.The corrosion potential increased with an increase in nitriding time and nitriding temperature.Corrosion rate followed the opposite trend and ultimately showed drastic reduction in corrosion rate after plasma-nitriding.Increase in corrosion potential with an increase in nitriding time could be possibly attributed to an increase in the amount of nitrides formed at a higher nitriding temperature and time.A thicker and dense nitrided layer would always enhance the corrosion resistance[21].The trend of corrosion potential in Fig.7was similar to the trend of the amount of total nitrides observed in Fig.4(b)further confirming the increase in corrosion resistance by coverage of nitride phase.Furthermore,improvement in corrosion resistance was found to be directly related to the increase in the amount of e-nitride at the surface(see Fig.4(a)); it had been suggested that the most important parameter for controlling the corrosion resistance of the nitrided specimens is the amount of e-nitride.The comparison of morphologies of corroded specimens after electrochemical behavior tests is shown in Fig.8.The unnitrided specimen corroded most severely,and an obvious general and uniform corrosion could be seen over the entire exposed surface,as shown in Fig.8(a).A high corrosion rate thus resulted under the anodic potential.In contrast,the surface of nitrided specimens was more resistant to general corrosion in comparison with unnitrided ones,but was subjected to pitting corrosion as shown in Fig.8(b)–(d).Therefore,the nitrided specimens showed lower corrosion rate than the unnitrided ones.The4708C nitrided specimen showed a higher pitting current density(4.21Â10À5A/cm2).After the test the corroded surface showed the highest pits(652/cm2)as shown in Fig.8(b).The high density of pits formed during the corrosion tests could be due to a very large number of sites in which local breakdown of the passivity can easily occur.The observed corrosion behavior could be ascribed to the fact that the amount of e-nitride(53.6%)was lower in4708C nitrided surface layer in comparison with the other nitrided specimens.Specimen nitrided at5008C showed a medium pitting current density(3.07Â10À5A/cm2).After linear polarization the corroded surface showed the presence of fewer and smaller pits(500/cm2) as shown in Fig.8(c).This behavior could be attributed to the moderate amount of e-nitride(57.0%)formed on the nitrided layer.In the5308C nitrided specimen,the pitting current density showed a significant reduction(6.13Â10À6A/cm2)in comparison with the other nitrided specimens.After the linear polarization,the specimen showed the lowest pits(239/cm2)on the corroded surface as shown in Fig.8(d).This indicated that very small number of sites in which the local breakdown of passivity occurred during the corrosion tests.This behavior was due to a dense protective nitrided layer with high amount of e-nitride(62.6%)formed on the surface.Fig.8(e)and(f)exhibits the corrosion mechanism of the unnitrided and plasma-nitrided specimens.The unnitrided speci-men showed large area damage caused by general corrosion.A large number of obvious grain boundaries and some intergranular corrosion were observed.In contrast,the nitrided specimen showed different surface morphology as illustrated in Fig.8(f). Due to the protection of nitrided layer,the surface had undergone typical pitting corrosion failure characterized by deep cavities, while only few grain boundaries could be observed within the cavities.Fig.9shows that the pits occurred during electrochemical testing increase almost linearly with the pitting current density, while the amount of e-nitride was found to increase with decreasing pitting current density,as expected for the amount of e-nitride controlling the corrosion resistance.3.2.2.Salt fog spray testsDuring the salt fog spray test,the specimens were regularly checked by visual means for any corrosion initiation.The results are also shown in Table3.As can be seen from Table3,the duration of corrosion initiation increased with an increase in nitriding temperature and nitriding time,an excellent improvement in corrosion resistance was observed for the nitrided specimens. Although the time for corrosion initiation of the unnitrided specimen is10min,it lasts much longer,more than10times for the nitrided specimens.Fig.10(a)–(d)shows the surface morphol-ogies of the unnitrided and plasma-nitrided specimens after salt spray tests for2.5h and followed by removing corrosion products. The unnitrided specimen corroded most severely,and a lotof Fig.9.Variations of pit density and amount of e-nitride with pitting current density for the specimens nitrided at different nitriding temperatures for12h.D.C.Wen/Applied Surface Science256(2009)797–804 802corrosion pits were seen on the surface.The nitrided specimens also suffered pitting corrosion.However,the corrosion pits on the nitrided specimens were much less than those on the unnitrided specimen.Furthermore,the corrosion attacks decreased with anincrease in nitriding temperature and nitriding time.The results of salt fog spray tests were in accord with the results of the electrochemical behavior tests.Fig.10(e)and (f)shows the corroded surface morphologies of the unnitrided and plasma-nitrided specimens after salt spray tests for 2.5h before removing corrosion products,respectively.The red rust and severe pitting corrosion could be seen on the unnitrided specimen.However,the red rust and pitting corrosion behavior of the nitrided specimen was improved significantly,as seen in Fig.10(f).The results indicated that the formation of nitrided layer with rich in e -nitride after plasma-nitriding is effective to enhance the corrosion resistance to localized corrosion.4.ConclusionsIn the present investigation,NAK80mold steel was plasma-nitrided at 470,500,and 5308C for 4,8,and 12h.The main conclusions from this investigation can be summarized asfollows:Fig.10.Corroded surface morphologies of the specimens after 2.5h salt spray tested:(a)unnitrided,(b)nitrided at 4708C for 12h,(c)nitrided at 5008C for 12h,(d)nitrided at 5308C for 12h,(e)unnitrided specimens before removing the corrosion products,and (f)specimen nitrided at 5308C for 12h before removing the corrosion products.Table 3Compound layer and diffusion layer thickness,core hardness,and corrosion initiation duration of unnitrided and plasma-nitrided specimens.Nitriding parameterCompound layerthickness (m m)Diffusion layerthickness (m m)Corehardness (HRC)Corrosion initiation duration (min)Unnitrided ––40.710At 4708C for 4h 1–225–4040.1106At 4708C for 8h 2–455–7539.8109At 4708C for 12h 3–690–11039.4113At 5008C for 4h 2–455–7040.5117At 5008C for 8h 5–795–12041.1121At 5008C for 12h 7–10155–18040.8132At 5308C for 4h 4–680–10541.3120At 5308C for 8h 6–9135–17041.7128At 5308C for12h8–12210–24041.2142D.C.Wen /Applied Surface Science 256(2009)797–8048031.Microstructure of core remains unaltered after plasma-nitrid-ing.The average core hardness of solubilized specimens increases from31.5to40.8HRC after the plasma-nitriding treatment,implying that the nitriding and aging can be achieved simultaneously in the same treatment cycle.2.The nitrided case was composed of an outer most compound layer constituting a mixture of e-nitride and g0-nitride,followed by the diffusion layer.The plasma parameters determined the thickness and composition of compound layer.The amounts of e-nitride and total nitrides increased with an increase in nitriding temperature and nitriding time.A maximum amount of total nitrides was formed when nitriding was carried out at 5308C for12h(95.1%).3.Corrosion study in3.5%NaCl solution revealed that plasma-nitriding significantly improved the corrosion resistance in terms of corrosion potential,corrosion and pitting current density,and corrosion rate.The improved corrosion resistance was found to be directly related to the increase in the amount of e-nitride at the surface,indicating the amount of e-nitride controlling the corrosion resistance.4.The corroded surface of unnitrided specimen showed an obviousgeneral and uniform corrosion over the entire exposed surface, while the corroded surface of nitrided specimens was replaced by pitting corrosion.The pit number and pitting current density progressively decreased with an increase in nitriding tempera-ture and nitriding time.References[1]http://www.daido.co.jp/english/products/tool/plasticmold.html.[2]Q.F.Peng,Improving abrasive wear by surface 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Corrosion and stress corrosion cracking in supercritical waterG.S.Was a,*,P.Ampornrat a ,G.Gupta a ,S.Teysseyre a ,E.A.West a ,T.R.Allen b ,K.Sridharan b ,L.Tan b ,Y.Chen b ,X.Ren b ,C.Pister baUniversity of Michigan,Nuclear Engineering,1911Cooley,Ann Arbor,MI 48109,United StatesbUniversity of Wisconsin,United StatesAbstractSupercritical water (SCW)has attracted increasing attention since SCW boiler power plants were implemented to increase the efficiency of fossil-based power plants.The SCW reactor (SCWR)design has been selected as one of the Generation IV reactor concepts because of its higher thermal efficiency and plant simplification as compared to current light water reactors (LWRs).Reactor operating conditions call for a core coolant temperature between 280°C and 620°C at a pressure of 25MPa and maximum expected neutron damage levels to any replaceable or permanent core com-ponent of 15dpa (thermal reactor design)and 100dpa (fast reactor design).Irradiation-induced changes in microstructure (swelling,radiation-induced segregation (RIS),hardening,phase stability)and mechanical properties (strength,thermal and irradiation-induced creep,fatigue)are also major concerns.Throughout the core,corrosion,stress corrosion cracking,and the effect of irradiation on these degradation modes are critical issues.This paper reviews the current understanding of the response of candidate materials for SCWR systems,focusing on the corrosion and stress corrosion cracking response,and highlights the design trade-offs associated with certain alloy systems.Ferritic–martensitic steels generally have the best resistance to stress corrosion cracking,but suffer from the worst oxidation.Austenitic stainless steels and Ni-base alloys have better oxidation resistance but are more susceptible to stress corrosion cracking.The promise of grain boundary engineering and surface modification in addressing corrosion and stress corrosion cracking performance is discussed.Ó2007Elsevier B.V.All rights reserved.1.IntroductionOne of the most promising advanced reactor concepts for Generation IV nuclear reactors is the Supercritical Water Reactor (SCWR).Operating above the thermodynamic critical point of water (374°C,22.1MPa),the SCWR offers many advan-tages compared to state-of-the-art LWRs including the use of a single phase coolant with high enthalpy,the elimination of components such as steam gener-ators and steam separators and dryers,a low coolant mass inventory resulting in smaller compo-nents,and a much higher efficiency ($45%vs.33%in current LWRs).Overall,the design provides a simplified,reduced volume system with high thermal efficiency.The challenge is provided by the substantial increase in operating temperature and pressure as compared to current BWR and PWR designs.The reference design for the SCWR [1,2]calls for an operating pressure of 25MPa and an outlet water temperature up to 620°C,Fig.1.Since supercritical water has never been used in nuclear power applications,there are numerous0022-3115/$-see front matter Ó2007Elsevier B.V.All rights reserved.doi:10.1016/j.jnucmat.2007.05.017*Corresponding author.Fax:+17347634540.E-mail address:gsw@ (G.S.Was).Journal of Nuclear Materials 371(2007)176–201potential problems,particularly with materials.Water in the supercritical phase exhibits properties significantly different from those of liquid water below the critical point.It acts like a dense gas and its density can vary with temperature and pres-sure from less than 0.1g/cc to values similar to that of water below the critical point.This allows one to ‘tune’the properties of SCW,such as the ion prod-uct,heat capacity,and dielectric constant,to fit the application of interest [1].The corrosive behavior of SCW over this range of densities varies widely depending upon the values of these properties [3,4].Supercritical water will,at the high reactor outlet temperatures,fall in the lower end of the density scale at around 0.2g/cc while having higher density at the reactor inlet.At the low density associated with the reactor outlet,water is a non-polar solvent and can dissolve gases like oxygen to complete mis-cibility.Depending upon what species are present and how much oxygen is present in the solution,SCW in this state can become a very aggressive oxi-dizing environment [3,5].On the other hand,the sol-ubility of ionic species is expected to be extremely low.This is a cause for concern for the general cor-rosion and stress corrosion cracking (SCC)suscepti-bility of the structural materials and fuel elements of reactors.While no experience exists with supercritical water reactors there is a significant operating history with supercritical fossil plants [6,7].In fact,as of 2004,there were some 268944MWe (462units)of installed capacity of coal-fired supercritical water power plants worldwide [7].Hence there is signifi-cant industry experience with supercritical water in power generation.However,a nuclear reactor coreis significantly different from a fossil-fired boiler.One key difference is geometry.A fossil-fired boiler consists of a large number of fire tubes that circulate water on the inside.These tubes have relatively thick walls,approximately 6–12mm in thickness.Further,the water sees a geometrically smooth and simple surface along its path through the boiler.Contrast this with the core of an SCWR,which con-sists of fuel rods,control rods,and water rods com-prising a fuel assembly with some 145assemblies forming the core.The wall thickness for the fuel rod cladding in the reference SCWR design is 0.63mm and the wall thickness for the water rods is 0.40mm.These very thin components do not pro-vide much margin for corrosion in the cores of supercritical water reactors.Oxide films of several hundred micrometer thickness are not unusual for boiler tubes with typical wall thicknesses in the 6–12mm range,but are unacceptable for water rods or fuel cladding.Fig.2illustrates the scale differ-ences between core components in an SCWR and the corresponding components in a fossil-fired SCW power plant.Reactor core components must also contend with irradiation that can affect both the water chemistry and the alloy microstructure.Radiolysis can result in an increase in the concentration of oxygen and other oxidizing species such as H 2O 2that raise the corrosion potential and increase susceptibility to processes such as stress corrosion cracking.Radiol-ysis is not at all understood in supercritical water and pioneering experiments are just now under-way.The very high-temperatures and significantly different properties of SCW compared to subcritical water make it difficult to estimate the effect of irradiation on this fluid.Initial measurements have shown that virtually no free radical reaction rates follow an Arrhenius law,and so the rate constants must be measured [8].Bartels et al.[8]have made measurements on the radiolysis yield of H 2as a function of temperature and have found that it increases steeply with temperature above about 350°C,Fig.3a.They have also found that rate constants for the reactions H 2+OH (Fig.3b)and OH +OH decrease at higher temperature indicating a slow down of the reaction rate at high-tempera-ture.Since hydrogen addition is a key strategy in controlling corrosion at high-temperature in current reactor designs,an understanding of the dependence of this reaction (Fig.3b)on temperature and radia-tion may be critical to controlling the corrosion potential in the core of an SCWR.A measureofFig.1.Pressure–temperature regime of SCWR operation com-pared to that for current BWRs and PWRs.G.S.Was et al./Journal of Nuclear Materials 371(2007)176–201177the corrosion potential in supercritical water will require development of a reference electrode that can withstand the SCW environment.Once the irra-diated water chemistry is understood so that the corrosion potential can be estimated or measured,then laboratory experiments can be carried out to more accurately simulate the environmental condi-tions expected for the core of a SCWR.Perhaps the most challenging problem is the role of irradiation on the microstructure and how these changes affect stress corrosion cracking.Irradiation assisted stress corrosion cracking has been a generic problem in light water reactors of all types and covering many austenitic and nickel-base alloys [9,10].This paper reviews the current understanding of the response of candidate materials for SCWR sys-tems,focusing on the corrosion and stress corrosion cracking response,and highlights the design trade-offs associated with certain alloy systems.WiththeFig.2.(a)A 1/8assembly model of the 21·21SCWR fuel assembly [2],and (b)comparison of typical fossil boiler tube dimensions to fuel rod and water rod dimensions in the reference SCWR design.178G.S.Was et al./Journal of Nuclear Materials 371(2007)176–201exception of the effect of irradiation on stress corro-sion cracking,the important issues of radiation response and radiolysis are not addressed in this review.2.Experimental programsIn support of the supercritical water reactor pro-gram,corrosion and stress corrosion cracking have been studied in pure supercritical water in ferritic-martenitic steels,austenitic stainless steels,Ni-base alloys,Zr-base alloys,and Ti-base alloys.Test temperatures have ranged from290to732°C. Dissolved oxygen concentration has ranged from <10ppb to8000ppb.Exposure times for corrosion tests have ranged from100to over1000h.In stress corrosion cracking studies,the effect of chemical additions have been examined,specifically H2SO4, HCl,H2O2,NaCl.Additionally,the affect of system pressure on SCC resistance has been studied.Some specific alloys have multiple designations. For example,HCM12A and T122are the same composition,as are NF616and T92.In thefigures, the alloys designation used is consistent with the reference from which the data was originally pub-lished.Tables1and3can be used to determine alloys with multiple designations.3.CorrosionThis section presents and summarizes the corro-sion in pure supercritical water.Specific details the corrosion test procedures are described in Ref.[11].The data and descriptions presented are dis-tinct from that on supercritical water oxidation (SCWO)experiments that contain halides for the purpose of breaking down organic materials.This latter database will not be covered in the review thatTable1Summary of experiments on corrosion in pure supercritical waterAlloy Class Alloy Temp.(°C)Water chemistry Exposuretime(h)Austenitic SS304,304L,316,316L,316+Zr,310,310S,310+Zr,TP347H,Sanicro28,D9,800H 290–650Deaerated(<10ppb)to8000ppb dissolved oxygen,100–1026Nickel-base600,625,690,718,825,C22,B2,C276,MAT21,MC 290–600Deaerated(<10ppb)to8000ppb dissolved oxygen,<0.1mS/cm100–575Ferritic–martensitic T91,T91a,T91b,HCM12A(T122),HCM12,HT-9(12Cr–1Mo–1WVNb),NF616(T92),MA956,2.25Cr–1Mo(T11),P2290–650Deaerated(<10ppb)to8000ppb dissolved oxygen,<0.1mS/cm100–1026ODS9Cr,12Cr,F–M,316,Inconel,Hastelloy G-30,19Cr,14Cr–4Al,16Cr–4Al,19Cr–4Al,22Cr–4Al360–60025ppb200–1026Zirconium Zr,Zr–Nb,Zr–Fe–Cr,Zr–Cr–Fe,Zr–Cu–Mo,Zr-2,Zr-4400–500Deaerated(<10ppbdissolved oxygen),<0.1mS/cm<2880Titanium Ti–3Al–2.5V,Ti–6Al–4V,Ti–15Mo–5Zr–3Al,Ti–15V–3Al–3Sn–3Cr 290–5508000ppb dissolved oxygen,0.1mS/cm500G.S.Was et al./Journal of Nuclear Materials371(2007)176–201179follows.The alloy systems,specific alloys and range of test conditions is summarized in Table1.3.1.Ferritic–martensitic steelsFerritic and martensitic(FM)steels were selected for possible use in supercritical water reactor systems because of their radiation resistance,high thermal conductivity,and low thermal expansion coefficients.To date,international programs have evaluated the following ferritic–martensitic steels: T22,P2,T91,HT9,HCM12,HCM12A(T122), NF616(T92),and numerous oxide dispersion strengthened steels including JAEA9Cr ODS,MA 956and(14–22)Cr–4Al versions[11–19].Surface modification,specifically the implantation of oxy-gen and yttrium,to reduce oxidation rates,has been performed on NF616and HCM12A.3.1.1.Oxide structureFor samples exposed to low dissolved oxygen concentration(less than300ppb),a dual-layer oxide is formed on ferritic–martensitic steels.A typ-ical structure and composition profile for an oxide grown on a ferritic–martensitic steel is shown in the SEM image and EDS profiles of Fig.4.Addi-tional information on oxide structure in ferritic–martensitic steels can be seen in Fig.5,an EBSD image.The image and the composition profilesshow two distinct layers.The oxygen content is sim-ilar in both outer and inner oxide layers.The outer oxide is predominantly iron oxide whereas the inner layer contains a significant amount of chromium. For300ppb and lower dissolved oxygen concentra-tion,XRD and EBSD have revealed that the outer oxide layer is magnetite(Fe3O4)and the inner layer is an iron–chromium spinel of composition (Fe,Cr)3O4.For samples exposed to high(2000ppb)dis-solved oxygen,an outer hematite layer also forms. The formation of a two-layer oxide in low oxygen concentration and a three-layer oxide in high oxy-gen concentration is consistent with predicted stable phases[20].In addition EDS analysis of the oxygen profile shown in Fig.4b indicates the presence of an internal oxidation transition zone underneath the inner spinel layer.Grain morphology in each oxide layer is revealed by the EBSD map.The spinel layer is composed of small equiaxed grains with a large aspect ratio($0.7 in average),while the magnetite layer is composed of large columnar grains with a small aspect ratio ($0.4in average).The grains in the magnetite layer are elongated along the direction parallel to the growth direction of oxide scale.Carbides retained on the prior austenite grain structure in the inner oxide layer are similar to that in the ferrite phase, indicating that the oxide has grown from the ferrite by solid-state oxidation.3.1.2.Oxide kineticsA strong correlation between increasing tempera-ture and increased oxidation has been found for all ferritic–martensitic steels[13,15,56].The typical effect of time and temperature on oxide growth in ferritic–martensitic steels can be seen in Figs.6and 7,using HCM12A as an example.At higher temper-atures(500°C and600°C),the growth of the oxide follows roughly parabolic kinetics(the growth expo-nent is$0.4for this small data set).At temperatures below the pseudo-critical point(371°C),very little oxide growth is seen out to1026h.The oxide growth increases significantly with temperature.Activation energies have been estimated from the oxidationof Fig. 4.Cross-section image(a)and EDS(b)of HCM12A exposed to600°C SCW containing25ppb dissolved oxygen for yers are magnetite,spinel,and metal from left to right.180G.S.Was et al./Journal of Nuclear Materials371(2007)176–201T91,HCM12A,and HT9exposed at temperatures from 400to 600°C and are 189,177,and 172kJ/mol,respectively.Extensive experiments by Graham and Hussey [21]have suggested that in Ni,Cr,Fe,and Fe–Cr alloys,the outer oxides grow predominantlybyFig.5.Cross-section EBSD scanning maps of SCW-exposed samples (a)500°C,25ppb dissolved oxygen,505h and (b)500°C,2000ppb dissolved oxygen,505h,where base metal,spinel,magnetite,and hematite are highlighted in red,blue,yellow,and magenta,respectively.Black areas are unindexed [20].(For interpretation of the references in colour in this figure legend,the reader is referred to the web version of thisarticle.)G.S.Was et al./Journal of Nuclear Materials 371(2007)176–201181outward diffusion of cations.In particular,the acti-vation energy for the diffusion of iron in Fe 3O 4is 230kJ/mol,for nickel in NiO the value is 234kJ/mol,and for chromium in Cr 2O 3,the value is 420kJ/mol [22].The activation energy for diffusion of oxygen in Fe 2O 3is 610kJ/mol [21],and while that for Fe 3O 4is considerably smaller,the diffusion coefficients cannot account for the measured oxide thickness.Nevertheless,evidence suggests that the inner oxide grows by the inward diffusion of oxygen.The diffusion is likely affected by short-circuit paths such as pores,cracks and grain boundaries,which would account for the higher rates.For times to 1026h,the oxides that develop on traditional (non-ODS versions)ferritic–martensitic steels are stable,maintain a constant average density,and do not spall.This is evident in the data presented in Fig.8that shows the weight gain is proportional to the oxide thickness.The oxide den-sity is greater for samples exposed at 2000ppb as compared to samples exposed at 25ppb dissolved oxygen.Plan view and cross-sectional images of the dual phase oxides also do not show any indica-tion of spallation.3.1.3.Effect of dissolved oxygenThe oxide growth rate and associated weight gain of ferritic–martensitic steels are dependent on the dissolved oxygen concentration.This dependence is shown in Fig.9using data from HCM12A.For dissolved oxygen concentrations between 10and 300ppb,the weight gain associated with oxidegrowth decreases slightly with dissolved oxygen concentration.However,the weight gain increases significantly for samples exposed to 2000ppb dis-solved oxygen,due to the development of a thicker oxide layer.In addition,the development of a den-ser oxide layer and the formation of a hematite layer formed at the higher dissolved oxygen content SCW may also play a role in the upsurge in weight gain observed for samples exposed to 2000ppb dissolved oxygen SCW.Combined water chemistry control in fossil plants [23,24]adds small amounts of oxygen to enhance the formation of hematite crystals between the magnetite grains,thus reducing the oxidation rate,perhaps by reducing the diffusion of oxygen through the multi-phase film.For the studies sup-porting SCWRs,significant hematite crystals have not been found between the magnetite grains at 10–300ppb oxygen where oxidation is the slowest.3.1.4.Effect of bulk chromium concentrationFor conventional ferritic–martensitic steels,increasing the bulk chromium concentration reduces the weight gain due to oxidation.An exam-ple of this correlation is shown in Fig.10where the 9at.%Cr alloy NF616has a greater weight gain than the 12at.%alloy HCM12A.The same trend is seen in the weight gain data for T91(9Cr)and HT9(12Cr)for all temperatures and times evalu-ated.These results agree with those from Jang [15],as well as Cho and Kimura [16,25]who exam-ined weight gain over a range of Cr contentandFig.8.Weight gain as a function of oxide thickness for HCM12A exposed to 25ppb and 2000ppb oxygen at 500°C.Average oxide density (the slope of the plotted lines)is 1526mg/cm 2for 2000ppb dissolved oxygen and 1338mg/cm 2for 25ppb dissolved oxygen [20].182G.S.Was et al./Journal of Nuclear Materials 371(2007)176–201found that increasing Cr led to decreasing weight gain.A different behavior is noted in oxide dispersion strengthened ferritic–martensitic steels.The weight gain data for the JAEA9Cr ODS alloy is also included in Fig.10.Even though this alloy only has9at.%Cr,it shows the lowest weight gain of any of the tested ferritic–martensitic steels.The ODS spinel that forms tends to become more porous at higher exposure times[26].3.1.5.Oxidation reduction methodsAs was mentioned in the introduction,the allow-able oxide thickness must be limited for thin-walled cladding or water rods.For a1026h exposure at 600°C,the oxide thickness in HCM12A is approx-imately65l m,approximately15%of the thickness of the original water rod wall.For ferritic–martens-itic steels to be acceptable for thin-walled compo-nents exposed to supercritical water,the thickness of the oxide(and the associated metal loss)must be reduced.One method for improvement is seen from the ODS data in Fig.10.The ODS steels typically show lower oxidation than conventional ferritic–martens-itic steels.The ODS steels have two significant differences from conventional ferritic–martensitic steels.First,they include nanometer sized Y–Ti–O particles,added for strengthening.A comparison of the9Cr ODS alloy with NF616,Figs.11and 12,shows that the ODS alloy has a much deeper internal oxidation layer.Thus the net conversion of metal to spinel appears to be driven by oxygen diffusion into the metal rather than cation diffusion to the oxide surface.Detailed microscopy indicates the formation of Y–Cr-rich oxides along grain boundaries in the steel near the metal–oxide inter-face may act to block cation diffusion[26].The inner spinel layer that forms in the ODS alloy is more por-ous,and the lower density,combined with a thinner oxide,leads to a smaller weight gain in the ODS material.Additionally,the ODS materials typically have smaller grain sizes than conventional ferritic–martensitic steels.The increased diffusion length along short-circuit diffusion paths may slow oxidation.A second method that shows promise in slowing oxidation in ferritic–martensitic steels is surfaceFig.11.(a)Cross-sectional SEM image of the9Cr ODS ferritic steel after exposure to supercritical water at500°C for1026h and (b)corresponding composition profile across the oxide thickness[26].G.S.Was et al./Journal of Nuclear Materials371(2007)176–201183composition modification.Both oxygen and yttrium have been implanted into the surface of various ferritic–martensitic steels.The change in weight gain with various surface implantation conditions is shown in Fig.13.At 500°C,oxygen implantation reduced the weight gain in T91,HCM12A,and HT9.The reduction in oxide thickness in HT9has an associated change in the oxide texture early in the development of the oxide,but at longer times,the textures are similar between samples with and without oxygen pre-implantation [27].One possible explanation for the texture is that,among randomly orientated initial grains,some with certain crystal orientations grow faster than the others to release the stress introduced by the ion implantation.A sec-ond possibility is that bombardment with oxygen ions may increase nucleation sites with certain pre-ferred orientations at the initial stage of oxide for-mation,resulting in a denser textured oxide layer [17].At 600°C,the oxygen pre-implantation was not effective in reducing weight gain due to oxidation,but pre-implantation of yttrium strongly reduced the oxide thickness in both NF616and HCM12A.A thin layer of yttrium is incorporated into the mag-netite layer that appears to slow cation diffusion through the magnetite layer,thus reducing the oxidation by approximately 50%,Fig.14.3.2.Austenitic steelsAustenitic Fe-base steels were selected for possi-ble use in supercritical water systems becauseofFig.12.(a)Cross-sectional SEM image of the 9Cr NF616ferritic steel after exposure to supercritical water at 500°C for 1026h and (b)corresponding composition profile across the oxide thickness [26].Fig.13.Effect of oxygen and yttrium surface implantation at 500°C,and 600°C on T91,HCM12A,HT9,NF616,and 9Cr ODS.184G.S.Was et al./Journal of Nuclear Materials 371(2007)176–201their corrosion resistance and relative radiation resistance compared to Ni-base alloys.To date,international programs have evaluated the follow-ing austenitic steels:304L,304,304H,316L,316,D9,310S,and 800H [11,13,14,28–30].Grain bound-ary engineering has been performed on Alloy 800H to reduce oxide spallation.3.2.1.Oxide structureMost results on austenitic stainless steels reveal that the oxide consists of a two-or three-layer struc-ture,consistent with that observed during exposure to air,vacuum or subcritical water [57–59].The outer layer generally consists of magnetite with an inner layer that is rich in chromium and is either an iron–chromium spinel or an iron chromium oxide with a hematite structure [11,50,59].As in the case of ferritic steels an internal oxidation layer is also observed between the inner oxide layer and the base metal as evidenced by a gradually decreas-ing oxygen diffusion profile in this transition zone.A typical structure and composition profile for an oxide grown on an austenitic steel (800H)is shown in the SEM image in Fig.15.The oxide structure is somewhat similar to those formed in ferritic–martensitic steels,except that the outer oxide layer is composed of both magnetite and hematite.An additional difference between austen-itic and ferritic–martensitic steels is the stability of the outer oxide layer.In many of the austenitic stainless steels with higher bulk concentrations of nickel and chromium,the outer oxide layer has a tendency to spall,as demonstrated in the plan view images from D9exposed to 500°C 2000ppb SCW as shown in Fig.16.3.2.2.Oxide kineticsAustenitic stainless steels have a smaller weight gain than ferritic–martensitic steels [11,13,56].An example of typical kinetics for tests to 1026hisFig.14.Oxide thickness and compositional profile in HCM12A (top)and HCM12A with surface implantation of yttrium (middle).A thin layer of yttrium becomes incorporated into the outer magnetite layer (bottom,area marked 1)with a resulting $50%thickness inoxide.Fig.15.Cross-section of oxide formed on Alloy 800H exposed to 500°C SCW with 25ppb dissolved oxygen concentration for 505h [31].G.S.Was et al./Journal of Nuclear Materials 371(2007)176–201185shown in Fig.17for alloy D9.For 500°C and 600°C,the oxide kinetics appear parabolic,but the data scatter is greater in austenitic alloys than in ferritic–martensitic steels.In some austenitic alloys,a portion of this scatter is attributed to spall-ation.The oxidation rate in austenitic alloys is smal-ler than in ferritic–martensitic paringFigs.7and 17,at 600°C the oxidation at 1026h is roughly a factor of 4larger in HCM12A than in D9.Like ferritic–martensitic steels,the oxidation rate increases dramatically with increasing temperature.The temperature dependence of alloy D9is shown in Fig.18.Because of spallation,as evidenced for D9in Fig.19,a simple relation between oxide thick-Fig.16.SEM morphologies of the D9samples after exposure to 2000ppb SCW at 500°C for (a)168h,(b)335h and (c,d)503h [60].ness and weight gain is not possible in many austen-itic alloys.Activation energy has been calculated for some austenitic steels,specifically 210kJ/mol for 304L and 214kJ/mol for 316L,higher than the energy measured in T91,HCM12A,and HT9.3.2.3.Effect of dissolved oxygenThe effect of dissolved oxygen is more complex in austenitic alloys than in ferritic–martensitic steels.This is demonstrated in Fig.20for D9and 21for 316and 316L.For D9exposed at 500°C,at short times,the weight gain is smaller at very high (2000ppb)dissolved oxygen than at low (25ppb dissolved oxygen).At longer times (between 333and 505h),the oxidation rate of the D9exposedto 2000ppb dissolved oxygen increases significantly.It is noted for the D9samples exposed to 25ppb oxygen content,particularly for longer exposure times that the weight gain remains nearly constant,an affect that is attributed to oxide spallation in these alloys.The weight gain measurements shown in Fig.20represent the cumulative effects of weight gain due to oxidation as well as spallation.The weight loss due to spallation counterbalances the weight gain due to the growth of the oxide layer.This observation is supported by our SEM observa-tions of oxide spallation in D9steels.It is speculated that this effect would be observed for the samples exposed to higher oxygen SCW if the samples had been exposed to longer durations.Weight gain rate for 316and 316L as a function of temperature from this work is compared to that of Kasahara et al.[13],who exposed 316L in a static autoclave at 8000ppb dissolved oxygen,in Fig.21.At low temperatures,the oxidation is minimal and no significant difference is noted between 316and 316L or between exposures at different oxygen con-centration.At 500°C,higher dissolved oxygen leads to greater weight gain in 316.At 550°C,the weight gain rate is actually higher for the deoxygenated (<10ppb)case than for the 8000ppb case.The effect of oxygen is not as straightforward in the austenitic steels as was seen in the ferritic–mar-tensitic steels and greater study is needed toidentifyFig.19.Oxide thickness for D9exposed to 500°C SCW at 25ppb dissolved oxygen for 1026h [60].。