铬-钼- v钢的回火脆性 外文翻译

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铬-钼-V钢的回火脆性可逆回火脆性（RTB）－在500-650℃温度范围内回火或缓慢冷却的钢的脆性，它被认为是造成在前奥氏体晶粒边界形成杂质（P，锑，锡，砷）[1-4]的原因。然而，有资料显示[5]，不仅有这些进程，而且有其他进程，像在500-600℃淬火钢有助于回火脆性在发生。在这项工作的关注是淬火铬钼合金，以防止在钒和磷的含量RTB的钢回火脆化。热件1（0.35％的V，0.015％P）重四十一吨是伪造的，以一个550毫米的大小,15毫米厚的板材被切断他们。另外，在一个100公斤的感应炉加热熔化。钒被添加到在一个重达16公斤的铸块中，它被锻造、轧制成10毫米厚的金属板。在所有的10-15毫米厚加热板放在油中980℃（1小时）淬火。从板淬火和在100-760℃回火10小时后制备样品。经过淬火和低温回火，所有加热件呈现一个典型的马氏体结构，但在高（超过600摄氏度）的温度回火后得到回火索氏体结构。在实验室加热的奥氏体晶粒尺寸较小（8-9级）比商业热（4级），它们的结构部件具有更大的分散性和均匀性的。脆性是从Tso和Ttemper变化中来确定，Ttemper是淬火钢的回火温度，Tso是韧-脆转变温度，它具有最完整的脆化特征。Tso的确定是通过5×5×27.5毫米的切口1毫米深（根半径0.25毫米）冲击试验样品。Tso被认为是在断裂50％纤维下的测试温度。拉伸强度通过5个直径为3毫米的样品确定的为20°。图1显示了热件1回火温度与力学性能的变化。这种热铸块的力学性能在100-600℃回火后几乎不变。当回火温度从600-760度提高时，强度特性急剧下降，而韧性增加。更改影响回火钢的稳定的情况下钒浓度，因为不含钒的钢的强度会下降至约500°。热件1开始时回火温度大约300℃，Tso的增加达到500-600度的最高值。如同淬火条件对峰值相比增加了100℃。回火温度进一步提高，Tso会下降，这与削弱开始相吻合。对于热件2-7的Tso和Ttemper变化的总体特点是相同的，尽管在较低温达到峰巅。这类钢在是Tso〜90℃条件冷却，低于热1大概-70±10℃。在730℃后的Tso锻炼价值与热2-7（-110至-130°）实际上是相同。它的峰高和立场取决于钒钢的内容。当钒浓度变化从0到0.55％的高峰上升％60°，同时转移〜100°。比较热件1和3，6，在史前的冶金学上有区别的，但相同钒含量（〜0.3％）是相同的，人们可以看到，在Tso峰地区的增长几乎是与淬火条件相同。在0.005和0.022％的磷的热处理后的测试中，磷对Tso没有什么影响。在回火期间的钢的力学性能改变显然取决于精细结构的变化。图3显示了在热淬火Tso 显微条件后，在600℃回火，对应于对强度特性极值边缘与Tso高峰在760°的回火之后，强度和Tso达到最低的。淬火后的结构由带有位错的板条马氏体组成。在板条彼此间略有错误导向，平均宽度〜0.3μ和长度〜5μ，并归为〜5×5μ。板条充满均匀分布的密度位错〜011cm-2。硬质合金阶段没有观察到淬火钢。在600℃回火10小时的后混乱阵列的整体性质和晶体的碎片将被保留。木板条的平均大小（宽度和长度）它的平均规模保持不变。唯一的变化是明显的沉淀分散的碳化物阶段。粒子的大小是150至200A，平均密度为~1015cm-2。这些沉淀在混乱通过板条马氏体空间分布。带长条形直径为〜250A和〜2•103A的较大沉淀物位于深的晶体的边界附近（板条马氏体）。这种类型的技术被作为MTCa电石鉴定。在760℃回火后急剧变化的钢的结构：位错是不规则的。大部分分布在整提失调，改建后更适合了，与位错领域相吻合。单元格的平均规模为〜0.5μ。该阶段碳化物形态的也有变化。微细分散的大小圆形的沉淀〜250A的位错在位于网络的连接处。随着他们有更大的圆形沉淀〜2•103A位于边界的十字路口。析出的碳化物回火后的存在也证实了热物理化学相分析热件1。当回火温度由原来〜300-600度的渗碳体碳化物的数量逐渐减少，M7C3增加。强有力的碳化物含量形成的残留增加元素：铬，钼，和钒.，这个过程在500-600度是特别活跃。这很可能是VC和MO2C膜类型不理化分析检测，在高度分散的碳化物阶段也有更多的热力学稳定相[6]。因此，氢脆回火过程中铬钼钒钢，作为对Tso与Ttemper出现明显的高峰表现，从我们的实验结果。它的峰值并没有关联的磷在钢存在，但与钒的浓度各不相同，在碳化物的形成，与继承发展脱位阵列的温度范围内位置。Tso达到高峰时的温度在Fe3C的变换更稳，与M7C3类似的现象已在铬钢观察[5]。脆化可能是由于体积和边界的影响。在铬钼钒钢的的情况二次硬化，体现在在Ttemper=300-500℃性能的增强可能与这些相关的影响有关。据了解，[7]，随着晶粒尺寸不变（或板条马氏体晶粒常量大小）Tso通过εTso=σy+C与低屈服强度σy（σ0.2相关的关系）同期相联系，其中C和ε为常数。热件1Tso呈线性关系中的屈服强度为σo.2大的变化范围，并只有在400-600度，特殊的碳化物开始制定改变，有偏离线性的关系。当然，这些过程首先影响边界地区，那里的条件是为硬质合金阶段准备，由于晶体结构的几何缺陷有利合金相的形成。可以设想，Tso峰值是由于在边界条件的变化。这个假设是证实了电子显微镜分析结果：在矩阵（位错密度，碎片大小条件）引起材料的[8，9]实力的增强，保持不变直至Ttemper=600°。这也证实了结构进行检查和对回火温度与屈服强度的变化。因此，如果发生任何更改的边界，增加脆化，那么，Tso同Ttemper变化将有虚线的形式。所有解释RTB[1，2，10]的理论都是建立在对边界的影响起主导作用基础上的。然而，前奥氏体晶粒边界损害并非我们的典型调查（热件1纹裂在淬火、回火后显现出来 ），没有磷的影响，这显然是对钼的存在钢铁中的解释。由此可见，由实验数据C–Mo–V钢回火过程中的脆性主要受碳化物形成的影响，更加准确地重建碳化物和的这些过程影响因素。这原种影响的理可以概括介绍如下，随回火温度的增加渗碳体开始凝聚，在250-350℃开始沉淀。更有效地遏制电石，在钢中容纳更多的碳化物成形元素比铁中350℃左右开始的特殊碳化物的原子核钢的成形，尤其是M7C3[8]。沉淀物均匀分布在整个体积的位错位置，低角度的板条马氏体的界限，高角度的板条殖民地边界，他们加强了矩阵和削弱（脆化）边界。在二次硬化温度时，很可能最大程度的削弱碳化物与基体或与沉淀、连贯的最大密度的边界，i.e.。如元素钒，促进碳化物细化，从而增加二次硬化[8]，增加脆化。在沉淀和聚结硬质合金阶段，出现了混乱的位错的同时，边界变得更加完善和矩阵被削弱。这两种效应导致韧性、脆性转变温度下降，这是从实际观察的结构。结论1。15Kh3MFA类型的钢都容易脆化，在给定的回火温度下达到高峰。在高峰期（Tso）的上限，在正值温度下材料开始削弱。2。该峰的高度和回火温度也相应增加时，几乎线性钒浓度从0提高到0.55％，但他们是在0.005-0.022％的磷浓度限制下表现出来的。3。回火钢的脆性取决于在碳化物相变（从渗碳特殊碳化物），保留的位错优先发生在碎片的边界。文献引用1.L.M.Utevskii,TemperBrittlenessofSteels[inRussian],Metallurgizdat,Moscow(1961),p.138.2.P.B.Mikhailov-Mikheev,ThermalEmbrittlementofSteels[inRussian],Mashgiz,Moscow--Leningrad(1956),p.56.3.J.Hollomon,Trans.ASM,36,473(1946).4.E.Houdremont,SpecialSteels[Russiantranslation],Vol.I,Metallurgiya,Moscow(1966),p.455.5.V.A.Korablev,Yu.I.Ustinovshchikov,andI.G.Khatskelevich,"Embrittlementofchromiumsteelswithformationofspecialcarbides,"Metalloved.Term.Obrab.Met.,No.I,16(1975).6.A.P.Gulyaev,I.K.Kupalova,andV.A.Landa,"Methodandresultsofphaseanalysisofhlgh-speedsteels,"Zavod.Lab.,No.3,298(1965). 7.J.HeslopandN.Perch,Phil.Mag.,~,No.34,1128(1958).8.V.V.Rybinetal.,"Themechanismofhardeningofsorbite-hardeningsteelandthepossibilityofdeterminingittheoreticallyandexperimentally,"in:MetalScience[inRussian],No.17,Sudostroenie,Leningrad(1973),p.105.9．L.K.Gordienko,SubstrucutralHardeningofMetalsandAlloys[inRussian],Nauka,Moscow(1973),p.64.10．E.E.Glikmanetal.,"Natureofreversibletemperbrittleness,"Fiz.Met.Metalloved.,36,365(1973).11.OliverWC,PharrGM(2004)JMaterRes19:312.KimJY,LeeBW,ReadDT,KwonD(2005)ScrMater52:35313.KimJY,LeeJS,LeeKW,KimKH,KwonD(2006)KeyEngMater326–328:48714.KimJY,LeeJJ,LeeYH,JangJI,KwonD(2006)JMaterRes21:297515.KimJY,KangSK,LeeJJ,JangJI,LeeYH,KwonD(2007)ActaMater55:355516.DowlingNE(1993)Mechanicalbehaviorofmaterials.PrenticeHall,EnglewoodCliffs17.KimJY,LeeKW,LeeJS,KwonD(2006)SurfCoatTechnol201:427818.DIN17175-79(1979)Seamlesssteeltubesforelevatedtemperatures19.AhnJH,KwonD(2001)JMaterRes16:317020.DieterGE(1988)Mechanicalmetallurgy.McGraw-Hill,Singapore外文原文TEMPERBRITTLENESSOFCr--Mo--VSTEELReversibletemperbrittleness(RTB)–embrittlementofsteelsduringtemperingorslowcoolinginthetemperaturerangeof500-650℃isconsideredtoresultfromtheformationofimpuritysegregates(P,Sb,Sn,As)inprioraustenitegrainboundaries[1-4].However,therearedataindicating[5]thatnotonlytheseprocessesbutotherprocessesfavoringembrittlementoccurinquenchedsteelduringtemperingat500-600°. ThisworkconcernsembrittlementduringtemperingofquenchedchromiumsteelsalloyedwithmolybdenumtopreventRTBinrelationtothevanadiumandphosphorusconcentrations.Ingotsofheat1(0.35%V,0.015%P)weighing41tonswereforgedtoasizeof550mm.Plates15mmthickwerecutfromthem.Theotherheatsweremeltedina100-kginductionfurnacewithuseofZhS-0ironinthecharge.Vanadiumwasaddedtothesteelduringpouringofaningotweighing16kg,whichwasforgedandrolledtoaplate10mmthick.Allheatsintheformofplates10-15mmthickwereoilquenchedfrom980°(1h).Sampleswerepreparedfromtheplatesafterquenchingandaftertemperingat100-760°for10h.Afterquenchingandlow-temperaturetempering,allheatshadatypicalmartensiticstructure,butafterhigh-temperaturetempering(above600°)asorbitestructure.Theaustenitegrainsizeofthelaboratoryheatswassmaller(grade8-9)thaninthecommercialheat(grade4),withgreaterdispersityandhomogeneityofthestructuralcomponents.EmbrittlementwasdeterminedfromthevariationofTsowithTtemper,whereTtemperisthetemperingtemperatureofthequenchedsteelandTsoistheductile--brittletransitiontemperature,whichmostcompletelycharacterizesembrittlement.Tsowasdeterminedonimpacttestsamples5×5×27.5mmwithanotch1mmdeep(rootradius0.25mm).Tsowastakenasthetestingtemperatureatwhichthefracturewas50%fibrous.Thetensilestrengthwasdeterminedat20°onfivesamples3mmindiameter.Figure1showsthevariationofthemechanicalpropertiesofheat1withthetemperingtemperature.Themechanicalpropertiesofthisheatarealmostconstantaftertemperingat100-600°.Whenthetemperingtemperatureisraisedfrom600-760°thestrengthcharacteristicsdecreasesharply,whiletheductilecharacteristicsincrease.Changingthevanadiumconcentrationaffectsthestabilityofthesteelduringtempering.Forthesteelwithoutvanadiumthestrengthbeginstodecreasearound500°.Forheat1,beginningwithtemperingattemperaturesaround300°,Tsoincreasesandreachesamaximumvalueat500-600°.AscomparedwiththequenchedconditiontheincreaseofTsoatthepeakis100°.WithfurtherincreaseofthetemperingtemperatureTsodecreases,whichcoincideswiththebeginningofweakening.Forheats2-7theoverallcharacterofthevariationofTsowithTtemperisthesame,although thepeaksoccuratlowertemperatures.ForthesteelsinthequenchedconditionTsois~90°lowerthanforheat1,andamountsto-70±10°.Aftertemperingat730°thevalueofTsoispracticallythesameforheats2-7(-110to-130°).Theheightofthepeakanditspositiondependsonthevanadiumcontentofthesteel.Whenthevanadiumconcentrationischangedfrom0to0.55%thepeakrises%60°andatthesametimeshifts~100°.Comparingheats1and3,6,differingintheirmetallurgicalprehistorybutsimilarinvanadiumcontent(~0.3%),onecanseethattheincreaseofTsointheregionofthepeakisalmostthesameasforthequenchedcondition.Inamountsof0.005and0.022%,phosphorushasnoeffectonTsoaftertheheattreatmentstested.Thechangesinthemechanicalpropertiesofthesteelduringtemperingevidentlydependonchangesinfinestructure.Figure3showsthemicrostructureofheatinthequenchedcondition,aftertemperingat600°,correspondingtotheedgeoftheplateauofthestrengthcharacteristicsandthepeakofTso,andaftertemperingat760°,wherethestrengthandTsoarelowest.Afterquenching,thestructureconsistsoflathmartensitewithwell-developeddislocationarrays.Thelathsareslightlymisorientedwithrespecttoeachother,withanaveragewidthof~0.3μandlength~5μ,andaregroupedincolonies~5×5μ.Thelathsarefilledwithevenlydistributeddislocationswithadensity~1011cm-2.Nocarbidephasewasobservedinthequenchedsteel.Aftertemperingat600°for10htheoverallcharacterofthedislocationarraysandthefragmentationofthecrystalsareretained.Theaveragesizeofthelaths(widthandlength)andtheaveragesizeofthecoloniesremainunchanged.Theonlynoticeablechangeistheprecipitationoffinelydispersedcarbidephases.Thesizeoftheparticlesis150-200Aandtheaveragedensity~1015cm-2.Theyareprecipitatedondislocationsandeventlydistributedthroughthebulkofthemartensitelaths.Largerprecipitateswiththeshapeofneedles(platelets)~250Aindiameterand~2·103Alongarelocatednearhighlymisorientedboundaries(grains,coloniesofmartensitelaths,themostmisorientedlaths,etc.)andintheboundariesthemselves.PrecipitatesofthistypewereidentifiedbymicrodiffractiontechniquesasMTCacarbide.Thestructureofthesteelchangessharplyaftertemperingat760°:thedislocationsarepolygonized.Thedislocationsdistributedthroughoutthebulkarerebuiltintoenergeticallymore suitableconfigurations–cellwalls–notcoincidingwithlong-rangefields.Theaveragesizeofthecellsis~0.5µ.Themorphologyofthecarbidephasesalsochanges.Finelydispersedroundedprecipitateswithasizeof~250Aarelocatedinthejunctionsofthedislocationnetwork.Alongwiththemtherearelargerroundedprecipitates~2·103locatedattheintersectionsofcellboundaries.Thepresenceofcarbideprecipitatesaftertemperingwasalsoconfirmedbyphysicochemicalphaseanalysisofheat1.Whenthetemperingtemperatureisraisedfrom~300-600°thequantityofFe3CcarbidegraduallydecreasesandthequantityofM7C3increases.Theconcentrationofstrongcarbide-formingelementsintheresiduesincreases:Cr,Mo,andV.Thisprocessisparticularlywelldevelopedat500-600°.ItishighlyprobablethatamongthehighlydispersedcarbidephasestherearealsomorethermodynamicallystablephasesoftheMo2CandVCtypenotdetectedbyphysicochemicalanalysis[6].ItfollowsfromourexperimentalresultsthatembrittlementoccursduringtemperingofCr–Mo–Vsteels,manifestasadistinctpeakonthecurveofTsovsTtemper.Thepeakisnotassociatedwiththepresenceofphosphorusinthesteelbutvarieswiththeconcentrationofvanadiumandislocatedinthetemperaturerangeofcarbideformation,whichdevelopswithinheritanceofthedislocationarrays.TsoreachesapeakatthosetemperatureswhereFe3CtransformstomorestableM7C3.Asimilarphenomenonhasbeenobservedinchromiumsteels[5].Embrittlementmaybeduetobothvolumeandboundaryeffects.InthecaseofCr–Mo–VsteelsthesecondaryhardeningmanifestintheincreaseofthestrengthatTtemper=300–500°maybeassociatedwiththefirstoftheseeffects.Itisknown[7]thatwithaconstantgrainsize(orconstantsizeofcoloniesofmartensitelaths)Tsoisassociatedwithalowyieldstrengthσy(σ0.2infirstapproximation)bytherelationshipεTso=σy+C,whereCandεareconstants.Forheat1Tsovarieslinearlywithσo.2inabroadrangeofchangesinyieldstrength,andonlywithtemperingat400-600°,wherespecialcarbidesbegintodevelop,isthereadeviationfromthelinearrelationship.Ofcourse,theseprocessesaffecttheboundaryzonesfirst,whereconditionsarefavorablefortheformationofcarbidephaseduetothegeometricimperfectionofthecrystalstructure.ItcanbeassumedthattheTsopeakisduetoachangeintheconditionoftheboundaries.Thisassumptionisconfirmedbytheresultsofelectronmicroscopicanalysis:Theconditionofthematrix(dislocationdensity,sizeoffragments),causinganincreaseinthestrengthofthematerial [8,9],remainsunchangeduptoTtemper=600°.Thisisconfirmedbyexaminationofthestructureandbythevariationoftheyieldstrengthwiththetemperingtemperature.Thus,ifnochangesoccurredintheboundarytoincreaseembrittlement,thenthevariationofTsowithTtemperwouldhavetheformofthedashedline.AlltheoriesexplainingRTB[1,2,10]arebasedontherecognitionofthedominantroleofboundaryeffects.However,damageintheboundariesofprioraustenitegrainswasnottypicalinourinvestigation(thefractureofheat1afterquenchingandaftertemperingwasquasibrittle),andnoeffectofphosphoruswasobserved,whichisevidentlyexplainedbythepresenceofmolybdenuminthesteel.ItfollowsfromtheexperimentaldatathattheembrittlementofC–Mo–Vsteelsduringtemperingisaffectedmainlybycarbideformation,morepreciselytherebuildingofcarbidesandfactorsaffectingtheseprocesses.Themechanismofthiseffectcanbepresentedingeneraltermsasfollows.Withincreasingtemperingtemperaturescementitebeginstocoalesce,precipitatingat250-350°.Insteelscontainingmoreeffectivecarbide-formingelementsthanirontheformationofnucleiofspecialcarbidesbeginsaround350°,especiallyM>C3[8].Precipitatingevenlythroughoutthevolumeondislocations,low-angleboundariesofmartensitelaths,andhigh-angleboundariesofcoloniesoflaths,theystrengthenthematrixandweaken(embrittle)theboundaries.Itisprobablethattheboundariesareweakenedmostwiththemaximumdensityofcarbidescoherentwiththematrixorwiththeirprecipitation,i.e.,atsecondaryhardeningtemperatures.Suchelementsasvanadium,promotingrefiningofcarbidesandthusincreasingsecondaryhardening[8],increasetheembrittlement.Withprecipitationandcoalescenceofcarbidephase,occurringsimultaneouslywithpolygonizationofdislocations,theboundariesbecomemoreperfectandthematrixisweakened.Boththeseeffectsleadtoadropoftheductile-brittletransitiontemperature,whichisinfactobserved.CONCLUSIONS1.Steelsofthe15Kh3MFAtypearesusceptibletoembrittlement,whichreachesapeakatagiventemperingtemperature.Theupperlimitofthepeak(Tso)coincideswiththetemperatureat whichthematerialbeginstoweaken.2.Theheightofthepeakandthetemperingtemperaturecorrespondingtoitincreasealmostlinearlywhenthevanadiumconcentrationisraisedfrom0to0.55%,buttheyareindependentofthephosphorusconcentrationwithinlimitsof0.005-0.022%.3.Temperbrittlenessofthesteelinvestigateddependsonthechangeinthecarbidephase(fromcementitetospecialcarbides)thatoccurswithretentionofthedislocationarrayspreferentiallyintheboundariesoffragments.LITERATURECITED1.L.M.Utevskii,TemperBrittlenessofSteels[inRussian],Metallurgizdat,Moscow(1961),p.138.2.P.B.Mikhailov-Mikheev,ThermalEmbrittlementofSteels[inRussian],Mashgiz,Moscow--Leningrad(1956),p.56.3.J.Hollomon,Trans.ASM,36,473(1946).4.E.Houdremont,SpecialSteels[Russiantranslation],Vol.I,Metallurgiya,Moscow(1966),p.455.5.V.A.Korablev,Yu.I.Ustinovshchikov,andI.G.Khatskelevich,"Embrittlementofchromiumsteelswithformationofspecialcarbides,"Metalloved.Term.Obrab.Met.,No.I,16(1975).6.A.P.Gulyaev,I.K.Kupalova,andV.A.Landa,"Methodandresultsofphaseanalysisofhlgh-speedsteels,"Zavod.Lab.,No.3,298(1965).7.J.HeslopandN.Perch,Phil.Mag.,~,No.34,1128(1958).8.V.V.Rybinetal.,"Themechanismofhardeningofsorbite-hardeningsteelandthepossibilityofdeterminingittheoreticallyandexperimentally,"in:MetalScience[inRussian],No.17,Sudostroenie,Leningrad(1973),p.105.9．L.K.Gordienko,SubstrucutralHardeningofMetalsandAlloys[inRussian],Nauka, Moscow(1973),p.64.10．E.E.Glikmanetal.,"Natureofreversibletemperbrittleness,"Fiz.Met.Metalloved.,36,365(1973).11.OliverWC,PharrGM(2004)JMaterRes19:312.KimJY,LeeBW,ReadDT,KwonD(2005)ScrMater52:35313.KimJY,LeeJS,LeeKW,KimKH,KwonD(2006)KeyEngMater326–328:48714.KimJY,LeeJJ,LeeYH,JangJI,KwonD(2006)JMaterRes21:297515.KimJY,KangSK,LeeJJ,JangJI,LeeYH,KwonD(2007)ActaMater55:355516.DowlingNE(1993)Mechanicalbehaviorofmaterials.PrenticeHall,EnglewoodCliffs17.KimJY,LeeKW,LeeJS,KwonD(2006)SurfCoatTechnol201:427818.DIN17175-79(1979)Seamlesssteeltubesforelevatedtemperatures19.AhnJH,KwonD(2001)JMaterRes16:317020.DieterGE(1988)Mechanicalmetallurgy.McGraw-Hill,Singapore