Weldability of Ferritic Alloys for Future Power Plants

REVIEW

Welding and weldability of candidate ferriticalloys for future advanced ultrasupercriticalfossil power plants

S. A. David*1, J. A. Siefert2 and Z. Feng1

Fossil fuels continue to be the primary source of energy in the world. The worldwide demand for

clean and affordable energy will continue to grow, and a strong emphasis has been placed on

increasing the efficiency and reducing the carbon footprint of new and existing fossil fired power

plants. Throughout Asia, Europe and the USA, this demand is being met with programmes to

develop advanced materials that have enhanced high temperature creep and corrosion

properties. A new class of ferritic alloys, known as creep strength enhanced ferritic steels, has

been developed to meet these requirements. This article focuses on the weldability of the

advanced ferritic alloys used in boilers and boiler components of ultrasupercritical coal fired

power plants. This review focuses on alloy selection; welding and weldability issues, including in

service weld failure such as type IV cracking; welding of dissimilar metals; and weld repair. Future

articles will address the welding and weldability issues of two other classes of materials, namely

austenitic stainless steels and nickel base superalloys.

Keywords: Ferritic alloys, Type IV cracking, Fossil energy, Ultrasupercritical, Dissimilar metal welds, Weld repair, Friction stir welding

IntroductionEfforts are under way to meet the world demand forclean and affordable energy, particularly in fossil powerindustries, to increase the plant operating efficiencies andto reduce carbon footprints of conventional coal firedpower plants and cycling heat recovery steam generatorapplications. This requires operating the plant at highersteam temperatures and pressures. Accomplishing thisgoal requires development and utilisation of advancedmaterials. In addition, successful development and utili-sation of these materials in advanced power systemsdepend on their ability to be welded or joined. Therefore,welding and weldability of these new materials becomekey issues. Weldability plays a critical role in the selectionof materials, in particular for the advanced coal firedpower plants.

Over the last 50 years, coal fired power plants haveevolved from a subcritical plant, which operates at a steamtemperature of 540uC and at a steam pressure of 16?5 MPawith a normal efficiency of 35% HHV [the higher heatingvalue includes latent heat of vaporization of water in thecombustion of coal. In this paper, plant efficiency is alwaysnoted in terms of HHV.], to the present ultrasupercritical(USC) plant, which is capable of operating at temperatures

.593uC, a steam pressure .24?8 MPa and an efficiency upto 42%. Now, efforts are under way in the USA, Europeand Japan to increase the efficiency of USC power plantsto .45% by increasing the steam temperature andpressure; these parameters are representative of so calledadvanced ultrasupercritical (AUSC) technology. Effortsfor alloy development for higher efficiency USC powerplants have focused on the optimisation of current creepstrength enhanced ferritic (CSEF) alloys and the modifica-tion of traditional chromium–molybdenum (Cr–Mo)steels. These alloys have excellent high temperature creepstrength and excellent physical and corrosion properties.Conventional alloy design principles are used to enhancecreep strength.1–3 Recently, within the realm of alloydesign, Bhadeshia4 has reviewed a currently availablequantitative methodology to attempt to design new heatresistant steels and welding consumables.

In addition to alloy development, many welding andjoining issues require attention to address the materialneeds of USC power plants. Unlike the other structuralmaterials, CSEF steels require special care and attentionduring welding if they are to retain their high temperatureproperties. They need a very precisely controlled weldingprocedure, and susceptibility of these alloys to type IVcracking [type IV cracking is creep cracking (also knownas ‘‘midlife cracking’’) that occurs during service in theheat affected zone of ferritic steel weldments] needs to beunderstood and mitigated. In this paper, power plantclassification, materials selection, welding and weldabilityissues, weldability of USC power plant materials andtopics including weld repair and welding of dissimilar

1Oak Ridge National Laboratory, One Bethel Valley Road, Oak Ridge, TN37831, USA2Electric Power Research Institute, 1300 West W T Harris Blvd, Charlotte,NC 28262, USA

*Corresponding author, email [email protected]

� 2013 Institute of Materials, Minerals and MiningPublished by Maney on behalf of the InstituteMORE OpenChoice articles are open access and distributed under the terms of the Creative Commons Attribution Non-Commercial License 3.0Received 30 April 2013; accepted 19 June 2013DOI 10.1179/1362171813Y.0000000152 Science and Technology of Welding and Joining 2013 VOL 18 NO 8 631

metals are reviewed, and a prognosis for welding themsuccessfully is offered.

Welding and weldabilityWelding is a critical and complex fabrication technologyused in the construction of energy systems. The term‘weldability’ has no universal definition. The Americanwelding society handbook defines weldability as ‘thecapacity of a material to be welded under the imposedfabrication conditions into a specific, suitably designedstructure and to perform satisfactorily in the intendedservice’.5 Although adequate, that definition does notaddress all of the issues related to material behaviourwhen exposed to the weld thermal cycle.

Recent discussions of the authors with others in thefield have led to a new way of looking at weldability. Toput it simply, ‘weldability is a measure of the ease withwhich a metal or an alloy can be welded or joined withoutdegradation that is detrimental to the weldment micro-structure or properties during or after welding and for theduration of intended service’.6 The definition is inclusiveof most of the problems encountered by the weldedstructures (e.g. hot and crater cracking, cold or hydrogencracking, reheat cracking and in service problems such astype IV cracking). Undeniably, the term ‘weldability’remains a highly subjective term, whose definition may bespecific to an individual or institution.

During welding, the original microstructure of thematerial is destroyed by melting, which creates a liquidpool that solidifies upon cooling. This region is calledthe fusion zone (FZ). Beside the FZ is the heat affectedzone (HAZ), a region of the metal that experiences heatbut does not melt. The rest of the material, which is notaffected by the welding heat, is the base metal.

Figure 1 is the schematic diagram showing the devel-opment of various subzones in the HAZ of P(T)91weldment during welding and their approximate correla-tion to the calculated P(T)91 equilibrium phase diagram7

(see Table 1 for alloy compositions of CSEF steels and thedesignations used and Table 2 for CSEF boiler materialapplications).8,9 Similar regions exist for all of the CSEFsteels. In particular, the figure shows the various sub-microstructural zones due to the thermal gradient experi-enced by the HAZ. Often, most of the welding andweldability problems encountered during welding are not

1 Schematic diagram of various subzones in HAZ of P(T)91 weld and its approximate relation to calculated equilibrium

phase diagram7 (courtesy of P. Mayr)

Table 1 Composition in of candidate CSEF steels/wt-%

Steel alloy Code case C Si Mn Cr Ni Mo W V Nb B N Others

T23 2199 0?06 0?20 0?45 2?25 … 0?1 1?6 0?25 0?05 0?003 0?010 Ti: 0?010T24 2540 0?08 0?30 0?50 2?25 … 1?0 … 0?25 … 0?004 Ti: 0?70P(T)91 Incorporated 0?10 0?40 0?40 9?0 … 1?0 … 0?20 0?08 … 0?050P(T)92 2179 0?07 0?06 0?45 9?0 … 0?50 1?8 0?20 0?05 0?004 0?060E911 2327 0?10 0?20 0?40 9?0 0?20 1?0 1?0 0?20 0?08 … 0?070P(T)122 2180 0?10 0?25 0?60 12?0 0?25 0?35 2?0 0?20 0?06 0?003 0?060 Cu: 0?80VM12 N/A 0?12 0?50 0?30 11?4 0?25 0?25 1?5 0?25 0?05 0?005 0?050 Co: 1?5

Table 2 Potential CSEF steel boiler materials andapplications for SC, USC and AUSC powerplants*8,9

Material

Application

P5pipe/headers

SH/RH5superheat/reheat tubing

Ww5waterwall tubing

Subcritical (SC) USC AUSC

T23 P, SH/RH P, SH/RH, ww wwT24 SH/RH, ww wwP(T)91 P, SH/RH P, SH/RH P, SH/RH, ww*P(T)92 P, SH/RH P, SH/RH P, SH/RH, ww*E911 P P ww*P(T)122 P P ww*VM12 SH/RH SH/RH ww*

*ww*: postweld heat treatment (PWHT) required. Note that therequirement of PWHT in waterwall applications introducescomplexity that may be difficult to address.

David et al. Welding and weldability of candidate ferritic alloys

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only associated with the structures of the FZ and the HAZ,but also due to the gradients in structure, which are ofcritical importance in focusing creep strain, for example inthe type IV cracking phenomenon. Also during welding,thermal and solidification shrinkage stresses develop, andmost of the stresses remain in the weldment as residualstresses and may affect the performance of CSEFweldments in the following ways:

(i) residual stresses in the weldment in the presenceof a corrosive atmosphere may increase thepotential for stress corrosion cracking (SCC).As welded grade 91 components that wereexposed to uncontrolled corrosive atmosphericconditions were found to be affected.10 Additionalissues, attributed to SCC, have been reportedduring start-up of new supercritical (SC) unitsemploying grades 23 and 24 in waterwallpanels11,12

(ii) the presence of residual stress and a triaxial stressstate may induce reheat cracking in componentsmade from grade 2313–16

(iii) the presence of residual stress, combined withimproper preheat in all CSEF materials, maylead to issues resulting in hydrogen inducedcracking.17–20

Several other factors must be considered before welding[e.g. joint design, selection of a joining process, jointrestraint, postweld heat treatment (PWHT), procedurequalification]. Typically, commercial alloys are selectedon the basis of mechanical properties or corrosionresistance. Generally, welding and weldability of theselected material are considered or evaluated at the endof the material selection cycle; such negligence ofwelding considerations may result in significant pro-blems in the construction of a component.

Weldability of a material is often determined by anumber of metallurgical and non-metallurgical factors,such as composition, microstructure, properties, processparameters and restraint. In actuality, to determine theweldability of an alloy, a testing procedure that duplicatesthe conditions that the material experiences during weldfabrication should be employed. In most cases, it is notfeasible to demonstrate weldability with a full size mock-up component with the appropriate materials, weldingconsumables and restraint. Therefore, a number of welda-bility tests have been developed to appropriately char-acterise the behaviour of the materials during welding andare described in the literature.21–23 Weldability tests havebeen designed to evaluate the weldability of metals andalloys during fabrication (e.g. hot cracking tests, implanttests, laminar tearing tests) and during service (e.g. tensileand ductility tests, fracture toughness tests, creep resis-tance, fatigue and corrosion resistance tests).23 Despite theabundance of weldability tests, no single universal testexists that can determine the weldability of materials. One

possible solution would be to develop an integratedcomputational materials engineering model with the indus-trial software tools that are currently available (Table 3).24

Coal fired power plantsThe development of advanced fossil energy technologiesfor increased efficiency and a reduced carbon footprintrequires the development of materials and structuresthat can withstand an increase in operating steamtemperature and pressure. In its most simplistic descrip-tion, a boiler carrying high pressure steam consists oftubes (i.e. small outside diameter and thin walls) andpipes (i.e. large outside diameter and thick walls). Pipescan be subdivided into headers that collect steam andpiping that transports hot steam to the turbine. Tubesconsist of furnace wall tubes and superheater/reheatertubes. Figure 2 shows a typical header consisting ofthick walled pipe penetrated by a number of tubes.25

The terms ‘subcritical’, ‘supercritical’ and ‘ultrasuper-critical’ refer to the temperature and pressure of a powerplant’s working fluid (steam). Steam conditions can beillustrated in a phase diagram for water.26 Figure 3 (notto scale) shows phase equilibria for water with variousphase fields (steam, water and ice). Solid lines A, B andC show where two adjacent phases coexist. It also showsthe values of pressure and temperature at which waterand steam are in equilibrium with each other. Point D inFig. 3 is the triple point (i.e. the temperature andpressure at which the three phases, namely water, iceand steam coexist). The normal boiling point of water inFig. 3 is at 100uC and 0?1 MPa; at that point on curveC, water and steam coexist. Increasing pressure andtemperature above that point causes water and steam tocoexist as two phases in the subcritical range. Point E oncurve C, which occurs at 374uC and 22?1 MPa, is calledthe critical point of water (Fig. 3), above which thephase boundary between water and steam terminatesand distinct steam and water phases do not exist.

Table 3 Software tools for integrated modelling24

Software Internet Link Platform

SYSWELD http://www.esi-group.com/products/welding DesktopVr-Weld http://www.goldaktec.com/vrweld.html DesktopVFT http://www.battelle.org/ DesktopWELDSIM http://www.aws.org/wj/2008/05/wj200805/wj0508-36.pdf DesktopSORPAS http://www.swantec.com/sorpas.htm DesktopE-WeldPredictor http://calculations.ewi.org/VJP/ Internet

2 Photograph of typical header25 (with permission of

Springer Science)



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Increasing pressure and temperature above that pointpushes steam into SC range. Further increase in eitherpressure or temperature would produce USC steam. Theterm ‘advanced ultrasupercritical’ has been recently usedto delineate more aggressive steam conditions fromthose representative of ‘USC’ although there is no phasechange in the steam.

Coal fired power plants are classified according tosteam conditions and plant efficiency.27 Increasing steamtemperature and pressure is desirable because doing soincreases efficiency (i.e. it reduces the amount of coalrequired to achieve an equivalent amount of energy) andtherefore reduces the emission of carbon dioxide. Table 4shows different types of power plants, their operatingconditions and their efficiencies. The subcritical powerplants typically operate with a steam temperature of,540uC and a pressure of 16?5 MPa, resulting in athermal efficiency of 35%. An increase in temperature andpressure moves the operating conditions into the SCrange. Typical operating conditions for SC power plantsare a steam temperature of 565 to 580uC and a pressure of24?8 MPa, resulting in an efficiency of 38%. The USCplants have an operating temperature in the range of 593to 620uC and a pressure .24?8 MPa. These state of theart plants operate at an efficiency .42%. Future AUSCplants will operate at temperatures .700uC and atpressures .27?5 MPa, resulting in efficiencies of .45%.

Candidate ferritic alloys for AUSC boilersTo achieve the efficiencies given in Table 4, AUSCpower plants will require the use of materials with high

creep strength, high corrosion resistance and adequatefabricability. In the opinion of materials experts, thematerials required for constructing an AUSC power plantare commercially available.25 An assessment of weldabi-lity is of paramount importance and necessary for theimplementation of candidate materials. In a US pro-gramme during the last decade, the Electric PowerResearch Institute (EPRI) has been leading an effort toapply existing materials and to develop new materials forAUSC boilers. Under the programme, the US Departmentof Energy and the Ohio Coal Development Office havefunded a consortium of major US boiler manufacturers toevaluate and develop materials for advanced steam cyclecoal fired power plants.27,28 The programme has set targetsteam temperature of .760uC and a plant efficiency of.45%. Similar programmes to develop the necessarytechnology to achieve AUSC conditions of 700uC and aplant efficiency .45% are under way in Europe, Japan,China and India.28–31 Regardless of the particular devel-opment effort, the singular goal of the worldwide interestin AUSC conditions is to increase plant efficiency throughthe utilisation of existing materials fabrication routes and/or development of new materials.

In advancing the AUSC fossil power plant technol-ogy, the first task was to identify and/or developmaterials with adequate high temperature strength andcorrosion resistance to withstand the operating condi-tions of the AUSC boiler components. The alloysselected for making AUSC boiler components such asheaders and tubes have to satisfy the high temperaturecreep strength requirement and must have excellentresistance to fireside corrosion and high temperaturesteam oxidation. The materials should also have therequired long term strength at the desired temperatureand pressure. Selection of materials for the headers andpiping may be based on the criterion of the temperatureto cause rupture in 105 h at 100 MPa. Figure 4 shows

Table 4 Classification of coal fired power plants9

Nomenclature Conditions (main steam/hot reheat) Net plant efficiency/HHV

Subcritical 16?5 MPa (2400 lb in22) 35565uC (1050uF)/565uC/1050uF

SC §24?8 MPa (3600 lb in22) 38565 (1050uF)/579uC (1075uF)

USC §24?8 MPa (3600 lb in22) §42593uC (1100uF)/620uC (1150uF)

AUSC 27?6–34?5 MPa (4000–5000 lb in22) §45704–760uC (1300–1400uF)

4 Average temperature for rupture in 105 h for various

boiler materials32

3 Phase diagram for water system26





the temperatures for various AUSC alloys;32,33 Fig. 5shows allowable stress derived by taking two-thirds ofthe extrapolated 105 creep rupture strength for varioussteels.32,34

Until a decade ago, power plants depended on lowalloy ferritic steels for metal temperatures to ,580uC. Attemperatures exceeding that threshold, oxidation resis-tance was not adequate for low chromium Cr–Momaterials. Therefore, austenitic steels were used withtemperatures .580uC. Because of their higher coefficientof thermal expansion and low thermal conductivity,austenitic stainless steels are susceptible to thermalfatigue problems in headers, main steam piping, valvesand other thick section components. Furthermore,compared with ferritic materials, austenitic steels are alsocostly. For those reasons, with a few exceptions, the useof austenitic stainless steels has been limited to tubing.

A number of alloys have been considered for heavysection components such as pipes and headers. TheCSEFs, a new family of Cr–Mo steels, have beendeveloped to extend the operating temperature range to620uC. A number of 2?25 wt-%Cr and 9–12 wt-%Crsteels have been developed to fill the needs of current

USC coal fired power plants. Alloy P(T)91 has been inthe American Society of Mechanical Engineers (ASME)code for nearly three decades, and its first commercialuse was in the late 1980s. It was developed by Oak RidgeNational Laboratory and Combustion Engineering inthe 1970s35 for applications in nuclear pressure vesselsfor the fast breeder reactor programme. Although theintended application never came to fruition, P(T)91 wasapproved by ASTM and ASME in 1984 and sawprimary interest in the fossil fired power plant industryfor utilisation in tubing, piping and headers. AlloyP(T)92, a modification of P(T)91, was produced bysubstituting W for Mo and by the addition of controlledamounts of B and N. P(T)92 alloys are actively beingutilised up to 620uC. Ultimately, the use of 9 wt-%CrCSEF steels may be limited by oxidation resistance. Toaddress that problem, a new 12 wt-%Cr steel has beendeveloped and is being commercialised. That material,called VM12SHC, is limited to a wall thickness of12?7 mm (0?50 in.) and has an advertised strengthequivalent to P(T)91. Other CSEF steels have beendeveloped to address a wide range of issues, includingP(T)122, E911, P(T)23 and T24 (see Table 1). Figure 6shows the evolution of these steels.1,32

Elements such as W, V, Nb, B and N have been addedto mainstay Cr–Mo steels to improve high temperaturecreep strength. The modification of the base alloy hasresulted in the development of several new advancedferritic steels that have a tempered bainitic microstruc-ture [T(P)23 and T24] or martensitic lath structure[P(T)91, P(T)92, P(T)122 and VM12]. The bainitic ormartensitic lath microstructure is further strengthenedby a combination of M23C6 type carbides, MX typecarbonitride precipitates and Mo and W solute atoms insolid solution.2

Creep strength enhanced ferritic steels have excellenthigh temperature creep strength, high thermal conduc-tivity and a low coefficient of thermal expansion. Theyare also less expensive than austenitic stainless steels andthus are looked at favourably and are used extensivelyworldwide as the material for pipes and headers.25 At

5 Allowable stress for various alloys comparing tempera-

ture capabilities32,34

6 Evolution of ferritic steels1,32





a given temperature, pressure and allowable stress, areduction in pipe wall thickness of .2 : 1 is possible ifP(T)91 is used instead of 2?25Cr–1Mo material. Figure 7shows relative wall thickness for P(T)91, P(T)92 andother ferritic alloys.36 For the same design conditions,less thermal stress will develop in thinner walled pipingand tubing made of P(T)91. Reduced wall thickness alsoallows for reduced joint thickness and hence less weldmetal deposited and less time to weld. All of thesefactors increase productivity. Although these alloys haveexcellent properties required for the USC power plantapplications, failures have been reported recently inpower plants after a few years in service in the HAZ ofweldments.37 This premature failure is known as type IVcracking. One of the main reasons for this failure is thatCSEF steels, when exposed to fabrication or repairprocedures, are not capable of regaining the originalmicrostructure that is responsible for their excellentcreep properties.38 The new alloys are being utilised insteam applications up to 620uC and/or pressure up to27?6 MPa.39

The basic P(T)91 steel has a normalised and temperedmartensitic structure. The creep strength of the steel ismainly derived from the solid solution effects and fromthe precipitation of stable carbides. The principles bywhich high temperature, high strength steel can bedesigned are the following:3

(i) precipitation hardening by alloy additions of W,Nb, Ti, C and N

(ii) solid solution hardening by additions of W andMo

(iii) dislocation hardening

(iv) hardening due to the presence of boundariesand sub-boundaries that impede dislocationmotion.

The strengthening of Cr–Mo steels is mainly due to amicrostructure consisting of high dislocation density andprecipitation of M23C6 carbides and MX carbonitridesalong the sub-boundaries within the grains.

Figure 8 shows micrographs produced by scanningelectron microscopy and transmission electron micro-scopy of a normalised and tempered P(T)91 steel.40

Figure 8a shows a scanning electron microscopy (SEM)image of prior austenite grain boundaries (PAGBs), andhigher magnification reveals well defined laths oftempered martensite structure with the carbide precipi-tates visible. During long creep exposure (thousands ofhours at 600uC), the tempered martensitic structurechanges by precipitation of a Laves phase Fe2 (W, Mo)and a Z phase that is a complex nitride of the type Cr(Nb, V) N. The concentration of W and Mo in the alloydetermines the amount of Laves phase that forms. Theprecipitation of Laves phase and Z phase causes the lossof long term creep strength of the base metal. This is dueto the precipitation of Laves phase, which reduces theamount of matrix W and Mo and thus reduces theirsolid solution strengthening effects. Large particles of Zphase also form at the expense of finely dispersed VN, aphenomenon that has been largely associated with thedegradation of creep properties of 10 to 12 wt-%CrCSEF steels such as grade 122. This results in theformation of a VN free zone around the carbide. Alsoduring long time creep exposure, the carbides and thecarbonitrides coarsen along the PAGBs by the Ostwaldcoarsening mechanism. All of these changes lead to theloss of creep strength of the alloy after long term creepexposure.

Most of the improvements in strength of ferritic alloyshave come from adding or subtracting elements to thebase composition and manipulating the microstructure.Recently, Abe2 has shown that it is possible to producesteels containing nanoscale VN that have better proper-ties than P(T)91 by reducing carbon levels in 9 wt-%Cralloys to 0?002 wt-%. He has also shown that theaddition of boron to conventional 9 to 12 wt-%Cr alloysstabilises the M23C6 carbide, thus preventing coarseningof the carbide and the degradation in the properties.

The normalising temperature for P(T)91 alloy isbetween 1040 and 1080uC. Normalising is carried out

7 Relative wall thickness for piping and tubing made of

P(T)91, P22 and other ferritic steels for same design

conditions36

8 Images (a SEM and b TEM) of normalised and tempered P(T)91 steel40 (with permission of Springer Science)





to take carbides and carbonitrides into solution com-pletely. The material is then cooled from the normalisingtemperature to room temperature to produce a fullymartensitic structure. It is important for all of theaustenite to be transformed to martensite. Otherwise,any retained austenite will have an effect on theproperties. Table 5 shows transformation temperatures,normalisation and PWHT ranges for CSEF steels.32

After normalisation, the fully martensitic structure inthe alloy is tempered. Ductility and toughness increaseduring tempering. Care must be exercised in all heattreatments for these steels. In particular, the temperingtemperature for P(T)91 must not exceed the A1

temperature. If it does exceed the temperature, austenitewill reform, which, upon cooling, will transform tountempered martensite (with a high cooling rate) or softferrite (with a low cooling rate). Retention of either oneof these transformation products in the material isdamaging to the toughness and creep properties.41

Furthermore, at temperatures near the A1 or after theaccumulation of many successive heat treatments overthe life of a component or improper processing, over-tempering will occur and will result in degradationthrough precipitate coarsening and recovery at thesubgrain boundaries. This will further degrade the creepstrength of the material.

The chemical composition of the alloy determines themaximum temperature for tempering and ferrite forma-tion. Recently, it was pointed out that the ASTMchemical composition specification may be wide enoughto exceed the critical temperature A1 duringPWHT.38,42,43 Figure 9 shows graphically how the lowertransformation temperature A1 could vary for sixdifferent 9 to 12 wt-%Cr alloys within the composition

range specified by ASTM. Chalk et al.43 have investi-gated the effect of variations in composition on the A1

temperature of P(T) 92 steel. Using thermodynamicmodelling, they have developed equations to predict theA1 temperature for welding consumables as a functionof composition.

Knowing the microstructure of a given steel isimportant in understanding the influence of a weldthermal cycle on the material and ultimately itsperformance. The tempered martensitic microstructureof P(T)91 base metal changes drastically when thematerial is exposed to a weld thermal cycle. The weldthermal cycle alters the microstructure significantly,resulting in a weldment with a large compositiongradient in the FZ due to solidification as well asregions in the HAZ containing coarse grained and finegrained microstructures. The HAZ is also associatedwith a soft zone. The resulting microstructure in theHAZ has been the source of problems and is the primaryreason for the structural component failure. Thepremature type IV failure in the HAZ is unique toCSEF steels and thus needs special care and attention.44

Welding and weldability of AUSCcandidate ferritic alloysConsideration of the appropriate welding technologies(i.e. filler metal, welding process, PWHT method)required during the fabrication and field constructionof fossil fired power plants is critical because arepresentative plant will contain millions of pounds ofmaterial and will need tens of thousands of welds to fullyerect. Welds may be required between similar and/ordissimilar materials. The process used may be a single

Table 5 Transformation temperatures in uC, Normalization and PWHT ranges for CSEF steels*32

Material MS MF BS A1 A3 Normalisation Tempering/PWHT

T23 y540 y325 y620 800–820 960–990 1040–1080 730–800/720–800 (PWHT)T24 y460 y260 y550 815 950 980–1020 720–775P(T)91{

Base (Alex.) 372–393 159–196 N/A 788–813 833–860 1050–1080 730–780/730–770 (PWHT)Base (V&M) 372–393 159–196 N/A 800–830 890–940 1050–1080 730–780/730–770 (PWHT)Weld 390–418 200–237 N/A 770–801 832–861 1050–1080 730–780/730–770 (PWHT)

P(T)92{Base (Alex.) 366–396 135–198 N/A 795–817 854–867 1040–1080 730–800Base (V&M) 366–396 135–198 N/A 795–835 900–925 1040–1080 730–800Weld 376–423 185–229 N/A 759–800 832–861 1040–1080 730–800

E911 360 160 N/A y805 y980 1040–1080 740–780/730–800 (PWHT)VM12 300 175 N/A y810–825 y890 1040–1080 750–800

*MS: martensite start temperature; MF: martensite finish temperature; BS: bainite start temperature; A1: lower critical transformationtemperature; A3: upper critical transformation temperature.{VM: Vallourec and Mannesmann; Alex.: Alexandrov.

9 Graph representation of possible ranges of A1 tempera-

ture for CSEF steels38

10 Combination of materials and welding processes

involved in fabrication of stream boiler component45





process, or multiple processes may be applied, depend-ing on the requirements for fabrication or construction.Figure 10 shows the combination of materials used andthe fabrication techniques and joining processes involvedin the fabrication of a steam boiler component.45

Utilisation of a particular material in a fossil firedplant requires extensive characterisation. Weldabilityevaluation of the materials for boilers, pipes and tubeshas been studied using product forms representative ofactual components. For example, in addition to creepstrength requirements for the base metal, weldmentstrength is also a determining factor for materialselection for boilers, tubes and pipes. Since the creepstrength of the weld metal is less than that of the basemetal, it is recommended that a ‘weld strength reductionfactor’ used in the design of high temperature nuclearpower plants be used in designing fossil power plantboiler components.46,47 The term ‘weld strength reduc-tion factor’ is defined as the ratio of the stress to rupturefor weld metal to the stress to rupture for boilermaterials for a given time and temperature. The ratiois particularly important in the construction of longseam welded piping and may also be important inregions of a component or system where a weld may beinfluenced by large bending stresses.

Welding of alloy P(T)91 and other CSEF steels canappear to be very difficult compared with welding ofother steels, but it is relatively straightforward if thewelding procedures and specifications are followed.Creep strength enhanced ferritic steels are highlyhardenable, and upon cooling after welding, theyundergo phase transformation that results in a fullymartensitic structure. To obtain the desired properties,the welds must undergo PWHT, and an interpasstemperature must be maintained.41 It is vital to ensurethat CSEF steels are not treated like their lower alloyCr–Mo cousins. Many failures in the fossil fired powerplants are attributable to faulty heat treatments thatresulted from the lack of training and supervision.Therefore, adequate training and supervision are criticalto obtain the properties that can be realised with theCSEF class of steels. Because of the increased use ofCSEF steels in current power plants and their antici-pated use in future power plants, the American WeldingSociety (AWS) and ASME are defining the codes andguidelines for welding CSEF steel, piping and tubes,

either by writing new documents or by revisiting existingdocuments (D10?10, D10?21; ASME’s VIII, B31?1, andB31?3).41

Filler materials for AUSC power plantsCreep strength enhanced ferritic alloys can be imple-mented in coal fired AUSC plants if they can be weldedsuccessfully. As indicated in Table 2, CSEF steels arebeing proposed in atypical applications, such as the useof T91/T92 in waterwall panels. Many of the establishedwelding processes can be used to weld alloy P(T)91successfully.20 However, development of consumablesthat provide the required weld metal creep, corrosionand toughness properties is a major concern. Manyconsumables are available to weld P(T)91, and theyare made to AWS specifications.48–52 Tables 6 showsvarious consumables available for welding CSEFsteels.20,44 The composition of a given CSEF fillermaterial is generally developed to match the base metalcomposition. Filler metal composition should bebalanced to avoid ferrite formation and to ensure afully martensitic structure in the weld metal. Presence offerrite in the weld metal leads to poor toughness andcreep properties. Work has shown that the filler metalwith reduced Nb, N and Si improved toughness of theweld metal.53 Having investigated some experimentalalloys, Sireesha et al.54 have shown that the Si and Nbcontent in the consumable to weld P(T)91 should belower than that in the base metal (higher levels of Si andNb promote possible retention of ferrite and reduce weldmetal toughness). The Si should be in the range 0?15 to0?25 wt-%. Nb should be in the range 0?03 to 0?05 wt-%.In addition, an increase in the Ni content has beenshown to have a beneficial effect on toughness.48,53,55 Nilowers the A1 temperature and should be maintained inthe range of 0?4–1?0 wt-% to prevent the A1 temperaturefrom falling below the PWHT temperature, leading toaustenite formation, which upon cooling transforms tountempered martensite. Hence, Ni content higher than1 wt-% may promote austenite formation in the weldmetal, and it is detrimental to its toughness properties.Additionally, the weld metal composition should bebalanced such that the weld metal is fully martensiticwithout any austenite. Specific guidelines must befollowed when filler metal is being developed forP(T)91. An extremely well controlled preheat, interpass

Table 6 Examples of welding consumables20,44

Welding process Spec., A/SFA T/P91* T/P23{ T/P24{ T/P921 E911" VM12SHC#

SMAW 5?5 E9015-B9 E9015-GGTAW GMAW 5?28 ER90S-B9 ER90S-G ER110S-GFCAW 5?29 E91T1-B9 N/A E91T1-GM N/A N/ASAW 5?23 EB9 EG

*Note that AWS will formally adopt ‘-B91’ to replace ‘-B9’ in order to be more specific in identifying CSEF filler materials.{Because a ‘G’ designation is specified, the CSEF steel grade must be specified to avoid confusion between materials. In this case,‘grade 23’ may be sufficient until AWS formally adopts a ‘-B23’ designation.{Because a ‘G’ designation is specified, the CSEF steel grade must be specified to avoid confusion between materials. In this case,‘grade 24’ may be sufficient until AWS formally adopts a ‘-B24’ designation.1Because a ‘G’ designation is specified, the CSEF steel grade must be specified to avoid confusion between materials. In this case,‘grade 92’ may be sufficient until AWS formally adopts a ‘-B92’ designation."Because a ‘G’ designation is specified, the CSEF steel grade must be specified to avoid confusion between materials. In this case,‘grade E911’ may be sufficient. There are no formal plans for AWS to adopt a specific 9%Cr designation for E911.#Because a ‘G’ designation is specified, the CSEF steel grade must be specified to avoid confusion between materials. In this case,‘grade VM12SHC’ may be sufficient. There are no formal plans for AWS to adopt a specific 12%Cr designation for VM12SHC or other12%Cr CSEF alloy.



temperature and PWHT are mandatory to ensure thatthe required creep rupture properties and toughness areobtained in the weldment.

During the welding of CSEF steels, it is also importantto preheat the base metal and to maintain an interpasstemperature in order to mitigate hydrogen crackingproblems. A preheat temperature of 150–200uC and aninterpass of y350uC are recommended; the requiredpreheat is material and restraint dependent. Guidelinesregarding preheat and interpass for CSEF steels are givenin Table 7.32 When required by governing code bodies,proper PWHT of all CSEF steels is critical to achieveproper tempering of martensite.56–58 The time at tem-perature for PWHT depends on the thickness of the part.The PWHT temperature should also be controlled so thatthe formation of austenite, which would result in the

formation of untempered martensite, is avoided. For atypical PWHT, the allowable ASME range for holdingtemperature is given in Table 7. If thickness is ,5 in., it is1 h in.21 of hold time; if thickness is .5 in., then it is5 hz15 min per inch of thickness. Figure 11 shows atypical weld thermal cycle experienced by the P(T)91 steelcomponent during welding and PWHT.20 Control overthermal gradients is also critical.48 Existence of a thermalgradient in a heavy section component due to the methodof heating (soak band) during PWHT may cause defectiveheat treatment. In such a case, a non-uniform PWHTmay result if the outside surface temperature reachesPWHT temperature well before the inside of thecomponent does. Practices such as increasing the size ofthe heating band, using properly sized heating pads toprevent large gaps or overlap through the circumferenceof the tube or pipe and using a control thermocoupleunder each heating pad are recommended in addition tothe guidelines specified by AWS D10?10.59

In the bainitic CSEF steels, interbead and layer tolayer tempering can play a particularly important role indetermining the properties of a weldment with respect tohardness, toughness and creep.60,61 The propertiesdepend on a number of variables, including heat input,fluxes, microlloying agents and the interaction ofresidual elements. In a multilayer weld region, varyingmicrostructures can form due to the thermal cycleexperienced by each bead from the subsequent weldpasses. Figure 12 shows macrostructure of a multipassweld and possible complexity of microstructures that

Table 7 Recommended preheat and interpass temperatures for common CSEF steels32

Alloy

Minimum preheat* Maximum interpass ASME allowable PWHT range

uF uC uF uC uC

T/P23 350–400{ 175–200 600 315 720–800T24 400 200 600 315 720–775T/P91 350–400{ 175–200 660 350 730–800T/P92 350–400 175–200 660 350

*The minimum preheat recommendations in this table do not supersede existing code requirements, such as those present in ASMEB31?1 for grade 91. Section 131?4?5, which addressed P5A and P5B (some now classified as P15C/E materials) requires 400uF (200uC)for material that has either a specified minimum tensile strength in excess of 60 000 lb in22 (413?8 MPa) or has both a specifiedminimum chromium content .6% and a thickness at the joint in excess of 0?50 in. (13?0 mm); 300uF (150uC) for all other materialshaving this P number.{If there is concern to have minimal hardness (,350 Hv) without PWHT (to minimise susceptibility to SCC), then a 400uF (204uC)preheat is advised.{Note that welding has been successfully performed using the gas tungsten arc welding (GTAW) process and reduced preheattemperatures.

11 Typical weld thermal cycle for P(T)91 steel component during PWHT20

12 Macrostructure of multipass weldment62 (courtesy of

P. Mayr)





may exist due to thermal effects and associated HAZsthat form during a multipass weld.62 Depending on thesequence in which the weld beads are laid and theproperties of the multiple HAZs (orientation, whetherthey are continuous or discontinuous) could promotetype I cracking within the weld metal. In a multipassweld, the bead shape and the welding position also havea significant effect on the weld properties.48 When theweld deposit is thin and wide, it allows for some degreeof tempering of the previous deposit from the heat ofwelding and promotes grain refinement in the previouslydeposited weld metal.48 Thicker beads limit the degree ofgrain refinement and lead to loss of toughness in theweld metal. As to the welding position, most of theobservations made are on welds produced in the flat(1G) position. A loss of toughness is observed inshielded metal arc welds (SMAW) made in vertical(3G) and overhead (4G) positions due to bead controlproblems and lack of tempering and grain refinement.48

Microstructure development in the FZ depends on thesolidification behaviour of the weld pool. The solidifica-tion behaviour controls the size and shape of the grains,the segregation and distribution of inclusions andporosity. Solidification behaviour and associated segre-gation are also key factors in the hot cracking behav-iour of the weld metal. Sometimes, it is convenientto consider weld metal as a miniature casting. There-fore, parameters important in determining microstruc-ture in casting, such as growth rate, temperaturegradient, undercooling and composition, also determineits character. A comprehensive review of the funda-mentals of weld metal solidification is available in theliterature.63

In the case of CSEF steels, segregation of low meltingresidual elements such as S and P during solidificationcan cause hot cracking (specifically crater cracking) inextreme cases.48 This has been reported in root passwelds when welding the bainitic CSEF steel T24.Segregation of alloying elements (e.g. Cr, C, N, Nb)could further affect the subsequent solid state transfor-mation in the steel. In thick section welding, in which theweld groove requires multiple overlapping weld passes,each subsequent weld pass subjects the earlier pass to anumber of weld thermal cycles, leading to a number ofHAZs within the weld. This is very important in CSEFsteels, which undergo solid state transformation uponheating and cooling. Owing to the short time attemperature, a limited amount of solute redistributionmay occur by diffusion.64

Important aspects of weld pool solidification are weldpool dynamics, development and steady state geometry.Weld pool shape is important in the development ofgrain structure and weld bead shape, size and penetra-tion. In predicting the weld pool shape and penetration,recent theoretical developments include formation of afree surface computational model with a coupledconduction and convection heat transfer model.65,66

The models are capable of predicting weld pool shapeand penetration as well as thermal gradient and coolingrate. A good knowledge of heat and mass transfer willaid in the optimisation of weld bead size and shapecontrol during temper bead procedure development.Solidification of weld metal proceeds in a dendritic orcellular dendritic fashion, depending on the thermal andconstitutional conditions at the interface.

Solidification of CSEF alloys appears straightfor-ward with LRd (ferrite) and the ferrite subsequentlytransforming to austenite (c), which then transforms tomartensite (for T/P91, T/P92 or VM12SHC) or bainite(for T23/24) upon cooling. Owing to the partitioning ofalloying elements between sold and liquid during weldpool solidification of a complex alloy such as P(T)91,extensive solute redistribution of alloying elements occurs.This partitioning of elements, also called segregation,leads to heterogeneity in composition within the weldmetal. Segregation on a fine scale of the order of dendritearm spacing (10 mm) is called microsegregation. Thisaspect of weld metal solidification is described in theliterature.63 When the scale is larger than several dendritespacings, it is called macrosegregation. Local variations incomposition can change the local transformation char-acteristics of the material.

Recent characterisation and modelling of non-equili-brium microstructure evolution during solidification ofP(T)91 alloy have revealed the presence of austenite in thehighly segregated (C and N) interdentritic region ofthe weld metal.67 In contrast, in the base metal, wherethe composition is homogenous after normalisation, theaustenite is completely transformed into martensite. Thisdifference in the microstructure was attributed to micro-segregation in P(T)91 during solidification. Although thesegregation effects are present, it is difficult to reveal theweld metal solidification substructure or the dendriticstructure, mainly because the weld metal goes throughsubsequent martensitic transformation after solidification,masking the solidification substructure. It is the variationin elemental concentrations and its reaction with a selectedetchant that reveal the dendritic pattern during metallo-graphic examination of the weldment structure. Theunderlying dendritic microstructure has been revealed byapplying the special metallographic technique of tintetching. Figure 13 shows the dendritic solidificationsubstructure in P(T)91 weld metal after tint etching.Solidification simulation was carried out using the Schiel–Gulliver model64,68 and Thermocalc.69 The analysis showsthat the first phase to solidify is d-ferrite (Fig. 1) and thattoward the end of solidification, the interdendritic regionswere enriched in elements C and N. The enrichment by Cand N leads to the stability of austenite c in theinterdendritic region. The variations in Cr concentrationin the solid phase were predicted using Schiel–Gullivermethod and are shown in Fig. 14. The calculations show

13 Dendritic solidification substructure observed in

P(T)91 after tint etching67




the Cr level from the core of the dendrite to theinterdendritic region varying from 8?9 to 11?3 wt-% andcorresponding C level varying from 0?02 to 0?4 wt-%.Based on these composition levels, the weld metal regionshould show a large temperature range over which thetransformation from austenite to martensite may occur.This indicates possible retention of austenite in the weldmetal, leading to an additional martensitic transformationduring cooling of the weld metal from the temperingtemperature. The impact of microsegregation on thetransformation behaviour was confirmed using thermo-mechanical simulations and time resolved X-ray diffrac-tion experiments.67

Welding processesSeveral welding processes are used to weld CSEF steels.Processes such as SMAW, submerged arc welding(SAW), gas tungsten arc welding (GTAW) and fluxcored arc welding (FCAW) are commonly used.48,70 Theprocess selected may vary depending on the specifics foreach weld [e.g. the number of welds to be made, therequired deposition rate, the thickness of the material,access to the weld joint and the welding position (flat,overhead or vertical)]. Care should be taken to see thatthe weld metal contains very low concentrations oftramp elements such as As, Sn, Sb, Pb, S and P. If S or Pis at a maximum of 0?010%, crater cracking and grainboundary cracking can be avoided.70 The variousprocesses used to weld 9–12 wt-%Cr steels are alsodescribed in the literature.48,70 Although it is a slowprocess, GTAW produces the highest quality welds. Theweld metal produced by GTAW has been found to havea higher toughness than the welds produced by fluxbased processes. Weld quality using the gas metal arcwelding (GMAW) process can be difficult to consistentlyreproduce, and only a few fabricators are qualified touse the process. Flux cored arc welding is a highdeposition rate process exceeding all other processesexcept for SAW. The use of the SMAW process to weldCSEF steels is well established and is the mostcommonly utilised process. The SAW process, used forthick section welding and in seam welding waterwalltubes to membranes, can be automatic or semiauto-matic. It is a high deposition rate process, and the beadshape and size should be controlled to achieve acomplete interbead tempering. In general, the beadthickness for SAW and FCAW should be no greater

than 3?8 mm. This allows for tempering of thepreviously deposited layer.48

Basic flux is recommended for SAW of P(T)91.70

Other fluxes will reduce the C and increase O and N inthe weld metal, affecting the strength level. For weldingprocesses in which elements may be lost across the arc(such as in GMAW, SMAW or SAW), maintaining theweld metal composition control is critical for all weldingprocesses. In particular, the levels of Nb, Ni, C and Nmust be controlled to result in good creep strength in theweld metal. For grade 91 steel, the weld metal shouldcontain a minimum of 0?09 wt-%C, 0?02 wt-%Nb and0?02 wt-%N.70 When Ti is present in the weld metal,lower N levels can be tolerated. However, Ti contentshould not exceed 0?01 wt-%, or TiN will form, and thusN, which is essential for high temperature strength, willbe depleted in the matrix.

The composition of the weld metal should becontrolled to maintain the A1 temperature so thatPWHT can be carried out safely without increasing thepossibility of forming retained austenite. For example,the NizMn level should not exceed 1?0 wt-%. NizMn inT/P91 weld metal at NizMn contents .1?0 wt-% maysuppress the A1 temperature ,790uC (the maximumallowable PWHT temperature), and narrow the range inwhich PWHT can be performed safely.71 Alloy P(T)91 isa highly hardenable alloy and is subject to hydrogencracking. In order to overcome this potential problem, alow hydrogen electrode (e.g. with a -H4 designation)should be used.70 The SAW and FCAW wires should beused with a -H4 designation and should be stored in aheated chamber. Efforts are under way to use computa-tion modelling for optimising the process parameters.Vasudevan et al.72 have used a genetic algorithm (GA) forthe optimisation of the A-TiG [tungsten inert gas (TIG)welding is an unconventional term used for GTAWprocess] process to weld modified 9Cr–1Mo alloys. A-TiG welding is same as GTAW process where anactivated flux is used on the surface of the steel to bewelded to improve the depth of weld penetration.73,74 Thechange in weld penetration is mainly due to the change inweld pool dynamics and the arc plasma characteristicsbrought about by surface active elements in the activatedflux.65,66 Genetic algorithm is a family of computationalmodels based on evolutionary strategies. It is a field ofartificial intelligence often used to generate usefulsolution to optimisation and search problems. A non-linear and complex process such as welding is ideallysuited for optimisation of the process using GA to achievedesired weld quality. The authors have used a GA basedmodel to determine the weld bead shape, weldingparameters, depth and penetration and HAZ width.The model has been used successfully to generateoptimised process parameters that agree well withobserved values of weld bead geometry and HAZ width.

A potential new process that could join CSEF steelswhile maintaining their creep strength, fatigue resistanceand corrosion properties is friction stir welding (FSW).It is a new and innovative process developed in the 1990sat The Welding Institute in the UK.75 The processdetails are shown schematically in Fig. 15 and aredescribed in detail in the literature.76

Friction stir welding involves plunging a rotating toolinto the material to be joined and translating the toolalong the joint line. Owing to the frictional heat

14 Calculated variations in chromium and carbon con-

centrations from core of dendrite to interdendritic

region67




generated by the rotating tool and the metal, the metal isplasticised and a metallurgical bond is created. Frictionstir welding, a solid state, low heat input process, couldbe a solution to the problems encountered by the Cr–Mosteels during welding, in particular, type IV cracking.Thomas et al.77 have demonstrated the feasibility ofFSW to join 12 wt-%Cr steels with parent metalproperties such as cross-weld tensile properties andbend tests. The FSW process offers numerous advan-tages for most materials; for example, it is an energyefficient solid state process, it does not use shielding gasor filler metal, the parent material composition isretained, steel thickness up to 25 mm can be weldedfrom two sides and the weldment undergoes very little orno distortion. At the same time, it has severaldisadvantages. For steels and other high temperaturematerials, tool wear is a major problem. Welding travelspeed is not comparable to conventional processes andthus not very economical. Bhadeshia and DebRoy78

have highlighted a number of issues and problemsassociated with the process.

Weldability issues in CSEF alloys

Type IV crackingFailures due to cracking in weldments are classified fromtype I to IV, depending on the location of the crack inthe weldment.79 Figure 16 schematically shows varioustypes of cracking. Type I cracking (e.g. solidificationcracks) occurs and remains in the weld metal. Type IIoriginates in the weld metal and can extend into theHAZ or base metal. Type III is located in the HAZ, inthe coarse grain region; an example would be reheat

cracking. Type IV is in the intercritical HAZ (ICHAZ)region.

The precise location of failures in the HAZ attributedto type IV cracking is the question and subject of manystudies and has yet to be defined. A well designedmicrostructure and stress simulation and characterisa-tion using neutron scattering may shed more light onthis issue. The effect of type IV failure with regard tolong term properties is shown in Fig. 1780 and comparedto base material using Keith Bell’s analysis of type IVfailures. The data for grade 91 and type IV failures areplotted using a Larson–Miller parameter as a functionof stress. The use of the Keith Bell – 20% line in Fig. 17bounds the minimum for type IV failure.81 An exampleof a type IV failure in a long term cross-weld creep testin E911 is shown in Figs. 18 and 19.82 In Fig. 18, thedamage and failure are clearly in the HAZ, specifically ina region consistent with either the fine grain HAZ(FGHAZ) or the ICHAZ (i.e. type IV failure). Thisobservation is further supported by the electron back-scatter diffraction image shown in Fig. 19, where themaximum damage (white areas absent of grain orienta-tion) is clearly concentrated in an FGHAZ. In addition,

15 Schematic of friction welding process showing inter-

action of tool pin with material76 (courtesy of The

Welding Institute)

16 Classification and types of cracking in weldments79

17 Grade 91 base material creep rupture comparison to

Keith Bell type IV failure analysis80

18 Example of type IV failure location in a cross-weld

creep specimen from E911 after 13 945 h at 600uC and

120 MPa82

19 Example of type IV failure, location and electron back-

scatter diffraction image across HAZ; creep damage

indicated by white regions in colour map82








the failure is characterised by low ductility, with thehighest degree of damage occurring subsurface near themidwall, as shown by comparison with the oppositeHAZ, which remains intact. These characteristics in thissimulated large specimen creep test are characteristic ofthose encountered in service.

The 9–12 wt-%Cr CSEF steels are prone to type IVcracking during long time service. The susceptibility ofcross-weld bainitic CSEF steel to type IV failure is lessclearly defined and is not as well researched.83,84 Themain reason for a type IV failure is believed to be that,once exposed to the weld thermal cycle, the alloys arenot able to retain their original microstructure, which iswhat gives them their excellent creep properties. In orderto understand this phenomenon, one has to very closelylook at the complex microstructure that develops in theHAZ of CSEF steels.

The microstructure development in the HAZ of thewelded CSEF steels P(T)91, P(T)92, P(T)122 steels andothers has been discussed by many authors.85–88

Figure 1 shows schematically the details of the micro-structure development in the HAZ during welding as itrelates to the phase diagram for P(T)91 steel. As shownin Fig. 1, the HAZ is made up of different zones withunique microstructures, depending on the peak tem-peratures Tp experienced by various regions during theweld thermal cycle. The different zones are as follows:

(i) a coarse grained HAZ (CGHAZ) next to thefusion line, where Tp reaches much higher thanA3

(ii) an FGHAZ, where Tp reaches just above A3

(iii) an ICHAZ, where Tp is between A1 and A3

(iv) an overtempered region, where Tp is below A1.

Figure 20 shows a hardness traverse across a P(T)91weldment in three conditions: as welded, PWHT andcreep tested.86 Figure 21 shows hardness map across anex service grade 91 weldment.32 The hardness profileshows a dip in hardness in the HAZ, indicating a softzone. In Fig. 21, the dark blue region in the centre of thethickness and white area near the top of the figureindicate softening consistent with the observations madein the ICHAZ. The location of the soft zone is still openfor discussion because of the complex microstructuralfeatures make it difficult to find its exact location. It canbe anywhere in the overtempered region, the intercriticalregion or the fine grain region. The soft zone in the HAZregion is the weakest link in the weldment. Premature

creep failure (type IV cracking) has been observed tooriginate in that region and is common to all CSEFalloys.

The evolution of the microstructures in these regionscan be rationalised based on the transformation thattakes place in each zone and on the stability of theprecipitates in the HAZ. In the CGHAZ, the peaktemperature is high enough to dissolve the carbideprecipitates that effectively pin the austenite grainboundary. In the absence of these precipitates, theaustenite grains can grow. In the FGHAZ, the peaktemperature is just above A3, and the carbide precipitatedissolution is incomplete, leaving the undissolvedparticles to stifle the growth of austenite grains in theFGHAZ. In addition, a partial reversion of themicrostructure to austenite occurs, as shown in Fig. 1.On cooling, the austenite transforms to untemperedmartensite. The temperature in the ICHAZ is not highenough to dissolve the carbides, but the temperaturemight be sufficient to coarsen them by Ostwald ripening.The austenite that forms in the ICHAZ is a low carbonaustenite. On subsequent tempering at temperaturesbelow A1, no secondary precipitation in the martensite isobtained from the low carbon austenite or from theovertempered ferrite. The microstructure in this regionundergoes enhanced recovery to form a soft zone. Theextent of these various zones in the HAZ depends on thethermal gradient in the HAZ. The thermal gradient isinfluenced by variables such as the material propertiesand heat input.

During welding, the region of the HAZ experiencingthe intercritical temperature range undergoes a changein the microstructure of the original base metal.

The original normalised and tempered base metalmicrostructure, consisting of high dislocation densityM23C6 carbide and carbonitride precipitates along thesubgrain boundaries, is replaced by large subgrains,coarsening of the M23C6 carbides, and a change in theshape of the carbonitrides from needles to spheres.

The main problem that needs to be solved to addresstype IV cracking is to determine which of the zones in theweldment is the weakest and to locate the place where thetype IV cracking occurs. To get an understanding of

20 Microhardness profile across P(T)91 weldment in as

welded, PWHT and creep tested conditions86 (with

permission of Springer Science)21 Hardness across ex service Grade 91w. Dark blue

through centre of thickness and white areas near top

of figure indicate softening, consistent with observa-

tions made in ICHAZ. Representative weld was

exposed at 580uC for y75 k h32





the phenomenon, it would be helpful if the properties ofeach zone with its unique microstructural features couldbe determined.

It is difficult to determine the properties of each zonebecause the HAZ in a weldment is narrow. To overcomethis problem, various investigators have carried outcreep rupture studies on simulated HAZ specimens,89,90

cross-weld specimen testing and small specimen creeptesting of localised regions in the HAZ.91–93 Simulatedspecimens were made by isothermally heating largespecimens to the required peak temperatures to repro-duce microstructures of the various zones in eachspecimen. The samples were also made by thermo-mechanical simulator (Gleeble) by subjecting the sam-ples to a predetermined thermal cycle. The creep resultswere useful but did not enable the researchers to resolvethe issue of which zone of the HAZ is responsible fortype IV cracking. Laha et al.86 have shown by theirinvestigation involving similar regions of HAZ micro-structures that the soft zone observed in the grade 91steels during welding is located in the ICHAZ. Othercreep rupture studies on cross-weld specimens show thatthe type IV cracking occurs in the FGHAZ, which is notthe same region as the soft zone. Finite element analysisof stress in a cross-weld specimen during creep testingpredicts a triaxial stress state in the FGHAZ, which willaccelerate creep void formation that could lead to typeIV cracking.85 The stress state in the specimen could bealtered by changing the width of the HAZ or weldgroove angle.94 Doing so can alter the creep behaviourof the weld. An additional finite element analysis modelof the cross-weld specimen also places the origin of typeIV cracking in the FGHAZ by allowing for grainboundary sliding.95 Such grain boundary sliding relaxesthe constraint from the surrounding heterogeneousmicrostructure, thus promoting failure in the FGHAZregion.

Although the above analysis places the type IVcracking in the FGHAZ, Laha et al.86,87 came to adifferent conclusion and place the type IV failure in theICHAZ. Within a Cr–Mo steel weld HAZ, each zone issurrounded by regions that have different creep flowbehaviour, and because of this inhomogeneity, a triaxialstress distribution develops across the weldment duringcreep testing. Using finite element modelling analysis,stress accumulation has been predicted in both theFGHAZ and the ICHAZ.85,96,97 The triaxial state ofstress that develops in the specimen is the main reasonfor the development of creep cavitation in the ICHAZ,which leads to type IV cracking. This is furtherconfirmed by the metallurgical observation of morefrequently formed creep cavitation in the ICHAZ.86 Inaddition, metallurgical features such as coarse carbideprecipitates in the ICHAZ may increase the possibilityof creep void nucleation.

Creep cavities have always been found to beassociated with large carbide particles. This has beenattributed to the higher stress concentration at theinterface between the carbide particles and the softmatrix in the ICHAZ region,96 leading to nucleation ofcreep cavities. Such an event and coalescence of cavitiesinto cracks during deformation aided by the triaxialstress in the ICHAZ leads to type IV cracking.

Abe et al.98 have investigated creep rupture propertiesof Gr 92, Gr 92N, and 9Cr–B (9 wt-%Cr–3 wt-%W–3

wt-%Co–0?02 wt-%V–0?05 wt-%Nb) steels using simu-lated HAZ specimens. Grade 92 exhibited a significantdecrease in time to rupture after a thermal cycle to a peaktemperature close to A3; the other two steels showedproperties that were the same as those of the base metaland no type IV cracking in welds. Characterisation of themicrostructure revealed that the Gr 92 steel after thermalcycle had a fine grain structure with a very littleprecipitation of M23C6 carbide along the PAGBs andsub-boundaries, whereas the other two steels had a grainsize similar to that of the base metal and extensiveprecipitation of carbides along the PAGBs and also alongthe lath and block boundaries. The authors attributed thedegradation in creep rupture properties in Gr 92 afterthermal cycle to A3 temperature to reduced precipitationof carbides and reduction in grain boundary and sub-boundary hardening, not to the fine grain structure.

The assessment of type IV cracking in different regionsof the HAZ has been investigated through the use of smallpunch creep testing, whereby discs measuring 3 mm indiameter and 0?35 mm in thickness were removed fromspecific regions of the electron beam weld HAZ in a 8Cr–2W–VTa experimental CSEF steel.92 In that study, discstaken from the tempered HAZ, FGHAZ, CGHAZ,fusion line and base metal were compared. The creep lifeof the FGHAZ was markedly lower than those of theother regions. A disc taken from the ICHAZ was notspecifically identified, and it is not clear whether theauthors used the tempered HAZ as a synonymous termwith the FGHAZ. In addition, the specimens wereextracted from a very narrow HAZ of an electron beamweld, and it is unclear by how much the specimens wererepresentative of the various zones of the HAZ. However,the small punch creep testing method has great potentialfor testing and understanding creep behaviour of HAZswith microstructural gradients. A similar study93 com-pared the HAZ of the base metal and the HAZ of theweld metal with specimens extracted from the weld metaland base material. The HAZ of the base metal showed apronounced decrease in life as compared to otherexamined constituents. These results offer an intriguingpossibility in the examination of the comparative lifebetween HAZ regions by testing small specimensextracted from real weldments.

Mitigation of type IV crackingCreep strength enhanced ferritic steel welds have a verycomplex HAZ that results in a decrease in creep rupturestrength and failure of the weldment during service. Thisis a major concern for the utilities. Premature failure isdue to the influence of the heterogeneity in the HAZmicrostructure, a triaxial stress state in service andthe accumulation of residual stress during welding. Inparticular, there is a concentration of stress and coarsen-ing of carbide in the ICHAZ and FGHAZ, leading tonucleation and growth of voids that lead to type IVcracking and premature failure. It is yet to be determinedwhether the crack initiates in the FGHAZ or ICHAZ. Anexample of a type IV failure in a 9 wt-%Cr alloy (in thisinstance E911) is shown in Fig. 18.

An approach that could eliminate type IV cracking inthese alloys and make them very useful would be to subjectthe whole weldment after welding to a normalisation andtempering treatment and to eliminate the complex HAZmicrostructures and stresses. However, that approachwould be impractical and costly. Therefore, a number of



investigators have focused on increasing the crackingresistance of the alloys.94,97,99–101 The methodologies toaddress resistance to cracking have been summarised byLaha et al.87 The cracking resistance of the steel has beenfound to increase by alloying and modifying the steel withC and B and by altering the carbide precipitation andstability. In addition, efforts have been made to alter thestress state of the weldment by joint design, weld processand varying HAZ width by changing the heat input.94

Several investigators102–104 have studied the effect ofwelding parameters on type IV cracking. Both jointgeometry and preheat have been found to improve creeplife significantly.

Boron has been found to be an effective agent inincreasing the cracking resistance of the alloy. Kondoet al.99 have investigated Cr steels (9Cr–3W–3Co–VNbB)with boron levels of 47 to 180 ppm and with low Nlevels. The N level should be controlled carefully so thatBN does not form and reduce the amount of B availablefor stabilising the carbides. A creep rupture test revealedthat a weld in high Cr steel containing B had superiorproperties compared with a weld in the same steel butwithout B. In addition, there was no evidence of type IVcracking in the steel. This improvement in properties andincrease in cracking resistance have been found to be dueto the absence of FGHAZ in weldments of B containingsteel, unlike weldments in high Cr steel that does notcontain B. In addition, B effectively suppresses theM23C6 carbide coarsening by replacing some of the Cwith B in the carbide.105,106 Figure 22 shows a transmis-sion electron microscopy (TEM) micrograph of nor-malised and tempered P(T)91 steel with and withoutthe addition of B.107 The steel without B has amuch coarser M23C6 carbide precipitate than the steeladded B. The C content of the alloy has been found toinfluence its creep properties and its susceptibility totype IV cracking.

Hirata et al.108 have found a correlation between Clevel and the resistance of the steel to cracking. Theyinvestigated 9Cr–3Co–3W–VNb alloy with C rangingfrom 0?005 wt-% to 0?1 wt-%. As the C contentdecreased, the creep strength of both base metal andthe simulated HAZ increased. The normal decrease inHAZ creep strength was eliminated. The increase inresistance to cracking has been attributed to the decreasein coarsening rate of carbide and the volume fraction ofcarbide as the carbon level decreases; longer term testsare still needed to definitively prove this claim.

An approach not explored actively to eliminate typeIV cracking is to alter the heat treatment schedule ratherthan change composition. Abe et al.98 have found thatGr 92N subjected to normalising treatment and notempering had the same creep rupture life as the basemetal. However it is not very conclusive if this, andabsence of type IV cracking is due to a change in theheat treatment procedure or composition since Gr 92Ncontained nitrogen.

Dissimilar metal weldsAustenitic to ferritic and ferritic to ferritic steeltransition joints are often used in fossil fired powerplants. In conventional boiler designs, austenitic steelsare used in the final stages of superheaters and reheaterpendants, where the temperature is high and whereconventional ferritic and CSEF materials would not lastthe anticipated design life. Hence, there is a need to havedissimilar metal weld (DMW) joints between austeniticand ferritic steels. Some practical insight into such jointsis available in the literature.48,70 Failures in the powerplants involving dissimilar metal joints brought about anincrease in the research activities in the 1950s.109 Thefailure of the joints was attributed to differences in thecoefficient of thermal expansion between the two steels,namely, 2JCr–1Mo ferritic steel and 316 austeniticstainless steel, and accelerated creep in a carbon depletedzone due to carbon migration in the ferritic matrixadjacent to the interface. Exhaustive studies have beenconducted to understand and provide better engineeringguidance to increase the performance of austenitic toferritic DMWs.110–113 Recent studies of DMWs betweenP(T)91 and alloy 800 using INCONEL 182 filler metalhave shown P(T)91 to be susceptible to a similar failuremechanism observed in P(T)22 weldments using nickelbase filler materials.114

The earliest elastic stress analysis in the joint geometrywas performed jointly by General Electric Company andOak Ridge National Laboratory in the mid-1970s.115

Their analysis showed that introducing a piece oftransition metal (alloy 800H), whose coefficient ofthermal expansion is between those of 2JCr–1Mo steeland 316 austenitic steel, reduced the stresses signifi-cantly. Further significant work in this area is describedin the literature.116,117

When dissimilar joints are made between P(T)91,P(T)11 and P(T)22 steels (ferritic to ferritic), stresses

22 Images (TEM) of normalised and tempered P(T)91 steel a without B and b with B107 (with permission of Springer

Science)




developed at the joint interface are low because theircoefficients of thermal expansion are low and their thermalconductivities are high. When exposed to PWHT or servicetemperatures, such joints undergo significant redistribu-tion of interstitial C and N and modification of themicrostructure.118,119 Carbon migrates from the low CrCr–Mo steel to the high Cr Cr–Mo steel. Such a migrationof C leads to a C depleted zone, a soft zone characterisedby low volume fraction of carbide precipitates in the lowCr steel. Similarly, on the high Cr steel side, a C enrichedzone develops, which is a hard zone characterised by ahigher volume fraction of carbide precipitates. Themigration of C is driven by a gradient in chemical potentialof C between the two alloys.120,121 The low carbon zonetends to dissolve carbides and therefore experiences nograin boundary pinning effects. This leads to recrystallisa-tion of grains in the region, the development andconcentration of a triaxial stress state and accumulationof deformation, resulting in premature failure of theweldment.122,123 The evaluation of cross-weld CSEFDMWs made between P(T)91 and P(T)23 in creep hasdemonstrated a similar propensity for premature failuredue to carbon migration.124,125

Recent efforts in DMW research has focused onmitigating carbon diffusion and the formation of hardand soft zones in dissimilar ferritic/ferritic joints. Oneapproach is to have a nickel based diffusion barrierbetween 2JCr–1Mo and 9Cr–1Mo steels.119 The pre-sence of such a diffusion barrier has been found toreduce the propensity to form hard and soft zones. Anumerical analysis to model and simulate the formationof soft and hard zones in the DMW joint with adiffusion barrier has been developed and has been usedto optimise the thickness of the diffusion barrierinterlayer.

Another approach has led to the development of HFS6,a new nickel based filler metal with similar thermalexpansion properties to the low alloy ferritic steel andlower Cr content to avoid carbon migration.126 The alloywas never used because of microfissuring problems. Afterfurther development and extensive characterisation, anew filler metal electrode was developed and has beencommercialised as filler metal EPRI P87(Ni–0?01C–0?3Si–1?5Mn–9?0Cr–2?0Mo–1?0Nb–38Fe–0?008S–0?008P, inwt-%). EPRI P87 has demonstrated excellent elevatedtemperature properties and promises an improvementin cross-weld rupture life between CSEF steels andaustenitic stainless steels. No tendency for C migrationor detrimental carbide formation has been documentedwhen EPRI P87 was welded against a 9 wt-%Cr CSEFsteel and austenitic stainless steel.127,128

Recently, a novel concept of graded transition material(a functionally graded steel whose composition andcoefficient of thermal expansion change gradually fromthe ferritic to austenitic) has been proposed.129 It couldreplace the traditional transition piece of alloy 800.Current DMW practice mostly involves buttering oneend of the base metal with a suitable filler metal, PWHT,and joining to the other.

Weld repairTo meet power demands and to keep costs low, utilitieswould like to keep the time required for the maintenanceand forced outages due to repair to a minimum. Toaddress this need, utilities require a sound and well

engineered repair strategy.130 Such an approach couldextend the operating life of a power plant and couldreduce cost and the replacement of damaged components.In the USA, most of the power plants are .25 years old,and a survey of utilities by EPRI found that the utilitieshave no standard guidelines for repair welding proce-dures for extending the life of aging power plants.131

Electric Power Research Institute’s survey of power plantequipment manufacturers, power producers and usersrevealed that utilities were experiencing repeat cracking atlocations where repairs had been made, revealing anurgent need for improved repair welding practices. Basedon these surveys, EPRI is developing procedures andguidelines for performing reliable weld repair of fossilfired power plant components fabricated from CSEFsteels.

P(T)91 is the steel of choice for heavy sections such aspipes and headers in the advanced fossil fired powerplants today. It has a martensitic microstructure withhigh hardness and requires tempering or PWHT, duringwhich the hardness of the steel decreases gradually.131

The PWHT operation is expensive and time consuming,and repeated PWHT thermal cycles (such as in weldrepair) may have a negative effect on the life of a givencomponent. Utilities recognise the potential benefit in adeveloped, suitable repair technology that would elim-inate PWHT after repair in CSEF materials such asP(T)91. It is known that when ferritic steels used in fossilfired power plants are exposed to the weld thermal cycle,the steel generates a range of microstructures in theHAZ (Fig. 1). The coarse grained region may besusceptible to cracking (e.g. hydrogen cracking, reheatcracking) soon after welding; some CSEF materials aremore susceptible to these mechanisms. The coarsegrained region also has high hardenability and willlikely form martensite upon cooling. Any weld repairprocess that can temper and/or eliminate the coarsegrained region is highly desirable to reduce risk to thepreviously mentioned cracking mechanisms.

The repair technique developed and most commonlyused is the temper bead welding technique, also knownas cold welding. The use of temper bead welding in fossilfired power plants has its roots in materials highlysusceptible to reheat cracking, such as CrMoV steels andin low alloy Cr–Mo steels.132,133 The process has beenadded for qualification to section IX of the ASMEBoiler and Pressure Vessel Code. In the temper beadprocess, weld metal is placed in a location to alter theproperties of the previous weld bead (see Fig. 23a).134

Once the top layer has been deposited, a second layer isdeposited, and the heat from the second layer penetratesinto the layer below it, where it tempers the weld bead inthe first layer and the HAZ (Fig. 23b). Sometimestempering of the first layer is improved by grinding offpart of the first layer bead and its HAZ before thesecond layer is applied. This process is called the halfbead tempering technique.

Bhaduri et al.135 have investigated two variations ofthe temper bead process using artificially aged 2?25Cr–1Mo and 9Cr–1Mo steels. The process variationinvolves the half bead temper bead and butter beadtemper bead techniques. In butter bead temper beadtechnique, the heat input to the weld is altered bysuitably changing the size of the electrode. Repair weldswere made on artificially aged specimens; the SMAW



process and matching filler metal were used according tothe ASME Boiler and Pressure Vessel Code. Of the twoprocesses, the butter bead temper bead process wasfound to be more suitable for both 2?25Cr–1Mo and9Cr–1Mo steels.

Significant advances, such as variations to the temperbead procedures, narrow groove welding and highdeposition rate processes such as FCAW,136 are beingmade to develop other, more effective welding and weldrepair technologies. In the past decade, a number ofinvestigators have examined the potential of using atemper bead weld repair for grade 91 components.137–140

To eliminate PWHT after welding grade T91 steel,Siefert et al.141 have investigated two temper bead repairwelding processes, the consistent layer process and thecontrolled deposition process, in which EPRI P87 nickelbase filler metal was used in tubing applications. In theconsistent layer process, the heat energy dissipation iscontrolled to develop a tempered martensitic structure inthe first few millimetres of the HAZ. In the controlleddeposition process, the heat input is varied by controllingthe electrode size. Of the two, the consistent layer processwas found to be more effective in tempering the repairwelds. P87 was found to be beneficial in a number of waysover conventional nickel base materials (ERNiFe-2 andERNr Cr-3), namely in that it prevents the formation ofdetrimental type I carbides in the CSEF steel at the fusionline of the weld. The application of such weldingtechniques was specifically examined and reported fortubing applications; because the potential for prematurefailure of the repair is internal to the furnace, it has far lessrisk in causing injuries to plant personnel.

Continuing research at EPRI has been initiated inexamining the potential of applying temper bead

welding techniques to thick section grade 91 componentsusing a variety of filler materials and techniques(Table 8).32 The application of a temper bead weldrepair may pose fewer issues in martensitic steels thanwith the implementation of similar techniques in legacybainitic Cr–Mo steels or CrMoV steels because martensi-tic steels have been experimentally shown to accumulateless residual stress during welding. This was observed inexperiments employing the Satoh test (see Fig. 24).142,143.The Satoh test is a simple constrained bar test, wherebythe material of interest is heated to various peaktemperatures or profiles consistent with manufacturingprocesses (such as welding, normalising, or tempering).Upon cooling, the thermal contraction leads to anaccumulation of stress, which can be plotted as a functionof temperature. As can be seen in Fig. 24, transformationplasticity during the cooling of a uniaxially constrainedsample from the austenite phase field acts to relieve thebuild-up of thermal stress as the sample cools. Residualstresses were observed to accumulate only after transfor-mation was completed and the specimens approachedambient temperature.144–146

By contrast, the non-transforming austenitic steelexhibited a continuous increase in residual stress withdecreasing temperature, as might be expected fromthe thermal contraction of a constrained sample. Thisobservation is profound in that it may be entirelypossible to utilise a temper bead procedure for grade 91,and it may be the case that martensitic materials aremore accepting of such procedures than their bainiticcounterparts. This observation may account for some ofthe issues reported in bainitic CSEF materials T/P23 andT/P24, which were not of concern during the develop-ment and widespread introduction of T/P91 and T/P92.

a remelting of weld metal and HAZ of first bead due to thermal effects of second bead; b effect of second layer ofweld metal on first layer134

23 Schematic diagram of temper bead welding process (with permission of Welding Journal)




OutlookCoal will continue to be a major source of energy in theworld. Coal fired power plant steam conditions havesuccessfully evolved from subcritical to USC andcontinue to extend into the AUSC regime. The prog-nosis of continued improvement in plant efficiency and areduced carbon footprint looks promising as significantadvances are being made in a number of technicaldisciplines, such as alloy design, microstructure–prop-erty relationships, fabrication techniques (welding inparticular) and welding power supplies, sensors, andcontrols. New alloy design methodologies can be used toengineer, design and develop advanced structural alloysand welding consumables for boilers and boiler compo-nents. Ideally, these alloys should be easily weldable andconsistently reproducible. As alloys with better hightemperature creep properties, corrosion resistance andother thermophysical properties are developed, weldingof such advanced alloys inherently may become morechallenging. Therefore, attention should be paid to theweldability issues early in the alloy design stage of futurematerials. Such attention to weldability is well placedbecause a boiler requires tens of thousands of welds as itprogresses from construction in a fabrication facility tofield erection and commissioning.

In the fabrication of boilers and other components,greater understanding is needed of the metallurgy ofmaterials subjected to the weld thermal cycle. Animmediate need is to better characterise existing complexCSEF steels gaining traction in industry (specifically,grades 23, 24 and 92). Such characterisation, as it relatesto material behaviour, the evolution of microstructureand the microstructure–property relationship that devel-ops when the alloys are welded, will directly enable moreprecise life assessment and more applicable weld repairprocedures. This undertaking must be enhanced andfurther refined using applicable computational model-ling and realistic weld simulation studies with propermaterial property input and understanding of theresidual stress states induced by welding. In this regard,future studies will involve an integration of materials,welding processes, microstructure, the microstructure–property relationship, sensors and controls and theirmodels to develop reliable materials and welds(Fig. 25).147

24 Satoh test shows how residual stresses develop on

cooling for various steels142,143

25 Integration of process, microstructure, properties, pro-

cess control and automation and their models147

Table 8 Key aspects of variables considered in evaluation of repair techniques for grade 91 piping and components32

Weld Base material

Weld metal

Welding procedure{ PWHTAWS design Trade name*,{

1A 1 E9015-B91 Thermanit Chromo9V Mod.

Normal procedureztype PWHT 1375uF (746uC)/2 h

2A, 2B ** E9015-B91 Thermanit Chromo9V Mod.

Normal procedurezlow PWHT 1250uF (676uC)/2 h


Temper bead None


Poor practice temper bead None

5A, 5B ** E8015-B8 9Cr–1Mo Temper bead None6A, 6B ** E9015-G (-B23) Thermanit P23 Temper bead None7A, 7B ** E9015-G (-B23) Thermanit P23 Normal procedureztype PWHT 1375uF (746uC)/2 h8A, 8B ** E9018-B3 2?25Cr–1Mo Temper bead None9A, 9B ** ENiCrFe-4

(in code case)EPRI P87 Temper bead None

10A, 10B ** ENiCrFe-2 INCO-WELD A Temper bead None

*Thermanit is a registered trademark of Bohler Thyssen Schweisstechnik Deutschland.{INCO-WELD is a registered trademark of the Special Metals Corporation family of companies.{The letter after each temper bead procedure signifies a specific welding procedure, detailed below it.1Service exposed grade 91 (‘A’).**Service exposed grade 91 (‘A’) renormalised and tempered Grade 91 (‘b’).





Current high deposition rate welding processes used inthe construction of boilers are SAW and SMAW. Otheravailable processes that show promise encompassadvanced arc welding processes such as FCAW, narrowgroove arc welding, hot wire GTAW to high energydensity welding processes such as reduced pressureelectron beam welding and laser welding. The feasibilityand benefits of these processes need to be evaluated onboth a material performance basis and a cost basis.Other advanced processes such as FSW and hybrid(laser gas metal arc) welding have the potential toimprove both productivity and quality of the welds. Inaddition, they are low heat input processes, which havethe advantages such as the reduction in residual stressesand distortion and the potential elimination of problemssuch as type IV cracking.

Welding transition joints between dissimilar alloys posean unavoidable challenge in the construction of currentand future coal fired power plant boilers. The design andimplementation of new materials that will be subjected toDMWs must be thoroughly evaluated and understood sothat costly outages from unanticipated issues in materialcompatibility can be avoided. For this issue to beaddressed, existing filler materials used for similarapplications in other industries will need to be evaluated,new filler metals may need to be designed, and new joiningprocesses may have to be investigated to ensure acceptableperformance of the DMWs. The new concept of a gradedcomposition transition joint has the potential to replacethe current technology and to mitigate problems such asthermal fatigue due to carbon migration and mismatchesin coefficients of thermal expansion.

With the worldwide power grid transforming at arapid pace, it is inevitable that current fossil fired powerplants will have to move from a baseload operation to acyclic operation to accommodate the addition ofrenewable energy sources. Thus, higher strength materi-als will be required for thinner sections. The emphasiswill be on components and welds being able towithstand thermal fatigue and other damage mechan-isms specific to cycling conditions. This realisation mayalter the way in which materials are designed (currentlystrength dominated) to a focus on other high tempera-ture material requirements, such as fatigue and ductility.

Regardless of the impetus for new material develop-ment, weldability assessment or weld process develop-ment, future research in CSEF steels must provideadequate understanding to develop well engineeredsolutions. It is inevitable that as the complexity ofthe base material, filler material or welding processincreases, the amount of data needed to adequatelyassess a given material or process multiplies dramati-cally. This observation was not fully appreciated beforethe implementation of longitudinally seam weldedpiping, application of T23 and T24 in waterwall panels,the realisation of improper PWHT cycles and the lack ofdetailed education for fabricators, erectors and end usersregarding these complex materials. Well engineeredsolutions demand a thorough understanding of pastfailures and limitations as well as an intimate interactionbetween researchers and industry to understand andultimately solve today’s most pressing challenges. Theunfortunate reality is that in the forum of CSEF steels,there are countless examples where this has not been thecase. As such, the opportunity for research surrounding

these materials remains fluid, extremely intriguing andripe with opportunity.

Acknowledgements

The authors wish to acknowledge Dr J. Chen (OakRidge National Laboratory) for his valuable time andassistance in preparing this manuscript. We thank Dr W.F. Newell, Euroweld, Newell associates and K. Coleman(EPRI) for valuable comments. We also would like tothank Professors H. K. D. H. Bhadeshia (CambridgeUniversity), T. DebRoy (Pennsylvania State University)and Dr J. Shingledecker (EPRI) for their review andvaluable comments. Research is sponsored by theOffice of Fossil Energy Advanced Research MaterialsProgram, US Department of Energy, under contractDE-AC05-00OR22725 with UT-Battelle, LLC.

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