Review Article Physical Metallurgy of Modern Creep...

Review ArticlePhysical Metallurgy of Modern Creep-Resistant Steel for SteamPower Plants Microstructure and Phase Transformations

V C Igwemezie12 C C Ugwuegbu2 and U Mark2

1Renewable Energy Marine Structures Centre for Doctoral Training (REMS CDT) Cranfield University CranfieldBedfordshire MK43 0AL UK2Department of Materials amp Metallurgical Engineering Federal University of Technology Owerri Ihiagwa Nigeria

Correspondence should be addressed to V C Igwemezie victorigwemeziecranfieldacuk

Received 26 June 2016 Accepted 11 October 2016

Academic Editor Sunghak Lee

Copyright copy 2016 V C Igwemezie et alThis is an open access article distributed under theCreative CommonsAttribution Licensewhich permits unrestricted use distribution and reproduction in any medium provided the original work is properly cited

The fact that the microstructure of steel depends on its composition and the heat treatment given to it has been heavily exploitedin the design of steel for power plant applications To obtain a steel that can function at the higher temperature where power plantsoperate without failure for extended life heat treatment is needed to produce fine and highly stable dispersion of carbides nitridesand intermetallic compounds in the microstructure of the material A significant contribution also comes from solid solutionstrengthening by substitutional solutes We review here various types of phases microstructures functions and interacting effectsof the various alloying elements in the design of steel for modern power plant application

1 Dimensions of Steel

The critical role that steel has played in the construction ofour society is largely due to its superior adaptation to thewideranges of mechanical requirements The strength and tough-ness of steel can vary with changes in its microstructure thatis the size and shape of grains and the detailed arrangementof the atoms [1]

There are behaviours of iron which are unseen fromordinary view but ultimately affect their physical behaviourThe macroscopic shape of steel crystal depends on theenvironment processing andmany other factors determinedmostly by the fact that a piece of steel is generally poly-crystalline and is filled completely with millions of one ormore kinds of crystals The fact that a piece of steel containsvarieties of crystals is frequently not apparent to many designengineers making them treat the metal as a continuumespecially in modelling approaches

The electrons form the bonds between atoms Thestrength of these bonds is a major factor which affects thenature and properties of materials Another consideration isthatmicrostructure of steel by nature is imperfect In practiceit is found that equilibrium in engineering steel demandsthat crystals should be imperfect or away from equilibrium

Steels may not be very useful as they are presently withoutcertain defects There is always an equilibrium concentrationof point defects and more significant are the nonequilibriumunintended defects introduced in the process of manufacture[15]

In the design of steel for power plant or for hightemperature application there are challenging engineeringrequirements for the materials engineer The summary ofthese requirements is as follows [6]

(1) High creep strength at high temperature(2) High toughness and resistance to embrittlement(3) Resistance to steam oxidation and corrosion(4) Ease of fabrication and weldability

Not all steels are candidates for high temperature applicationStudy has shown that austenitic steels have higher coefficientof thermal expansion and low conductivity than ferritic steelsIn other words ferrite has much smaller thermal expansioncoefficient and high thermal conductivity than austenite asshown in Table 1

When considered for elevated temperature service mate-rials with large thermal expansivities become susceptibleto thermal fatigue especially in thick sections of the kind

Hindawi Publishing CorporationJournal of MetallurgyVolume 2016 Article ID 5468292 19 pageshttpdxdoiorg10115520165468292

2 Journal of Metallurgy

Table 1 Comparison of the coefficient of thermal expansion (CTE) and thermal conductivity for ferritic and austenitic steel at 100∘C Thecompositions (wt) of the steels are also given [2 3]

(a)

Steel CTE120583mminus1Kminus1 Thermal conductivityWmminus1Kminus1

Ferritic 99 249Austenitic 173 163

(b)

Steel AISI type C Mn P S Si Cr NiFerritic 403 015 100 0040 003 050 115ndash130 mdashAustenitic 305 012 200 0045 003 100 17ndash19 105ndash130

frequently employed in power generation This becomes amajor problem due to the fact that steam turbines are oftenturned off and on during their service life This thermalcycling can cause fatigue in austenitic steels This has beenthemain reasonwhy austenitic steels in spite of their superiorcreep strength are rejected in favour of ferritic steels forthe construction of power plant [6] Martensitic steels havesmaller thermal expansion and larger thermal conductivitythan austenitic steels and Ni base superalloys and are nowoffering highest potential to designing thick section steels forpower plant application

In the development of modern steam power plants thegoal is centred around increasing the steam parameters fromthe current optimum of about 620∘C at 28MPa up to 700∘Cat 35MPa in order to improve efficiency and thereby reduceCO2emission [19] Further advanced steam conditions of

700∘C and above have been already initiated to gain netefficiency higher than 50 at Thermie AD700 project inEurope and at DOE vision 21 project [13 20 21] In thelight of available knowledge so far these steam conditionscannot be reached with ferritic-martensitic alloys which arecurrently the state-of-the-art material for steam power plants[19 22] Hence these projects involve the replacement of 9ndash12Cr martensitic steels by Ni base superalloys for the highesttemperature regimes Tominimize the requirement of expen-sive Ni base superalloys as Ni superalloy is about ten timescostlier than high temperature steels 9ndash12Cr martensiticsteels are planned to be used in the next highest temperaturecomponents of such very high temperature plantsTherefore9ndash12Cr martensitic steels are still strongly desired to expandthe present temperature range up to 650∘C and more Itshould be noted that research is also going back to austeniticsteels to see their possible use for 700∘C USC plants in Japan

2 Phase Transformation in Creep Steel forPower Plant

The general rule for determining if creep will occur in metalsor not is [23 24]

119879 gt 03ndash04 (119879119872) (1)

Pure iron has a melting temperature 119879119872 of about 1536∘C

and the equation above suggests that pure iron would creep

at temperature where traditional steam power plant operatesthat is in temperature range of 450ndash565∘C

Since service temperature depends on the melting tem-perature of the steel a direct approach to solving the creepproblem is to use metals with high melting temperatures sothat the temperature at which creep occurs is increased Anexample of suchmetals is tungsten with119879

119872of 3407∘C In this

case creep becomes important above about 1363∘C which isfar above the aforementioned temperature range for powerplants This apparent advantage comes with considerableconsequences Firstly W is several times more expensivethan steel which could make affordability of the power plantvery unlikely for most countries This cost implication andthe attendant low demand will have a negative impact onthe huge investment in the technology Secondly the oxideof tungsten (WO

3) sublimes and so forth is not considered

as a practical solution Another approach could be to usenickel-based superalloys Study has shown that the creeprupture strength of the nickel-based superalloys is verygood at high temperatures and they have a sufficiently lowexpansion coefficient Unfortunately they are about 6ndash10times more expensive than low alloy steels depending on thealloy composition and market prices of alloying elements Inaddition nickel-based superalloys aremore expensive toweldand machine

Almost all steels rely on the transformation betweenaustenite and ferrite for obtaining the desiredmicrostructureWhen the high temperature 119891119888119888-phase austenite in steeldecomposes to the less dense 119887119888119888-phase ferrite a number ofdifferent microstructures and morphologies can form whichdepend on the cooling rate the presence of alloying elementsand the conditions and availability of lower energy nucleationsites for heterogeneous nucleation The transformation tem-perature between austenite and ferrite can be greatly affectedby the presence of solutes such as C Si Mn Ni Mo Cr WTi Nb andV For example 08 wt carbon lowers the secondsolid-state transformation temperature from 910∘C to 723∘C

The transformation temperature can also be altered bychanging the cooling rate from elevated temperatures Arapidly cooled steel transforms at a lower temperature andthere is shorter amount of time available to accomplish thenecessary changes in atomic positions before the materialfinally cools down to room temperature The products ofsuch transformation are metastable phases The advantage of

Journal of Metallurgy 3

Table 2 Stored energy as function of microstructure defined relative to microstructure for which stored energy is arbitrarily set to zero[adapted from [4]]

Case Microstructure Stored energy JmolPhase mixture inFe-02C-15Mn wt at 27∘C1 Ferrite graphite and cementite 02 Ferrite and cementite 703 Paraequilibrium ferrite and paraequilibrium cementite 3854 Bainite and paraequilibrium cementite 7855 Martensite 1214Phase mixture inFe-01C-22Cr-1Mn wt at600∘C6 Ferrite and paraequilibrium cementiterarr ferrite and M

2X (eg when steel enters service) 63

7 Ferrite and paraequilibrium cementiterarr ferrite and M23C6(eg after long service time) 0

fast cooling is that the structural scale is reduced which isadvantageous in achieving greater strength and toughness [1]

To obtain a steel that can function at the higher tem-perature where power plants operate without failure forextended life pure iron is alloyed with other elements suchas C Cr Ni Mn Mg W Ti V Mo and Nb and thenheat-treated to produce steel with required creep propertiesCreep steels are able to survive for such long periods as 30years because the operating temperature is only about half ofthe absolute melting temperature making the migration ofatoms very slow to cause any significant change in the steelmicrostructure

In power plants the longevity of iron alloys relies on thefact that the diffusivities are incredibly small [4 25] Notwith-standing this slow but significant changes are observed tooccur over the long service life Another consequence ofthe small diffusion coefficients is that the dominant creepmechanism is the climb of dislocations over obstacles withthe help of thermal energy The obstacles are mainly carbideparticles which are dispersed throughout the microstructure[4]

Another reason why the changes that are expected overthe years must occur at a slow rate is that the steels aregenerally tempered before they enter service producing avery stablemicrostructure and relieving stressesThe stabilitycan be discussed in terms of the free energy stored in thematerial as illustrated in Table 2 [4]

The stored energy of the plain carbon steel (Phasemixture in Fe-02C-15Mn wt at 27∘C) is described relativeto the equilibrium state which is amixture of ferrite graphiteand cementite without any defects The phases in cases 1 and2 involve the equilibrium partitioning of all elements so as tominimize free energy that is minimize the tendency to reactfurther In cases 3ndash5 the substitutional solutes are frozen intheir position

It can be seen that bainite and martensite have higherstored energies relative to ferrite because of strain energycontributions and solute trapping The term paraequilibriummeans that the phase has the same ratio of iron to substitu-tional solute atoms as the parent phase fromwhich it formed

Cases 6 and 7 are for a classical power plant alloyFe-01C-22Cr-1Mo wt at 600∘C Here the reduction instored energy when the microstructure changes from M

2X

to M23C6microstructure is very small This means that

this transformation will occur at a corresponding slow rateIn summary the reaction rate of a transformation can becontrolled by the available Gibbs free energy also referred toas the driving force and by the diffusivity of carbon

3 The Matrix Microstructure of Creep Steel

The ability of power plant steels to resist creep deformationdepends on the presence in the microstructure of fineand highly stable dispersion of carbides and intermetalliccompounds which precipitate during tempering or duringelevated temperature service [4] The carbides interfere withthe climb and glide of dislocations and retard the coarseningrate of themicrostructure as awhole for example the size andshape of martensite or bainite plates In other words theseparticles not only interfere with the progress of dislocationsbut also stabilize the microstructure so that features such aslath boundaries change very slowly during long-term serviceat elevated temperatures [7] A significant contribution alsocomes from solid solution strengthening by substitutionalsolutesThe solid solution strengthening component is foundto becomemore important after prolonged service at elevatedtemperature as the microstructural contribution to strength-ening diminishes due to annealing effects

The heat treatments given to the steels before service areusually so severe that the precipitates are in an overagedcondition The microstructures are therefore relatively sta-ble before entering service though further microstructuralchanges are inevitable given that service involves manydecades at temperatures where substitutional atoms havesignificant mobility and this change is sometimes used tomonitor remaining life of the steel components

Majority of steels used in fossil power plants in the wholeworld are constructed from low alloy steels containing CrMo and V as the major alloying elements The total alloyconcentration is usually less than 5wt [4] These include


Table 3 Chemical compositions (wt) of some low-alloy creep-resistant steels for power plant construction

Designation C Si Mn Ni Mo Cr V1Cr05Mo 015 025 050 mdash 060 095 mdash025CrMoV 015 025 050 005 050 030 02505Cr05Mo025V 012 025 050 mdash 060 045 0251CrMoV 025 025 075 070 100 110 035225Cr1Mo 015 025 050 010 100 230 000

Table 4 Heat treatment conditions matrix phase and precipitates for steels of varying chromium content [5]

Steel Water quenching treatment Temperingair cooled treatment Phases present Precipitates after tempering2Cr-2W 930∘Chr 730∘Chr Bainite M

23C6 M6C M2C

5Cr-2W 930∘Chr 750∘Chr Bainite M23C6

9Cr-2W 950∘Chr 750∘Chr Martensite M23C6

12Cr-2W 1000∘Chr 750∘Chr Martensite + 120575-ferrite (016) M23C6


steels such as 05Cr05Mo025V 1CrMoV and 225Cr1Mo thecompositions of which are given in Table 3 along with someother common creep-resistant steels

These alloys have formed the backbone of the powergeneration and petrochemical industries for over 50 years foroperating temperatures of 565∘C or less 225Cr1Mo steel iswidely used for superheater tubing in power plant and as fillermaterials for joining 05Cr05Mo025V steam piping [26]

The primary microstructures of these low alloy creep-resistant steels are essential for leaner steels allotriomorphicferrite and pearlite [6] In other words about a decade agoor so for the classical or common creep steels the bestestablished steels relied on allotriomorphic ferrite or bainiteas the startingmicrostructureThesemicrostructures containvery small grains of the order of a quarter of a micron whichprovide additional opposition to the movement of disloca-tions If diffusion creep was prevalent in these materialsa large grain size would be needed to increase the energyrequired for diffusion between and through grains [6]

To mitigate oxidation and corrosion in these creepsteels deliberate increase in the concentration of chromiumwas made Consequently 120575-ferrite formed which is detri-mental to creep strength To eliminate this ferrite fromthe microstructure balancing additions of other alloyingelements were added leading to sufficient increase in thehardenability of the steel On normalization heat treatmentit is found that the microstructure of this new steel systembecame martensite instead of bainite for identical coolingconditions

Abe and Nakazawa (1992) [5] presented a study on theeffect of varying chromium content on creep steel Thesummary of their study is shown in Table 4

By varying the chromium content of simple 2wt Wsteel the microstructure changes from bainitic to martensiticand then to a dual-phase matrix of martensite and 120575-ferriteof about 016 volume fraction 120575-ferrite is coarse and soft andconsequently detrimental to creep strength The phase canform during the solidification of ferritic steels [27ndash29] Inother words higher alloy content steel such as 9ndash15wt Cr

200

300

400

500

600

700

800

0 2 4 6 8log (times)

Z

Y

X

Ferrite

Bainite

Martensite

Tem

pera

ture

(∘C)

Figure 1 Calculated TTT diagrams for steels of composition 119883(015C-02Si-05Mn-1Mo-23Cr wt) 119884 (43Cr wt) and 119885 (93Crwt) [9ndash12]

steels tends to be martensitic on cooling rather than bainiticThis is because as alreadymentioned the larger alloy contentcauses an increase in the hardenability of the steel to the pointwhere it cannot transform to bainiteThis is illustrated by thecalculated time-temperature transformation (TTT) diagramsfor 23 (119883) 43 (119884) and 93 (119885) wt chromium steels shownin Figure 1

In Figure 1 the upper curve for each case representsdiffusion transformation whereas the lower curve representsbainite In general 9ndash15 wt Cr steels are fully air hardenedmeaning that the microstructure formed on air cooling fromnormalizing is completely martensitic Therefore marten-sitic matrix is seen as arising by chance The martensitemicrostructure is reported to turn out to be the most exciting


Table 5 Nominal chemical compositions (wt) of ferritic steels

Steels C Si Mn Cr Mo W Co V Nb B N Others2Cr

T22 012 03 045 225 10 mdash mdash mdash mdash mdash mdash mdashHCM2S 006 020 045 225 01 160 mdash 025 005 0003 mdash mdash

9CrT9 012 060 045 900 10 mdash mdash mdash mdash mdash mdash mdashHCM9M 007 030 045 900 20 mdash mdash mdash mdash mdash mdash mdashT91 010 040 045 900 10 mdash mdash 020 008 mdash 005 mdashNF616 007 006 045 900 05 180 mdash 020 005 0004 006 mdashTempaloy F-9 006 050 060 900 10 mdash mdash 025 040 0005 mdash mdashEM12 010 040 010 900 20 mdash mdash 030 040 mdash mdash mdash

12CrHT91 020 040 060 120 10 mdash mdash 025 mdash mdash mdash 05NiHT9 020 040 060 120 10 05 mdash 025 mdash mdash mdash 05NiHCM12 010 030 055 120 10 10 mdash 025 005 mdash 003 mdashHCM12A 011 010 060 120 04 20 mdash 020 005 0003 006 10CuNF12 008 020 050 110 02 26 25 020 007 0004 005 mdash

SAVE12 010 030 020 110 mdash 30 30 020 007 mdash 004 007Ta004Nd

ferritic creep-resistant steels Consequently modern heat-resistant steels are based onmartensite microstructureThesehigher alloy content steels for higher temperatures andgreater corrosionoxidation resistance are generally referredto as 9ndash12wt Cr steels [6 27ndash29] for the well-establishedones

Martensitic steels usually have to be tempered in order toobtain an optimum combination of strength and toughnessThe tempering process usually leads to a decrease in strengthdue to the precipitation of iron carbides from carbon thatwas originally in solid solution in the martensite Howeverwhen the steel contains strong carbide forming elements suchas Mo Cr V Ti and W it becomes possible to recover thestrength during prolonged tempering at elevated tempera-tures This is called secondary hardening it occurs becausesubstitutional alloying elements (such as Cr) can diffuseduring prolonged annealing to precipitate finely dispersedalloy carbides

The alloy carbides in secondary-hardened steels are foundto be more thermodynamically stable than iron carbides andshow little tendency to coarsen However many of the alloycarbides find it difficult to nucleate so that the equilibriumcarbide is not necessarily the first to precipitate Insteadthere is a sequence of alloy carbides precipitated in an orderdetermined by a combination of kinetic and thermodynamicconsiderationswhich leads ultimately towards an equilibriummicrostructure

Many of the precipitates found in power plant steelshave crystal structures and compositions which are quitedifferent from those of the ferrite matrix The precipi-tatematrix interfacial energy can therefore be expected tobe large making it difficult for the equilibrium precipitateto nucleate Consequently decomposition often starts withthe formation of one or more metastable phases which are

kinetically favoured These initial phases eventually dissolveas equilibrium is approached The phases normally appearand reappear in the course of the microstructure attain-ing equilibrium This progression towards equilibrium atservice temperatures makes creep steels useful It is wellestablished that the fracture toughness of many power plantsteels deteriorates during service at elevated temperatures fortwo reasons Firstly the carbide particles particularly thoselocated at the prior austenite boundaries coarsen and henceprovide easier sites for crack or void nucleation Secondlythe elevated temperatures permit the impurities to diffuserelatively rapidly and saturate the boundaries

Table 5 shows the composition of a selection of ferriticpower plant steels

One can see that there are three main well-establishedgroups of ferritic power plant steels as shown in Table 5usually designated 2 9 and 12wt Cr steels The 9ndash12wtCr steels are currently undergoing the most research anddevelopment due to their high strength and good corrosionresistance However research has been done [8 17 30 31]and currently going on in the development of 15Cr1Mo creep-resistant steels for power plant temperatures up to 700∘C

4 Heat Treatment of Power Plant Steels

In general the primary microstructures of power plant alloysoften consist of 120575-ferrite martensite bainite allotriomorphicferrite and retained austenite as the major phases obtainedfollowing a normalizing heat treatment [32] The phasesformed in these alloy steels are dependent not only on theelements that make up their composition but also on the heattreatments applied to the components before they go intoservice Studies have demonstrated how easy it is to changethe microstructure of the steels by making little change in


the alloy content The situation can be further complicatedby heat treatments applied to the steel which can have a largeeffect on the microstructure and the mechanical propertiesThe heat treatment produces carbides and precipitates thathave a direct effect on the properties of the steel The mainintention of the heat treatment is to optimize the mechanicalproperties of the alloy and to relieve stresses Generally twotypes of heat treatment are applied a normalization treatmentand a tempering treatment

Normalization involves austenitization at a high temper-ature in order to dissolve many of the alloying elementsinto solid solution The temperature is usually high enoughto dissolve stable precipitates already in the steel The pre-cipitate may be carbides and nitrides so that on coolingfrom austenitizing temperature they precipitate as a finedispersion of particles The austenitization temperature thatgives complete dissolution of precipitates varies for differentalloy compositions A too low temperature gives incompletedissolution of carbide and this leads to loss of creep rupturestrength in 12Cr-Mo-V steel A too high temperature givesundesirable formation of 120575-ferrite phase Recent publicationshave emphasised that a high Ac

1temperature (where austen-

ite formation first begins during heating) leads to good creepstrength [33ndash35]

After austenitization the cooling rate will determine thephases that are formed from austenite and depends on alloycomposition In some low alloy steels slow cooling may formundesirable ferrite Formation of ferrite or low bainite if thecooling is too fast will result in an unsatisfactory distributionof carbides and this is found to affect creep properties Tocontrol cooling rates modifying composition to change thehardenability or using a suitable coolingmedium is necessary

The normalization is usually followed by tempering heattreatment to obtain the required mechanical properties Thistempering process nucleates precipitates and relieves stressIt is well established that the creep properties improve withsubsequent tempering and then eventually degrade afterthe precipitation of the larger equilibrium alloy carbidesMost creep steels are subjected to very severe tempering(sim700∘C for several hours) causing general coarsening andthe precipitation of ever more stable alloy carbides andintermetallic compounds The precipitates phases and com-position of a steel and the heat treatments applied to it are allimportant and interacting factors affecting creep propertiesIt is these solid-state reactions which ultimately determinethe mechanical stability of the steels and hence their usefuldesign lives [32]

In practice two tempering heat treatments are oftenapplied to the alloys after normalization The first is meantto eliminate any retained austenite in the microstructure asthis could influence dimensional stability An initial temperof 500∘C can induce the decomposition of the austenite somachining can be carried out accurately The second is toproduce a ldquostablerdquo microstructure consisting of a variety ofalloy carbides in a ferritic matrix This tempering producesthe required mechanical properties

Tempering is carried out below the austenite trans-formation temperature so that there is no transformationto austenite Temperature may be in excess of 500∘C in

Table 6 Typical heat treatments schedule for power plant steelswhere Y0336 Z1092 A and B are names for some developed creepsteels [6]

Steel Y0336 Z1092 A BNormalizationTemperature ∘C 1053 1048 1200 1180Duration hr 1135 116 2 2Cooling rate AC AC AC ACTemperingTemperature ∘C 750 742 800 800Duration hr 1267 125 4 4Cooling rate AC AC AC ACAnnealingTemperature ∘C mdash mdash 740 740Duration hr mdash mdash 4 4Cooling rate mdash mdash AC ACAC air cool

order to allow the diffusion of substitutional atoms over thedimensions of precipitate growth in a reasonable period oftime 225Cr1Mo steel is usually given a stress-relief heattreatment at 700∘C for several hours before going into serviceat a temperature of approximately 565∘C [4] The conditionschosen affect the precipitate size and creep properties Theinitial precipitations formed are those for which nucleationis easiest At longer times mores table phases may form withthe complimentary dissolution of the less stable carbidesAlthough the nucleation of these latter phases is moredifficult their formation leads to a reduction in the freeenergy of the system and so is thermodynamically preferableTypical heat treatment schedule in the design of novel creepsteel by Cole (2000) [6] is shown in Table 6

The heat treatments applied to these steels have tobe carefully chosen and controlled to produce a suitablemicrostructure The choice has to take into account thecomposition of the alloy and the phases whichwill be formedNormalizing at a too low temperature or tempering at a toohigh temperature could have a deleterious effect on the creeprupture strength

Study has shown clearly that carbide particle sizesincrease with tempering temperatureThis growth is undesir-able for creep properties since a dispersion of fine particles isdesired Sikka et al (1984) [36] reported that tempering 9Cr-1Mo alloy above a nominal temperature of 760∘C producedinsignificant effects on creep rupture strength but temperingbelow this temperature significantly improved it

5 Description of Carbides in PowerPlant Steels

The evolution of microstructure in steel is a complex anddynamic process involving simultaneous transformations ofdifferent types of phases All the products compete for thelimited nucleation sites resources and space A fast coolingcondition can preserve the austenite into lower temperature


regions so that finer structures prevail When many trans-formations occur at the same time the fast reaction mayconsume much of the austenite The migration of carbonthrough interstices the creation rate of nuclei on the austenitegrain boundaries the growth rates of individual phases andso forth all react and interact with each other [1]

The task in designing power plant steels is usually tomodel the evolution of these precipitates and dissolutionreactions In dealing with precipitations in creep steels asimple notation system is used to describe the chemicalformula of the precipitate Here ldquoMrdquo is used to represent themetal content (mixture of metal atoms) and ldquoXrdquo representsthe carbon and nitrogen content

There are varieties of phases that could be found in thesepower plant steels The following list shows the well-knownprecipitates that determine themicrostructure of power plantsteels and are crucial in the development of creep strain [15]

Types of Precipitates That Can Be Found in Power Plant Steels[15 20] They are listed as follows

Possible phasesGraphiteEpsilon Fe

24C

Cementite Fe3C

Chi Fe2C

M2X M6C M23C6M7C3Laves

M5C2Z-phase 120583-Phase 120594-phase

M5C2is among the latest discovery in 1Cr-05Mo steels 120583-

phase has also been discoveredSince these precipitates are often metastable they do not

represent the lowest free energy state They form becausethey nucleate easily and they may dissolve as the materialapproaches equilibrium with time Therefore the first alloycarbide to form is generally not the equilibrium carbideleading to precipitation sequences that produce more stablecarbides [6 15] Sometimes carbides which are metastableat first become stable during service when the compositionof the steel changes This is reported to happen whentwo different kinds of steels are joined together leading todiffusion fluxes across the junction

Cementite is iron-rich carbide that forms very rapidlywhereas graphite despite an equilibriumphase forms incred-ibly slowly since it is difficult to nucleate Cementite con-taining no alloying element additions is seen to have anapproximately hexagonal close-packed (hcp) arrangementof metal atoms with localized distortions to accommodatethe carbon atoms or simply has an orthorhombic structureThe structure of cementite containing no alloying elementadditions has the structure Fe

3C

M3C denotes predominantly an iron-rich carbide with

the same orthorhombic structure as Fe3C but here several

other alloying elements can be found to partition into thiscarbide in significant quantities Other metal elements cansubstitute for iron in alloy steels The unit cell dimensionsare found to be affected by partial substitution of alloyingelements Mn is found to dissolve in large quantities as can

Cr with up to one-fifth of the Fe atoms being replaced byCr (specific to low alloy steels) Ni and Co can also dissolveas they form metastable orthorhombic carbides although itis found that they usually partition to ferrite Mo W and Vhave also been found to have limited solubilities in M

3C In

general M3C carbides can be referred to using the general

formula (FeCrMnMo)3C [26]

Tempering causes the cementite to enrich towards itsequilibrium composition It is found that the rate of enrich-ment is fastest when the particles are of small size and theferrite is saturated in carbide forming solute atoms [26 32]In low Cr steels it is found that the cementite will enrichslowly and is unlikely to reach its equilibrium compositionin the component lifetime but in high Cr and high alloysteels the cementite enriches rapidly and then dissolves asother alloy carbides precipitateThis phenomenon is believedto be the basis of secondary hardening in these steels Thesteel softens as the cementite forms and removes carbon fromsolid solution but is then hardened by the alloy carbideswhich interact with dislocations thus giving the so calledsecondary hardening It is possible to form more stablecarbides nitrides and borides than cementiteThese are seento form by the dissolution of a relatively coarse cementitedispersion which is replaced by finer alloy carbide dispersion[32]

It is found that when strong carbide forming elementsare present alloy carbides would form in preference tocementite However alloy carbides do not form until thesteels are tempered in the range of 500ndash600∘C because themetallic alloying elements cannot diffuse sufficiently rapidlyto allow the alloy carbides to nucleate Carbon like nitrogencan diffuse much more rapidly since it moves interstitiallythrough the iron lattice thus iron carbide tends to be the firstprecipitates formed in these steels

In many cases M2X type precipitates are found to be the

first to form after cementite and they have a hexagonal crystalstructure (hcp)HereMdenotes one ormore of these alloyingelements V CrMn FeNi NbMo Ta andWX represents Cor N It has been determined that in complex steels the phaseis best regarded as (CrMoV)

2(CN) Study found that in

Mo steel containing low Cr content and no nitrogen suchas 225Cr-1Mo steels M

2X is often close to Mo

2C [37] In 9ndash

12wtCr steels containing nitrogen the composition ofM2X

is found to be closer to Cr2N

M2X is found to exist as fine needle-like particles dis-

persed throughout the matrix and generally considered tonucleate on dislocations and martensite lath boundaries Itsprecipitation is found to often cause secondary hardeningwhich is particularly seen to be important for low alloysteels where M

2X plays a major role in the creep strength

of the alloy The precipitation of Mo2C is usually seen to be

the major factor in conferring creep resistance on low alloyferritic steels Study found that increasing Ni contents in steeldecreases the stability of M

2X phase [26]

Stabilization of a Cr-rich M2X phase in 12 wt Cr steels

can also cause secondary hardening [38] Unfortunately theelements which stabilize M

2X such as Cr and Mo are seen

to cause the formation of 120575-ferrite a phase detrimental to


creep properties therefore limiting the quantities which maybe added

In their study Baker and Nutting (1959) [37] observedthat M

7C3forms after the formation of M

2X (ie M

2C)

M7C3is a Cr-rich carbide with the trigonal structure of

Cr7C3 having a solubility of Fe up to 60 This phase can

form after cementite formation (M3C) without any M

2X

being observed In Thompsonrsquos review Titchmarsh (1978)[39] is said to have found that the Cr Fe ratio can be greaterthan 1 in 225Cr1Mo steel Mn V and Mo can also dissolvewith decreasing probabilities respectively

It has been found that M7C3tends to nucleate in the

vicinity of cementite or at the cementiteferrite interface of a225Cr-1Mo steel M

7C3seemed to be observed if a sufficient

Cr concentration is available The dissolution of cementite isseen to provide an area sufficiently rich in chromium inwhichM7C3could precipitate since thesewill be themost favourable

sites [32 40] This phase has been found to coarsen rapidlyand gave no beneficial contribution to creep rupture strength

M23C6is chromium-rich carbide which precipitates after

either M7C3or M2X in 9ndash12wt Cr It is observed to possess

complex fcc structure of Cr23C6[26] and the precipitate

may also contain W Mo V and Ni [5 6 41] In steelscontaining significant amounts of Mo it is believed that theformula Fe

21Mo2C6can result and that in steels containing

both Cr and Mo its composition can be anywhere withinthe above This carbide is often found to precipitate withequilibrium composition and is predominant after tempering[42] The phase is found to nucleate on the prior austenitegrain boundaries or martensite lath boundaries or adjacentto M7C3

Overaging is accompanied by the formation of M23C6

and dissolution of M2X It is then often found that M

2X

dissolves after long-term exposure at high temperature in 9ndash12wtCr steels and so does not improve the long-term creepresistance It was also found that vanadium stabilizes M

7C3

and so decreases the rate of release into the matrix of carbonand chromium for the growth of M

23C6

Various studies show that the composition of M23C6

varies with time The carbide is found to enrich in Fe after5 years in 225Cr-1Mo steels with a high Cr content withsome low Mo content and after 18 years in service the phasecontained much more chromium-rich M

23C6 The carbide

M23C6is frequently observed in the form of large or coarse

particles and is not thought to contribute directly to creepstrength That is M

23C6is less effective in hindering creep

deformation Delaying its precipitation would have the effectof stabilizing the finer dispersions of M

2X and MX to longer

times with a possible enhancement of creep strength Studyhas also shown that the tendency forM

23C6to coarsen should

also be greatest in the Cr-rich alloys [6]However it has been suggested that M

23C6carbides may

stabilize lath boundaries and delay grain growth In his studyAbe (1999) [13] made a conclusion that M

23C6precipitates

are the most effective in 9Cr-W system for pinning lathboundaries and reducing the coarsening rate of the lathsThisstudy shows that coarsening occurs by sideways movementof lath boundaries and the coalescence of two boundaries

M23C

Y

Y

6

Figure 2 Schematic of coarsening of martensite laths by themovement of Y-junction in 9Cr-Wsteels during creep testing at 550ndash650∘C for up to 15 000 hours with varying tungsten content [6 13]

Figure 2 illustrates the movement of Y-junction of triplepoints of lath boundaries to cause coarsening

Figure 3 shows an image obtained using TEM from athin foil specimen of 9Cr steel given severe heat treatments

The microstructure in Figure 3(a) shows temperedmartensite which retains the essential lath-like morphol-ogy and carbides The absence of coarsening is said to bedue to stabilization of the microstructure by the carbideswhich nucleated heterogeneously at the lath boundaries [6]Figure 3(b) shows carbides aligned at the martensite lathsboundaries

9Cr-Mo steels after creep test have shown preferentialrecovery of the microstructure in the vicinity of prior austen-ite grain boundariesThis recovery phenomenon is associatedwith carbides forming and coarsening preferentially at theprior austenite grain boundaries removing solutes and dis-solving smaller precipitates in the surrounding regions Theloss of pining effect by the precipitates allows the recovery ofthe martensitic microstructure [6] This process is depictedschematically as in Figure 4 for a 12Cr alloy

These mechanisms are reported to account for the loss inlong-term creep strength often experienced by 9ndash12Crwtsteels caused by the coarsening of the microstructure Thisprocess has been reported to have occurred formodified 9Cr-1Mo steel at 600 and 650∘C for 10 000 hours [14] Figure 5shows degradation in creep rupture strength for a modified9Cr-1Mo steel by this process

For the improvement of long-term creep rupturestrength research has shown that the dissolution of smallprecipitates and the coarsening of large ones need to beretarded Recovery of the microstructure also needs to bedelayed especially in the vicinity of prior austenite grainboundaries

The coarsening of M23C6carbides in 12CrMo(W)VNbN

steel is reported to be accompanied by dissolution of fineMXcarbonitrides Increase inNi content in 12CrMoV steel resultsin accelerated microstructure degradation with more rapidcoarsening of M

23C6 dissolution of MX and precipitation of

coarse M6X and Fe

2Mo Increase in nitrogen content results

in accelerated microstructure degradation during creep at650∘C with more rapid precipitation of a phase referred to asZ-phase


1 120583m

(a)

1 120583m

(b)

Figure 3 TEMmicrograph of Z1092 (9Cr steel) showing the general microstructure [6]

(a) (b)

Figure 4 Illustration of the degradation mechanism of a 12Cr-04Mo-2W-Cu-V-Nb steel (a) shows the microstructure prior to creep testingand (b) illustrates the microstructure after long-term creep deformation [6]

One of the important observations in the recent literatureis that many ferritic steels are strong in creep in the earlystages of creep life but the strength then deteriorates sharplyas a consequence of the development of local regions ofrecovery where damage is then intensified The observationsreported in [6] confirm that recovery is not usually homoge-neous A possible solution is suggested of the use of a dualheat treatment to obtain more homogeneous dispersion ofprecipitates and retard localized recovery The observationsalso support the design philosophy that the concentrationof Ni should be minimized since it causes acceleration ofmicrostructure degradation with more rapid coarsening ofM23C6

In Fe-9Cr-W steels W content has been found to affectthe coarsening of M

23C6during creep testing at 600∘C [13]

as shown in Figure 6 However this information is usuallyinterpreted with caution because the addition of tungstenalso causes the precipitation of Laves phase (intermetalliccompound) [13]

This study found that the coarsening of M23C6and the

martensite laths were reduced by increasing the amount oftungsten The increase in W also caused the precipitation ofLaves phase which is said to be beneficial in decreasing creeprate and suppressing the recovery of the microstructure

When steels contain Mo and relatively low level of CrM6C precipitates which seems to be the equilibrium carbide

[35 43] Therefore M6C is found to be essentially Mo-

rich carbide with fcc [26] M6C is found to nucleate on

M2X andM

23C6interphase boundaries prior austenite grain

boundaries or martensite lath growing rapidly at the expenseof all surrounding carbides nucleating at existing particles[6]The transformation ofMo

2C toM

6C is said to occurmore

rapidly in bainite than in ferriteWork reported by Nutting (1999) [44] indicates that this

phase may not nucleate from existing carbides but as thecarbides dissolve solutes are transferred toM

6C by diffusion

The carbide is often coarse and since it causes the dissolutionof less stable fine carbides M

6C can cause a reduction in creep

strength It has been shown that the coarsening rate of M6C is

greater than that of M23C6and so is a particularly undesirable

phase in creep steel design This carbide is found to bemainly rich in Mo in 05Cr05Mo025V steel and precipitatein 12CrMoV steels with high Ni content [44]

MX denotes some carbides which may be observed inpower plant steels and often has a cubic structure like NaCl[41] It was observed to be often rich in vanadium althoughNbC and TiC have also been found to form The vanadiumcarbide structure is basically VC but extends to V

4C3 There


10

100

1000

Stre

ss (M

Pa)

100 1000 10000 10000010Time to rupture (hr)

750∘C700∘C650∘C

600∘C550∘C

Figure 5 Degradation in creep rupture strength for amodified 9Cr-1Mo steel adapted from [14] or stress versus time to rupture curvesfor 9Cr-1Mo (or T91)

Laves phase

1 2 3 4 50Tungsten (wt)

M23C6

0

0005

001

0015

002

0025

003

Mol

e fra

ctio

n of

pha

se

Figure 6 Coarsening data from Fe-01C-9Cr-05Mn-03Si wtsteel containing a variety of tungsten additions tested in creep at600∘C [13]

is also some solubility of Fe Cr and Mo in these MC Thecarbides MoC and WC were observed to have hexagonalstructures Other precipitates of this form such as VN werefound to be desirable because they form a fine distributionof small and stable particles which are beneficial to creepstrength VN and NbC could be found in a 12 wt Cr steelcontaining V andNb VNwas found as a plate-like phase andNbC in a more spheroidal form

After long exposure to elevated temperatures Laves phaseprecipitates and it has the general composition Fe

2M where

the alloy content may be W Mo or a combination of bothLaves phase is found to have a higher interfacial energy It

has often been observed in 9ndash12wt Cr steels containing Wor Mo and in 9Cr1Mo steel of approximate composition

11 wt Si 44wt Mo 17 wt Cr and 28wt Fe

It was also found that this phase tends to form preferentiallyin Cr-depleted regions These regions were found to becomparatively rich in Mo indicating that a principal kineticconstraint on Laves phase formation is the diffusion of Moatoms It has been established that increasing the Si contentof 9Cr-2Mo steel accelerates and increases Laves phaseprecipitation It was found that Mn additions were useful inretarding the formation of Laves phase [45] Precipitationof Laves phase on lath boundaries decreases creep rateand suppressed the recovery of the microstructure A Sicontent of 03 wt was found to provide the best balancebetween recovery of the microstructure and coarsening ofthe precipitate of some creep steel Laves phase has beenfound to initially form on prior austenite grain boundariesfollowed by precipitation on lath boundaries and finally someprecipitation within the laths [45] It may also benefit fromnucleation on existing phases such as M

23C6 Study found

formation of Laves phase at grain boundaries andmartensiticlath interfaces were an important factor in maintaining creeprupture strength For tungsten containing steels Laves phaseprecipitation is found to be beneficial for creep strengththan for solid solution strengthening by tungsten Lavesphase containing W is found to be thermally stable butthe coarsening rate of the particles after precipitation ishigh [45] This gives large Laves phase precipitates afterseveral thousand hours of creep exposure indicating that thebeneficial effects for creep strength would not last into thelong-term

Laves phase is seen to be the equilibrium precipitate attemperatures below sim600∘C in the 10CrMoV steel The vari-ation of the equilibrium volume fraction with temperatureis shown in Figure 7 From the graph Laves phase is notobserved beyond about 600∘C

There have been a number of studies of Laves phaseprecipitation in 9CrMoWV steel (NF616) [46] a com-mercial alloy designated for service at temperatures above600∘C in power plant The steel (wt Fe-0106C-896Cr-047Mo-0051N-0069Nb-020V-183W) contains tungstenwhich makes it more susceptible to Laves formation whichoccurs over a greater range of temperatures than in the10CrMoV steel In the study in [47] Laves phase (Fe

2Mo) and

progressive coarsening of (CrFe)23C6carbides with respect

to increasing test duration and temperature were observedThe formation and growth of Fe

2Mo are reported to occur at

Cr2Nmatrix interface caused by the depletion of Cr solutes

in the vicinity of Cr2N precipitates [47]

Another phase usually found in 12CrMoVNb steels attemperature greater than 550∘C is Z-phase This phase isa complex nitride phase containing Cr V and niobium(Nb) (of the form Cr(Nb V)N and it is found to form inclose association with dissolving NbX particles [47 48] Asthis phase grows it is seen to gradually replace MX andM2X nitrogen-rich phases which improve the creep rupture

strength [49] Therefore Z-phase directly contributes to areduction in the creep rupture strength of the observed


0

0002

0004

0006

0008

001

500 550 600 650

Equi

libriu

m v

olum

e fra

ctio

n

Temperature (∘C)

Figure 7 Variation of equilibrium volume fraction of Laves phasewith temperature in 10CrMoV steel calculated using MTDATA [7]

12CrMoVNb steel Precipitation of large Z-phase particlesCr(V Nb)N which depletes the matrix of fine MX nitrides(V Nb)N was observed to be the major cause of earlybreakdown in long-term creep strength of a number of new9ndash12Cr martensitic steels [50]

Study shows that large concentration of niobium(09 wt) enhances the precipitation of 119885-phase In 12Crsteel used by [51] precipitation of Z-phase was observed to beaccomplished by Cr atoms diffusing into (V Nb)N particlesand later this structure transforms to cubic or tetragonalZ-phase It was also reported that Z-phase formed rapidlyat prior austenite grain boundaries Also the formation wasfastest at 650∘C

In general the precipitation of Z-phase M6X and

Fe2(WMo) Laves phase during creep causes a loss of creep

strength at long times because they consume existing fineM2X andMX and orM

23C6precipitates [42 50] Retardation

of the Z-phase formation is necessary to improve strength ofcreep-resistant steels [49]

Figure 8 is the sigmoidal behaviour in creep rupture datafor four creep steels at 650∘C showing the effect of Z-phaseformation [16]

The synergetic effect of Z-phase precipitation and tung-sten depletion of solid solution due to Fe

2W Laves phase

formation is believed to be the reason for the sigmoidalshape of creep rupture strength curve of TAF 650 steel Theloss of creep strength in the modified 9Cr-1Mo (or T91)is reported [16] to be due to preferential recovery of themicrostructure in the vicinity of prior austenite boundariesas shown in Figure 9 The preferential recovery promotesthe onset of acceleration creep and hence causes prematurerupture The dissolution of MX and the precipitation of Z-phase are reported to promote the preferential recovery

Mod 9Cr-1MoTAF 650

TAFTAF 650

40

400

Stre

ss (M

Pa)

100 1000 10000 10000010Time to fracture (h)

650∘C

Figure 8 Comparison of stress versus time to rupture curves forvarious martensitic steels at 650∘C adapted from [16]

1 120583m

Figure 9 TEM micrograph after creep rupture testing for 34141 hrat 600∘C and 100 MPa [16]

Generally it is believed that the loss of creep strengthin 9ndash12Cr steels at 550∘C and above is due mainly to thedissolution of fine M

2X and MX carbonitrides during creep

exposure which is driven by the precipitation of coarse Z-phase and or M

6X carbonitrides which are thermodynami-

cally more stable than M2X and MX [16 52ndash54]

A report has been made on the discovery of M5C2in ex-

service 1Cr-05Mo creep steels [55] This phase is said to havea monoclinic structure and has not been identified beforein creep-resistant Cr-No steels The precipitates were foundto be rod-shaped and appeared to nucleate heterogeneouslyon M2C and remain in ferrite regions from which M

2C had

disappeared suggesting that M5C2is more stable than M

2X

The discoveries of M5C2and Z-phase are more recent than

the other phases in ferritic power plant steels


400 nm

(a) (10min)

400 nm

(b) (30min)

400 nm

(c) (1 hr)

400 nm

(d) (16 hr)

Figure 10 This figure shows micrographs which are examples of all the transitions observed at 600∘C [7]

Table 7 Heat treatments of 10CrMoV steel for carbide stabilitystudy [7]

Temperature ∘C Time h600 016 05 1 4 16650 016 05 2 8800 005 016

A visual picture of some of these carbides can be seenfrom the study carried out to understand the evolution ofmicrostructure in the early stages of tempering in a bidto establish carbide stability diagram for 10CrMoV steelThe starting material was in quenched condition The heattreatment given to the test piece was as listed in Table 7

The result of series ofmicrographs examined after a giventime using TEM of samples held at 600∘C is as shown inFigure 10

In Figure 10(a) analysis shows that the small needle-shaped particles are M

2X precipitated along with M

3C

After 30 minutes the M2X precipitate intensified while

the M3C volume fraction remained approximately constant

(Figure 10(b)) After 1 hour there were additional blockyprecipitates seen which were identified as M

23C6 M3C was

present as well as M2X (Figure 10(c)) Finally after 16 hours

the microstructure consisted of relatively large M23C6parti-

cles localized to lath boundaries along with M2X occurring

within the laths (Figure 10(d))The sequence observed in this test was summarized as

followsM3C +M

2X 997888rarr M

3C +M

2X +M

23C6

997888rarr M2X +M

23C6

(2)

It was reported that some fine MX phase could be presentalthough none was said to be found in this particularinvestigation The same sequence was reported for test at650∘C but the precipitation kinetics was faster for all phasesat this temperature

At 800∘C M23C6was observed after only 3min by

which time M3C had dissolved The approximate carbide


stability diagram of this experiment can be found in thesepublications [26 45] At 600∘C the equilibrium precipitatefor the 225Cr1Mo steel isM

23C6 while for the 10CrMoV steel

the equilibrium precipitates are M23C6and M

2X [7]

A common observation reported here is that once ametastable phase has completed precipitating (or enriching)it will start to dissolve to provide solute for all the more stablephases whose extent of reaction is less So in the case of the10CrMoV steel M

3C and Laves phase eventually dissolve to

provide solute toM2X andM

23C6which are the stable phases

For the 10CrMoV steel the order of events is first the rapidcompletion of M

3C enrichment as the steel has a very large

Cr concentrationThis is followed by the completion of Lavesprecipitation for this particular temperature (600∘C) Lavesphase is ametastable (brittle) phase It has been explained thatbecause the precipitation of Laves phase is extremely sluggishat the time it stops precipitating (sim30min) the volumefraction which has formed is very small (sim2 times 10minus9) Afterthis time it dissolves rapidly and so is not experimentallyobserved Finally M

2X and M

23C6form as the equilibrium

phasesThis order of precipitation and enrichment is also

observed in 225Cr1Mo steel but due to low Cr concentra-tion and higher C content in this steel it is reported thatM2C precipitation occurs and is finished well before M

23C6

precipitation and the volume fractions of M2C and M

3C

are greater Therefore precipitation or enrichment of thesephases removes a greater proportion of the available solutefor the matrix than is removed in the 10CrMoV steel Thereduction in the solute available for M

23C6suppresses its

formation to later times [7]

6 Effect of Alloying Elements of Creep Steel onPhase Stability

The large number of alloying elements in creep steel producesa huge variety of microstructure where the matrix can differas well as the composition The composition influences thephasesrsquo precipitation sequences and distribution

The elements that make the composition of these steelshave specific functions even though theremay be interactionamong them Increasing or decreasing their amount in searchof better creep steel hinges on the understanding of theirspecific roles in the alloy system These elements are usually

C Mn Ni Co Bo Cu Cr Si W V Mo Ti Nb Ta(and Nd) Al and N

Elements such as Mn Mo Cr V W Ti and Nb can bothform stable carbides and be found in solid solution withinthe ferrite [26]The studies on power steels show that Cr WandMoprovide solid solution strengtheningTheprecipitatesthat commonly form in this alloy system include M

23C6 as

well as MX (VN and NbC) The drop in the strengtheningeffect of precipitate hardening is due to the coarsening ofthese precipitates [6]

The carbon content of the steel determines the amount ofcarbon in the carbide and the excess remains in solid solutionin ferrite Nitrogen is reported to enter the carbide phase

readily forming carbonitrides with iron and many alloyingelements and separate alloy nitride phases in the presence ofTi or Al [26]

The known characteristics of some of the elements are asfollows

Carbon (C) Carbon stabilizes austenite relative to ferrite Itis essential for the formation of carbides which cause thesecondary hardening of power plant steels It usually haspeak value because high levels of carbon lead to unacceptablereductions in mechanical properties such as toughness andmay cause cracking after normalization and also after weld-ing These effects have been studied by [38 45 56]

Manganese (Mn) Manganese stabilizes austenite but is oftenfound to have an adverse effect on the creep strength of powerplant steel This phenomenon was attributed to the retentionof austenite which will be rich in carbon and nitrogen soreducing the effects of secondary hardening Study showsthat increasing Mn content may increase the growth rate ofM6X an undesirable and coarse phase which can remove W

from solid solution and cause the dissolution of other moredesirable precipitates [57]

Nickel (Ni) Nickel stabilizes austenite and is found to accel-erate the precipitation and coarsening of carbide phase andLaves phaseThis coarsening reduces the creep properties andso is detrimental to creep properties of power plant alloysNi concentrations greater than 04wt have been shownto deteriorate the creep strength of high chromium steelappreciably [54] This is because Ni increased the rate ofdissolution of M

2X as a consequence of precipitation of M

6X

phase It increases the coarsening rate of M6X and M

23C6

[54]In the search for a creep-resistant steel capable of func-

tioning at temperatures of 650ndash750∘C in a very recent study15Cr1Mo ferritic steel was alloyed with increasing Ni contentas shown in Table 8

The result of that endeavour is shown in Figure 1115Cr1Mo steel gave ferritic and martensitic phases upon

quenching with volume fraction of martensite increasingwith Ni addition The precipitates observed were reported tobe Laves phase and 120594-phase (which are intermetallic com-pounds) Cr

23C6(chromium carbide) and Z-phase (Cr(V

Nb)N) [8 31] Ni was seen to cause preferential precipitationofCr23C6at grain boundaries between the ferrite andmarten-

site Once again Ni is seen to accelerate the precipitation ofLaves phase This study seems to show that at 650∘C no effectin creep strength is seen by increasing Ni content of 15Cr1Mosteel but there was huge creep strength improvement incomparison with T92 creep steel whose composition is givenas 9Cr-05Mo-18W-V-Nb at each temperature This hugedifference is said to be due to increased precipitation of Lavephase and chromium carbide The coarsening effect of Ni oncarbide and Lave phases was not reported

Cobalt (Co) Cobalt is austenite stabilizer and is found notto coarsen the steel microstructure like Mn and Ni Studyshows that cobalt strengthens the martensitic matrix and


Table 8 15Cr1Mo steel studied [8]

C Cr Ni Co Mo W V Nb N B08Ni 0051 1516 086 196 103 603 021 006 0037 0003114Ni 0048 1513 141 199 103 602 021 0058 0037 0003121Ni 0051 1507 215 204 098 603 021 0054 0036 00031

100 1000 10000 100000 100000010Time to rupture (h)

20

100

500

Stre

ss (M

Pa)

T92 750∘C T92 700∘CT92 650∘C 15Cr-21Ni 750∘C15Cr-21Ni 700∘C 15Cr-21Ni 650∘C

Figure 11 Comparison of stress versus creep rupture strengthof 15Cr steels with increasing Ni content and T92 at 650ndash750∘Cadapted from [8] also see [17]

enhances the stability of fine precipitates such as MX and120583-phase Cobalt was found to remain in solid solution in12Cr steels evenwith concentrations of up to 10wtThis isbeneficial effect for the creep rupture strength of creep steelsIrvine et al (1960) [38] and Igarashi and Swaragi (1997) [58]have studied this beneficial effect It is thought that cobaltpromotes the formation of a stable oxide layer Howeverthe effects of cobalt additions on oxidation resistance arenot well known Major disadvantage in the use of Co asalloying element is that it makes steel difficult to recycle Inother words cobalt-containing steels have a very low scrapvalue In recent study additions of Co and Ni were observedto strongly influence the creep strength and precipitationbehaviour of intermetallic compounds (Laves phase120594-phaseand 120583-phase) during creep exposure at high temperaturefor 15Cr1Mo6W steel [31] They found that creep strength isincreased by precipitation of fine particles and homogenousdispersion of intermetallic particles in the ferritic matrix

Boron Study on boron additions shows beneficial effect onhigh temperature strength of 10Cr steel It is found that long-term creep rupture strength was improved by B additions upto approximately 001 wt This increase in rupture strengthby boron has been attributed to stabilization of M

23C6at

the grain boundaries retardation of the coarsening of theprecipitate prevention of grain boundary sliding and thereby

dynamic recrystallization Some research work on this can befound in [14 57 59 60]

However there appears to be no benefit to creep ruptureproperties at higher boron percentages In fact excessiveboron is detrimental to strength due to the formation ofundissolved boride phases [45] Increased levels of boronare also found to lead to problems with weldability andforgeability Excessive levels of boron can also combine withnitrogen to form quite coarse BN precipitates This is seen toreduce the amount of finer nitrides in the alloy such as VNwhich can have beneficial effects on creep rupture propertiessuch as ductility [61]

Copper Copper is an austenite stabilizing elements and isoften added purely to counteract the undesirable formation of120575-ferrite Hence copper is often used as an addition to powerplant steels to retard 120575-ferrite formation since it stabilizesthe austenite phase High copper additions up to 2wt havebeen found to cause the precipitation of fine Cu particles onmartensite lath boundaries which pins them and retardingthe recovery of the martensite It was also reported thatCu additions caused refinement of the prior austenite grainsize One concern with large copper additions is that it mayprecipitate and become removed from the alloy during heattreatment Studies in this area have been reported by [18 5362]

Chromium Chromium stabilizes ferrite and is carbide for-mer The addition of chromium to ferritic steels has beenshown to improve the oxidation resistance The high Crsteels are extremely corrosion-resistant but the lower Crsteels are susceptible to corrosion [6] Studies have longestablished that Cr improves the steamoxidation resistance offerrous alloys [63 64] Large chromium additions of 9-10wtprovide hardenability good creep strength and resistance tocorrosion An increase of 11 wt Cr was found to markedlyincrease corrosion resistance at creep steel at 650∘C hence atleast 11 wtCr has been recommended in the design of a steelfor use at 650∘C Disadvantage of high concentrations of Cr isthe promotion of undesirable 120575-ferrite formation Ennis et al(1998) [65] and Abe and Nakazawa (1992) [5] have reportedthis study

Silicon Silicon is a ferrite stabilizer and it is found to influencethe rate at which carbide precipitates It is also found toaccelerate the precipitation and coarsening of Laves phase in9Cr steels Silicon additions likeCr promote the formation ofundesirable 120575-ferrite phase [66] However silicon can be veryimportant in the formation of protective oxidation layersIt was reported that about 08 wt Si is required to form aprotective layer at 650∘C for 9wt Cr steels while 04 wt Si


9 10 11 12 138Cr (wt)

T = 1100∘C

0

10

20

30

40

50

60

70

80

120575-f

errit

e (w

t)

05 wt Si1wt Si

0wt Si

Figure 12 120575-ferrite volume fraction as a function of Cr wt fora steel of composition 008C-01Si-05Mn 813Cr-3W-02V-0005B-0005N wt [18]

is needed in the case of 12 wt Cr steels [6] However thereis optimum silicon content of approximately 03 wt Si togive acceptable corrosion resistance and mechanical strengthsince it causes acceleration of Laves phase precipitation andcoarsening Study has shown that silicon content above04 wt leads to deterioration in tensile strength creeprupture strength and toughness due to 120575-ferrite formationTherefore the formation of 120575-ferrite has been the majorproblem with silicon additions [18]

Figure 12 shows calculation using Thermo-Calc for amodel steel comparing the amount of 120575-ferrite which formsduring austenitization with Cr and Si content

The figure clearly shows that the addition of high Cr or Sicontent will cause 120575-ferrite formation during austenitizationaround 1100∘C To mitigate against this austenite stabiliz-ing element is commonly added These characteristics arebalanced with the effect on the mechanical properties whendesigning ferritic alloys with significant silicon contentsOther reports on Si study can be found in [41 66 67]

Tungsten Tungsten is a ferrite stabilizer and strong carbideformer It has been found to stabilize M

2X though less

effective Tungsten promotes the formation of Laves phaseLike Cr and Si high contents of tungsten may lead to theformation of undesirable 120575-ferrite [45]

In general tungsten addition has been found to improve thecreep rupture strength of ferritic heat-resistant steelsThis effectwas attributed to solid solution hardening and precipitationstrengthening due to M

23C6and Laves phase M

23C6and

Laves phase were found to precipitate at grain boundariesand this is believed to be the main factor responsible for theimproved creep rupture strength Study shows that increasedW concentrations in a 9Cr alloy reduced the coarsening rateof M23C6precipitates In turn this reduced the coarsening

of martensite laths due to the pining effects of M23C6 thus

delaying recovery Report also has it that the coarsening rateof Laves phase was very large in these alloys Reports oneffects of tungsten addition to creep steel can be found readilyin [29 46 68]

Vanadium Vanadium is a ferrite stabilizer and strong carbideformer Vanadium is a key element due to the stability of itsnitride and this nitride is found to improve the creep rupturestrength of steel Like W it stabilizes M

2X phase and also

MX Vanadium additions to 10ndash12wt Cr steels have beenfound to significantly improve the long-term creep strengthThe strengthening effect of vanadium may also be enhancedby the removal of available carbon into vanadium carbidesthus leaving Mo in the matrix where it contributes to solidsolution strengthening Reports on this study can be foundin [6 69ndash72]

Molybdenum Molybdenum forms stable carbides andremains in ferrite contributing to solid solution strengthen-ing Like W and Va Mo has been shown to stabilize M

2X

phase and a phase given as M3C6phase Increasing Mo

concentration above 1 wt has been found to promoteformation of M

6C Laves phase and 120575-ferrite in 9wt Cr

steels Study in this area can be found in [69 71 73]

Titanium Titanium is strong carbide and nitride formerand improves the creep rupture strength of ferritic steelsHowever increase in Ti may result in a large reductionin rupture ductility due to Ti based inclusions preventingdeformation at grain boundaries

Niobium Niobium addition is found to give MX carbidesand nitrides which are small and stable These precipitateshelp to pin grain boundaries and prevent grain growthThe effectiveness of this mechanism relies on dissolving Nbduring austenitization otherwise insoluble NbC remainscausing coalescence of the precipitates coarsening them andaccelerating recovery Increased level of Nb may also causethe formation of undesirable Z-phase Further information onthis can be found in [69]

Tantalum (and Neodymium) The additions of Ta and Ndto creep steel have been found to be beneficial to the creeprupture strength of weld joints in a 01C-11Cr-3W-3Co-V-Nb-N steel Igarashi and Swaragi (1997) [58] reported thatthis effect was attributed to the formation of TaN and NdNprecipitates These nitrides were found to be extremely stableand do not coarsen appreciably during welding and thesubsequent heat treatment

Aluminum The addition of Al is thought to be detrimentalto the creep rupture strength since it leads to the formationof AlN instead of VN or NbN AlN is a coarse phase andnot useful to creep strength while the others are fine anddispersed and so improve creep property

Nitrogen Nitrogen like carbon is small enough to occupyinterstitial sites in the iron lattice and stabilizes austeniterelative to ferrite It has been found to stabilizeMXprecipitate


in Cr steels probably of the composition CrN and thishas been one of the reasons suggested why N additionsmay increase creep strength [38] Nitrogen addition is foundto aid the formation of MX type precipitate such as VNand NbN which are fine and generally desirable for creepstrength Potential disadvantage militating against full uti-lization of nitrogen is said to be the difficulty associatedwith increasing nitrogen content of ferritic steels aboveapproximately 008wt without pressurizing the cast in anitrogen atmosphere This adds to the cost of designing

However study shows that addition of nitrogen to boroncontaining steels may precipitate BN which can offset thebeneficial effects of both B andN Report on effect of nitrogenon creep steel can also be found in [38 61 74]

Phosphorus The presence of phosphorus in substantialamount is found to lead to segregation to grain boundarieswhich embrittles the alloy Study shows that it reduces thevolume fraction of beneficial M

2X precipitates and increases

creep rate Phosphorus is said to interact and drive thesegregation of Mo at prior austenite grain boundaries andresults in loss of ductility in 9Cr1Mo steel Pilling et al (1982)[70] andNoble et al (1990) [75] have some report on the effectof this element on creep steel

Sulphur Formation of sulphides such as MnS can pro-vide preferential nucleation sites for cavitation Precipitationoccurs on prior austenite grain boundaries when coolingfrom the austenite The cavitation is caused by the lowinterfacial adhesion betweenMnS and the ferrite matrix [44]

It is usually useful to examine the combined effectof elements that behave in a similar manner rather thanindividual effects For example 120575-ferrite promoters are Cr SiW and Mo For the effect of W and Mo concentrations on9Cr alloys it is found that in the range 05 gtMo wt gt 08Mo improved the creep rupture strength [76]This is just oneexample of the many interactions between variables whichmake alloy design extremely difficult

From above discussions one can see that the elementsthat make up the creep steel composition are added for avariety of reasons The carbide formers are Ti W Va Cr NbMo Si and the undesirable 120575-ferrite promoters tend to beferrite stabilizers like Si Mo Cr W V and the counteractingelements the austenite stabilizers that is C Mn Ni Co Cuand N

7 Conclusion

We can see that power plant steels are designed by con-sidering the known effects of each single element and theirinteractions This interaction makes the design of this classof steel extremely challenging venture These elements areadded either to stabilize phases which are beneficial to creepresistance or to suppress phases which are detrimental Thenature of steel has created opportunity to design steels thathave served man in the high temperature domain up to650∘C for decadesThere are emerging possibilities of alloyingwith nickel and boron to produce steel that functions at 700∘Cand above where the alloys are independent of any carbides

as strengthening factors The current studies on 15Cr steelare keeping the hope alive The success of this is expected toincrease tremendously the efficiency of power plants by 50and the concomitant reduction in the emissions of gases suchas CO

2 NO119909 and SO

119909 The immense possibility of altering

steel microstructure to create wide ranges of mechanicalproperties underscores the critical role that steel has playedin the construction of our society It just appears that thepotential of steel is limitless

Competing Interests

The authors declare that there are no competing interestsregarding the publication of this paper

Acknowledgments

The authors are grateful to HKDHBhadeshia of University ofCambridge for his numerous insightful and inspiring papersand his willingness to be open in information disseminationin phase transformation study They relied heavily on hisresearch activities They also wish to commend David Gra-ham Cole whose PhD thesis is a mine of information Hiswork aided in no small measure in this review

References

[1] J Chen Modelling of simultaneous transformations in steels[PhD thesis] Department ofMaterials Science andMetallurgyFitzwilliam College University of Cambridge 2009

[2] H K D H Bhadeshia ldquoNovel aspects of ferritic steels for thegeneration of electricityrdquo in Proceedings of the 4th InternationalSymposium onMechanical Science based on Nanotechnology pp143ndash146 Sendai Japan 2007

[3] M Murugananth and H K D H Bhadeshia ldquoComponents ofthe creep strength of weldsrdquo inMathematical Modelling of WeldPhenomenamdashVI H Cerjak andH K DH Bhadeshia Eds pp243ndash260 Institute of Materials 2002

[4] H K D H Bhadeshia A Strang and D J Gooch ldquoFerriticpower plant steels remanent life assessment and approach toequilibriumrdquo International Materials Reviews vol 43 no 2 pp45ndash69 1998

[5] F Abe and S Nakazawa ldquoMicrostructural evolution and creepbehaviour of bainitic martensitic and martensitendashferrite dualphase Cr-2W steelsrdquo Materials Science and Technology vol 8no 12 pp 1063ndash1069 1992

[6] D G Cole Designs of heat-resistant steels for small powerplant [PhD thesis] Materials Science and Metallurgy DarwinCollege University of Cambridge 2000

[7] J D Robson and H K D H Bhadeshia ldquoModelling precip-itation sequences in powerplant steels Part 2mdashapplication ofkinetic theoryrdquoMaterials Science and Technology vol 13 no 8pp 640ndash644 1997

[8] M Shibuya Y Toda K Sawada H Kushima and K KimuraldquoImproving the high-temperature creep strength of 15Cr ferriticcreep-resistant steels at temperatures of 923ndash1023Krdquo MaterialsScience and Engineering A vol 652 pp 1ndash6 2016

[9] H K D H Bhadeshia ldquoEstimation of the microstructure andproperties of ferritic creep-resistant steelsrdquo in Modelling ofMicrostructural Evolution in Creep ResistantMaterials A Strang


and M McLean Eds pp 15ndash38 The Institute of MaterialsLondon UK 1999

[10] H K D H Bhadeshia ldquoThermodynamic analysis of isothermaltransformation diagramsrdquoMetal Science vol 16 no 3 pp 159ndash165 1982

[11] F Brun T Yoshida J D Robson V Narayan H K DH Bhadeshia and D J C MacKay ldquoTheoretical design offerritic creep resistant steels using neural network kinetic andthermodynamicmodelsrdquoMaterials Science and Technology vol15 no 5 pp 547ndash554 1999

[12] D Cole CMartin-Moran A G Sheard H K D H Bhadeshiaand D J C MacKay ldquoModelling creep rupture strength offerritic steel weldsrdquo Science and Technology of Welding andJoining vol 5 no 2 pp 81ndash89 2000

[13] F Abe ldquoCoarsening behaviour of martensite lathes in temperedmartensitic 9Cr-Wsteels during creep deformationrdquo inProceed-ings of the 4th International Conference on Recrystallization andRelated Phenomena pp 289ndash294The Japan Institute of Metals1999

[14] F Abe ldquoResearch and development of advanced ferritic steelfor 650∘C USC boilersrdquo in Proceedings of the at 4th Ultra-SteelsWorkshop that took place in 2000 at Tsukuba InternationalCongress CenterTitle Ultra-Steel 2000 Innovative StructuralMaterials for Infrastructure in 21st Century Sponsor pp 119ndash129National Research Institute for Metals Tsukuba Japan 2000

[15] H K D H Bhadeshia ldquoDesign of ferritic creep-resistant steelsrdquoISIJ International vol 41 no 6 pp 626ndash640 2001

[16] F Abe High Performance Creep Resistant Steels for 21st CenturyPower Plants National Institute for Materials Science (NIM)Tsukuba Japan 1999

[17] M Shibuya Y Toda K Sawada H Kushima and K KimuraldquoEffect of precipitation behavior on creep strength of 15Crferritic steels at high temperature between 923 and 1023KrdquoMaterials Science and Engineering A vol 592 pp 1ndash5 2014

[18] F Abe M Igarashi N Fujitsuna and K Kimura ldquoResearchand development of advanced ferritic steels for 650∘C USCboilersrdquo in Proceedings of the 6th Liege Conference in Materialsfor Advanced Power Engineering vol 5 part 1 pp 259ndash268 1998

[19] F Kauffmann T Klein A Klenk and KMaile ldquoCreep behaviorand in-depth microstructural characterization of dissimilarjointsrdquo Science and Technology of Advanced Materials vol 14no 1 Article ID 014203 2013

[20] H K D H Bhadeshia ldquoCreep-resistant steel-case studyrdquo inKinetics and Microstructure Modelling Lecture Note for CourseMP6 pp 1ndash6 Department ofMaterials ampMetallurgy Universityof Cambridge 2013

[21] R Viswanathan J F Henry J Tanzosh G Stanko J Shin-gledecker and B Vitalis ldquoMaterials for ultrasupercritical coalfired power plant boilersrdquo in Proceedings of the 29th Interna-tional Technical Conference onCoalUtilization and Fuel SystemsClearwater Fla USA 2004

[22] A Helmrich Q Chen G Stamatelopoulos and B ScarlinldquoMaterials development for advanced steamboilersrdquo inProceed-ings of the 8th Liege Conference onMaterials for Advanced PowerEngineering J Lecomte-Beckers M Carton F Schubert and PJ Ennis Eds vol 2 pp 931ndash948 September 2006

[23] W D CallisterMaterials Science and Engineering An Introduc-tion JohnWiley amp Sons New York NY USA 7th edition 2007

[24] J Rosler H Harders and M Baeker Mechanical Behaviour ofEngineeringMaterialsMetals Ceramics Polymers andCompos-ites Springer Berlin Germany 2007

[25] H K D H Bhadeshia Bainite in Steels TransformationsMicrostructure and Properties The Institute of Materials Lon-don UK 2nd edition 2001

[26] R C Thomson Carbide composition changes in power plantsteels as a method of remanent creep life prediction [PhD thesis]Department of Materials Science and Metallurgy NewnhamCollege University of Cambridge 1992

[27] FMasuyama andT Yokoyama ldquoHigh temperaturematerials forpower engineeringrdquo inMaterials for Advanced Power Engineer-ing D Coutsouradis J H Davidson J Ewald et al Eds pp301ndash308 Kluwer Academic Dordrecht The Netherlands 1994

[28] K Kimura H Kushima F Abe K Yagi and H IriesldquoAssessment of creep strength properties of 9ndash12Cr steelsrdquo inAdvances in Turbine Materials Design and Manufacturing 4thInternational Charles Parsons Turbine Conference pp 257ndash269The Institute of Materials London UK 1997

[29] H K D H Bhadeshia Mechanical Properties of Martensite inHeat-resistant Steels Proceedings of Ultra-Steel 2000 NationalResearch Institute for Metals Tsukuba Japan 2000

[30] K Kimura Y Toda H Kushima and K Sawada ldquoCreepstrength of high chromium steel with ferrite matrixrdquo Interna-tional Journal of Pressure Vessels and Piping vol 87 no 6 pp282ndash288 2010

[31] M Shibuya Y Toda K Sawada H Kushima and K KimuraldquoEffect of nickel and cobalt addition on the precipitation-strength of 15Cr ferritic steelsrdquoMaterials Science and Engineer-ing A vol 528 no 16-17 pp 5387ndash5393 2011

[32] J D Robson and H K D H Bhadeshia ldquoModelling precipi-tation sequences in power plant steels part 1mdashKinetic theoryrdquoMaterials Science and Technology A vol 13 no 8 pp 631ndash6391997

[33] Y Murata M Morinaga and R Hashizume ldquoDevelopmentof ferritic steels for steam turbine rotors with the aid of amolecular orbital methodrdquo in Advances in Turbine MaterialsDesign and Manufacturing 4th International Charles ParsonsTurbine Conference A Strang W M Banks and R D ConroyEds pp 270ndash282The Institute of Materials London UK 1997

[34] M Morinaga R Hashizume Y Murata et al ldquoElectronicapproach to the design of ferritic steels for power plantsrdquo inMaterials for Advanced Power Engineering D Coutsouradis JH Davidson J Ewald et al Eds part 1 pp 319ndash328 KluwerAcademic Dordrecht The Netherlands 1994

[35] C J Tillman and D V Edmonds ldquoAlloy carbide precipitationand aging during high-temperature isothermal decompositionof an Fe-4Mo-02C alloy steelrdquoMetals Technology vol 1 no 10pp 456ndash461 1974

[36] V K SikkaM G Cowgill and BW Roberts ldquoCreep propertiesof modified 9Cr-1Mo steelrdquo in Proceedings of the Topical Confer-ence on Ferritic Alloys for Use inNuclear Energy Technologies pp412ndash423 TMS-AIME Warrendale Pa USA 1984

[37] R G Baker and J Nutting ldquoThe tempering of 225Cr1Mo steelafter quenching and normalizingrdquo Journal of the Iron and SteelInstitute vol 191 no 3 pp 257ndash268 1959

[38] K J Irvine and F B Pickering ldquoThe temperature characteristicsof low-carbon low-alloy steelsrdquo Journal of Iron and Steel Insti-tute vol 194 no 2 pp 137ndash153 1960

[39] J M Titchmarsh ldquoThe identification of second-phase particlesin steels using an analytical transmission electron microscoperdquoin Proceedings of the 9th International Conference on ElectronMicroscopy JM Sturgess Ed vol 1 pp 618ndash619MicroscopicalSociety of Canada Toronto Canada 1978


[40] J Beech andDHWarrington ldquoM7C3toM23C6transformation

in chromium containing alloysrdquo Journal of the Iron and SteelInstitute vol 204 no 5 pp 460ndash468 1966

[41] J H Woodhead and A G Quarrell ldquoRole of carbides in low-alloy creep resisting steelsrdquo Journal of the Iron and Steel Institutevol 203 pp 605ndash620 1965

[42] F Abe ldquoBainitic and martensitic creep-resistant steelsrdquo CurrentOpinion in Solid State and Materials Science vol 8 no 3-4 pp305ndash311 2004

[43] D V Edmonds and R W K Honeycombe ldquoStructure andproperties of an isothermally transformed Fe-4Mo-02C alloyrdquoJournal of the Iron and Steel Institute vol 211 pp 209ndash216 1973

[44] J Nutting ldquoThe structural stability of low alloy steels for powergeneration applicationrdquo in Advanced Heat Resistant Steel forPower Generation R Viswanathan and J Nutting Eds pp 12ndash30 Institute of Materials Minerals and Mining 1999

[45] J D RobsonModelling of carbide and Laves phase precipitationin 9ndash12 wt chromium steels [PhD thesis] Materials Scienceand Metallurgy Clare College University of Cambridge 1996

[46] M Ohgami HMinura andH Naoi ldquoCreep rupture propertiesand microstructures of a new ferritic W containing steelrdquo inProceedings of the 5th International Conference on Creep ofMaterials pp 69ndash73 Lake Buena Vista Fla USA 1992

[47] B K Choudhary ldquoTertiary creep behaviour of 9Cr-1Mo ferriticsteelrdquo Materials Science and Engineering A vol 585 pp 1ndash92013

[48] A Strang and V Vodarek ldquoZ phase formation in martensitic12CrMoVNb steelrdquoMaterials Science andTechnology vol 12 no7 pp 552ndash556 1996

[49] K Sawada K Suzuki H Kushima M Tabuchi and K KimuraldquoEffect of tempering temperature on Z-phase formation andcreep strength in 9Cr-1MondashVndashNbndashN steelrdquo Materials Scienceand Engineering A vol 480 no 1-2 pp 558ndash563 2008

[50] H K Danielsen and J Hald ldquoOn the nucleation and dissolutionprocess of Z-phase Cr(VNb)N in martensitic 12Cr steelsrdquoMaterials Science and Engineering A vol 505 no 1-2 pp 169ndash177 2009

[51] L Cipolla H K Danielsen D Venditti P E Di Nunzio J Haldand M A J Somers ldquoConversion of MX nitrides to Z-phase ina martensitic 12 Cr steelrdquo Acta Materialia vol 58 no 2 pp669ndash679 2010

[52] J Pilling and N Ridley ldquoTempering of 225 Pct Cr-1 Pct Mo lowcarbon steelsrdquo Metallurgical Transactions A vol 13 no 4 pp557ndash563 1982

[53] A Iseda Y Sawaragi S Kato and F Masuyama ldquoDevelopmentof a new 01C-11Cr-2W-04Mo-1Cu steel for large diameter andthick wall pipe for boilersrdquo in Creep Characterization Damageand Life Assessments Proceedings of the 5th International Con-ference on Creep of Materials Lake Buena Vista Florida USAASM International 1992

[54] V Vodarek and A Strang ldquoEffect of nickel on the precipitationprocesses in 12CrMoV steels during creep at 550∘Crdquo ScriptaMaterialia vol 38 no 1 pp 101ndash106 1997

[55] S DMann D GMcCulloch and B CMuddle ldquoIdentificationof M5C2carbides in ex- service 1Cr-05Mo steelsrdquoMetallurgical

and Materials Transactions A vol 26 article 509 1995[56] T Fujita and K Asakura ldquoEffects of carbon on creep rupture

strength and toughness of high Cr-Mo heat-resisting steelscontainingV andNbrdquoTransactions of the Iron and Steel Instituteof Japan vol 26 no 12 pp 1073ndash1079 1986

[57] K Miyata M Igarashi and Y Sawaragi ldquoEffect of trace ele-ments on creep properties of 006C-225Cr-16W-01Mo-025V-005Nb steelrdquo ISIJ International vol 39 no 9 pp 947ndash954 1999

[58] M Igarashi and Y Swaragi ldquoDevelopment of 01C-11Cr-3W-3Co-V-Nb-Ta-Nd-N ferritic steel for USC boilersrdquo in Proceed-ings of the International Conference on Power Engineering pp107ndash112 Japan Society of Mechanical Engineers (JSME) TokyoJapan 1997

[59] I M Park and T Fujita ldquoLong-term creep rupture propertiesandmicrostructure of 12Cr heat resisting steelsrdquo Transactionsof the Iron and Steel Institute of Japan vol 22 no 11 pp 830ndash8371982

[60] M Hattestrand and H-O Andren ldquoBoron distribution in 9ndash12 chromium steelsrdquoMaterials Science and Engineering A vol270 no 1 pp 33ndash37 1999

[61] F Abe ldquoEffect of boron on microstructure and creep strengthofadvanced ferritic power plant steelsrdquo Procedia Engineeringvol 10 pp 94ndash99 2011

[62] T Tsuchiyama Y Futamura and S Takaki ldquoStrengthening ofheat resistant martensitic steel by Cu additionrdquo in Creep andFracture of EngineeringMaterials and Structures T Sakuma andK Yagi Eds vol 171ndash174 of Key Engineering Materials pp 411ndash418 2000

[63] H L Solberg G A Hawkins and A A Potter ldquoCorrosionof unstressed steel specimens by high temperature steamrdquoTransactions of the American Society of Mechanical Engineersvol 64 pp 303ndash316 1942

[64] I A Rohrig R M Van Duzer and C H Fellows ldquoHightemperature steam corrosion studies at Detroitrdquo Transactions ofthe American Society of Mechanical Engineers vol 66 pp 277ndash290 1944

[65] P J Ennis Y Wouters and W J Quadakkers ldquoThe effects ofoxidation on the service life of 9ndash12wtCr steelsrdquo in AdvancedHeat Resistant Steels for Power Generation San SabastianSpain R Viswanathan and J Nutting Eds pp 457ndash467 IoMCommunications 1998

[66] Y Kato M Ito Y Kato and O Furukim ldquoEffect of Si onprecipitation behaviour of Nb-laves phase and amount of Nbin Solid solution at elevated temperature in high purity 17Cr-05Nb Steelsrdquo Materials Transactions vol 51 no 9 pp 1531ndash1535 1986

[67] N Fujitsuna M Igarashi and G R Booker ldquoAccelerationof Fe3W precipitation and its effect on creep deformationbehaviour of 85Cr-2W-VNb steels with Sirdquo in Creep andFracture of EngineeringMaterials and Structures T Sakuma andK Yagi Eds vol 171ndash174 ofKey EngineeringMaterials pp 469ndash476 2000

[68] J Vanaja K LahaM DMathew T Jayakumar and E RajendraKumar ldquoEffects of tungsten and tantalum on creep defor-mation and rupture properties of reduced activation ferritic-martensiticrdquo Procedia Engineering vol 55 pp 271ndash276 2013

[69] T Fujita K Asakura T Sawada T Takamatsu and Y OtoguroldquoCreep rupture strength andmicrostructure of lowC-10Cr-2Moheat-resisting steels with V and NbrdquoMetallurgical TransactionsA vol 12 no 6 pp 1071ndash1079 1981

[70] J Pilling andN Ridley ldquoTempering of 225 Pct Cr-1 PctMo LowCarbon Steelsrdquo Metallurgical Transactions A vol 13 no 4 pp557ndash563 1982

[71] D J Gooch ldquoCreep fracture of 12CrMoV steelrdquo Metal sciencevol 16 no 2 pp 79ndash89 1982

[72] T Fujita and N Takahashi ldquoThe effects of V and Nb on the longperiod creep rupture strength of 12 wtCr heat-resisting steel


containingMo and BrdquoTransactions of the Iron and Steel Instituteof Japan vol 18 pp 269ndash278 1979

[73] V Foldyna A Jakobova R Riman and A Gemperle ldquoEffectof structural factors on the creep properties of modifiedchromium steelsrdquo Steel Research vol 62 no 10 pp 453ndash4581991

[74] F Masuyama and M Ohgami ldquoCreep behaviour of highnitrogen ferritic steelrdquo in Proceedings of the 7th InternationalSymposium on Aspects of High Temperature Deformation andFracture in Crystalline Materials Y Hosoi H Yoshinaga HOikawa and K Maruyama Eds pp 325ndash332 Japan Instituteof Metals Sendai Japan 1993

[75] F W Noble B A Senior and B L Eyre ldquoThe effect of phos-phorus on the ductility of 9Cr-1Mo steelsrdquo Acta Metallurgica etMaterialia vol 38 no 5 pp 709ndash717 1990

[76] H NaoiM Ohgami Y Hasegawa et al ldquoMechanical propertiesof 12Cr-W-Co ferritic steels with high creep rupture strengthrdquoin Materials for Advanced Power Engineeringrsquo Part 1 D Cout-souradis J J Davidson J Ewald and T Fujita Eds vol 1 pp425ndash434 Kluwer Academic DordrechtThe Netherlands 1994

Submit your manuscripts athttpwwwhindawicom

ScientificaHindawi Publishing Corporationhttpwwwhindawicom Volume 2014

CorrosionInternational Journal of

Hindawi Publishing Corporationhttpwwwhindawicom Volume 2014

Polymer ScienceInternational Journal of



CeramicsJournal of


CompositesJournal of

NanoparticlesJournal of



International Journal of

Biomaterials


NanoscienceJournal of

TextilesHindawi Publishing Corporation httpwwwhindawicom Volume 2014

Journal of

NanotechnologyHindawi Publishing Corporationhttpwwwhindawicom Volume 2014

Journal of

CrystallographyJournal of


The Scientific World JournalHindawi Publishing Corporation httpwwwhindawicom Volume 2014


CoatingsJournal of

Advances in

Materials Science and EngineeringHindawi Publishing Corporationhttpwwwhindawicom Volume 2014

Smart Materials Research



MetallurgyJournal of


BioMed Research International

MaterialsJournal of


Nano

materials


Journal ofNanomaterials


Table 1 Comparison of the coefficient of thermal expansion (CTE) and thermal conductivity for ferritic and austenitic steel at 100∘C Thecompositions (wt) of the steels are also given [2 3]

(a)

Steel CTE120583mminus1Kminus1 Thermal conductivityWmminus1Kminus1

Ferritic 99 249Austenitic 173 163

(b)

Steel AISI type C Mn P S Si Cr NiFerritic 403 015 100 0040 003 050 115ndash130 mdashAustenitic 305 012 200 0045 003 100 17ndash19 105ndash130

frequently employed in power generation This becomes amajor problem due to the fact that steam turbines are oftenturned off and on during their service life This thermalcycling can cause fatigue in austenitic steels This has beenthemain reasonwhy austenitic steels in spite of their superiorcreep strength are rejected in favour of ferritic steels forthe construction of power plant [6] Martensitic steels havesmaller thermal expansion and larger thermal conductivitythan austenitic steels and Ni base superalloys and are nowoffering highest potential to designing thick section steels forpower plant application

In the development of modern steam power plants thegoal is centred around increasing the steam parameters fromthe current optimum of about 620∘C at 28MPa up to 700∘Cat 35MPa in order to improve efficiency and thereby reduceCO2emission [19] Further advanced steam conditions of

700∘C and above have been already initiated to gain netefficiency higher than 50 at Thermie AD700 project inEurope and at DOE vision 21 project [13 20 21] In thelight of available knowledge so far these steam conditionscannot be reached with ferritic-martensitic alloys which arecurrently the state-of-the-art material for steam power plants[19 22] Hence these projects involve the replacement of 9ndash12Cr martensitic steels by Ni base superalloys for the highesttemperature regimes Tominimize the requirement of expen-sive Ni base superalloys as Ni superalloy is about ten timescostlier than high temperature steels 9ndash12Cr martensiticsteels are planned to be used in the next highest temperaturecomponents of such very high temperature plantsTherefore9ndash12Cr martensitic steels are still strongly desired to expandthe present temperature range up to 650∘C and more Itshould be noted that research is also going back to austeniticsteels to see their possible use for 700∘C USC plants in Japan

2 Phase Transformation in Creep Steel forPower Plant

The general rule for determining if creep will occur in metalsor not is [23 24]

119879 gt 03ndash04 (119879119872) (1)

Pure iron has a melting temperature 119879119872 of about 1536∘C

and the equation above suggests that pure iron would creep

at temperature where traditional steam power plant operatesthat is in temperature range of 450ndash565∘C

Since service temperature depends on the melting tem-perature of the steel a direct approach to solving the creepproblem is to use metals with high melting temperatures sothat the temperature at which creep occurs is increased Anexample of suchmetals is tungsten with119879

119872of 3407∘C In this

case creep becomes important above about 1363∘C which isfar above the aforementioned temperature range for powerplants This apparent advantage comes with considerableconsequences Firstly W is several times more expensivethan steel which could make affordability of the power plantvery unlikely for most countries This cost implication andthe attendant low demand will have a negative impact onthe huge investment in the technology Secondly the oxideof tungsten (WO

3) sublimes and so forth is not considered

as a practical solution Another approach could be to usenickel-based superalloys Study has shown that the creeprupture strength of the nickel-based superalloys is verygood at high temperatures and they have a sufficiently lowexpansion coefficient Unfortunately they are about 6ndash10times more expensive than low alloy steels depending on thealloy composition and market prices of alloying elements Inaddition nickel-based superalloys aremore expensive toweldand machine

Almost all steels rely on the transformation betweenaustenite and ferrite for obtaining the desiredmicrostructureWhen the high temperature 119891119888119888-phase austenite in steeldecomposes to the less dense 119887119888119888-phase ferrite a number ofdifferent microstructures and morphologies can form whichdepend on the cooling rate the presence of alloying elementsand the conditions and availability of lower energy nucleationsites for heterogeneous nucleation The transformation tem-perature between austenite and ferrite can be greatly affectedby the presence of solutes such as C Si Mn Ni Mo Cr WTi Nb andV For example 08 wt carbon lowers the secondsolid-state transformation temperature from 910∘C to 723∘C

The transformation temperature can also be altered bychanging the cooling rate from elevated temperatures Arapidly cooled steel transforms at a lower temperature andthere is shorter amount of time available to accomplish thenecessary changes in atomic positions before the materialfinally cools down to room temperature The products ofsuch transformation are metastable phases The advantage of













2X












23C6 M6C M2C











200

300

400

500

600

700

800


Z

Y

X

Ferrite

Bainite

Martensite

Tem

pera

ture

(∘C)










































24C

Cementite Fe3C

Chi Fe2C







3C





3C In










2C [37] In 9ndash








2X phase [26]









2X (ie M

2C)




2X















2X


7C3


23C6



23C6 The carbide





2X and MX to longer




23C6carbides may


23C6precipitates


M23C

Y

Y

6











coarse M6X and Fe




1 120583m

(a)

1 120583m

(b)


(a) (b)











M2X andM



2C toM




6C by diffusion







4C3 There


10

100

1000

Stre

ss (M

Pa)


750∘C700∘C650∘C

600∘C550∘C


Laves phase


M23C6

0

0005

001

0015

002

0025

003

Mol

e fra

ctio

n of

pha

se




2M where





23C6 Study found




2Mo) and









0

0002

0004

0006

0008

001

500 550 600 650

Equi

libriu

m v

olum

e fra

ctio

n

Temperature (∘C)











2W Laves phase


Mod 9Cr-1MoTAF 650

TAFTAF 650

40

400

Stre

ss (M

Pa)


650∘C


1 120583m









2C had


2X




400 nm

(a) (10min)

400 nm

(b) (30min)

400 nm

(c) (1 hr)

400 nm

(d) (16 hr)








3C




23C6 M3C was





followsM3C +M

2X 997888rarr M

3C +M

2X +M

23C6

997888rarr M2X +M

23C6

(2)








2X [7]








2X and M




23C6


3C


23C6suppresses its







23C6 as










6X


23C6












100 1000 10000 100000 100000010Time to rupture (h)

20

100

500

Stre

ss (M

Pa)





23C6at








9 10 11 12 138Cr (wt)

T = 1100∘C

0

10

20

30

40

50

60

70

80

120575-f

errit

e (w

t)

05 wt Si1wt Si

0wt Si






2X though less




23C6and





2X phase and also



2X



















7 Conclusion



2 NO119909 and SO



Competing Interests


Acknowledgments


References



























































































CeramicsJournal of







Biomaterials




Journal of


Journal of





CoatingsJournal of

Advances in








MaterialsJournal of


Nano

materials















2X












23C6 M6C M2C











200

300

400

500

600

700

800


Z

Y

X

Ferrite

Bainite

Martensite

Tem

pera

ture

(∘C)










































24C

Cementite Fe3C

Chi Fe2C







3C





3C In










2C [37] In 9ndash








2X phase [26]









2X (ie M

2C)




2X















2X


7C3


23C6



23C6 The carbide





2X and MX to longer




23C6carbides may


23C6precipitates


M23C

Y

Y

6











coarse M6X and Fe




1 120583m

(a)

1 120583m

(b)


(a) (b)











M2X andM



2C toM




6C by diffusion







4C3 There


10

100

1000

Stre

ss (M

Pa)


750∘C700∘C650∘C

600∘C550∘C


Laves phase


M23C6

0

0005

001

0015

002

0025

003

Mol

e fra

ctio

n of

pha

se




2M where





23C6 Study found




2Mo) and









0

0002

0004

0006

0008

001

500 550 600 650

Equi

libriu

m v

olum

e fra

ctio

n

Temperature (∘C)











2W Laves phase


Mod 9Cr-1MoTAF 650

TAFTAF 650

40

400

Stre

ss (M

Pa)


650∘C


1 120583m









2C had


2X




400 nm

(a) (10min)

400 nm

(b) (30min)

400 nm

(c) (1 hr)

400 nm

(d) (16 hr)








3C




23C6 M3C was





followsM3C +M

2X 997888rarr M

3C +M

2X +M

23C6

997888rarr M2X +M

23C6

(2)








2X [7]








2X and M




23C6


3C


23C6suppresses its







23C6 as










6X


23C6












100 1000 10000 100000 100000010Time to rupture (h)

20

100

500

Stre

ss (M

Pa)





23C6at








9 10 11 12 138Cr (wt)

T = 1100∘C

0

10

20

30

40

50

60

70

80

120575-f

errit

e (w

t)

05 wt Si1wt Si

0wt Si






2X though less




23C6and





2X phase and also



2X



















7 Conclusion



2 NO119909 and SO



Competing Interests


Acknowledgments


References



























































































CeramicsJournal of







Biomaterials




Journal of


Journal of





CoatingsJournal of

Advances in








MaterialsJournal of


Nano

materials








23C6 M6C M2C











200

300

400

500

600

700

800


Z

Y

X

Ferrite

Bainite

Martensite

Tem

pera

ture

(∘C)










































24C

Cementite Fe3C

Chi Fe2C







3C





3C In










2C [37] In 9ndash








2X phase [26]









2X (ie M

2C)




2X















2X


7C3


23C6



23C6 The carbide





2X and MX to longer




23C6carbides may


23C6precipitates


M23C

Y

Y

6











coarse M6X and Fe




1 120583m

(a)

1 120583m

(b)


(a) (b)











M2X andM



2C toM




6C by diffusion







4C3 There


10

100

1000

Stre

ss (M

Pa)


750∘C700∘C650∘C

600∘C550∘C


Laves phase


M23C6

0

0005

001

0015

002

0025

003

Mol

e fra

ctio

n of

pha

se




2M where





23C6 Study found




2Mo) and









0

0002

0004

0006

0008

001

500 550 600 650

Equi

libriu

m v

olum

e fra

ctio

n

Temperature (∘C)











2W Laves phase


Mod 9Cr-1MoTAF 650

TAFTAF 650

40

400

Stre

ss (M

Pa)


650∘C


1 120583m









2C had


2X




400 nm

(a) (10min)

400 nm

(b) (30min)

400 nm

(c) (1 hr)

400 nm

(d) (16 hr)








3C




23C6 M3C was





followsM3C +M

2X 997888rarr M

3C +M

2X +M

23C6

997888rarr M2X +M

23C6

(2)








2X [7]








2X and M




23C6


3C


23C6suppresses its







23C6 as










6X


23C6












100 1000 10000 100000 100000010Time to rupture (h)

20

100

500

Stre

ss (M

Pa)





23C6at








9 10 11 12 138Cr (wt)

T = 1100∘C

0

10

20

30

40

50

60

70

80

120575-f

errit

e (w

t)

05 wt Si1wt Si

0wt Si






2X though less




23C6and





2X phase and also



2X



















7 Conclusion



2 NO119909 and SO



Competing Interests


Acknowledgments


References



























































































CeramicsJournal of







Biomaterials




Journal of


Journal of





CoatingsJournal of

Advances in








MaterialsJournal of


Nano

materials









































24C

Cementite Fe3C

Chi Fe2C







3C





3C In










2C [37] In 9ndash








2X phase [26]









2X (ie M

2C)




2X















2X


7C3


23C6



23C6 The carbide





2X and MX to longer




23C6carbides may


23C6precipitates


M23C

Y

Y

6











coarse M6X and Fe




1 120583m

(a)

1 120583m

(b)


(a) (b)











M2X andM



2C toM




6C by diffusion







4C3 There


10

100

1000

Stre

ss (M

Pa)


750∘C700∘C650∘C

600∘C550∘C


Laves phase


M23C6

0

0005

001

0015

002

0025

003

Mol

e fra

ctio

n of

pha

se




2M where





23C6 Study found




2Mo) and









0

0002

0004

0006

0008

001

500 550 600 650

Equi

libriu

m v

olum

e fra

ctio

n

Temperature (∘C)











2W Laves phase


Mod 9Cr-1MoTAF 650

TAFTAF 650

40

400

Stre

ss (M

Pa)


650∘C


1 120583m









2C had


2X




400 nm

(a) (10min)

400 nm

(b) (30min)

400 nm

(c) (1 hr)

400 nm

(d) (16 hr)








3C




23C6 M3C was





followsM3C +M

2X 997888rarr M

3C +M

2X +M

23C6

997888rarr M2X +M

23C6

(2)








2X [7]








2X and M




23C6


3C


23C6suppresses its







23C6 as










6X


23C6












100 1000 10000 100000 100000010Time to rupture (h)

20

100

500

Stre

ss (M

Pa)





23C6at








9 10 11 12 138Cr (wt)

T = 1100∘C

0

10

20

30

40

50

60

70

80

120575-f

errit

e (w

t)

05 wt Si1wt Si

0wt Si






2X though less




23C6and





2X phase and also



2X



















7 Conclusion



2 NO119909 and SO



Competing Interests


Acknowledgments


References



























































































CeramicsJournal of







Biomaterials




Journal of


Journal of





CoatingsJournal of

Advances in








MaterialsJournal of


Nano

materials







2X (ie M

2C)




2X















2X


7C3


23C6



23C6 The carbide





2X and MX to longer




23C6carbides may


23C6precipitates


M23C

Y

Y

6











coarse M6X and Fe




1 120583m

(a)

1 120583m

(b)


(a) (b)











M2X andM



2C toM




6C by diffusion







4C3 There


10

100

1000

Stre

ss (M

Pa)


750∘C700∘C650∘C

600∘C550∘C


Laves phase


M23C6

0

0005

001

0015

002

0025

003

Mol

e fra

ctio

n of

pha

se




2M where





23C6 Study found




2Mo) and









0

0002

0004

0006

0008

001

500 550 600 650

Equi

libriu

m v

olum

e fra

ctio

n

Temperature (∘C)











2W Laves phase


Mod 9Cr-1MoTAF 650

TAFTAF 650

40

400

Stre

ss (M

Pa)


650∘C


1 120583m









2C had


2X




400 nm

(a) (10min)

400 nm

(b) (30min)

400 nm

(c) (1 hr)

400 nm

(d) (16 hr)








3C




23C6 M3C was





followsM3C +M

2X 997888rarr M

3C +M

2X +M

23C6

997888rarr M2X +M

23C6

(2)








2X [7]








2X and M




23C6


3C


23C6suppresses its







23C6 as










6X


23C6












100 1000 10000 100000 100000010Time to rupture (h)

20

100

500

Stre

ss (M

Pa)





23C6at








9 10 11 12 138Cr (wt)

T = 1100∘C

0

10

20

30

40

50

60

70

80

120575-f

errit

e (w

t)

05 wt Si1wt Si

0wt Si






2X though less




23C6and





2X phase and also



2X



















7 Conclusion



2 NO119909 and SO



Competing Interests


Acknowledgments


References



























































































CeramicsJournal of







Biomaterials




Journal of


Journal of





CoatingsJournal of

Advances in








MaterialsJournal of


Nano

materials




1 120583m

(a)

1 120583m

(b)


(a) (b)











M2X andM



2C toM




6C by diffusion







4C3 There


10

100

1000

Stre

ss (M

Pa)


750∘C700∘C650∘C

600∘C550∘C


Laves phase


M23C6

0

0005

001

0015

002

0025

003

Mol

e fra

ctio

n of

pha

se




2M where





23C6 Study found




2Mo) and









0

0002

0004

0006

0008

001

500 550 600 650

Equi

libriu

m v

olum

e fra

ctio

n

Temperature (∘C)











2W Laves phase


Mod 9Cr-1MoTAF 650

TAFTAF 650

40

400

Stre

ss (M

Pa)


650∘C


1 120583m









2C had


2X




400 nm

(a) (10min)

400 nm

(b) (30min)

400 nm

(c) (1 hr)

400 nm

(d) (16 hr)








3C




23C6 M3C was





followsM3C +M

2X 997888rarr M

3C +M

2X +M

23C6

997888rarr M2X +M

23C6

(2)








2X [7]








2X and M




23C6


3C


23C6suppresses its







23C6 as










6X


23C6












100 1000 10000 100000 100000010Time to rupture (h)

20

100

500

Stre

ss (M

Pa)





23C6at








9 10 11 12 138Cr (wt)

T = 1100∘C

0

10

20

30

40

50

60

70

80

120575-f

errit

e (w

t)

05 wt Si1wt Si

0wt Si






2X though less




23C6and





2X phase and also



2X



















7 Conclusion



2 NO119909 and SO



Competing Interests


Acknowledgments


References



























































































CeramicsJournal of







Biomaterials




Journal of


Journal of





CoatingsJournal of

Advances in








MaterialsJournal of


Nano

materials




10

100

1000

Stre

ss (M

Pa)


750∘C700∘C650∘C

600∘C550∘C


Laves phase


M23C6

0

0005

001

0015

002

0025

003

Mol

e fra

ctio

n of

pha

se




2M where





23C6 Study found




2Mo) and









0

0002

0004

0006

0008

001

500 550 600 650

Equi

libriu

m v

olum

e fra

ctio

n

Temperature (∘C)











2W Laves phase


Mod 9Cr-1MoTAF 650

TAFTAF 650

40

400

Stre

ss (M

Pa)


650∘C


1 120583m









2C had


2X




400 nm

(a) (10min)

400 nm

(b) (30min)

400 nm

(c) (1 hr)

400 nm

(d) (16 hr)








3C




23C6 M3C was





followsM3C +M

2X 997888rarr M

3C +M

2X +M

23C6

997888rarr M2X +M

23C6

(2)








2X [7]








2X and M




23C6


3C


23C6suppresses its







23C6 as










6X


23C6












100 1000 10000 100000 100000010Time to rupture (h)

20

100

500

Stre

ss (M

Pa)





23C6at








9 10 11 12 138Cr (wt)

T = 1100∘C

0

10

20

30

40

50

60

70

80

120575-f

errit

e (w

t)

05 wt Si1wt Si

0wt Si






2X though less




23C6and





2X phase and also



2X



















7 Conclusion



2 NO119909 and SO



Competing Interests


Acknowledgments


References



























































































CeramicsJournal of







Biomaterials




Journal of


Journal of





CoatingsJournal of

Advances in








MaterialsJournal of


Nano

materials




0

0002

0004

0006

0008

001

500 550 600 650

Equi

libriu

m v

olum

e fra

ctio

n

Temperature (∘C)











2W Laves phase


Mod 9Cr-1MoTAF 650

TAFTAF 650

40

400

Stre

ss (M

Pa)


650∘C


1 120583m









2C had


2X




400 nm

(a) (10min)

400 nm

(b) (30min)

400 nm

(c) (1 hr)

400 nm

(d) (16 hr)








3C




23C6 M3C was





followsM3C +M

2X 997888rarr M

3C +M

2X +M

23C6

997888rarr M2X +M

23C6

(2)








2X [7]








2X and M




23C6


3C


23C6suppresses its







23C6 as










6X


23C6












100 1000 10000 100000 100000010Time to rupture (h)

20

100

500

Stre

ss (M

Pa)





23C6at








9 10 11 12 138Cr (wt)

T = 1100∘C

0

10

20

30

40

50

60

70

80

120575-f

errit

e (w

t)

05 wt Si1wt Si

0wt Si






2X though less




23C6and





2X phase and also



2X



















7 Conclusion



2 NO119909 and SO



Competing Interests


Acknowledgments


References



























































































CeramicsJournal of







Biomaterials




Journal of


Journal of





CoatingsJournal of

Advances in








MaterialsJournal of


Nano

materials




400 nm

(a) (10min)

400 nm

(b) (30min)

400 nm

(c) (1 hr)

400 nm

(d) (16 hr)








3C




23C6 M3C was





followsM3C +M

2X 997888rarr M

3C +M

2X +M

23C6

997888rarr M2X +M

23C6

(2)








2X [7]








2X and M




23C6


3C


23C6suppresses its







23C6 as










6X


23C6












100 1000 10000 100000 100000010Time to rupture (h)

20

100

500

Stre

ss (M

Pa)





23C6at








9 10 11 12 138Cr (wt)

T = 1100∘C

0

10

20

30

40

50

60

70

80

120575-f

errit

e (w

t)

05 wt Si1wt Si

0wt Si






2X though less




23C6and





2X phase and also



2X



















7 Conclusion



2 NO119909 and SO



Competing Interests


Acknowledgments


References



























































































CeramicsJournal of







Biomaterials




Journal of


Journal of





CoatingsJournal of

Advances in








MaterialsJournal of


Nano

materials







2X [7]








2X and M




23C6


3C


23C6suppresses its







23C6 as










6X


23C6












100 1000 10000 100000 100000010Time to rupture (h)

20

100

500

Stre

ss (M

Pa)





23C6at








9 10 11 12 138Cr (wt)

T = 1100∘C

0

10

20

30

40

50

60

70

80

120575-f

errit

e (w

t)

05 wt Si1wt Si

0wt Si






2X though less




23C6and





2X phase and also



2X



















7 Conclusion



2 NO119909 and SO



Competing Interests


Acknowledgments


References



























































































CeramicsJournal of







Biomaterials




Journal of


Journal of





CoatingsJournal of

Advances in








MaterialsJournal of


Nano

materials






100 1000 10000 100000 100000010Time to rupture (h)

20

100

500

Stre

ss (M

Pa)





23C6at








9 10 11 12 138Cr (wt)

T = 1100∘C

0

10

20

30

40

50

60

70

80

120575-f

errit

e (w

t)

05 wt Si1wt Si

0wt Si






2X though less




23C6and





2X phase and also



2X



















7 Conclusion



2 NO119909 and SO



Competing Interests


Acknowledgments


References



























































































CeramicsJournal of







Biomaterials




Journal of


Journal of





CoatingsJournal of

Advances in








MaterialsJournal of


Nano

materials




9 10 11 12 138Cr (wt)

T = 1100∘C

0

10

20

30

40

50

60

70

80

120575-f

errit

e (w

t)

05 wt Si1wt Si

0wt Si






2X though less




23C6and





2X phase and also



2X



















7 Conclusion



2 NO119909 and SO



Competing Interests


Acknowledgments


References



























































































CeramicsJournal of







Biomaterials




Journal of


Journal of





CoatingsJournal of

Advances in








MaterialsJournal of


Nano

materials












7 Conclusion



2 NO119909 and SO



Competing Interests


Acknowledgments


References



























































































CeramicsJournal of







Biomaterials




Journal of


Journal of





CoatingsJournal of

Advances in








MaterialsJournal of


Nano

materials




















































































CeramicsJournal of







Biomaterials




Journal of


Journal of





CoatingsJournal of

Advances in








MaterialsJournal of


Nano

materials



Review Article Physical Metallurgy of Modern Creep...

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