Reduction of CO2 emission is an urgent task for powerindustries, especially for fossil power plants. The amount ofCO2 emission decreases with increasing operation tempera-ture of the plants. Plant operation at higher temperature re-quires heat resistant materials with higher creep strength.Thermal fatigue of the structural components is a graveconcern of the fossil power plants subjected to daily orweekly start-and-stop operations. Ferritic steel has a highthermal conductivity and a low thermal expansion coeffi-cient, and is less susceptible to the thermal fatigue thanaustenitic stainless steel. Because of these attractive proper-ties ferritic heat resistant steel has been used extensively inthe fossil power plants. However, operation temperature ofpower plant has been limited owing to the poor creepstrength of ferritic steel.
Alloy development of the advanced ferritic heat resistantsteel started from Cr–Mo steel. V, Nb and N were added toincrease its creep strength, and Cr content was raised tohave better oxidation resistance necessary for the plant op-
erations at higher temperature. The alloy development hasresulted in several advanced ferritic steels, such asMod.9Cr–1Mo steel (T91/P91: 9Cr–1Mo–VNb), NF616steel (T92/P92: 9Cr–1.8W–0.5Mo–VNb), and HCM12Asteel (T122/P122: 11Cr–2W–0.5Mo–CuVNb).1,2) The creepstrength of NF616 and HCM12A is as good as austeniticstainless steel. These high Cr ferritic steels are character-ized by a tempered martensitic lath structure. They com-monly contain MX type carbo-nitrides (M5V, Nb andX5C, N) and M23C6 (M5Cr) type carbides.
The high Cr ferritic steel has been used successfullyunder ultra super critical (USC) steam conditions at 600°C,and may be used at 625°C. USC power plants of the nextgeneration are to be operated at 650°C. Further improve-ment of creep strength is necessary for this advancement.To establish the alloy design concept for this alloy develop-ment, creep deformation and fracture behavior of the highCr ferritic steel have been studied extensively in the lastdecade. Several conferences on the steel and related sub-jects have been held in the last few years. The present paperwill provide an overview of the strengthening mechanisms
ISIJ International, Vol. 41 (2001), No. 6, pp. 641–653
Strengthening Mechanisms of Creep Resistant TemperedMartensitic Steel
Kouichi MARUYAMA, Kota SAWADA2) and Jun-ichi KOIKE
Department of Materials Science, Graduate School of Engineering, Tohoku University, Aobayama, Aoba-ku, Sendai 980-8579Japan. 2) National Research Institute for Metals, Sengen, Tsukuba 305-0047 Japan.
(Received on November 29, 2000; accepted in final form on February 19, 2001 )
The creep deformation resistance and rupture life of high Cr ferritic steel with a tempered martensitic lathstructure are critically reviewed on the basis of experimental data. Special attention is directed to the follow-ing three subjects: creep mechanism of the ferritic steel, its alloy design for further strengthening, and lossof its creep rupture strength after long-term use.
The high Cr ferritic steel is characterized by its fine subgrain structure with a high density of free disloca-tions within the subgrains. The dislocation substructure is the most densely distributed obstacle to disloca-tion motion in the steel. Its recovery controls creep rate and rupture life at elevated temperatures.Improvement of creep strength of the steel requires a fine subgrain structure with a high density of free dis-locations. A sufficient number of pinning particles (MX particles in subgrain interior and M23C6 particles onsub-boundaries) are necessary to cancel a large driving force for recovery due to the high dislocation densi-ty. Coarsening and agglomeration of the pinning particles have to be delayed by an appropriate alloy designof the steel.
Creep rupture strength of the high Cr ferritic steel decreases quickly after long-term use. A significant im-provement of creep rupture strength can be achieved if we can prevent the loss of rupture strength. In thesteel tempered at high temperature, enhanced recovery of the subgrain structure along grain boundaries isthe cause of the premature failure and the consequent loss of rupture strength. However, the scenario isnot always applicable. Further studies are needed to solve this important problem of high Cr ferritic steel.MX particles are necessary to retain a fine subgrain structure and to achieve the excellent creep strength ofthe high Cr ferritic steel. Strengthening mechanism of the MX particles is another important problem left un-solved.
KEY WORDS: steel for elevated temperature service; creep; strengthening mechanism; alloy design; microstructure; microstructural degradation.
of the high Cr ferritic steel with a tempered martensitic lathstructure under creep condition.
2. Heat Treatment and Microstructural Features ofHigh Cr Ferritic Steel
The high Cr ferritic steel is normalized in an austeniteregime (usually at 1 040–1 100°C) and then cooled to roomtemperature. A high Cr concentration (9–12 mass% Cr) en-ables martensitic transformation during air cooling, and amartensitic lath structure with a high density of dislocationsis introduced into the steel after normalizing treatment. Thesteel is then subjected to tempering at a temperature be-low Ae1. During tempering, the martensitic lath structurechanges into a subgrain structure.3)
Figure 1 shows an optical micrograph of high Cr ferriticsteel. Prior austenite grains are divided into packets andfurther into blocks. There are many elongated subgrains inthe blocks, and each subgrain contains a high density offree dislocations in it. High Cr ferritic steel for steam tur-bine rotors is tempered at lower temperature (usually650–730°C) to keep its higher yield stress. It contains ahigh density of free dislocations within subgrains. Steel forboiler tubes and pipes is tempered at higher temperature(usually 750–780°C), and has a lower density of free dislo-cations. Because of the different tempering temperatures,the precipitation sequence is slightly different between therotor and tube/pipe steels.
The following three types of precipitates appear in highCr ferritic steel4): M23C6 carbide, MX carbo-nitride andFe2M Laves phase. In the Cr carbide M23C6 a part of Cratoms are replaced with Fe and other alloying elements.5,6)
The M23C6 particles are formed during tempering, and lo-cated on grain boundaries and sub-boundaries as describedin Fig. 2. Their precipitation may continue during creeptests of the rotor steel tempered at low temperature.7)
The MX carbo-nitrides are classified into VN and NbC.Parts of N atoms in VN and C atoms in NbC are substitutedwith C and N, respectively. There are four types of MX par-ticles. Coarse primary NbX particles remain after normaliz-ing.8) Fine platelet VX and fine spherical NbX form duringtempering.9,10) VX wings are formed on fine NbX particlesduring creep tests.9,10) Precipitation of MX may continueduring the creep tests of the rotor steel.5,8,11,12) The MX par-ticles except the primary NbX are fine and distributed uni-formly within subgrains5,9) as well as on sub-boundaries.
The Laves phase (Fe2M) is an intermetallic compound:Fe2W and Fe2Mo in W and Mo containing steels, respec-tively. It is not present after tempering, but precipitates ongrain boundaries and sub-boundaries during creep tests.8,13)
The rotor steel often contains fine M2X (mainly Cr2N)particles within subgrains.5,6,8,14,15) HCM12A steel has Cuparticles on sub-boundaries.11)
3. Assessment of Strength Mechanisms
As mentioned in Sec. 2, the high Cr ferritic steel containsseveral types of obstacles to dislocation motion. They aresub-boundaries, free dislocations within subgrains, andM23C6, Fe2M and MX particles. Solute atoms such as Moand W can contribute to creep strength. The strengthening
mechanisms attributed to these obstacles are classified intothe following three categories:
Dislocation hardening: sub-boundaries and free disloca-tions
Particle hardening: M23C6 and Fe2M on sub-boundariesand MX within subgrains
Solution hardening: W and Mo in solutionIn this section we assess the contribution of each type ofobstacles to creep strength of ferritic steel.
3.1. Solution Hardening
Minimum creep rates of a iron (0.001mass%C) and a Fe (0.001mass%C)–2.3mass%W alloy are compared in Fig. 3.16) The alloy concentration is expressed in mass%throughout this paper. Both materials do not have the parti-
Fig. 1. A typical optical microstructure of high Cr ferritic steel.
Fig. 2. A schematic drawing of precipitates in high Cr ferriticsteel.
Fig. 3. Minimum creep rates of ferritic steels at 600°C.
cles and the subgrain structure with the free dislocations(referred to as dislocation substructure hereafter). The Waddition reduces the creep rate by three orders of magni-tude. This fact proves that the solid solution hardening iseffective if the dislocation hardening and the particle hard-ening are absent. Iwanaga et al.17) have confirmed the solu-tion hardening in a Fe–9Cr–2Ni alloy with a dislocationsubstructure. It is to be noted in Fig. 3 that the reduction increep rate by the addition of 2.3% W is less significant inthe Fe–9Cr–VNbCN steel with the particles and the dislo-cation substructure than in the a iron without these obsta-cles.
3.2. Particle Hardening by MX within Subgrains
As is evident in Fig. 3, the creep rate of a iron is reducedtremendously by introducing MX particles into the Fe–VNbC steel. This result confirms the particle hardening byMX particles. Two explanations have been proposed on theparticle hardening. One is that MX particles themselves actas obstacles to dislocation motion,18,19) and the other is thatthey slow down recovery of the dislocation substructure andretain the dislocation hardening for longer duration.8)
3.3. Dislocation Hardening
A comparison is made in Fig. 420) between creep rates ofbainitic (B) steel with the dislocation substructure and fer-ritic (F) steel without the dislocation substructure. The dis-location substructure of bainite is similar to that of tem-pered martensite. When MX particles and solute atoms(0.6W10.15Mo) are present, the introduction of the dislo-cation substructure (from the ferritic (F1P1S) steel to thebainitic (B1P1S) steel) lowers the creep rate to 1/70,demonstrating the effectiveness of the dislocation harden-ing in creep. The dislocation hardening is confirmed in Fig. 3 between Fe–9Cr–VNbCN with the dislocation sub-structure and Fe–VNbC without the dislocation substruc-ture, and also in 9Cr–1Mo steel.21)
As is evident in Fig. 4, the dislocation substructure can-not reduce the minimum creep rate (F→B) without parti-cles and solute atoms, since the dislocation substructurecannot be retained during creep without these obstacles.Therefore, particles and/or solute atoms are necessary to re-veal the strengthening by the dislocation substructure. Onthe other hand, a similar amount of particles (F1P vs.B1P) or solute atoms (F1P1S vs. B1P1S) can reducecreep rate more effectively when the dislocation substruc-ture is present. These facts suggest that the dispersed parti-cles and the solute atoms assist the dislocation hardening inaddition to their own roles in strengthening.
3.4. Particle Hardening Due to Particles on Sub-boundaries
Figure 517) shows creep curves of 9Cr–0.1C steel havingM23C6 particles on sub-boundaries and of 9Cr–2Ni steelwithout particle. Both steels have the dislocation substruc-ture typical of tempered martensite. The M23C6 particlesobviously improve creep resistance, since they slow downthe recovery of the dislocation substructure.17–19,22)
Igarashi and Sawaragi23) have studied creep of ferriticsteel with Fe2M particles but without the dislocation sub-structure and MX particles, and have confirmed an increasein creep resistance by the Fe2M particles. Tsuchiyama et
al.24) have demonstrated significant decrease in creep ratedue to stabilization of the dislocation substructure by Cuparticles.
4. Strengthening Mechanism of Advanced High CrFerritic Steel
As proved in Sec. 3, each strengthening mechanism,namely the solution hardening, the dislocation hardening orthe particle hardening, is effective in creep when the othertwo mechanisms are absent. Practical steels have all the ori-gins of the three strengthening mechanisms, i.e. soluteatoms, dislocation substructure and particles. Furthermore,several kinds of particles, for example MX, M23C6 andFe2M, coexist in the steels. We have to specify the strength-ening mechanism of such practical steels in this section. Itis reasonable to assume that the obstacles determining theathermal yield stress at intermediate temperatures also con-trol creep deformation and rupture life. As mentioned inSecs. 3.1 and 3.3, a simple additive rule does not holdamong the three strengthening mechanisms. As shown inSec. 4.1, the most closely spaced obstacles in a material de-termine the athermal yield stress of the material. We haveto discuss which obstacles determine the athermal yieldstress of the high Cr ferritic steel.
4.1. Orowan Stress of a Material Containing SeveralKinds of Point Obstacles
Suppose there are three kinds of point obstacles A, B andC in a material as listed in Table 1. The average spacing of
Fig. 4. Minimum creep rates of ferritic (F) and bainitic (B) steelsat 600°C and 98 MPa. 1P and 1S mean that MX parti-cles and solute atoms (0.6W10.15Mo) were added to thesteels.
Fig. 5. Creep curves of tempered martensitic steels with(Fe–9Cr–2Ni) and without (Fe–9Cr–0.1C) M23C6 parti-cles.
the ith obstacles on a slip plane is denoted as l i. The fol-lowing equation correlates l i to the number density Ni ofthe obstacles on the slip plane:
Ni51/l i2 ....................................(1)
Orowan stress s i of the ith obstacles is defined by the fol-lowing equation and is given in the table:
s i50.8MGb/l i...............................(2)
where M is the Taylor factor (53), G is the shear modulus(64 GPa at 650°C), and b is the length of Burgers vector(0.25 nm). The total number density of the obstacles Nt andthe average spacing l t of the whole obstacles determine theactual Orowan stress of the material containing the threekinds of obstacles. They are given by
.................................(3)
.............................(4)
The Orowan stress sor of the material is given in Table 1(see “total”). It is obvious that the additive rule does nothold between sor of the material and s i due to each kind ofpoint obstacles.
.................................(5)
The value of sor is primarily determined by obstacle A withthe highest density, when the average spacing of the otherobstacles, B1C, is sufficiently greater than that of obstacleA: say more than twice. In such a case, we may neglect theother obstacles B and C when discussing particle strength-ening of the material.
4.2. Obstacle Determining Athermal Yield Stress
Typical values of the volume fraction V, diameter dp andspacing lp of the major particles contained in high Cr fer-ritic steels are listed in Table 2. They were estimated on thebasis of experimental data reported in literature.2,6,8,11) TheOrowan stress of each kind of particles was evaluated byEq. (2). The Orowan stresses due to the particles in thetable are candidates that may give rise to athermal yield
stress s a. The value of s a of the high Cr ferritic steel withtempered martensitic lath structure is in the range of 400 to500 MPa. The M23C6 particles give the largest value ofOrowan stress in the table, but the value is too small to ex-plain the experimental level of s a.
The tempered martensitic lath structure, namely the dis-location substructure, is another plausible candidate deter-mining s a. The athermal yield stress sr due to the free dis-locations within subgrains is given by
sr50.5MGb√—r f .............................(6)
Free dislocation density s f decreases with increasing tem-pering temperature and is in the range of 1–1031014
m22.6,7,25,26) The dislocation density corresponds to s a of240 to 760 MPa. The subgrain structure also contributes tos a. The athermal yield stress s sg due to the subgrain struc-ture is given by
s sg510Gb/l s ................................(7)
where l s is the subgrain width (a short width in the case of elongated subgrains)8,25–29) and in the range of 350 to500 nm in high Cr ferritic steels.7,25,26) The values of l s
give s a of 300 to 450 MPa. The dislocation substructuregives the greater values of s a than the particles and can ex-plain the experimental value of s a. Therefore it can be con-cluded that the dislocation substructure is a major obstaclewhich determine the athermal yield stress of the high Crferritic steel with the tempered martensitic lath structure.
4.3. Obstacle Controlling Creep Rate
At intermediate temperatures, moving dislocations passthrough the dislocation substructure athermally by the aidof applied stress. At elevated temperatures, on the otherhand, diffusion assists dislocations in passing through theobstacles, and yield stress goes down below s a in tensiletests. In this creep regime, the dislocation substructure re-mains to be the major obstacle controlling creep rate, if itdoes not significantly recover during creep tests. Figure 626)
Table 1. Orowan stress s i estimated from the density Ni andspacing l i of each type of obstacles. The obstacles A,B and C are supposed to be free dislocations withinsubgrains, M23C6 and Fe2M particles, respectively.
Table 2. Volume fraction, diameter and spacing of each kindof particles in high Cr ferritic steel, together withOrowan stress estimated form the values of interpar-ticle spacing.
Fig. 6. Recovery of dislocation substructure during creep of11Cr–2.6W–0.1Mo–CoVNb steel at 650°C under 98MPa, together with the creep curve under the creep con-dition and athermal yield stresses estimated from the dis-location substructure.
shows an example of changes in the dislocation substruc-ture which took place during a creep test of 11Cr–2.6W–0.1Mo–CoVNb steel. The free dislocation density withinsubgrains decreases and the subgrain width increases withthe progress of creep deformation. These changes are thetypical recovery process of the dislocation substructure.The athermal yield stresses corresponding to the free dislo-cation density and the subgrain width were estimated byEqs. (6) and (7), respectively, and are drawn in the figure.They are kept at the significantly higher level than the creepstress (98 MPa) over the whole period of creep test. On theother hand, the Orowan stress due to the particles in Table 2is about 150 MPa at the beginning of creep and decreasesduring creep. These facts point out that the dislocation sub-structure is the main obstacle that controls dislocation mo-tion during the whole period of creep tests. The creep strainaccumulation in Fig. 6(a) is accelerated at the late stage ofcreep. This acceleration corresponds to the accelerated re-covery of the dislocation substructure.
5. Recovery Process during Creep
5.1. Recovery of Dislocation Substructure
The dislocation substructure of tempered martensiticsteel is essentially stable even at 650°C, but plastic defor-mation promotes its recovery.19,30) Figure 726,31) shows atypical example of such recovery processes observed inthree high Cr ferritic steels during creep at 650°C under98 MPa. The subgrain width increases and the free disloca-tion density within subgrains decreases with the progress ofcreep deformation.6–8,19,25,26,30,32) The elongated subgrainsrelated to martensitic transformation become equiaxed sub-grains typical of creep deformation.
The evolution of subgrain width given in Fig. 7 is repre-sented as a function of creep strain in Fig. 825,33,34) togetherwith other experimental results. The subgrain width l s in-creases with creep strain and then reaches a stationaryvalue above a critical strain e c. Similar saturation occurs inthe free dislocation density within subgrains. The values ofe c are about 0.1 in many cases.17,20,25,33–35) It should benoted that the difference among the subgrain width versustime curves in Fig. 7(b) disappears in Fig. 8. The relationbetween (l s2l so)/(l s*2l so) and creep strain e is represent-ed by a straight line irrespective of creep conditions andmaterials33–35):
...............................(8)
where l so and l s* are the initial and stationary values ofsubgrain width, and a is a constant independent of creepconditions and materials tested. On the other hand, Blumand Götz8,29) have proposed the following equation relatingx (subgrain width or free dislocation spacing) to creepstrain e :
..............(9)
where xo and x* are the initial and stationary values of x,and k(s c) is a constant depending on creep stress s c. The
stationary values of subgrain width l s* and free dislocationdensity r f* are given by the following equations8,26–28,36):
l s*510Gb/s c ..............................(10)
r f*5(s c/0.5MGb)2 ..........................(11)
where b is the length of Burgers vector, G the shear modu-lus, and M the Taylor factor. These values are primarily de-termined by s c and independent of materials and initialstates of the dislocation substructure.25,33–35)
5.2. Effects of Recovery Speed on Creep DeformationResistance
As is evident in Fig. 7, the recovery speed of the disloca-tion substructure is different among the three steels. It isslower in the 11Cr–2.6W–0.1Mo–CoVNb and 9Cr–1.8W–0.4Mo–VNb steels containing W than the 9Cr–1Mo–VNbsteel without W. Creep curves corresponding to the recov-ery processes are shown in Fig. 7(a). The following rela-
Fig. 7. (a) Creep curves, and changes in (b) subgrain width and(c) free dislocation density within subgrains during creepof three high Cr ferritic steels at 650°C under 98 MPa.
Fig. 8. Evolution of subgrain width as a funtion of creep strain.
tionship is well known between the creep rate e and thesubgrain width37):
e~l s3.....................................(12)
This equation explains that the slow growth of the subgrainwidth in the W containing steel reduces its creep rate andconsequently extends its creep rupture life. The good corre-lation between the recovery speed and the creep rate con-firms the decisive role of the dislocation substructure increep.
6. Microstructure Design of High Cr Ferritic Steel
In the high Cr steel with tempered martensitic lath struc-ture, recovery of the dislocation substructure controls creepdeformation and subsequent creep failure. Eq. (12) suggeststhat a material can have high creep deformation resistanceif a fine subgrain structure is introduced into the material.High resistance to recovery of the dislocation substructureis also essential to keep a good creep resistance for long du-ration as demonstrated in Fig. 7. In this section we discusshow to design the microstructure of the high Cr ferriticsteel for improving creep strength.
6.1. Dislocation Substructure
Creep curves of four steels are compared in Fig. 9 underthe same creep condition. Characteristics of their disloca-tion substructures are summarized in Table 3. The 11Crsteels have smaller subgrain width than the 9Cr steels. Thehigh W concentration of the 11Cr steels may help the for-mation of the fine subgrain structure. The creep resistanceof the 11Cr–2W–0.3Mo–CuVNb steel in Fig. 9(a) is worsethan that of the 9Cr–1.8W–0.4Mo–VNb steel despite thesmaller subgrain width of the former. This result point outthat the other parameter of dislocation substructure, namelythe free dislocation density within subgrains may also havea significant effect on creep resistance of the steel. On thebasis of a result of 1Cr–1Mo–0.25V bainitic steel,74) Abe38)
has proposed that the extent of primary creep strain increas-es with increasing the free dislocation density within sub-grains. We have to assess this proposal. The free dislocationdensity within subgrains increases with decreasing temper-ing temperature. The 11Cr–2.6W–0.1Mo–CoVNb steel de-forms at a lower speed than the 9Cr–1Mo–VNb steel, be-cause of the finer subgrain width and the higher free dislo-cation density of the former. As seen in Fig. 9(b), the densi-ty of free dislocations has a significant effect on the extentof primary creep strain defined as the strain to the mini-mum creep rate. The strain is large in the 9Cr–1Mo–VNband 11Cr–2W–0.3Mo–CuVNb steels with the low densityof free dislocations, and is small in the 11Cr–2.6W–0.1Mo–CoVNb steel with the high density of free disloca-tions. As a result, the specimens reach 1% strain at 10% ofrupture life tr in the former and at 70% of tr in the latter.The result of Fig. 9 is opposite to the proposal of Abe, but aresult similar to Fig. 9 has been reported on a predeformedCu alloy.39) It can be concluded that a high density of freedislocations in addition to a fine subgrain width is better tohave high creep resistance.
It is easy to control the free dislocation density withinsubgrains by changing tempering temperature. The density
of dislocations introduced during martensitic transforma-tion primarily determines the subsequent subgrain size.Singh and Bhadeshia40) have proposed that a large drivingforce for martensitic transformation and a large deforma-tion resistance of the austenitic phase surrounding marten-site domains are required for making a fine subgrain struc-ture. Low transformation temperature of martensite is alsogood to make fine subgrains.41)
A comparison is made in Fig. 1042) between stress–rup-ture curves of 12Cr–1W–1Mo–VNb steel tempered at the
Fig. 9. Creep curves of four high Cr ferritic steels at 650°Cunder 98 MPa. The time scale in (b) is normalized bytime to rupture tr of each material.
Fig. 10. Effect of tempering temperature (750 and 800°C) oncreep rupture strength at 600 and 650°C.
Table 3. Heat treatment and the initial states of dislocationsubstructure of the high Cr ferritic steels tested.
two temperatures. The specimen tempered at the lower tem-perature (750°C) has higher dislocation density and longerrupture life at high stress.42,43) Its superior creep rupturestrength usually disappears at longer rupture life.18,19,42) Thelow temperature tempering sometime gives shorter rupturelife at low stress as seen in Fig. 10, since a high density ofdislocations enhances quick recovery.42,44) It should benoted that a material with a high density of dislocations isuseless unless its premature recovery of dislocation sub-structure is prevented by some means.
6.2. Pinning Particles
Two processes have been pointed out on the recovery ofsubgrain structure.38) (a) Disappearance of a sub-boundaryby annihilation of dislocations constructing the sub-bound-ary, and (b) mutual annihilation of two sub-boundaries withopposite signs. The first process is dominant at least in theearly stage of recovery, though the latter process may occurat the later stage. Dislocations constructing sub-boundariesare in their stable position after tempering, and it is not easyfor them to climb in the sub-boundaries. However, if dislo-cations (represented by the dotted line) glide into the sub-boundaries from the sub-grain interior as shown in Fig. 11,they disturb the balance in the sub-boundaries and bringabout annihilation of dislocations in the sub-boundaries. Asa result, the dislocation density decreases in the sub-bound-aries as well as within the subgrain. The number of disloca-tions dNd supplied to a unit length of sub-boundary duringa time period dt is given by
dNd5r f vgdt ...............................(13)
where r f and vg are the density and the velocity of free dis-locations within subgrains. The annihilation rate of disloca-tions at sub-boundaries is proportional to the product ofdNd and the climb velocity of dislocations. In order to re-duce the annihilation rate, vg is required to be small.Dispersion of fine MX particles is necessary to keep vg
small and retain fine subgrains.45)
The climb velocity vc of dislocations in sub-boundaries isgiven by
vc5AF ...................................(14)
where A is the mobility of dislocations and expressed as
A5Db/kT .................................(15)
The driving force F is given by
F5sdb2spb ..............................(16)
In these equations, D is the diffusion coefficient, k isBoltzmann’s constant, and T is the absolute temperature.The interaction force sdb between dislocations is inverselyproportional to dislocation spacing. Since the interactionforce increases with increasing dislocation density, Eq. (14)can explain the enhanced recovery of the steel tempered atthe lower temperature in Fig. 10. A large pinning force sdbof particles is necessary to cancel the large driving force ofrecovery. The pinning force sp is related to interparticlespacing by Eq. (2). A fine interparticle spacing of M23C6
and Fe2M is needed to keep fine subgrains for long dura-tion.
7. Alloy Design for Stable Dislocation Substructure
7.1. Self-diffusion Coefficient
Diffusion coefficient is the most important parameter increep, and is required to be as small as possible in order toretain fine particles for long duration and to slow down therecovery of the dislocation substructure. Ferritic steel trans-forms from a para- to a ferro-magnetic state below theCurie temperature Tc, and its self-diffusion coefficient is re-duced by the magnetic transformation.75,76) The high Cr fer-ritic steel is used in this temperature regime below Tc. Thediffusion coefficient of ferritic steel at 600°C is given inFig. 12.8) The diffusion coefficient decreases with increas-ing Tc of the material.5,8,75,76) We can control Tc by changingalloy concentration. The removal of Ni, Si and Mn, or addi-tion of Co decreases diffusion coefficient.
A result of COST project in Europe is given in Fig. 13.18)
The abscissa is a microstructural stability parameter definedby the following equation
.........................................(17)
where Ae1 and Ae3 are A1 and A3 temperatures in equilibri-um, and a1, a2, and a3 are experimental constants. The para-meter increases with increasing Ae1 and (Ae3–Ae1). It has
Fig. 11. A schematic drawing of the recovery process of disloca-tions at sub-boundaries.
Fig. 12. Diffusion coefficient of ferritic steel at 600°C as a func-tion of Curie temperature of the steel.
been claimed that ferritic steel has better creep rupturestrength when the parameter is large,5,18,19) though the phys-ical basis of the parameter has not been justified yet. It isknown that the best steel (B2) in the figure has low concen-tration of Si and Mn, leading to a low diffusion coeffi-cient.5,19)
7.2. Pining Particles
Major particles in the high Cr ferritic steel are MXcarbo-nitride within subgrains and M23C6 carbide and Fe2MLaves phase on sub-boundaries. The particles on sub-boundaries act as obstacles to dislocation climb motion andslow down the recovery of subgrain structure.8) Figure 1431)
shows changes in the density of M23C6 and Fe2W particlesduring creep tests of three high Cr ferritic steels at 650°C.The W containing steel includes 9Cr–1.8W–0.4Mo–VNband 11Cr–2.6W–0.1Mo–CoVNb steels. The density ofFe2W in the 11Cr–2.6W–0.1Mo–CoVNb steel is the high-est among the three steels. The density of M23C6 is at leastfive times higher than that of Fe2M, suggesting that M23C6
particles are the major obstacle to dislocation climb.26,31,38,46)
The density of M23C6 is higher in the W containing steel.This fact results in the slower recovery of the dislocationsubstructure in the W containing steel shown in Fig.7.26,31,46) The presence of B in the W containing steels alsocontributes to their high density of M23C6.
Precipitation sequence and particle size measured in10Cr–1W–1Mo–VNb steel (tempered at 730°C for 12 h)
are summarized in Fig. 15.10) The M23C6 and MX particleshave been formed already during tempering, and theircoarsening and agglomeration take place during creep ex-posure. Usual tempering temperature is too high to formFe2W and Fe2Mo. Fine particles of Fe2M precipitate duringthe creep exposure, but they grow faster than the M23C6
particles. A Z phase, Cr(V, Nb)N, also precipitates duringthe high temperature exposure. They are fine just after pre-cipitation, but grow quickly. On the other hand, the MXparticles grow very slowly and any appreciable change inthe diameter cannot be detected at the creep temperature upto a creep exposure time of 43104 h. In order to make gooduse of the pinning particles, we have to understand theirprecipitation and coarsening behavior during high tempera-ture exposure.
7.3. Precipitation of Particles
The fraction of particles w(t) which has precipitated at atime t is represented by the following Johnson–Mehl–Avrami equation.8)
w(t)512exp{2(t/t)3/2}......................(18)
..............(19)
where Co, Cea and Ce
b are the initial concentration of an al-loying element, its equilibrium concentration in matrix, andthe value in particles, respectively. In order to keep W andMo atoms in solution, the rate constant t for the precipita-tion of Fe2M should be large. On the contrary, the precipita-tion proceeds at a faster speed when diffusivity D, densityof precipitation sites Np and/or Co are larger. For example,Fe2W precipitates faster in HCM12A steel than in NF616steel, since Cu particles in the former provide a large valueof Np.
11) Faster precipitation of Fe2W in steel with a high Wconcentration can be explained by its increased value of Co.
7.4. Ostwald Ripening
Coarsening of pining particles is accompanied by theiragglomeration and the consequent reduction of pinningforce sp in Eq. (16). To keep the pinning force at a highlevel, the coarsening rate should be low. The coarsening isdescribed by the following equation:
dpn2do
n5Kdt ...............................(20)
where do and dp are the average diameter of particles before
Fig. 13. Creep strength which gives the rupture lives at 600°Cplotted against a microstructural stability parameter de-fined by Eq. (17).
Fig. 14. Number of particles counted in a unit area as a functionof exposure time at 650°C.
Fig. 15. Precipitation sequence and particle diameter in 10Cr–1W–1Mo–VNb steel.
testing and at a time t, and Kd is a constant. The exponent ndepends on the coarsening mechanism: n52 (coarseningcontrolled by interface diffusion), 3 (volume diffusion), 4(grain boundary diffusion), and 5 (pipe diffusion). The ex-ponent n often takes 3, and Kd is given by the followingequation in this case:
.....................(21)
where g is the interfacial energy of particles, Vma and Vm
b themolar volumes of ferrite matrix and particles, DM
a the diffu-sion coefficient of M atoms in ferrite, ca
M, cbM and c M
a /b theequilibrium mole fraction of M atoms in ferrite, particlesand at particle surface, respectively. The most important pa-rameters in the equation are Da
M and c Ma /b. ca
M is often usedas c M
a /b. Equation (21) points out that growth of particles isslow when DM
a and caM of the major elements in the particles
are low in the ferrite matrix.Coarsening of M23C6 particles on sub-boundaries is de-
scribed by Eq. (20) with n53, and Kd is given by Eq. (21).It has been reported that Kd decreases with decreasing Moconcentration.8) Reduction of Ni6) (see Fig. 19) and Mn47,48)
also lowers the coarsening rate of M23C6 particles due to thedecrease of D. Diameter of M23C6 particles measured in9Cr–W steel after creep rupture is plotted against creep ex-posure time in Fig. 16.38) The figure demonstrates the de-layed coarsening of M23C6 particles in the high W steel.However, the reduced coarsening rate of M23C6 particles inthe W containing steel cannot be explained simply by Eq.(21), and a multicomponent, multiphase coarsening theoryis needed to interpret the coarsening in engineeringsteels.77) It is widely accepted that B reduces the coarseningrate of M23C6.
7,14,49–51) The reduction is ascribed to the seg-regation of B to M23C6.
7,50,51) The segregation has been con-firmed with AP-FIM.11,52) The slow growth of M23C6 stabi-lizes fine subgrains and improves the creep strength of steelwith W and B.
The coarsening rate of MX particles is 1/10 of that ofM23C6.
8) Their coarsening is described by Eq. (20) withn55, suggesting the coarsening controlled by pipe diffu-sion.28) The value of Kd of MX particles is small due to thelow solubility of Nb and V in the ferrite matrix.8) Additionof W28) and B53,54) does not affect the coarsening of MXparticles. MX carbo-nitride is not an equilibrium phase, anda Z phase is formed by the consumption of MX particles.
The appearance of the Z phase has been reported in manyhigh Cr ferritic steels.5,6,8,10,53) The Z particles readily growand are not useful for strengthening.5,8) The formation ofthe Z phase has to be prevented to keep fine MX particlesfor long duration. The formation can be delayed by reduc-ing Nb6,8) and Ni6) concentration.
The coarsening of particles in ferritic steels is accelerat-ed by creep deformation. The acceleration of M23C6 parti-cles has been reported in literature.6,8,9,30,55) The formationof the Z phase is also promoted by creep deformation.10)
8. Loss of Creep Rupture Strength during Long-termService
8.1. Stress–Rupture Behavior
As shown in Fig. 17, creep rupture strength of 11Cr–2.6W–0.1Mo–CoVNb steel decreases quickly at rupturelives longer than 2 000 h. The steel is tempered at a lowertemperature (680°C) and has the higher creep rupturestrength than the 9Cr–1.8W–0.4Mo–VNb, 11Cr–2W–0.3Mo–CuVNb and 9Cr–1Mo–VNb steels tempered at ahigher temperature (760–780°C). However, its superiorityin rupture strength disappears at 100 MPa. If this loss ofrupture strength is prevented by some means as is the casein the 11Cr–2.6W–0.1Mo–CoReVNb steel (680°C temper-ing), then we can achieve a significant improvement of rup-ture strength. The loss of creep rupture strength is more ev-ident in the steel tempered at a lower temperature6,14,25,26,31)
or in the steel containing W of more than 2%.56) However,the loss has been reported on 9Cr–1Mo–VNb steel whichdoes not contain W and is tempered at a higher tempera-ture.57) Stress–rupture curves of the steel are drawn in Fig.18(a), and the solid curves were predicted from the short-term data at the higher temperatures. The long-term datapoints at 600 and 650°C deviate from the prediction toshorter rupture life. A similar loss of rupture strength hasbeen observed in 11Cr–2W–0.3Mo–CuVNb steel temperedat 780°C.58)
Figure 19(a)14) is an example of the loss of rupturestrength reported in 12Cr steel (without W and tempered at650°C (0.52% Ni) or 675°C (1.15% Ni)). This figure pointsout another important feature. The quick decrease in rup-ture strength ceases at further lower stresses, and the stress–rupture data show the sigmoidal curves. Similar sigmoidal
Fig. 16. Effect of W on the growth of M23C6 particles duringcreep of 9Cr–W steel at 600°C.
Fig. 17. Stress-rupture data to high Cr ferritic steel at 650°C.
behavior has been reported in other steel.6)
8.2. Causes of Loss of Creep Rupture Strength
Figure 18(b)57) shows a dislocation substructure of the9Cr–1Mo–VNb steel crept at 600°C under 100 MPa. Theloss of rupture strength is evident in Fig. 18(a) under thecreep condition. Small subgrains remain in the grain interi-or, but a completely recovered band without sub-boundaryis formed along the grain boundary. Such enhanced recov-ery along grain boundaries is not seen when the loss of rup-ture strength does not occur at high stresses. Kushima etal.57) have proposed that the enhanced recovery is the originof the loss of rupture strength. The loss of rupture strengthand the band of enhanced recovery have been confirmed in11Cr–2W–0.3Mo–CuVNb steels.58,59) Creep curves of thesteel tested under 100 MPa at two temperatures are drawnin Fig. 20(a).59) The time scale at each temperature is con-verted to 650°C according to the following equation:
t650°C5tT e m,T /e m,650°C........................(22)
where tT is the time at a temperature T, and e m,T and e m,650°C
are the minimum creep rate at T and 650°C. The reductionof area measured after creep fracture is given in the figure.The steel is brittle and the loss of rupture strength is evidentat 650°C, whereas it is fully ductile and the loss does notoccur at 700°C. The curve converted from 700°C suggeststhat the rupture life is extended by three times if the steel is
fully ductile at 650°C. The above findings propose the following scenario of the
loss of rupture strength. During long-term tests, the en-hanced recovery of subgrain structure takes place alonggrain boundaries. Strain concentration along the boundaryregions forms grain boundary cracks. This results in thelow ductility, the premature failure and the loss of rupture
Fig. 18. (a) Stress–rupture data of 9Cr–1Mo–VNb steel. (b) TEMmicrograph of the steel crept to rupture, tr534 141 h, at600°C under 100 MPa.
Fig. 19. (a) Stress–rupture data of 12Cr–0.5Mo–VNb steel at600°C. (b) Effect of Ni on the growth of M23C6 particlesin the steel at 600°C.
Fig. 20. Creep curves tested at several temperatures under thesame stress. (a) 11Cr–2W–0.3Mo–CuVNb steel, and (b)11Cr–2.6W–0.1Mo–CoVNb steel.
strength. In the 9Cr–1Mo–VNb steel of Fig. 18, MX parti-cles are consumed to form the Z phase, and the disappear-ance of the MX particles accords with the start of the lossof rupture strength.60)
The aforementioned scenario is not always true. Creepcurves of 11Cr–2.6W–0.1Mo–CoVNb steel under 98 MPaat 650°C (with the loss of rupture strength) and 700°C(without the loss) are given in Fig. 20(b). The time scale ofthe creep curve at 700°C is converted to 650°C based onEq. (22). The curve of 700°C suggests a creep curve we canexpect if the loss of rupture strength does not occur. Theloss of rupture strength corresponds to the sudden accelera-tion of creep rate at a low strain in both 11Cr–2W–0.3Mo–CuVNb and 11Cr–2.6W–0.1Mo–CoVNb steels. However,the reduction of area was always 90% and the enhanced re-covery along grain boundaries was not observed in the lat-ter at 650°C.59) This fact indicates that there are severalother causes that bring about the loss of rupture strength.
High Cr ferritic steel containing more than 2% W showthe loss of rupture strength.56) The drop of rupture strengthstarts when the precipitation of Fe2W has finished.28,56,61)
Coarsening of M23C6 carbides in the 12Cr–0.5Mo–VNbsteel is depicted in Fig. 19(b).14) Diameter dp of the parti-cles is plotted against creep exposure time t according toEq. (20) with n53. The dp
3–t curves change from quick toslow coarsening at a critical time: 13104 h in 1.15% Ni and33104 h in 0.52% Ni. The critical time accords with the appearance of the sigmoidal behavior in Fig. 19(a). In thesteel, therefore, the loss of rupture strength accords with thequick coarsening of M23C6.
8.3. How to Recover from Loss of Creep RuptureStrength
The quick drop of creep rupture strength in Fig. 19(a)shifts toward longer life when the coarsening of M23C6 par-ticles is slow in the low Ni steel. This fact suggests that theloss can be delayed when Ni, Mn and Si are removed to re-duce self-diffusion coefficient in the ferrite matrix.
Figure 17 includes the result of 11Cr–2.6W–0.1Mo–CoReVNb steel with Re.62) The steel contains similaramounts of alloying elements to the 11Cr–2.6W–0.1Mo–CoVNb steel except Re. The loss of rupture strength doesnot occur in the Re containing steel, suggesting an impor-tant role of Re. Reduction of impurities such as Al from the11–Cr–2.6W–0.1Mo–CoVNb steel is also effective to pre-vent the loss of rupture strength63) Re is an interesting ele-ment that can improve creep strength of high Cr ferriticsteel.22,62,64,65) Recovery of the dislocation substructure isdelayed in Re containing steel. Most of Re atoms are in so-lution,22,62) and they should contribute to the stable disloca-tion substructure. However, the detailed mechanism of thestabilization has not been clarified yet.
9. Summary of Roles of Major Alloying Elements
9.1. W
It is widely accepted that the addition of W to high Crferritic steel improves its creep strength.66,67) Part of Watoms in the steel form Fe2W particles and the rest are insolution. Spacing of the Fe2W particles is wider than thoseof M23C6 and MX particles, and the Fe2W particles con-
tribute little to creep strength. Cottrell atmosphere draggingdue to W atoms68) and the effect of W on self-diffusion co-efficient69) cannot explain the strengthening by W addition.Major roles of W in the strengthening mechanism are relat-ed to dislocation hardening. The W addition lowers Ms
point and makes a fine subgrain structure. The fine sub-grains remain for long duration,56,70) since the density ofM23C6 pinning particles is kept at a high level in W contain-ing steel.48,67) W by itself plays many important roles.However, each role can be played by other elements. TheB2 steel18) developed recently in Europe does not containW but provides similar creep strength to W containing steelsuch as NF616 and HCM12A.
9.2. B
The addition of a very small amount of B to high Cr fer-ritic steel significantly improves its creep strength.50,54,71) Bdelays the coarsening of M23C6 particles7,14,49–51) and keepsfine subgrains.14,72) Most of B atoms are in solution. Theymay be combined with vacancies and reduce self diffusioncoefficient of ferrite matrix.8) They may form Cottrell at-mosphere around dislocations and increase creep deforma-tion resistance.71) However the roles of B in solution havenot been fully understood yet.
9.3. MX Particles
V and Nb are added to disperse fine and thermally stableMX particles within grains. Although the density of MXparticles is low in the usual high Cr ferritic steel (see Table2), the particles are necessary to keep fine subgrains and toimprove creep strength of the steel. A comparison is madein Fig. 21 among stress–rupture curves of three ferriticsteels: 9Cr–1.8W–0.4Mo–VNb steel (0.16V, 0.06Nb, 0.11Cand 0.036N) with W and MX,31) 9Cr–1Mo–VNb steel(0.22V, 0.09Nb, 0.10C and 0.051N) with MX but withoutW,36) and 9Cr–2W steel (0.08C) with W but without MXparticles.73) Although the 9Cr–1Mo–VNb steel contains thelarger amounts of V, Nb and N than the 9Cr–1.8W–0.4Mo–VNb steel, the former is inferior to the latter in creep rup-ture strength because of the absence of W. However, therupture lives of 9Cr–2W steel is substantially shorter thanthose of the 9Cr–1Mo–VNb steel in spite of W addition.The important role of the MX particles is obvious, but theirstrengthening mechanism has not been solved in detail.
9.4. Other Elements
The roles of the major alloying elements in the high Crferritic steel are summarized in Table 4. A Cr content of9–12% has to be added to attain sufficient oxidation resis-tance above 600°C. B, W and Co suppress the coarseningof M23C6 pining particles, and keep fine subgrains for along duration. The reduction of Si, Mn, Ni, and Mo alsocontributes to retain the fine subgrain structure. Re preventsthe drop of creep rupture strength after long-term use.
The presence of a d ferrite phase is harmful both to creepstrength and toughness of the high Cr ferritic steel. Co orCu is added to suppress the formation of the d ferrite. Alloydesign based on the d-election state is useful for the reduc-tion of the d ferrite.47,48)
Igarashi et al.8) have proposed a new steel based on a dif-ferent strengthening concept from the conventional high Crferritic steel. They introduced coherent FePd precipitates
into the steel. The dislocation substructure and the MX par-ticles typical of the conventional steel are not necessary ifthe precipitates are thermally stable and their spacing is fineenough. Kimura et al.79) have pointed out recently that fer-ritic steel without the dislocation substructure is superior tomartensitic steel in long-term creep strength. This fact ad-vises us against taking the conventional strengthening con-cept based on the dislocation hardening. Further studies arerequired to assess these proposals.
10. Concluding Remarks
High Cr ferritic steel is characterized by its fine and elon-gated subgrain structure, and there is a high density of freedislocations within the subgrains. The steel contains M23C6
and Fe2M particles on sub-boundaries and MX particleswithin subgrains. The dislocation substructure is the mostdensely distributed obstacle to dislocation motion in thesteel. It determines the athermal yield stress of the steel atintermediate temperatures, and its recovery controls creeprate and rupture life of the steel at elevated temperatures.Improvement of creep strength of the high Cr ferritic steelrequires a fine subgrain structure with a high density of freedislocations. Because of a large driving force for recoverydue to the high dislocation density, pinning particles (MXin subgrain interior and M23C6 on sub-boundaries) andsolute atoms (such as W, Co and B) are necessary to keepthe fine subgrains. Coarsening and agglomeration of thepinning particles have to be delayed by reducing Si, Ni andMn concentration.
Creep rupture strength of the high Cr ferritic steel de-creases quickly after long-term tests. In the steel temperedat high temperature, the enhanced recovery of the subgrainstructure along grain boundaries is the cause of the prema-ture failure and of the consequent loss of rupture strength.Disappearance of the MX particles triggers the enhancedrecovery. However, the scenario is not always applicable. Asignificant improvement of creep rupture strength can beachieved if we can prevent the loss of rupture strength.Reduction of diffusion coefficient in the ferrite matrix, ad-dition of Re and removal of Al from the high Cr ferriticsteel have been proposed to prevent the loss of rupture
strength. However, many things are left unsolved on thiscrucial problem.
W, B and MX carbo-nitride are very important elementsin creep of the high Cr ferritic steel. The addition of W re-duces the initial subgrain size. W and B delay the coarsen-ing of M23C6 particles and consequently retain the fine sub-grains for a long duration. The MX particles also help tokeep the fine subgrains. As a result, the high Cr ferriticsteel has achieved the excellent creep strength.
Acknowledgments
The present research was supported by a grant-in-aidfrom the Ministry of Education, Science, Sports andCulture, Japan (No. 10555225). The authors express theirgratitude to Nippon Steel Corporation and Iron and SteelInstitute of Japan for their financial support.
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