Recent Advance in Creep for PP

Sadhana Vol. 28, Parts 3 & 4,June/August 2003, pp. 709–730. © Printed in India

Recent advances in creep-resistant steels for power plantapplications

P J ENNIS1 and A CZYRSKA-FILEMONOWICZ2

1Research Centre Jülich, Institute for Materials and Processes in Energy Systems,IWV-2, D-52425 Julich, Germany2University of Mining and Metallurgy, Faculty of Metallurgy and MaterialsScience, Al Mickiewicza 30, PL-30059 Kraków, Polande-mail: [email protected]

Abstract. The higher steam temperatures and pressures required to achieveincrease in thermal efficiency of fossil fuel-fired power-generation plants necessi-tate the use of steels with improved creep rupture strength. The 9% chromium steelsdeveloped during the last three decades are of great interest in such applications. Inthis report, the development of steels P91, P92 and E911 is described. It is shownthat the martensitic transformation in these three steels produces high dislocationdensity that confers significant transient hardening. However, the dislocation den-sity decreases during exposure at service temperatures due to recovery effects andfor long-term creep strength the sub-grain structure produced under different con-ditions is most important. The changes in the microstructure mean that great care isneeded in the extrapolation of experimental data to obtain design values. Only datafrom tests with rupture times above 3,000h provide reasonable extrapolated values.It is further shown that for the 9% chromium steels, oxidation resistance in steamis not sufficiently high for their use as thin-walled components at temperatures of600◦C and above. The potential for the development of steels of higher chromiumcontents (above 11%) to give an improvement in steam oxidation resistance whilstmaintaining creep resistance to the 9% chromium steels is discussed.

Keywords. Chromium steels; creep rupture strength; power/plant components;steam oxidation.

1. Introduction

The constraints that are currently placed on power generation plant in terms of environmentalimpact and economics have focussed attention on the development of high efficiency, lowemission systems. If thermal efficiencies of generating plants can be increased, fuel can beconserved (less fuel is required for a given power output) and emissions reduced (lowerfuel consumption means lower emissions of environmentally damaging gases). Increase inthe thermal efficiency of a power plant can be most effectively achieved by increasing thetemperature and, to a lesser extent, the pressure of the steam entering the turbine. Mostmodern steam power stations now in operation reach efficiencies of around 42% with steam

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710 P J Ennis and A Czyrska-Filemonowicz

Figure 1. Stress rupture strengths of the currently used and the newly developed power station steels.

temperatures of 600◦C and pressures of 25–30MPa. The next generation of steam powerplants should be capable of operating with steam at 625–650◦C, to enable thermal efficienciesof around 45% to be achieved. Further enhancement of the thermal efficiency may be obtainedby combining an advanced steam cycle plant with a gas turbine; in this way, efficiencies ofover 50% are possible. Of course, the increasing operating temperatures and pressures imposeincreasingly stringent requirements on the materials of construction. In the present paper wewill consider the developments that have taken place in the high chromium ferritic/martensiticsteels for advanced steam power plants.

In figure 1, the stress rupture strengths of the currently used and the new power stationsteels are compared on the basis of the maximum service temperature for a 1,00,000h stressrupture strength of 100MPa. It may be seen that the maximum service temperature increaseswith increasing complexity of the steel composition and the more highly alloyed steels havesufficient stress rupture strength to be considered for application at temperatures in excessof 600◦C. Indeed the new high chromium steels have similar stress rupture strengths toaustenitic stainless steels. There are several reasons for the reluctance to use the austeniticsteels; obviously the increased cost of the steel with the high chromium and nickel contentsis a disadvantage but there are technical problems because the thermal expansion coefficientof austenitic materials is at least 50% higher than that of ferritic steels. This means that carehas to be taken during cooling and heating to avoid excessive thermal stresses that can leadto fatigue failures.

2. Development of 9% chromium steels

In the 1970s, there was considerable interest in 9% chromium steels for components offast breeder nuclear reactors. On the basis of the familiar Fe9Cr1Mo steel used since the

Recent advances in creep-resistant steels for power plant applications 711

Table 1. Details of high chromium steels, chemical compositions in wt%.

Element P9 P91 P92 E911

C max. 0·15 0·10 0·124 0·105Si 0·20–0·65 0·38 0·02 0·20Mn 0·80–1·30 0·46 0·47 0·35P max. 0·030 0·020 0·011 0·007S max. 0·030 0·002 0·006 0·003Cr 8·5–10·5 8·10 9·07 9·16Mo 1·70–2·30 0·92 0·46 1·01W – – 1·78 1·00V 0·20–0·40 0·18 0·19 0·23Nb 0·30–0·45 0·073 0·063 0·068B – – 0·003 -N – 0·049 0·043 0·072Ni max. 0·30 0·33 0·06 0·07Al – 0·034 0·002 –Form and pipe, Ø159, 20 pipe, Ø300, 40 flat bar, 100× 16

dimensions, mm wall thickness wall thickness 1 h/1050◦C+Heat treatment 1 h/1050◦C+ 2 h/1070◦C+ 1 h/750◦C,

1 h/750◦C, 2 h/775◦C,1,00,000h stress

rupture strength at 35 94 115 110600◦C, MPa∗

∗Values for P91 from Canonico (1991), P92 from Wachteret al (1995), for E911 from Staubiet al(1998)

1950s in petrochemical plants, an improved steel was developed by the Oak Ridge NationalLaboratory (Sikkaet al 1981) and subsequently incorporated into the ASTM specificationsunder the designation P91 (ASTM 1986). A remarkable increase in the stress rupture strengthwas achieved by the addition of 0·2%V, 0·06Nb and 0·05N. In Japan, a steel developmentprogramme of Nippon Steel led to the steel NF616 (Nippon Steel 1991), which is nowdesignated P92 in the ASTM specification. With P92 a further increase in stress rupturestrength was obtained by an addition of 1·8%W and a reduction of the Mo content from 1to 0·5%. In the European COST (Co-operation in Science and Technology) Action 501, asimilar 9% chromium steel was developed; this steel is designated E911, contains 1%Mo and1%W, and offers similar stress rupture strength to P92 (Staubliet al1998)].

Chemical compositions and production details for the high chromium steels that wereinvestigated are given in table 1. Although much of the work described was carried out on thesteel P92, the findings are in principle also applicable to P91 and E911.The essential differencebetween the steels is the tungsten content and with increasing tungsten the propensity forLaves phase Fe2(Mo,W) formation increases.

3. Characteristics of high chromium steels

The Fe–Cr constitutional diagram is shown in figure 2. At compositions near to 9% Cr, thereis an extensive austenitic region from 820 to 1200◦C and the two-phase region betweenaustenite and ferrite has a very narrow temperature range. This means that it is possible toaustenitise the steel and on cooling to produce a practically fully martensitic structure, with


Figure 2. Fe–Cr constitutional dia-gram.

minimal amounts, if any, of delta ferrite, which is generally regarded as detrimental for hightemperature strength properties. The high creep rupture strength of the P91 steel relies onthe martensitic transformation hardening with additional strengthening by precipitation ofcarbides, nitrides and carbonitrides of Nb and V. In P92 and E911, the W additions contributeto additional solid solution strengthening of the martensitic matrix. The development line isillustrated in figure 3 which shows the 1,00,000h stress rupture strengths of P9 (data fromISO 1981), P91 (Canonico 1994), P92 (Wachteret al1995) and E911 (Staubliet al1998) at600 and 650◦C.

3.1 Microstructure

Chromium steels (9–12%) are heat-treated to produce a martensitic microstructure that issubsequently tempered to improve the ductility and impact strength at low temperatures. Thetreatment consists of austenitisation at temperatures around 1100◦C followed by temperingat around 750◦C.

3.1aMicrostructure of P92 after austenitising:Cooling in air is sufficient to initiate themartensitic transformation after austenitising treatment because of the high chromium content.Figure 4 shows the microstructure of the steel P92 after austenitising at 970◦C for 2h. The

Figure 3. The development of 9% chromium steels.


(b)

(a)

Figure 4. Microstructure of thesteel P92 after austenitising at970◦C for 2h: (a) optical micro-graph showing martensite lathswith a small amount of retainedaustenite;(b) transmission elec-tron micrograph showing needle-like Fe-rich M3C particles witha Widmanstätten structure withinmartensite laths.

treated steel exhibits a martensitic structure with high dislocation density and a small amountof retained austenite at the lath boundaries. Within the large martensite laths, needle-like Fe-rich M3C particles form a Widmanstätten structure with the usual Bagaryatskii orientationto the ferritic matrix. During austenitisation at 970◦C, not all M23C6 particles are dissolved,whereas austenitisation at 1070◦C and above leads to complete dissolution of this carbide type.Nb(C, N) precipitates are observed in all specimens after austenitisation and their presence inthe structure may inhibit austenite grain growth. The size of the particles suggests that theyare not dissolved during austenitisation.

The effect of variations in the austenitising temperature on the microstructure is shownin figure 5. The main difference in the microstructure is the increase in lath width (from0·38nm at 970◦C, to 0·42nm at 1070◦C and 0·58nm at 1145◦C) and in the prior austenitegrain size, which increases from 10µm at 970◦C to 20µm at 1070◦C and 60µm at 1145◦C.The microstructural changes caused by austenitising of P92 steel in different temperaturesare described by Enniset al (1997) and Zielinska-Lipiecet al (1997).

3.1b Microstructure of P92 after tempering:Figure 6 shows the microstructures of P92austenitised at 1070◦C and tempered at 715, 775 and 835◦C. During tempering, two main


Figure 5. TEM micrographs of P92austenitised for 2h at(a) 970◦C and(b) 1145◦C, and subsequently tempered2h/775◦C, showing the increased marten-site lath width at the higher austenitisingtemperature.

processes take place. First, recovery causes a reduction in the high dislocation density afteraustenitization and the formation of sub-grains and dislocation networks. These processesare accelerated at the higher tempering temperatures, so that tempering at 715◦C leads toslightly higher dislocation density than standard tempering at 775◦C. Tempering at 835◦Ccauses sharp reduction of about 75% in the dislocation density.

Second, the precipitation of carbides, nitrides or carbonitrides occurs during tempering(figure 7). The M3C precipitated after austenitisation dissolves as the more stable carbidesor nitrides of chromium, molybdenum, niobium and vanadium form. M23C6 is precipitatedon prior austenite grain boundaries, on subgrain boundaries and within the martensite laths.The precipitates that are important for the mechanical properties of P92 are the fine M(C,N):spheroidal Nb-rich carbonitrides and plate-likeV-rich nitrides. The larger, spheroidal particlesof Nb(C,N) appear to remain undissolved during austenitizing and during tempering actas nucleation sites for the plate-like V-rich nitrides, thus forming the V-wing complexes(Zielinska-Lipiecet al1997).The small number of such precipitates, however, make it unlikelythat they are of great significance for the mechanical properties of this steel.

The results of the microstructural parameter measurements (dislocation density, sub-grainwidth and precipitate dimensions) are summarised in figure 8, with the quantitative measure-ments normalised to the steel in the usual heat treatment condition(2 h/1070◦C+2 h/775◦C).The actual values are given in table 2. Data for creep specimens tested at 600◦C are includedfor comparison. It can be seen that the dislocation density decreases by a factor of around 3


Figure 6. TEM micrographs of P92austenitised at 1070◦C for 2h and tem-pered for 2h at(a) 715◦C, (b) 775◦Cand (c) 835◦C; the micrographs showincreased recovery of martensite anddecreased dislocation density as the tem-pering temperature is raised.

after high-temperature tempering and after creep exposure for a few thousand hours at 600◦C.The sub-grain width is substantially increased by creep deformation, being a factor of 3–4higher in the creep specimens tested at 600◦C for 10,000h or more than in the pre-test condi-tion. The precipitate dimensions are not so strongly affected, but significant increases in theM23C6 dimensions are seen after creep testing.

The other particle precipitated in 9–12% chromium steels is the Laves phase, Fe2(W, Mo),see figure 9. This takes place after long-term ageing or creep at 600 and 650◦C. It was observedthat the growth of the Laves phase occurred during the first 10,000h of ageing at 600◦C; andHattesrand & Andrén (2001) determined the volume fraction to be 1%. The final size of theLaves phase is much larger in the specimens aged at 650◦C in comparison to those aged at600◦C.


Figure 7. TEM micrograph of P92austenitised at 970◦C for 2h andtempered at 775◦C. Large M23C6on sub-grain boundaries and fineM.(C,N) within the sub-grains.

3.1c Comparison of the microstructure of P91, P92 and E922 steels:After the standardheat treatment (given in table 1), all three steels exhibit similar microstructures, as shown infigure 10. Austenitising produce a martensitic structure with high dislocation density withinthe martensite laths. During tempering, recovery causes the formation of sub-grains anddislocation networks. The creep strength of 9–12% Cr steels is correlated inversely with themartensite lath width and therefore with the sub-grain size. Measurements of the averagesub-grain width and of the dislocation density within the sub-grains, performed by means ofquantitative TEM, are presented in table 3.

It can be seen that sub-grain size is fairly similar in all the steels are investigated. Thesmall differences are connected with different prior austenite grain size. Dislocation den-

Figure 8. Results of quantitative measurements of P92 microstructural features after different heattreatments: 2 h/970◦C + 2 h/1070◦C, 2 h/1070◦C + 2 h/715◦C, 2 h/1070◦C + 775◦C (as received),2 h/1070◦C + 835◦C, 2 h/1145◦C + 2 h/775◦C and after creep exposure at 600◦C for 1,500, 10,000and 33,000h.


Table 2. Quantitative microstructural parameters for P92 in different heat treated conditions and aftercreep testing.

Taust Ttemp Creep Dislocation Sub-grain Mean dia Mean dia(◦C) (◦C) test density(1014m−2) width (µm) M23C6, (nm) (MX, nm)

970 775 – 8·7 ± 1·2 0·38± 0·1 – –1070 715 – 9·0 ± 1·0 0·37± 0·1 72± 16 14± 11070 775 – 7·5 ± 0·9 0·42± 0·1 89± 13 16± 11070 835 – 2·3 ± 0·6 0·50± 0·1 82± 12 16± 11145 775 – – 0·58± 0·1 68± 18 16± 11070 775 600◦C/1, 500 h 5·3 ± 0·6 0·70± 0·1 119± 8 18± 11070 775 600◦C/10, 000 h 2·5 ± 0·5 1·4 ± 0·1 125± 10 21± 21070 775 600◦C/33, 000 h 1·5 ± 0·4 1·5 ± 0·1 131± 12 30± 3

sities in P91 and P92 steels are similar, in both steels a little higher than in E911 steel.However, it must be borne in mind that only one heat of each steel has been examined indetail and the small differences observed may not be significant in the light of heat to heatvariations.

Besides recovery processes, the precipitation of carbides, carbonitrides and nitrides occursduring tempering. In all three steels examined, M23C6 carbides containing Cr, Fe, Mo (W)precipitate preferentially on the prior austenite grain boundaries and on the martensite lathboundaries. These precipitates retard the sub-grain growth and therefore increase the strengthof the materials. In P91 steel mainly spheroidal Nb-rich carbonitrides are observed withinthe martensite laths. In P92 and E911 steels, three types of MX; Nb(C,N), plate-like VN andsmall complex Nb(C,N)-VN, are found (Enniset al1997, 2002; Zielinska-Lipiecet al1997;Haldet al1998; Vanstone 1998; Hättestrand & Andrén 2001).

The microstructural evolution of P91, P92 and E911steels during ageing at service tem-peratures is discussed in detail by Haldet al (1998), and Ennis & Cyrska-Filemonowicz(2001). As an example, the microstructural evolution of P92 steel during very long term creepdeformation at 600 and 650◦C up to about 57,500h is discussed in detail in the followingsection.

Figure 9. Laves phase precipitatedin P92 after creep testing at 650◦Cfor 6,500h.


Figure 10. TEM micrographs of(a) as receivedP91; (b) P92; and(c) E911 steels, showing theelongated sub-grains of tempered martensite.

Table 3. Dislocation density and mean sub-grain size of as received P91, P92 andE911 steels.

Steel Dislocation density× 10−14(m−2) Mean sub-grain size,(µm)

P91 7·5 ± 0·8 0·4 ± 0·06P92 7·9 ± 0·8 0·4 ± 0·09E911 6·5 ± 0·6 0·5 ± 0·05


Figure 11. Iso-stress plots at120MPa of secondary creep ratefor P92 in as-received conditionand after different heat treat-ments: 2 h/970◦C + 2 h/775◦C;2 h1070◦C +2 h/775◦C; 2 h/1070◦C+2 h/835◦C; 2 h/1070◦C, furnacecool to 780◦C, 8 h/780◦C, furnacecool to RT (ferritic structure, nomartensite).

3.2 Creep rupture properties

From the creep curves, the secondary creep rate is derived by taking the slope of the straight-line portion of each curve. Figure 11 shows the creep strength, plotted as iso-stress curvesat 120MPa for the secondary creep rate, of the P92 material in different heat treatmentconditions. Decreasing the austenitising temperature from the usual 1070◦C to 970◦C doesnot significantly affect the creep strength. However, increasing the tempering temperaturefrom 775 to 835◦C, a temperature just above theAc1 transformation temperature of 825◦C,leads to a marked fall in the creep strength. The low initial dislocation density produced bythis heat treatment, with no other significant differences in the microstructural parameters, isthe reason for the low creep strength. This effect is particularly relevant for weldments, sincein the heat-affected zone there is always a region which has been exposed to temperaturesaround 850◦C and therefore this region is the most likely site for creep failure according toEggeleret al (1987).

The dislocation density decreases not only as a result of increase in tempering temperaturebut also during prolonged exposure at lower temperatures, typical of service conditions. Thiscan be seen in the quantitative TEM investigations carried out on tested creep specimens.After long testing times, the dislocation density is around 2× 1014 m−2, a decrease of 75%compared with the as received material. The microstructural changes that occur in the first3,000h of exposure mean that there is therefore a danger of overestimating the long-termcreep strength of P92 if there is a preponderance of short duration data in the evaluation, inwhich the dislocation density remains at a high level during the whole test. The extrapolationof the experimental data for estimation of the long-term creep rupture strength is discussedlater in more detail.

Figure 12 shows the secondary creep rates of P91, P92 and E911 at 600 and 650◦C plot-ted against the applied stress(σ ). Data taken from published literature for the same heat areincluded where available. It can be seen that at high stresses the differences in the secondarycreep rates of the three steels are relatively small. In the low stress region, however, the differ-ences between the steels becomes more pronounced. The creep deformation characteristicsmay be described by the Norton equation (minimum creep rate is proportional to the appliedstress to the powern) with two different values for the Norton stress exponentn. At highstresses, the value ofn is around 16 while at lower stresses, the data conform to ann value of6. The change inn is indicative of a change in the creep characteristics. Figure 12 also showsthat at high stresses the differences in the secondary creep rates of the three steels are relatively


Figure 12. Secondary creep rates for P91, P92 and E911.

small and the high dislocation densities resulting from the martensite transformation domi-nates the deformation. In the low stress region, however, the differences between the steelsbecome more pronounced. The evolution of the microstructure is examined to investigate thereason behind the two deformation regions.

3.3 Microstructural stability of P92 during long creep deformation

The degree of change in the microstructure is dependent on duration and the temperature ofcreep deformation. Well-developed sub-grains of low dislocation density in the interiors arecharacteristic features of long-term exposed specimens. Figure 13 shows the microstructuresof P92 steel creep deformed at 600◦C up to 32909h and at 650◦C up to 27433h.

The results of quantitative measurements of dislocation density and sub-grain width withincreasing creep deformation time for P92 specimens exposed at 600 and 650◦C are shown infigure 14. The quantitative evaluation of the microstructure shows that in the first 3,000h ofexposure there is a rapid reduction in the dislocation density and an increase of the sub-grainwidth. After long testing times, the dislocation density is around 2× 1014 m−2, a decrease of75% compared with the as-received material.

Similar effects were observed for the P91 and E911 steels. The microstructure of the P91specimens creep specimens tested at 600◦C for 9,200h and steel E911 tested for 17,500h(figures 15 a and b) exhibit polygonal sub-grains as a results of the extensive deformation. Thedislocation densities in the steels P91, P92 and E911 after creep testing at 600◦C are shownin table 4. As observed in P92 steel, the mean sub-grain widths of E911 and P91 increase andthe dislocation densities within the sub-grains decrease with increasing creep exposure.

Other important microstructural changes during creep deformation of the investigated steelare the size, morphology and distribution of the carbide, nitride and carbonitride precipitatesas well as the chemical composition of the precipitates and the matrix. The fine MX, mainlyplate-like VN and spheroidal Nb(C,N), precipitate intragranularly and act as obstacles formoving dislocations, thus contributing to increased creep strength of P92 steel (figure 15).Some complex carbonitrides, consisting of a spheroidal Nb-rich particle to which the V-richparticle is attached, are also present. Sub-grain boundaries and prior austenite grain boundaries


Figure 13. TEM micrograph ofP92 steel after creep deformation at(a) 600◦C for 33,000h and(b) 650◦C for 27,500h.

Figure 14. Dislocation densities andsub-grain widths for P92 after creeptesting at 600 and 650◦C.


Figure 15. Dislocations interac-tion with fine carbide precipitatesin steel P92 specimen after creeptesting at 650◦C for 6468h.

were pinned by M23C6. These precipitates retard sub-grain growth and therefore increase thestrength of the material.

In steel exposed for longer durations, Laves phase Fe2(W,Mo) (figure 16) and other inter-metallic phases are formed. The particles of Laves phase formed during creep at 650◦C weremuch larger than those precipitated at 600◦C. The stacking faults formed within the Lavesphases could be observed as characteristic streaks on the selected area electron diffractionpattern (figure 16b). This effect allowed easy distinction between the Laves phase and theM23C6, as both particles were precipitated very frequently in close vicinity, preferentially onsub-grain boundaries. Statistical measurements of the mean particle sizes revealed signifi-cant coarsening of the M23C6 and intermetallic phases. The mean particle diameter of M23C6

increased with increasing exposure time at 600 and 650◦C, while the MX revealed insignifi-cant change in size.

Precipitation processes in P92 are influenced not only by temperature but also by stress.Figure 17 shows the results of statistical measurements of the M23C6 and MX particles formedin P92 specimens (head and gauge lengths) after creep deformation at 600◦C. It can be seenthat the coarsening of M23C6 carbides is accelerated by stress while the effect of stress-inducedcoarsening of MX is insignificant. These findings concerning the coarsening of carbides inchromium steels are in good agreement with other studies (Eggeleret al 1987; Schaffernaket al1998; Hattesrand & Andrén 2001).

In addition to precipitate morphology and distribution, the chemical compositions of thephases that precipitate with increasing exposure duration are important. STEM/EDS analyses

Table 4. Dislocation density in steels P91, P92 and E 911 creep tested at 600◦C.

ca 1,000h ca 10,000h ca 1,70,000h 22,000h

P91 (4·8 ± 0·5) × 1014 (1·1 ± 0·4) × 1014 nd ndP92 (5·4 ± 0·5) × 1014 (2·5 ± 0·5) × 1014 (2·3 ± 0·5) × 1014 ndE911 (5·0 ± 0·5) × 1014 nd (2·2 ± 0·4) × 1014 (2·1 ± 0·4) × 1014

nd - not determined


Figure 16. Laves phases formed during creep deformation of P92 at 650◦C for 27500h(a) TEMmicrograph and(b) corresponding electron diffraction pattern.

showed that the M23C6 are enriched in Cr, Fe, Mo and W, whereas the Laves phase is enrichedin W and Mo (Czyrska-Filemonowiczet al 2001). An increase in the creep deformationtemperature from 600 to 650◦C results in precipitation of much larger particles of Lavesphase. With increasing creep duration, the precipitation of the Laves phase removes Mo andW from the matrix solid solution and the strengthening of the matrix is decreased. This willbe discussed in the next section.

The results of our investigations are in general agreement with other TEM investigationsdescribing the microstructure during creep deformation of 9% Cr steels (Foldynaet al1996;Hald et al 1998; Nowakowskiet al 1998; Vanstone 1998; Hättesrand & Andrén 2001b).Strang & Vodarek (1998), however, reported that there is another important precipitate thatcontributes to the softening of high chromium steels, namely the formation of large particlesof Cr(Nb,V)N nitride (Z phase), replacing to some extent the fine VN precipitates. In the P92steel the Z phase has not yet been observed in specimens exposed for up to 57,000h (Enniset al2000).

The question arises as to the relevance for long-term creep strength of the high initial dis-location density, which results from martensitic transformation and decreases during the firstfew thousand hours of exposure at service temperatures. To clarify this, creep tests were car-ried out on P92 specimens heat-treated so that martensitic transformation was suppressed.The microstructure consisted of ferrite with a very low dislocation density(9·8 × 1013 m−2)

Figure 17. Statistical measurements ofthe M23C6 and MX particles formed inP92 specimens (head and gauge lengths)after creep deformation at 600◦C.


Figure 18. Stress rupture curves at600◦C for P92, P91 and P91 com-pared with the values for a ferriticstructure P92.

with carbide (M23C6) and carbonitride precipitates. Some of the M23C6 particles were par-ticularly coarse, up to around 1µm diameter. The secondary creep rate at 600◦C, 120MPafor specimens with ferritic microstructure was found to be 10,000 times higher that that ofP92 in the usual austenitised and tempered condition. Figure 18 compares the 600◦C stressrupture strengths of martensitic and ferritic P92 specimens.

It is clear that the martensite transformation makes a considerable contribution to the creeprupture strength of the 9% Cr steels, even though the martensitic structure degenerates intoferrite and the initially high dislocation density decreases with exposure time. The initialmartensitic structure does allow the formation of a stable subgrain structure and it is thissubstructure together with the fine carbonitride precipitates which confers the high creeprupture strength.

A ferritic structure containing similar fine carbonitride precipitates to the martensitic P92cannot provide the creep rupture strength required.This means that developments to strengthenthe 9% Cr steels for applications at 600◦C and above will be restricted by the necessity for thesteels to exhibit the martensitic transformation. The maximum operating temperature must besufficiently below the tempering temperature to avoid too rapid a recovery of the martensite.

4. Assessment of long-term creep rupture strength

Prediction of the long-term strength using any extrapolation procedure is difficult owing toseveral factors. The scatter in the experimental data, especially between different heats withinthe compositional and heat treatment specifications for a given steel; the temperature andduration ranges of the available data; microstructural changes that occur in the materials duringtesting and influence deformation; and environmental effects, such as oxidation, that reducethe effective load-bearing cross-section may all lead to erroneous predictions of strength. Oneof the earliest attempts to carry out extrapolation of creep rupture data was made by Norton(1929), although this book is better known for the first publication of the creep power lawequation, now widely used in creep deformation modelling. By relating the secondary creeprate to the applied stress and then using for design a maximum tolerable creep rate, a hightemperature component could be dimensioned. Larson & Miller (1952) applied the Arrhenius


rate equation for thermally activated creep processes. By assuming that the secondary creeprate was inversely proportional to the rupture time,tR, that is, the faster the specimen creeps,the shorter the rupture time, a time–temperature parameter could then be derived:

PLM = T (C + log tR),

whereT is the temperature inK, tR the rupture time andC is a constant, related to theactivation energy for creep. This implies that long times at low temperatures are equivalentto shorter times at higher temperatures. This enables all stress rupture data to be plotted on asingle master curve and for estimation of the stress rupture strength, no extrapolation of thecurve beyond the experimental data points is needed. The use of this parameter does, however,dictate the form of the stress rupture data curves; if the log of the rupture time is plotted asa function of the reciprocal test temperature inK for constant stress, straight lines should beobtained which intersect at 1/T = 0, log tR = −C.

Microstructural studies and analysis of the stress dependence of the secondary creep rate(figure 12) shows that there seems to be a difference in creep behaviour at low stress andhigh stress. Using the Larson–miller time–temperature parameter, the experimental data wereevaluated and the 1,00,000h stress rupture strengths at 600 and 650◦C were calculated, firstby using all the data and second by using only data for rupture times above 3,000h. Theextrapolated 1,00,000h stress ruture strengths are shown in figure 19. After 3,000h at 600and 650◦C, the dislocation density of P92 reaches a more or less constant value. It can beseen from figure 19 that the extrapolation using only data from the longer duration tests givessignificantly lower 1,00,000h rupture strengths at both temperatures.

The Larson–Miller assumption that the secondary creep rateεs is inversely proportional tothe rupture timetR was further refined by Monkman & Grant (1956) to give the Monkman–Grant (MG) equation,

εms · tR = K1,

Figure 19. Extrapolated 1,00,000h rupture strengths at 600◦C for P92 calculated with data fromdifferent rupture time ranges.


Figure 20. Monkman–Grant equation plots for P91, P92 and E911 steels.

whereK1 andm are constants. Rearranging, we obtain

log εs = −(1/m) log tR + logK1.

Plotting log rupture time as a function of log secondary creep rate therefore gives a straightline from whichK1 andm can be determined. The constantm is usually in the range 0·8–1·2.The MG equation together with the Norton equation can be used for extrapolation. From theMG plot the secondary creep rate for a given rupture life can be read off and, from the Nortonequation plot, the stress which gives rise to this creep rate can be determined. The Monkman–Grant plots for P91, P92 and E911 are shown in figure 20 and indicate that a rupture life of1,00,000h at 600◦C corresponds to a secondary creep rate of 1·5×10−5%/h. From figure 12,this creep rate is obtained at a stress of 110MPa for P92, 105MPa for E911 and 85MPa forP91. The values obtained for P92 agree well with the Larson–Miller extrapolation made usingonly data of above 3,000h duration.

5. Steam oxidation resistance of high chromium steels

Almost without exception, high temperature materials rely on the selective oxidation of one ormore alloy constituents to form a protective oxide scale. Two conditions need to be fulfilled;first, there must be a sufficiently high concentration of the selectively oxidised elements inthe matrix and, second, the diffusion rate of these elements must be fast enough to ensure thatthey replenish the matrix below the growing scale, thus ensuring long-term protection. In thehigh Cr steels, clearly Cr is the most important constituent with regard to oxidation resistance.

In the development of modified 9Cr1Mo steels, the emphasis was on improvements inthe stress rupture strength. Long-term creep tests carried out at temperatures up to 650◦Cshow that the oxidation resistance in air is excellent due to the formation of tightly adherent,protective oxide scales. The protective scales formed in air on 9% Cr steels are identifiedas consisting of (Fe,Cr)2O3 and (Fe,Cr,Mn)3O4. However, in steam-containing atmospheres,the scales formed at 600 and 650◦C are thick and consist of an external Fe3O4 scale and aninternal two-phase scale of Fe3O4 and (Fe,Cr,Mn)3O4 (Ehlers & Quadakkers 2001). Belowthe oxide scale, internal oxidation of Cr to form (Fe,Cr,Mn)3O4 occurs at the martensite


Ni coating

Fe3O4

(Fe, Cr, Mn) spinel

FeO + Cr2O3 stringers

steel matrix

Figure 21. Scale microstructure ofP91 exposed for 1,000h at 650◦C inAr-50% H2O.

lath boundaries. Figure 21 shows typical oxide scales formed on the 9% Cr steels in steam-containing atmospheres at 650◦C and figure 22 summarises the mass changes that occur.

There are several concerns about the high oxidation rates seen in steam. First, the loss inload-bearing cross-section due to the oxide scale formation and the internally oxidised zoneleads to stress increases and therefore a reduction in service lives. The reduction in life is,of course, dependent on the initial wall thickness of the components under consideration.Calculations reported by Quadakkers & Ennis (1998) have shown that the life reduction at600◦C in steam is significant for components with wall thicknesses below about 6 mm. Asecond, and more difficult to quantify, effect is the spalling of the oxide. The presence ofspalled oxide particles in the steam entering the turbine can cause erosion problems andlocal blockages. Furthermore, thick oxide scales on heated tubes can lead to decrease in heattransfer across the tube walls, resulting in overheating and subsequent creep failure of thetubes.

In order to achieve acceptable oxidation resistance at 600 and 650◦C, according to Ehlerset al (2001) a Cr content of at least 11% is required, to enable the formation of a protectivespinel scale. The oxidation resistance may be enhanced further by the addition of selectively

Figure 22. Mass change data for Cr steels exposed in Ar-50% H2O and in air at 650◦C.


oxidisable elements, the main contenders being Si and Mn. Si is, however, a ferrite former,and when added to the steel in amounts that are clearly beneficial for the steam oxidationresistance can lead to toughness and fabrication difficulties.

There are some additional factors that need to be considered in assessing the steam oxi-dation resistance of the high Cr steels. The effect of steam test parameters, such as flowrate and pressure, are not sufficiently well established and there is a need for further inves-tigations to ensure that laboratory data reflects sufficiently well the expected behaviour in apower plant. Such tests are being carried out in the new COST 522 Action, see Allenet al(1998).

The steam oxidation resistance of the high Cr steels is enhanced by:

• a Cr content of at least 11%;• the addition of oxygen active elements, such as Si and Mn;• increasing the diffusion of Cr to the surface by suitable bulk alloying additions or by

surface deformation treatments

6. Potential for further development

The need to obtain the optimum microstructure for high creep rupture strength and the require-ment for improved steam oxidation resistance make contradictory demands on the steel com-position. For satisfactory creep and stress rupture strength, Cr contents of around 9–10%allow the desired fully martensitic microstructure to be obtained. For adequate resistance tosteam oxidation, Cr contents above 11% are necessary. The aim of current steel developmentis to raise the Cr content to 11–12%, and to add austenite-stabilising elements to produce thefully martensitic structure. In this way, it is hoped that stress rupture strength levels similarto 9% Cr steels can be reached at higher Cr contents in the range 11–12%. Some successhas been reported by Tsudaet al (1998), but long-term creep data are not yet available andextrapolations are therefore uncertain.

A promising line of development being followed in the new COST 522Action is the additionof Co, which enlarges the austenite field in the composition direction with no decrease inthe a/g transition temperature, to 11–12% Cr steels (Allenet al 1998). The disadvantage ofCo is, of course, the high cost of the raw material. Additions of V and Nb are necessary forprecipitation hardening of the matrix. Regarding the solid solution strengthening additions ofMo and W, there are two approaches; in the first, additions of Mo alone are being investigated,in order to prevent the formation of Laves phase. There is, however, sufficient evidence tosuggest that W additions do confer improved stress rupture strength, at least to test durationsof around 50,000h that have been achieved in creep testing.

There are other development routes towards a ferritic, high strength Cr steel with goodresistance to steam oxidation. The effectiveness of increased amounts of nitride precipitatesin conferring high stress rupture strength has being investigated by Goecmenet al(1998). Thesteels under development contain increased amounts of V and N and are given an appropriateheat treatment to produce finely dispersed nitrides. A similar development has been reportedby Pugh (1998), the strengthening particles in this case being TiN, introduced by solid statereaction between CrN and the base steel containing 12% or more Cr and 1–2% Ti. The Ospreyprocess and a powder metallurgical production route have been investigated. Another ideaunder investigation is the addition of FeWTi carbide powder to steel melts with the aim ofproducing uniform and fine dispersion of Ti and W carbides in the steel matrix, as describedby Nutting (1999).


The application of surface treatments and coatings is a promising, innovative line of deve-lopment. The coating of large components and small diameter tubes does, however, presenta considerable technological and economic challenge.

The aim of the current developments is to combine in one steel the strength of the 9%Cr steels and the steam oxidation resistance of the 12% Cr steels, for the temperature rangeup to 625◦C. This target appears to be achievable. However, the extension of the operatingtemperature range of the steels to 650◦C requires further increase in creep strength; it is byno means clear whether such an increase can be achieved either by enhancing the alreadyestablished mechanisms of strengthening or by the new concepts being considered.

The support of the Polish State Committee for Scientific Research and German BMBF isgratefully acknowledged. The authors also wish to thank their colleagues, H J Penkalla andW J Quadakkers, for their valuable contributions to the work described in this paper.

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