Recent advances in creep-resistant steels for power plant applications

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

Recent advances in creep-resistant steels for power plant applications P J ENNIS1 and A CZYRSKA-FILEMONOWICZ2 1

Research Centre Jülich, Institute for Materials and Processes in Energy Systems, IWV-2, D-52425 Jülich, Germany 2 University of Mining and Metallurgy, Faculty of Metallurgy and Materials Science, Al Mickiewicza 30, PL-30059 Kraków, Poland e-mail: [email protected] Abstract. The higher steam temperatures and pressures required to achieve increase in thermal efficiency of fossil fuel-fired power-generation plants necessitate the use of steels with improved creep rupture strength. The 9% chromium steels developed during the last three decades are of great interest in such applications. In this report, the development of steels P91, P92 and E911 is described. It is shown that the martensitic transformation in these three steels produces high dislocation density that confers significant transient hardening. However, the dislocation density decreases during exposure at service temperatures due to recovery effects and for long-term creep strength the sub-grain structure produced under different conditions is most important. The changes in the microstructure mean that great care is needed in the extrapolation of experimental data to obtain design values. Only data from tests with rupture times above 3,000 h provide reasonable extrapolated values. It is further shown that for the 9% chromium steels, oxidation resistance in steam is not sufficiently high for their use as thin-walled components at temperatures of 600◦ C and above. The potential for the development of steels of higher chromium contents (above 11%) to give an improvement in steam oxidation resistance whilst maintaining 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 environmental impact and economics have focussed attention on the development of high efficiency, low emission systems. If thermal efficiencies of generating plants can be increased, fuel can be conserved (less fuel is required for a given power output) and emissions reduced (lower fuel consumption means lower emissions of environmentally damaging gases). Increase in the thermal efficiency of a power plant can be most effectively achieved by increasing the temperature and, to a lesser extent, the pressure of the steam entering the turbine. Most modern steam power stations now in operation reach efficiencies of around 42% with steam 709

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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 25–30 MPa. The next generation of steam power plants should be capable of operating with steam at 625–650◦ C, to enable thermal efficiencies of around 45% to be achieved. Further enhancement of the thermal efficiency may be obtained by combining an advanced steam cycle plant with a gas turbine; in this way, efficiencies of over 50% are possible. Of course, the increasing operating temperatures and pressures impose increasingly stringent requirements on the materials of construction. In the present paper we will consider the developments that have taken place in the high chromium ferritic/martensitic steels for advanced steam power plants. In figure 1, the stress rupture strengths of the currently used and the new power station steels are compared on the basis of the maximum service temperature for a 1,00,000 h stress rupture strength of 100 MPa. It may be seen that the maximum service temperature increases with increasing complexity of the steel composition and the more highly alloyed steels have sufficient stress rupture strength to be considered for application at temperatures in excess of 600◦ C. Indeed the new high chromium steels have similar stress rupture strengths to austenitic stainless steels. There are several reasons for the reluctance to use the austenitic steels; obviously the increased cost of the steel with the high chromium and nickel contents is a disadvantage but there are technical problems because the thermal expansion coefficient of austenitic materials is at least 50% higher than that of ferritic steels. This means that care has to be taken during cooling and heating to avoid excessive thermal stresses that can lead to fatigue failures. 2. Development of 9% chromium steels In the 1970s, there was considerable interest in 9% chromium steels for components of fast breeder nuclear reactors. On the basis of the familiar Fe9Cr1Mo steel used since the

Recent advances in creep-resistant steels for power plant applications

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Table 1. Details of high chromium steels, chemical compositions in wt%. Element C Si Mn P S Cr Mo W V Nb B N Ni Al Form and dimensions, mm Heat treatment 1,00,000 h stress rupture strength at 600◦ C, MPa∗

P9 max. 0·15 0·20–0·65 0·80–1·30 max. 0·030 max. 0·030 8·5–10·5 1·70–2·30 – 0·20–0·40 0·30–0·45 – – max. 0·30 –

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P91 0·10 0·38 0·46 0·020 0·002 8·10 0·92 – 0·18 0·073 – 0·049 0·33 0·034 pipe, Ø159, 20 wall thickness 1 h/1050◦ C+ 1 h/750◦ C, 94

P92 0·124 0·02 0·47 0·011 0·006 9·07 0·46 1·78 0·19 0·063 0·003 0·043 0·06 0·002 pipe, Ø300, 40 wall thickness 2 h/1070◦ C+ 2 h/775◦ C, 115

E911 0·105 0·20 0·35 0·007 0·003 9·16 1·01 1·00 0·23 0·068 0·072 0·07 – flat bar, 100 × 16 1 h/1050◦ C+ 1 h/750◦ C, 110

∗

Values for P91 from Canonico (1991), P92 from Wachter et al (1995), for E911 from Staubi et al (1998)

1950s in petrochemical plants, an improved steel was developed by the Oak Ridge National Laboratory (Sikka et al 1981) and subsequently incorporated into the ASTM specifications under the designation P91 (ASTM 1986). A remarkable increase in the stress rupture strength was achieved by the addition of 0·2%V, 0·06 Nb and 0·05 N. In Japan, a steel development programme of Nippon Steel led to the steel NF616 (Nippon Steel 1991), which is now designated P92 in the ASTM specification. With P92 a further increase in stress rupture strength was obtained by an addition of 1·8%W and a reduction of the Mo content from 1 to 0·5%. In the European COST (Co-operation in Science and Technology) Action 501, a similar 9% chromium steel was developed; this steel is designated E911, contains 1% Mo and 1% W, and offers similar stress rupture strength to P92 (Staubli et al 1998)]. Chemical compositions and production details for the high chromium steels that were investigated are given in table 1. Although much of the work described was carried out on the steel P92, the findings are in principle also applicable to P91 and E911. The essential difference between the steels is the tungsten content and with increasing tungsten the propensity for Laves 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, there is an extensive austenitic region from 820 to 1200◦ C and the two-phase region between austenite and ferrite has a very narrow temperature range. This means that it is possible to austenitise the steel and on cooling to produce a practically fully martensitic structure, with

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Figure 2. Fe–Cr constitutional diagram.

minimal amounts, if any, of delta ferrite, which is generally regarded as detrimental for high temperature strength properties. The high creep rupture strength of the P91 steel relies on the martensitic transformation hardening with additional strengthening by precipitation of carbides, nitrides and carbonitrides of Nb and V. In P92 and E911, the W additions contribute to additional solid solution strengthening of the martensitic matrix. The development line is illustrated in figure 3 which shows the 1,00,000 h stress rupture strengths of P9 (data from ISO 1981), P91 (Canonico 1994), P92 (Wachter et al 1995) and E911 (Staubli et al 1998) at 600 and 650◦ C. 3.1 Microstructure Chromium steels (9–12%) are heat-treated to produce a martensitic microstructure that is subsequently tempered to improve the ductility and impact strength at low temperatures. The treatment consists of austenitisation at temperatures around 1100◦ C followed by tempering at around 750◦ C. 3.1a Microstructure of P92 after austenitising: Cooling in air is sufficient to initiate the martensitic 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 2 h. The

Figure 3. The development of 9% chromium steels.


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(a)

(b)

Figure 4. Microstructure of the steel P92 after austenitising at 970◦ C for 2 h: (a) optical micrograph showing martensite laths with a small amount of retained austenite; (b) transmission electron micrograph showing needlelike Fe-rich M3 C particles with a Widmanstätten structure within martensite laths.

treated steel exhibits a martensitic structure with high dislocation density and a small amount of retained austenite at the lath boundaries. Within the large martensite laths, needle-like Ferich M3 C particles form a Widmanstätten structure with the usual Bagaryatskii orientation to the ferritic matrix. During austenitisation at 970◦ C, not all M23 C6 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 in the structure may inhibit austenite grain growth. The size of the particles suggests that they are not dissolved during austenitisation. The effect of variations in the austenitising temperature on the microstructure is shown in figure 5. The main difference in the microstructure is the increase in lath width (from 0·38 nm at 970◦ C, to 0·42 nm at 1070◦ C and 0·58 nm at 1145◦ C) and in the prior austenite grain size, which increases from 10 µm at 970◦ C to 20 µm at 1070◦ C and 60 µm at 1145◦ C. The microstructural changes caused by austenitising of P92 steel in different temperatures are described by Ennis et al (1997) and Zielinska-Lipiec et al (1997). 3.1b Microstructure of P92 after tempering: Figure 6 shows the microstructures of P92 austenitised at 1070◦ C and tempered at 715, 775 and 835◦ C. During tempering, two main

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Figure 5. TEM micrographs of P92 austenitised for 2 h at (a) 970◦ C and (b) 1145◦ C, and subsequently tempered 2 h/775◦ C, showing the increased martensite lath width at the higher austenitising temperature.

processes take place. First, recovery causes a reduction in the high dislocation density after austenitization and the formation of sub-grains and dislocation networks. These processes are accelerated at the higher tempering temperatures, so that tempering at 715◦ C leads to slightly higher dislocation density than standard tempering at 775◦ C. Tempering at 835◦ C causes sharp reduction of about 75% in the dislocation density. Second, the precipitation of carbides, nitrides or carbonitrides occurs during tempering (figure 7). The M3 C precipitated after austenitisation dissolves as the more stable carbides or nitrides of chromium, molybdenum, niobium and vanadium form. M23 C6 is precipitated on 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-like V-rich nitrides. The larger, spheroidal particles of Nb(C,N) appear to remain undissolved during austenitizing and during tempering act as nucleation sites for the plate-like V-rich nitrides, thus forming the V-wing complexes (Zielinska-Lipiec et al 1997). The small number of such precipitates, however, make it unlikely that they are of great significance for the mechanical properties of this steel. The results of the microstructural parameter measurements (dislocation density, sub-grain width and precipitate dimensions) are summarised in figure 8, with the quantitative measurements 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 included for comparison. It can be seen that the dislocation density decreases by a factor of around 3


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Figure 6. TEM micrographs of P92 austenitised at 1070◦ C for 2 h and tempered for 2 h at (a) 715◦ C, (b) 775◦ C and (c) 835◦ C; the micrographs show increased recovery of martensite and decreased dislocation density as the tempering 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 3–4 higher in the creep specimens tested at 600◦ C for 10,000 h or more than in the pre-test condition. The precipitate dimensions are not so strongly affected, but significant increases in the M23 C6 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 observed that the growth of the Laves phase occurred during the first 10,000 h of ageing at 600◦ C; and Hättesrand & Andrén (2001) determined the volume fraction to be 1%. The final size of the Laves phase is much larger in the specimens aged at 650◦ C in comparison to those aged at 600◦ C.

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Figure 7. TEM micrograph of P92 austenitised at 970◦ C for 2 h and tempered at 775◦ C. Large M23 C6 on sub-grain boundaries and fine M.(C,N) within the sub-grains.

3.1c Comparison of the microstructure of P91, P92 and E922 steels: After the standard heat treatment (given in table 1), all three steels exhibit similar microstructures, as shown in figure 10. Austenitising produce a martensitic structure with high dislocation density within the martensite laths. During tempering, recovery causes the formation of sub-grains and dislocation networks. The creep strength of 9–12% Cr steels is correlated inversely with the martensite lath width and therefore with the sub-grain size. Measurements of the average sub-grain width and of the dislocation density within the sub-grains, performed by means of quantitative TEM, are presented in table 3. It can be seen that sub-grain size is fairly similar in all the steels are investigated. The small differences are connected with different prior austenite grain size. Dislocation den-

Figure 8. Results of quantitative measurements of P92 microstructural features after different heat treatments: 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, 10,000 and 33,000 h.


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Table 2. Quantitative microstructural parameters for P92 in different heat treated conditions and after creep testing. Taust (◦ C)

Ttemp (◦ C)

Creep test

Dislocation density (1014 m−2 )

Sub-grain width (µm)

Mean dia M23 C6 , (nm)

Mean dia (MX, nm)

970 1070 1070 1070 1145 1070 1070 1070

775 715 775 835 775 775 775 775

– – – – – 600◦ C/1, 500 h 600◦ C/10, 000 h 600◦ C/33, 000 h

8·7 ± 1·2 9·0 ± 1·0 7·5 ± 0·9 2·3 ± 0·6 – 5·3 ± 0·6 2·5 ± 0·5 1·5 ± 0·4

0·38 ± 0·1 0·37 ± 0·1 0·42 ± 0·1 0·50 ± 0·1 0·58 ± 0·1 0·70 ± 0·1 1·4 ± 0·1 1·5 ± 0·1

– 72 ± 16 89 ± 13 82 ± 12 68 ± 18 119 ± 8 125 ± 10 131 ± 12

– 14 ± 1 16 ± 1 16 ± 1 16 ± 1 18 ± 1 21 ± 2 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 in detail and the small differences observed may not be significant in the light of heat to heat variations. Besides recovery processes, the precipitation of carbides, carbonitrides and nitrides occurs during tempering. In all three steels examined, M23 C6 carbides containing Cr, Fe, Mo (W) precipitate preferentially on the prior austenite grain boundaries and on the martensite lath boundaries. These precipitates retard the sub-grain growth and therefore increase the strength of the materials. In P91 steel mainly spheroidal Nb-rich carbonitrides are observed within the martensite laths. In P92 and E911 steels, three types of MX; Nb(C,N), plate-like VN and small complex Nb(C,N)-VN, are found (Ennis et al 1997, 2002; Zielinska-Lipiec et al 1997; Hald et al 1998; Vanstone 1998; Hättestrand & Andrén 2001). The microstructural evolution of P91, P92 and E911steels during ageing at service temperatures is discussed in detail by Hald et al (1998), and Ennis & Cyrska-Filemonowicz (2001). As an example, the microstructural evolution of P92 steel during very long term creep deformation at 600 and 650◦ C up to about 57,500 h is discussed in detail in the following section.

Figure 9. Laves phase precipitated in P92 after creep testing at 650◦ C for 6,500 h.

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Figure 10. TEM micrographs of (a) as received P91; (b) P92; and (c) E911 steels, showing the elongated sub-grains of tempered martensite.

Table 3. Dislocation density and mean sub-grain size of as received P91, P92 and E911 steels. Steel

Dislocation density × 10−14 (m−2 )

Mean sub-grain size, (µm)

P91 P92 E911

7·5 ± 0·8 7·9 ± 0·8 6·5 ± 0·6

0·4 ± 0·06 0·4 ± 0·09 0·5 ± 0·05


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Figure 11. Iso-stress plots at 120 MPa of secondary creep rate for P92 in as-received condition and after different heat treatments: 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, furnace cool to 780◦ C, 8 h/780◦ C, furnace cool to RT (ferritic structure, no martensite).

3.2 Creep rupture properties From the creep curves, the secondary creep rate is derived by taking the slope of the straightline portion of each curve. Figure 11 shows the creep strength, plotted as iso-stress curves at 120 MPa for the secondary creep rate, of the P92 material in different heat treatment conditions. Decreasing the austenitising temperature from the usual 1070◦ C to 970◦ C does not significantly affect the creep strength. However, increasing the tempering temperature from 775 to 835◦ C, a temperature just above the Ac1 transformation temperature of 825◦ C, leads to a marked fall in the creep strength. The low initial dislocation density produced by this heat treatment, with no other significant differences in the microstructural parameters, is the reason for the low creep strength. This effect is particularly relevant for weldments, since in the heat-affected zone there is always a region which has been exposed to temperatures around 850◦ C and therefore this region is the most likely site for creep failure according to Eggeler et al (1987). The dislocation density decreases not only as a result of increase in tempering temperature but also during prolonged exposure at lower temperatures, typical of service conditions. This can 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 first 3,000 h of exposure mean that there is therefore a danger of overestimating the long-term creep strength of P92 if there is a preponderance of short duration data in the evaluation, in which the dislocation density remains at a high level during the whole test. The extrapolation of the experimental data for estimation of the long-term creep rupture strength is discussed later in more detail. Figure 12 shows the secondary creep rates of P91, P92 and E911 at 600 and 650◦ C plotted against the applied stress (σ ). Data taken from published literature for the same heat are included where available. It can be seen that at high stresses the differences in the secondary creep rates of the three steels are relatively small. In the low stress region, however, the differences between the steels becomes more pronounced. The creep deformation characteristics may be described by the Norton equation (minimum creep rate is proportional to the applied stress to the power n) with two different values for the Norton stress exponent n. At high stresses, the value of n is around 16 while at lower stresses, the data conform to an n value of 6. The change in n is indicative of a change in the creep characteristics. Figure 12 also shows that at high stresses the differences in the secondary creep rates of the three steels are relatively

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

small and the high dislocation densities resulting from the martensite transformation dominates the deformation. In the low stress region, however, the differences between the steels become more pronounced. The evolution of the microstructure is examined to investigate the reason 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 of creep deformation. Well-developed sub-grains of low dislocation density in the interiors are characteristic features of long-term exposed specimens. Figure 13 shows the microstructures of P92 steel creep deformed at 600◦ C up to 32909 h and at 650◦ C up to 27433 h. The results of quantitative measurements of dislocation density and sub-grain width with increasing creep deformation time for P92 specimens exposed at 600 and 650◦ C are shown in figure 14. The quantitative evaluation of the microstructure shows that in the first 3,000 h of exposure there is a rapid reduction in the dislocation density and an increase of the sub-grain width. After long testing times, the dislocation density is around 2 × 1014 m−2 , a decrease of 75% compared with the as-received material. Similar effects were observed for the P91 and E911 steels. The microstructure of the P91 specimens creep specimens tested at 600◦ C for 9,200 h and steel E911 tested for 17,500 h (figures 15 a and b) exhibit polygonal sub-grains as a results of the extensive deformation. The dislocation densities in the steels P91, P92 and E911 after creep testing at 600◦ C are shown in table 4. As observed in P92 steel, the mean sub-grain widths of E911 and P91 increase and the dislocation densities within the sub-grains decrease with increasing creep exposure. Other important microstructural changes during creep deformation of the investigated steel are the size, morphology and distribution of the carbide, nitride and carbonitride precipitates as well as the chemical composition of the precipitates and the matrix. The fine MX, mainly plate-like VN and spheroidal Nb(C,N), precipitate intragranularly and act as obstacles for moving 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-rich particle is attached, are also present. Sub-grain boundaries and prior austenite grain boundaries


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Figure 13. TEM micrograph of P92 steel after creep deformation at (a) 600◦ C for 33,000 h and (b) 650◦ C for 27,500 h.

Figure 14. Dislocation densities and sub-grain widths for P92 after creep testing at 600 and 650◦ C.

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Figure 15. Dislocations interaction with fine carbide precipitates in steel P92 specimen after creep testing at 650◦ C for 6468 h.

were pinned by M23 C6 . These precipitates retard sub-grain growth and therefore increase the strength of the material. In steel exposed for longer durations, Laves phase Fe2 (W,Mo) (figure 16) and other intermetallic phases are formed. The particles of Laves phase formed during creep at 650◦ C were much larger than those precipitated at 600◦ C. The stacking faults formed within the Laves phases could be observed as characteristic streaks on the selected area electron diffraction pattern (figure 16b). This effect allowed easy distinction between the Laves phase and the M23 C6 , as both particles were precipitated very frequently in close vicinity, preferentially on sub-grain boundaries. Statistical measurements of the mean particle sizes revealed significant coarsening of the M23 C6 and intermetallic phases. The mean particle diameter of M23 C6 increased with increasing exposure time at 600 and 650◦ C, while the MX revealed insignificant 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 M23 C6 and MX particles formed in P92 specimens (head and gauge lengths) after creep deformation at 600◦ C. It can be seen that the coarsening of M23 C6 carbides is accelerated by stress while the effect of stress-induced coarsening of MX is insignificant. These findings concerning the coarsening of carbides in chromium steels are in good agreement with other studies (Eggeler et al 1987; Schaffernak et al 1998; Hättesrand & Andrén 2001). In addition to precipitate morphology and distribution, the chemical compositions of the phases 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.

P91 P92 E911

ca 1,000 h

ca 10,000 h

ca 1,70,000 h

22,000 h

(4·8 ± 0·5) × 1014 (5·4 ± 0·5) × 1014 (5·0 ± 0·5) × 1014

(1·1 ± 0·4) × 1014 (2·5 ± 0·5) × 1014 nd

nd (2·3 ± 0·5) × 1014 (2·2 ± 0·4) × 1014

nd nd (2·1 ± 0·4) × 1014

nd - not determined


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

showed that the M23 C6 are enriched in Cr, Fe, Mo and W, whereas the Laves phase is enriched in W and Mo (Czyrska-Filemonowicz et al 2001). An increase in the creep deformation temperature from 600 to 650◦ C results in precipitation of much larger particles of Laves phase. With increasing creep duration, the precipitation of the Laves phase removes Mo and W from the matrix solid solution and the strengthening of the matrix is decreased. This will be discussed in the next section. The results of our investigations are in general agreement with other TEM investigations describing the microstructure during creep deformation of 9% Cr steels (Foldyna et al 1996; Hald et al 1998; Nowakowski et al 1998; Vanstone 1998; Hättesrand & Andrén 2001b). Strang & Vodarek (1998), however, reported that there is another important precipitate that contributes to the softening of high chromium steels, namely the formation of large particles of Cr(Nb,V)N nitride (Z phase), replacing to some extent the fine VN precipitates. In the P92 steel the Z phase has not yet been observed in specimens exposed for up to 57,000 h (Ennis et al 2000). The question arises as to the relevance for long-term creep strength of the high initial dislocation density, which results from martensitic transformation and decreases during the first few thousand hours of exposure at service temperatures. To clarify this, creep tests were carried 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 of the M23 C6 and MX particles formed in P92 specimens (head and gauge lengths) after creep deformation at 600◦ C.

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Figure 18. Stress rupture curves at 600◦ C for P92, P91 and P91 compared with the values for a ferritic structure P92.

with carbide (M23 C6 ) and carbonitride precipitates. Some of the M23 C6 particles were particularly coarse, up to around 1 µm diameter. The secondary creep rate at 600◦ C, 120 MPa for specimens with ferritic microstructure was found to be 10,000 times higher that that of P92 in the usual austenitised and tempered condition. Figure 18 compares the 600◦ C stress rupture strengths of martensitic and ferritic P92 specimens. It is clear that the martensite transformation makes a considerable contribution to the creep rupture strength of the 9% Cr steels, even though the martensitic structure degenerates into ferrite and the initially high dislocation density decreases with exposure time. The initial martensitic structure does allow the formation of a stable subgrain structure and it is this substructure together with the fine carbonitride precipitates which confers the high creep rupture strength. A ferritic structure containing similar fine carbonitride precipitates to the martensitic P92 cannot provide the creep rupture strength required. This means that developments to strengthen the 9% Cr steels for applications at 600◦ C and above will be restricted by the necessity for the steels to exhibit the martensitic transformation. The maximum operating temperature must be sufficiently 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 to several factors. The scatter in the experimental data, especially between different heats within the compositional and heat treatment specifications for a given steel; the temperature and duration ranges of the available data; microstructural changes that occur in the materials during testing and influence deformation; and environmental effects, such as oxidation, that reduce the effective load-bearing cross-section may all lead to erroneous predictions of strength. One of 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 law equation, now widely used in creep deformation modelling. By relating the secondary creep rate to the applied stress and then using for design a maximum tolerable creep rate, a high temperature component could be dimensioned. Larson & Miller (1952) applied the Arrhenius


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rate equation for thermally activated creep processes. By assuming that the secondary creep rate 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 ), where T is the temperature in K, tR the rupture time and C is a constant, related to the activation energy for creep. This implies that long times at low temperatures are equivalent to shorter times at higher temperatures. This enables all stress rupture data to be plotted on a single master curve and for estimation of the stress rupture strength, no extrapolation of the curve 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 as a function of the reciprocal test temperature in K for constant stress, straight lines should be obtained 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 and high stress. Using the Larson–miller time–temperature parameter, the experimental data were evaluated and the 1,00,000 h stress rupture strengths at 600 and 650◦ C were calculated, first by using all the data and second by using only data for rupture times above 3,000 h. The extrapolated 1,00,000 h stress ruture strengths are shown in figure 19. After 3,000 h at 600 and 650◦ C, the dislocation density of P92 reaches a more or less constant value. It can be seen from figure 19 that the extrapolation using only data from the longer duration tests gives significantly lower 1,00,000 h rupture strengths at both temperatures. The Larson–Miller assumption that the secondary creep rate ε˙ s is inversely proportional to the rupture time tR was further refined by Monkman & Grant (1956) to give the Monkman– Grant (MG) equation, ε˙ sm · tR = K1 ,

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

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

where K1 and m are constants. Rearranging, we obtain log ε˙ s = −(1/m) log tR + log K1 . Plotting log rupture time as a function of log secondary creep rate therefore gives a straight line from which K1 and m can be determined. The constant m is usually in the range 0·8–1·2. The MG equation together with the Norton equation can be used for extrapolation. From the MG plot the secondary creep rate for a given rupture life can be read off and, from the Norton equation 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 of 1,00,000 h 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 110 MPa for P92, 105 MPa for E911 and 85 MPa for P91. The values obtained for P92 agree well with the Larson–Miller extrapolation made using only data of above 3,000 h duration.

5. Steam oxidation resistance of high chromium steels Almost without exception, high temperature materials rely on the selective oxidation of one or more 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 in the matrix and, second, the diffusion rate of these elements must be fast enough to ensure that they replenish the matrix below the growing scale, thus ensuring long-term protection. In the high 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 in the stress rupture strength. Long-term creep tests carried out at temperatures up to 650◦ C show 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 identified as consisting of (Fe,Cr)2 O3 and (Fe,Cr,Mn)3 O4 . However, in steam-containing atmospheres, the scales formed at 600 and 650◦ C are thick and consist of an external Fe3 O4 scale and an internal two-phase scale of Fe3 O4 and (Fe,Cr,Mn)3 O4 (Ehlers & Quadakkers 2001). Below the oxide scale, internal oxidation of Cr to form (Fe,Cr,Mn)3 O4 occurs at the martensite


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Ni coating

Fe3O4

(Fe, Cr, Mn) spinel

FeO + Cr2 O3 stringers steel matrix

Figure 21. Scale microstructure of P91 exposed for 1,000 h at 650◦ C in Ar-50% H2 O.

lath boundaries. Figure 21 shows typical oxide scales formed on the 9% Cr steels in steamcontaining 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 in load-bearing cross-section due to the oxide scale formation and the internally oxidised zone leads 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 at 600◦ C in steam is significant for components with wall thicknesses below about 6 mm. A second, and more difficult to quantify, effect is the spalling of the oxide. The presence of spalled oxide particles in the steam entering the turbine can cause erosion problems and local blockages. Furthermore, thick oxide scales on heated tubes can lead to decrease in heat transfer across the tube walls, resulting in overheating and subsequent creep failure of the tubes. In order to achieve acceptable oxidation resistance at 600 and 650◦ C, according to Ehlers et al (2001) a Cr content of at least 11% is required, to enable the formation of a protective spinel 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% H2 O and in air at 650◦ C.

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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 oxidation resistance can lead to toughness and fabrication difficulties. There are some additional factors that need to be considered in assessing the steam oxidation resistance of the high Cr steels. The effect of steam test parameters, such as flow rate and pressure, are not sufficiently well established and there is a need for further investigations to ensure that laboratory data reflects sufficiently well the expected behaviour in a power plant. Such tests are being carried out in the new COST 522 Action, see Allen et 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 requirement for improved steam oxidation resistance make contradictory demands on the steel composition. 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 to steam oxidation, Cr contents above 11% are necessary. The aim of current steel development is to raise the Cr content to 11–12%, and to add austenite-stabilising elements to produce the fully martensitic structure. In this way, it is hoped that stress rupture strength levels similar to 9% Cr steels can be reached at higher Cr contents in the range 11–12%. Some success has been reported by Tsuda et al (1998), but long-term creep data are not yet available and extrapolations are therefore uncertain. A promising line of development being followed in the new COST 522Action is the addition of Co, which enlarges the austenite field in the composition direction with no decrease in the a/g transition temperature, to 11–12% Cr steels (Allen et al 1998). The disadvantage of Co is, of course, the high cost of the raw material. Additions of V and Nb are necessary for precipitation hardening of the matrix. Regarding the solid solution strengthening additions of Mo 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 to suggest that W additions do confer improved stress rupture strength, at least to test durations of around 50,000 h that have been achieved in creep testing. There are other development routes towards a ferritic, high strength Cr steel with good resistance to steam oxidation. The effectiveness of increased amounts of nitride precipitates in conferring high stress rupture strength has being investigated by Goecmen et al (1998). The steels under development contain increased amounts of V and N and are given an appropriate heat treatment to produce finely dispersed nitrides. A similar development has been reported by Pugh (1998), the strengthening particles in this case being TiN, introduced by solid state reaction between CrN and the base steel containing 12% or more Cr and 1–2% Ti. The Osprey process and a powder metallurgical production route have been investigated. Another idea under investigation is the addition of FeWTi carbide powder to steel melts with the aim of producing uniform and fine dispersion of Ti and W carbides in the steel matrix, as described by Nutting (1999).


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The application of surface treatments and coatings is a promising, innovative line of development. The coating of large components and small diameter tubes does, however, present a 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 range up to 625◦ C. This target appears to be achievable. However, the extension of the operating temperature range of the steels to 650◦ C requires further increase in creep strength; it is by no means clear whether such an increase can be achieved either by enhancing the already established mechanisms of strengthening or by the new concepts being considered.

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