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Alloy compositions and mechanical properties of 9–12% chromium steels with martensitic–austenitic microstructure
Materials Science and Engineering A272 (1999) 292 – 299 www.elsevier.com/locate/msea
Alloy compositions and mechanical properties of 9–12% chromium steels with martensitic–austenitic microstructure U.E. Klotz a,*, C. Solenthaler a, P. Ernst b, P.J. Uggowitzer a, M.O. Speidel a a
Institute of Metallurgy, Swiss Federal Institute of Technology, ETH-Zentrum, CH-8092 Zu¨rich, Switzerland b ABB Corporate Research Limited, CH-5405 Baden-Da¨ttwil, Switzerland Received 7 June 1999; received in revised form 24 June 1999
Abstract Nitrogen alloyed 9–12% chromium steels with high amounts of manganese and nickel have been investigated. The steels are of a lamellar duplex microstructure, consisting of tempered martensite and about 30 vol.% austenite. The martensite is precipitation hardened with fine and homogeneously distributed vanadium nitrides. A good combination of strength and toughness has been achieved, with a room temperature yield strength level of up to 1000 MPa. As will be shown even at temperatures up to 550°C the new steels exhibit outstanding strength and ductility properties. Two different steels have been investigated. One steel was solely nitrogen alloyed, while in the other one parts of the nitrogen have been replaced by carbon. The steel which was alloyed with nitrogen only showed marked embrittlement after long-term ageing. Small carbon additions prevented ageing embrittlement. This is explained by the different precipitation behaviours. © 1999 Elsevier Science S.A. All rights reserved. Keywords: 9 – 12 % Cr-steels; Martensitic–austenitic microstructure; Microduplex; Long-term aging; Embrittlement
1. Introduction Steels with 9–12% chromium are often used for gas turbine compressors or steam turbine rotors and casings [1–3]. These steels offer a good combination of strength, high-temperature strength, toughness, and creep strength. However, all such steels in use today suffer from insufficient hot strength and creep strength at temperatures above 450°C (see Fig. 5). Therefore, nitrogen alloyed steels become more and more useful in terms of improving strength and creep resistance in this class of steels [4– 6]. The improvement of properties is caused by the precipitation of fine and homogeneously distributed Cr–V nitrides instead of coarser carbide precipitates [7,8]. In order to meet the requirements for future turbine application a new alloy design has been developed at ETH Zurich in collaboration with ABB Corporate Research, Baden [9]. It is based on a duplex
* Corresponding author. Present address: EMPA Du¨bendorf, U8 berlandstraße 129, CH 8600 Du¨bendorf, Switzerland. Tel.: +41-1823-45-87; fax: + 41-1-823-40-34. E-mail address: [email protected] (U.E. Klotz)
structure consisting of tempered martensite and austenite, which was proposed by Greenfield [10] as ‘‘austenite strengthened steel’’. Such a microstructure is formed during controlled tempering treatment of a martensitic microstructure after normalising and quenching. Because of the fine distribution of the austenite, this microstructure was referred to as a ‘microduplex’ structure. In microduplex steels the austenite formed during tempering can be stabilised by a well-balanced manganese-to-nickel ratio [9]. Due to the partial substitution of carbon by nitrogen a dispersion of fine and stable vanadium nitrides or carbo-nitrides is obtained, while in conventional nitrogen free steels the chromium–vanadium carbides are much less temper resistant. Typical composition ranges for microduplexsteels are (wt.%): 8–14 Cr, 0–6 Co, 2–7 Mn, 1–5 Ni, 0.5–2 Mo and/or W, 0.2–0.7 V, 0–0.1 C and 0.05– 0.15 N. The steels were subjected to mechanical testing to determine the properties by tensile tests and Charpy V-notch impact tests. Furthermore, the austenite content was measured after the quench and temper treatment as well as after long-term ageing. In order to obtain an explanation for the measured properties on a
0921-5093/99/$ - see front matter © 1999 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 9 ) 0 0 4 9 0 - 6
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3. Results
Table 1 Chemical composition of the alloys investigated (wt.%)
917 919
293
Fe
Cr
Co
Mn
Ni
Mo
V
N
C
3.1. Heat treatment and microstructure
bal. bal.
10.1 10.1
6.2 6.1
5.2 5.0
2.1 2.0
1.2 1.2
0.47 0.48
0.11 0.09
0.01 0.05
The heat treatment was optimised in order to achieve maximum hardness after quenching and an impact toughness of at least 50 J at room temperature after tempering. For both alloys Table 2 illustrates the optimised heat treatment procedure. The tempering was performed between the Ac1 and the Ac3 temperature. Therefore, during tempering a two-phase microstructure of austenite and tempered martensite was formed. A typical microstructure after normalising and tempering is shown in Fig. 1. The light optical image of alloy 917 (Fig. 1a) shows a microstructure typical for martensitic 9–12% chromium steels. Due to the low resolution of the light microscope, tempered martensite and austenite cannot be distinguished, while at much higher magnification in the TEM the two-phase microstructure becomes obvious (Fig. 1b). It is characterised by a lamellar arrangement of austenite (g, dark areas) and tempered martensite (a%, bright areas), with fine, plate-like precipitates of vanadium nitride in the tempered martensite as shown in Fig. 2a. Because of the two phases in fine lamellar arrangement, it is referred to as a ‘microduplex structure’. During tempering, partitioning of the alloying elements takes place. Manganese and nickel become enriched in the austenite, stabilising it against martensitic transformation during cooling from tempering temperature to room temperature. Fig. 2b shows the distribution of nickel revealed by EDX analysis at the specimen section given in Fig. 1b. The bright areas represent high nickel content and are austenitic, while the dark areas contain small amounts of nickel and are of tempered martensite. Manganese shows a similar distribution as nickel. The austenite content after heat treatment-about 30 vol.% for both alloys, as shown in Table 2, was measured by X-ray diffraction. Most of the vanadium precipitates as vanadium nitride in the tempered martensite while the austenite contains nearly no precipitates as shown in Fig. 2a. The precipitates have the shape of thin plates or discs on the {100} planes of the tempered martensite with the Bain orientation relationship [11,12]. Calculations by means of ThermoCalc® for alloy 917 (Fig. 3a), show that the two-phase microstructure of
microstructural basis, the materials have been investigated by means of transmission electron microscopy (TEM).
2. Material and experimental procedure Two typical alloy compositions are given in Table 1. While alloy 917 is nitrogen-alloyed only, alloy 919 contains nitrogen plus carbon. Apart from the typical alloying elements in 9 – 12% chromium steels, both alloys contain high amounts of cobalt, which is necessary to secure an austenitic solidification, by which degassing can be avoided. Both alloys were cast as 10 kg ingots under a nitrogen pressure of 0.08 MPa, forged to bars of 70×30× 200 mm, and subsequently normalised, quenched, and tempered. Tensile tests, Charpy impact tests, and creep tests were carried out according to the respective DIN standards. For tensile tests, specimens of 6 mm in diameter with a gauge length of 30 mm were used. Testing temperature varied between 22 and 700°C. Charpy V-notch specimens with 10× 10 mm cross-section were tested in the temperature range between − 196 and 200°C. Specimens for dilatometric tests, to determine the MS temperature, were 5 mm in diameter with a length of 10 mm. They were induction heated in vacuum at 5 K s − 1 and cooled at 20 K s − 1. The tempering time was 20 h at 600°C, 4 h at 625°C, 2 h at 650°C, and 1 h at 675, 700 and 750°C, respectively. The specimens investigated by TEM were cut from Charpy impact specimens from different heat treatment conditions, and were thinned by a mixture of perchloric acid and acetic acid. The element distribution in the microstructure was determined in the STEMmode by energy dispersive X-ray (EDX) distribution maps.
Table 2 Heat treatment conditions, grain size, and austenite content (VV (g)) after heat treatment Alloy
Normalising
Tempering
Grain size (mm)
Austenite content (%)]
917 919
1175°C/1.5 h/air cooling 1125°C/1.5 h/air cooling
600°C/20 h/air cooling 600°C/20 h/air cooling
87 38
26.3 9 3.3 37.2 94.0
294
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austenite and ferrite is in thermal equilibrium below temperatures of about 700°C. With decreasing temperature the austenite content decreases, while the ferrite content increases. Other phases in thermal equilibrium are vanadium nitride below 1200°C, and s-phase below 570°C. Because the two alloys investigated differ in carbon content by only 0.05 wt.%, the ThermoCalc® calculations are applicable for both alloys. As shown, the measured austenite content after tempering at 600°C for 20 h is somewhat lower than the equilibrium content. Fig. 3b shows the equilibrium composition of austenite as a function of temperature, indicating a strong enrichment of manganese and nickel in the austenite, especially at low temperatures. With increasing man-
Fig. 2. (a) TEM EDX-analysis showing the distribution of nickel after quench and temper treatment. Same specimen area as in Fig. 1b. Bright areas represent high nickel-content, and are austenitic. (b) Small, plate-like precipitates of vanadium nitride in the tempered martensite. Dark regions represent austenite (g), bright regions tempered martensite (a%). TEM micrograph.
Fig. 1. Microstructure of alloy 917 after quench and temper treatment. (a) Optical micrograph, (b) TEM bright-field micrograph. Bright areas represent tempered martensite, dark areas represent austenite.
ganese and nickel content, the austenite becomes more stable against martensitic transformation, see Eq. (1) below. The formation of s-phase lowers the chromium content in the austenite; therefore the austenite approaches a maximum chromium content at 570°C. The austenite stability after different tempering treatments has been investigated dilatometrically by thermal expansion tests. Fig. 4 shows the martensite start-temperature (MS temperature) as a function of tempering temperature. The calculated MS temperature was derived by an empirical equation of Jaffe [13], using the calculated austenite composition of Fig. 3b:
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295
MS (°C)=550–350%C – 40%Mn – 35%V – 20%Cr –17%Ni– 10%Cu – 10%Mo –8%W +15%Co + 30%Al (wt.%)
(1)
The calculated MS temperatures are very similar for both alloys; therefore in Fig. 4 the calculated values are shown only for alloy 917. The measured MS temperature of alloy 917 is in good agreement to the calculated results (Fig. 4). At tempering temperatures below 625°C the MS temperature is below 0°C for alloy 917, implying that the austenite formed during tempering remains stable during cooling to room temperature. Tempering temperatures above 625°C will result in unstable austenite. Because these steels are planned to be used for turbine parts which have operating temperatures between 500 and 600°C, insufficient austenite stability in this temperature range will result in volume changes
Fig. 4. Martensite start temperature of the austenite formed during tempering as a function of tempering temperature.
due to martensitic transformation and reaustenitisation within each thermal cycle of the turbine. These permanent volume changes will cause severe fatigue damage and, therefore, have to be avoided by sufficient austenite stability up to temperatures of 600–650°C. At tempering temperatures below 675°C the carbonalloyed steel alloy 919 shows lower austenite stability compared to alloy 917, as indicated by a higher MS temperature under the same tempering condition. Therefore, alloy 919 has to be tempered at temperatures of maximum 600°C to obtain a stable microduplex structure. Tempering above 600°C results in the formation of new martensite after tempering, which causes an increase in yield strength and a decrease in fracture elongation.
3.2. Mechanical properties
Fig. 3. Calculation of phase-equilibrium for alloy 917 by means of ThermoCalc®. (a) Phase-fractions, and (b) austenite composition as a function of temperature.
The yield strength (0.2% proof stress) as a function of test temperature is illustrated for both alloys for different heat treatment conditions (Fig. 5). For alloy 917 the strength decreases with increasing tempering time, while the impact energy at room temperature increases. Alloy 919 shows a somewhat lower strength at nearly the same room temperature impact toughness. With increasing testing temperature the strength of alloy 919 decreases slightly more pronouncedly compared to alloy 917. Compared to the commercially available alloys ‘Cronidur15’ and ‘X12CrNiMo12’ which are in use today for gas turbine compressor discs, the newly developed alloys 917 and 919 show higher strength levels at the same or even higher room temperature impact energy.
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The impact energy versus test temperature of the alloys in the tempered condition is shown in Fig. 6. For alloy 917, sharp transition behaviour is observed, with an upper shelf energy of about 80 J, and a ductile brittle transition temperature (DBTT) of −50°C. Alloy 919 in comparison shows inferior impact properties, meaning a lower upper shelf energy and a higher DBTT. The upper shelf energy is 10 – 20 J lower, while the transition temperature is shifted to higher temperatures (−20°C).
Fig. 5. Yield strength of alloy 917 and alloy 919 as a function of test temperature, compared to two typical steels for use in gas turbine compressors.
Fig. 7. Impact properties before and after long-term ageing at 550°C reported by Charpy impact tests as a function of temperature. (a) Alloy 917, (b) alloy 919.
3.3. Properties after long-term ageing
Fig. 6. Impact energy as a function of test temperature.
The long-term stability, especially in terms of the DBTT, are of great importance for materials for turbine application. Long-term ageing tests have been performed at 550°C up to 10 000 h. The ductile brittle transition behaviour before and after ageing is compared in Fig. 7 for both alloys investigated. With alloy 917 a pronounced loss in impact toughness is observed. After 2809 h the upper shelf energy is about 40 J, compared to 80 J before ageing, and the DBTT increased from −50°C before ageing to about +60°C after ageing. At a longer ageing time of 10 000 h, no
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297
Table 3 Result of room temperature tensile tests and austenite content (VV (g)) before (quenched and tempered) and after long-term ageinga Quenched and tempered
Quenched and tempered +550°C/2809 h
Quenched and tempered +550°C/10 000 h
917 Rp0.2/Rm (MPa) A5/Z (%) VV (g) (%)
916/1118 24.2/66.0 26.3 9 3.3
944/1122 22.0/52.8 26.19 5.3
864/1099 21.3/49.0 29.7 9 0.8
919 Rp0.2/Rm (MPa) A5/Z (%) VV (g) (%)
756/1101 22.0/59.8 37.2 9 4.0
a
– – –
719/992 24.3/56.8 34.6 93.6
Rp0.2, 0.2% proof stress; Rm, ultimate tensile strength; A5, fracture elongation; Z, reduction of area.
further embrittlement takes place. Alloy 919 instead, shows nearly no loss in impact energy after ageing at 550°C for 10 000 h. The DBTT after ageing is − 10°C, compared to −20°C before ageing, and the upper shelf energy is 60–80 J in both cases. The effect of long-term ageing on tensile properties is small, compared to the effect on impact properties. The result of tensile tests before and after long-term ageing is shown in Table 3. After ageing for 2809 h at 550°C, alloy 917 shows a small increase in yield strength Rp0.2, coupled with a small decrease in fracture elongation A5. With longer ageing time (10 000 h), the yield strength decreases again, at nearly constant fracture elongation. Alloy 919 shows a decrease in yield strength and a small increase in fracture elongation after ageing at 550°C for 10 000 h, compared to the quenched and tempered condition. The austenite content before and after longterm ageing is also given in Table 3. For alloy 917 the austenite content decreases in the first 2809 h of ageing, but remains constant at further ageing up to 10 000 h. No data is available for 2809 h of ageing for alloy 919, however, after 10 000 h of ageing a smaller austenite content was also measured, compared to the quenched and tempered condition. Therefore, it is expected that alloy 919 shows a similar behaviour compared to alloy 917.
4.1. Quenched and tempered condition Investigations by means of TEM in the quenched and tempered condition showed a different chromium distribution in alloy 917 and alloy 919. In alloy 919 the formation of chromium-rich precipitates at the phase boundaries of austenite and tempered martensite during tempering takes place, as can be seen in the chromium distribution revealed by TEM-EDX analysis (Fig. 8). Those precipitates were shown to be chromium–vanadium carbonitrides. In alloy 917 instead, only fine precipitates of vanadium nitride are present in the tempered martensite (Fig. 2), while, as the TEM studies showed, chromium is evenly distributed. The formation of chromium–vanadium carbonitrides in alloy 919 is assumed to be responsible for the lower austenite stability after tempering (Fig. 4), because the amount of carbon and nitrogen in solid solution in the austenite is reduced. As shown in [13–15] the marten-
4. Discussion The experimental results revealed a great difference between steel No. 917 alloyed with nitrogen only, and steel No. 919 alloyed with carbon plus nitrogen. This was most obvious in the impact energy after long-term ageing, and in the stability of the austenite against martensitic transformation. In the ‘short-term’ properties instead, like yield strength or impact energy in the quenched and tempered condition, the effect of carbon addition is much lower.
Fig. 8. Chromium-distribution in alloy 919 in the quenched and tempered condition after tempering at 600°/4 h. Bright regions (arrows) represent high chromium content.
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4.2. Properties after long-term ageing
Fig. 9. Microstructure of alloy 917 after long-term ageing at 550°C/ 2809 h revealed by TEM. (a) Electron image with precipitates of molybdenum-rich s-phase at the phase boundaries (dark). (b) Element distribution of nickel and chromium revealed by X-ray diffraction in the STEM-mode. Grey areas have high nickel-content and represent austenite; bright areas are low in nickel and represent tempered martensite. The chromium-intensity is plotted in black.
site start temperature increases with decreasing nitrogen and/or carbon content. The chromium – vanadium carbonitrides are further assumed to be responsible for weaker tensile and impact properties of alloy 919. Because of their coarse and film-like nature at the phase boundaries, they reduce strength and impact toughness. The fine and evenly distributed vanadium nitrides in alloy 917 instead, are more suitable to achieve high strength and toughness.
TEM-EDX analysis of alloy 917 after long-term ageing (550°C/2809 h) revealed the formation of a molybdenum-rich phase and chromium–vanadium rich phase (Fig. 9) at the phase boundaries of austenite and tempered martensite. The molybdenum-rich phase was identified by X-ray diffraction as s-phase, the chromium–vanadium-rich phase as chromium–vanadium nitride. For two reasons the molybdenum-rich s-phase is supposed to play a minor role in embrittlement. First, because alloy 919 with the same molybdenum-content as alloy 917, showed no embrittlement, and second, because of the fine and homogeneous dispersion of the s-phase. The main difference between both alloys is the precipitation of nitrides in the case of alloy 917 or carbonitrides in the case of alloy 919. The precipitation differs by means of the amount of precipitates, the moment at which precipitation takes place, and the place where precipitates are formed. Therefore, as will be shown, this is assumed to be the main factor causing embrittlement. Due to the precipitation of nitrides in alloy 917 during long-term ageing, the amount of nitrogen and chromium in solid solution in austenite and tempered martensite is lowered, influencing the austenite composition. However, this influence is localised to the austenite lamellae close to the phase boundary where precipitates are formed. The MS temperature of austenite increases with decreasing nitrogen content in the austenite [14,15]. A higher MS temperature, i.e. a reduced stability of the austenite, enhances its susceptibility to martensitic transformation during plastic deformation, for example in an impact test. The formation of virgin martensite, however, may lead to embrittlement. In alloy 919 chromium–vanadium carbonitrides are formed during tempering (Fig. 8) together with the austenite. This is different from alloy 917, which showed no coarse nitrides at the phase boundaries of austenite and tempered martensite after tempering, but only very fine vanadium nitrides within the tempered martensite (Fig. 2a). The formation of carbonitrides during tempering lowers impact energy and strength in the quenched and tempered condition because of their coarse and slim-like nature at the phase boundaries. However, the formation during tempering has also consequences on the stability of the austenite after tempering as well as after long-term ageing. After tempering the austenite will show a higher MS temperature, because of its lower carbon and nitrogen contents, which is clearly shown in Fig. 4 for tempering temperatures below 675°C. During long-term ageing the tendency to further precipitation of carbonitrides is much lower, and, therefore, practically no embrittle-
U.E. Klotz et al. / Materials Science and Engineering A272 (1999) 292–299
ment occurs in alloy 919 after ageing at 550°C for 10 000 h.
5. Conclusions The newly developed 9 – 12% chromium-steels with microduplex structure are characterised by a good combination of strength, and toughness even at elevated temperatures up to 550°C. The two steels investigated showed differences in their short-term and longterm properties. While the steel alloyed with nitrogen only is characterised by high strength and high toughness, it suffers from embrittlement after long-term ageing. The carbon – nitrogen-alloyed steel shows lower strength and toughness. However, it shows much better long-term properties, because no embrittlement took place. The reason for embrittlement is assumed to be the reduced stability of the austenite due to the formation of chromium – vanadium nitrides during ageing and the coarse and slim-like nature at the phase boundaries. The reduced nitrogen and chromium content in the austenite increases its proneness to martensitic transformation. Therefore, due to plastic deformation during mechanical testing, virgin martensite is formed, causing brittle fracture in impact tests.
.
299
Acknowledgements The authors would like to acknowledge ABB Corporate Research for financial support in this work and carrying-out the ageing experiments. Drs Reiner Steins and Alkan Go¨cmen are acknowledged for helpful discussions. References [1] K.J. Irvine, D.J. Crowe, F.B. Pickering, J. Iron Steel Inst. 195 (1960) 386. [2] V. Foldyna, Z. Kubon, M. Filip, K.-H. Mayer, C. Berger, Steel Res. 67 (1996) 375. [3] W. Bendick, M. Ring, Steel Res. 67 (1996) 382. [4] J. Hald, Steel Res. 67 (1996) 369. [5] A. Goecmen, R. Steins, C. Solenthaler, P.J. Uggowitzer, M.O. Speidel, ISIJ Int. 36 (1996) 768. [6] M. Shiga, K. Hidaka, S. Nakamura, Y. Fukui, ISIJ Int. 35 (1995) 1400. [7] B. Anthamatten, PhD-thesis, ETH Zurich, No. 9047, 1990. [8] F. Krafft, Fortschritt-Berichte VDI Reihe 5 Nr. 222, VDI, Dusseldorf, 1991. [9] R. Steins, PhD-thesis, ETH Zurich, No. 11997, 1996. [10] P. Greenfield, J.B. Marriott, K. Pithan (Eds.), A Review of the Properties of 9 – 12% Chromium Steels for Use as HP/IP Rotors in Advanced Steam Turbines, 1988, p.61. [11] D.H. Jack, Acta Metall. 24 (1976) 137. [12] M. Pope, P. Grieveson, K.H. Jack, Scand. J. Metall. 3 (1973) 29. [13] L.D. Jaffe, J.H. Hollomann, Trans. AIME 167 (1946) 617. [14] J. Kunze, B. Beyer, Z. Metallk. 88 (1997) 722. [15] U.R. Lenel, B.R. Knott, Metall. Trans 18A (1987) 767.
Alloy compositions and mechanical properties of 9–12% chromium steels with martensitic–austenitic microstructure U.E. Klotz a,*, C. Solenthaler a, P. Ernst b, P.J. Uggowitzer a, M.O. Speidel a a
Institute of Metallurgy, Swiss Federal Institute of Technology, ETH-Zentrum, CH-8092 Zu¨rich, Switzerland b ABB Corporate Research Limited, CH-5405 Baden-Da¨ttwil, Switzerland Received 7 June 1999; received in revised form 24 June 1999
Abstract Nitrogen alloyed 9–12% chromium steels with high amounts of manganese and nickel have been investigated. The steels are of a lamellar duplex microstructure, consisting of tempered martensite and about 30 vol.% austenite. The martensite is precipitation hardened with fine and homogeneously distributed vanadium nitrides. A good combination of strength and toughness has been achieved, with a room temperature yield strength level of up to 1000 MPa. As will be shown even at temperatures up to 550°C the new steels exhibit outstanding strength and ductility properties. Two different steels have been investigated. One steel was solely nitrogen alloyed, while in the other one parts of the nitrogen have been replaced by carbon. The steel which was alloyed with nitrogen only showed marked embrittlement after long-term ageing. Small carbon additions prevented ageing embrittlement. This is explained by the different precipitation behaviours. © 1999 Elsevier Science S.A. All rights reserved. Keywords: 9 – 12 % Cr-steels; Martensitic–austenitic microstructure; Microduplex; Long-term aging; Embrittlement
1. Introduction Steels with 9–12% chromium are often used for gas turbine compressors or steam turbine rotors and casings [1–3]. These steels offer a good combination of strength, high-temperature strength, toughness, and creep strength. However, all such steels in use today suffer from insufficient hot strength and creep strength at temperatures above 450°C (see Fig. 5). Therefore, nitrogen alloyed steels become more and more useful in terms of improving strength and creep resistance in this class of steels [4– 6]. The improvement of properties is caused by the precipitation of fine and homogeneously distributed Cr–V nitrides instead of coarser carbide precipitates [7,8]. In order to meet the requirements for future turbine application a new alloy design has been developed at ETH Zurich in collaboration with ABB Corporate Research, Baden [9]. It is based on a duplex
* Corresponding author. Present address: EMPA Du¨bendorf, U8 berlandstraße 129, CH 8600 Du¨bendorf, Switzerland. Tel.: +41-1823-45-87; fax: + 41-1-823-40-34. E-mail address: [email protected] (U.E. Klotz)
structure consisting of tempered martensite and austenite, which was proposed by Greenfield [10] as ‘‘austenite strengthened steel’’. Such a microstructure is formed during controlled tempering treatment of a martensitic microstructure after normalising and quenching. Because of the fine distribution of the austenite, this microstructure was referred to as a ‘microduplex’ structure. In microduplex steels the austenite formed during tempering can be stabilised by a well-balanced manganese-to-nickel ratio [9]. Due to the partial substitution of carbon by nitrogen a dispersion of fine and stable vanadium nitrides or carbo-nitrides is obtained, while in conventional nitrogen free steels the chromium–vanadium carbides are much less temper resistant. Typical composition ranges for microduplexsteels are (wt.%): 8–14 Cr, 0–6 Co, 2–7 Mn, 1–5 Ni, 0.5–2 Mo and/or W, 0.2–0.7 V, 0–0.1 C and 0.05– 0.15 N. The steels were subjected to mechanical testing to determine the properties by tensile tests and Charpy V-notch impact tests. Furthermore, the austenite content was measured after the quench and temper treatment as well as after long-term ageing. In order to obtain an explanation for the measured properties on a
0921-5093/99/$ - see front matter © 1999 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 9 ) 0 0 4 9 0 - 6
U.E. Klotz et al. / Materials Science and Engineering A272 (1999) 292–299
3. Results
Table 1 Chemical composition of the alloys investigated (wt.%)
917 919
293
Fe
Cr
Co
Mn
Ni
Mo
V
N
C
3.1. Heat treatment and microstructure
bal. bal.
10.1 10.1
6.2 6.1
5.2 5.0
2.1 2.0
1.2 1.2
0.47 0.48
0.11 0.09
0.01 0.05
The heat treatment was optimised in order to achieve maximum hardness after quenching and an impact toughness of at least 50 J at room temperature after tempering. For both alloys Table 2 illustrates the optimised heat treatment procedure. The tempering was performed between the Ac1 and the Ac3 temperature. Therefore, during tempering a two-phase microstructure of austenite and tempered martensite was formed. A typical microstructure after normalising and tempering is shown in Fig. 1. The light optical image of alloy 917 (Fig. 1a) shows a microstructure typical for martensitic 9–12% chromium steels. Due to the low resolution of the light microscope, tempered martensite and austenite cannot be distinguished, while at much higher magnification in the TEM the two-phase microstructure becomes obvious (Fig. 1b). It is characterised by a lamellar arrangement of austenite (g, dark areas) and tempered martensite (a%, bright areas), with fine, plate-like precipitates of vanadium nitride in the tempered martensite as shown in Fig. 2a. Because of the two phases in fine lamellar arrangement, it is referred to as a ‘microduplex structure’. During tempering, partitioning of the alloying elements takes place. Manganese and nickel become enriched in the austenite, stabilising it against martensitic transformation during cooling from tempering temperature to room temperature. Fig. 2b shows the distribution of nickel revealed by EDX analysis at the specimen section given in Fig. 1b. The bright areas represent high nickel content and are austenitic, while the dark areas contain small amounts of nickel and are of tempered martensite. Manganese shows a similar distribution as nickel. The austenite content after heat treatment-about 30 vol.% for both alloys, as shown in Table 2, was measured by X-ray diffraction. Most of the vanadium precipitates as vanadium nitride in the tempered martensite while the austenite contains nearly no precipitates as shown in Fig. 2a. The precipitates have the shape of thin plates or discs on the {100} planes of the tempered martensite with the Bain orientation relationship [11,12]. Calculations by means of ThermoCalc® for alloy 917 (Fig. 3a), show that the two-phase microstructure of
microstructural basis, the materials have been investigated by means of transmission electron microscopy (TEM).
2. Material and experimental procedure Two typical alloy compositions are given in Table 1. While alloy 917 is nitrogen-alloyed only, alloy 919 contains nitrogen plus carbon. Apart from the typical alloying elements in 9 – 12% chromium steels, both alloys contain high amounts of cobalt, which is necessary to secure an austenitic solidification, by which degassing can be avoided. Both alloys were cast as 10 kg ingots under a nitrogen pressure of 0.08 MPa, forged to bars of 70×30× 200 mm, and subsequently normalised, quenched, and tempered. Tensile tests, Charpy impact tests, and creep tests were carried out according to the respective DIN standards. For tensile tests, specimens of 6 mm in diameter with a gauge length of 30 mm were used. Testing temperature varied between 22 and 700°C. Charpy V-notch specimens with 10× 10 mm cross-section were tested in the temperature range between − 196 and 200°C. Specimens for dilatometric tests, to determine the MS temperature, were 5 mm in diameter with a length of 10 mm. They were induction heated in vacuum at 5 K s − 1 and cooled at 20 K s − 1. The tempering time was 20 h at 600°C, 4 h at 625°C, 2 h at 650°C, and 1 h at 675, 700 and 750°C, respectively. The specimens investigated by TEM were cut from Charpy impact specimens from different heat treatment conditions, and were thinned by a mixture of perchloric acid and acetic acid. The element distribution in the microstructure was determined in the STEMmode by energy dispersive X-ray (EDX) distribution maps.
Table 2 Heat treatment conditions, grain size, and austenite content (VV (g)) after heat treatment Alloy
Normalising
Tempering
Grain size (mm)
Austenite content (%)]
917 919
1175°C/1.5 h/air cooling 1125°C/1.5 h/air cooling
600°C/20 h/air cooling 600°C/20 h/air cooling
87 38
26.3 9 3.3 37.2 94.0
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austenite and ferrite is in thermal equilibrium below temperatures of about 700°C. With decreasing temperature the austenite content decreases, while the ferrite content increases. Other phases in thermal equilibrium are vanadium nitride below 1200°C, and s-phase below 570°C. Because the two alloys investigated differ in carbon content by only 0.05 wt.%, the ThermoCalc® calculations are applicable for both alloys. As shown, the measured austenite content after tempering at 600°C for 20 h is somewhat lower than the equilibrium content. Fig. 3b shows the equilibrium composition of austenite as a function of temperature, indicating a strong enrichment of manganese and nickel in the austenite, especially at low temperatures. With increasing man-
Fig. 2. (a) TEM EDX-analysis showing the distribution of nickel after quench and temper treatment. Same specimen area as in Fig. 1b. Bright areas represent high nickel-content, and are austenitic. (b) Small, plate-like precipitates of vanadium nitride in the tempered martensite. Dark regions represent austenite (g), bright regions tempered martensite (a%). TEM micrograph.
Fig. 1. Microstructure of alloy 917 after quench and temper treatment. (a) Optical micrograph, (b) TEM bright-field micrograph. Bright areas represent tempered martensite, dark areas represent austenite.
ganese and nickel content, the austenite becomes more stable against martensitic transformation, see Eq. (1) below. The formation of s-phase lowers the chromium content in the austenite; therefore the austenite approaches a maximum chromium content at 570°C. The austenite stability after different tempering treatments has been investigated dilatometrically by thermal expansion tests. Fig. 4 shows the martensite start-temperature (MS temperature) as a function of tempering temperature. The calculated MS temperature was derived by an empirical equation of Jaffe [13], using the calculated austenite composition of Fig. 3b:
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MS (°C)=550–350%C – 40%Mn – 35%V – 20%Cr –17%Ni– 10%Cu – 10%Mo –8%W +15%Co + 30%Al (wt.%)
(1)
The calculated MS temperatures are very similar for both alloys; therefore in Fig. 4 the calculated values are shown only for alloy 917. The measured MS temperature of alloy 917 is in good agreement to the calculated results (Fig. 4). At tempering temperatures below 625°C the MS temperature is below 0°C for alloy 917, implying that the austenite formed during tempering remains stable during cooling to room temperature. Tempering temperatures above 625°C will result in unstable austenite. Because these steels are planned to be used for turbine parts which have operating temperatures between 500 and 600°C, insufficient austenite stability in this temperature range will result in volume changes
Fig. 4. Martensite start temperature of the austenite formed during tempering as a function of tempering temperature.
due to martensitic transformation and reaustenitisation within each thermal cycle of the turbine. These permanent volume changes will cause severe fatigue damage and, therefore, have to be avoided by sufficient austenite stability up to temperatures of 600–650°C. At tempering temperatures below 675°C the carbonalloyed steel alloy 919 shows lower austenite stability compared to alloy 917, as indicated by a higher MS temperature under the same tempering condition. Therefore, alloy 919 has to be tempered at temperatures of maximum 600°C to obtain a stable microduplex structure. Tempering above 600°C results in the formation of new martensite after tempering, which causes an increase in yield strength and a decrease in fracture elongation.
3.2. Mechanical properties
Fig. 3. Calculation of phase-equilibrium for alloy 917 by means of ThermoCalc®. (a) Phase-fractions, and (b) austenite composition as a function of temperature.
The yield strength (0.2% proof stress) as a function of test temperature is illustrated for both alloys for different heat treatment conditions (Fig. 5). For alloy 917 the strength decreases with increasing tempering time, while the impact energy at room temperature increases. Alloy 919 shows a somewhat lower strength at nearly the same room temperature impact toughness. With increasing testing temperature the strength of alloy 919 decreases slightly more pronouncedly compared to alloy 917. Compared to the commercially available alloys ‘Cronidur15’ and ‘X12CrNiMo12’ which are in use today for gas turbine compressor discs, the newly developed alloys 917 and 919 show higher strength levels at the same or even higher room temperature impact energy.
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The impact energy versus test temperature of the alloys in the tempered condition is shown in Fig. 6. For alloy 917, sharp transition behaviour is observed, with an upper shelf energy of about 80 J, and a ductile brittle transition temperature (DBTT) of −50°C. Alloy 919 in comparison shows inferior impact properties, meaning a lower upper shelf energy and a higher DBTT. The upper shelf energy is 10 – 20 J lower, while the transition temperature is shifted to higher temperatures (−20°C).
Fig. 5. Yield strength of alloy 917 and alloy 919 as a function of test temperature, compared to two typical steels for use in gas turbine compressors.
Fig. 7. Impact properties before and after long-term ageing at 550°C reported by Charpy impact tests as a function of temperature. (a) Alloy 917, (b) alloy 919.
3.3. Properties after long-term ageing
Fig. 6. Impact energy as a function of test temperature.
The long-term stability, especially in terms of the DBTT, are of great importance for materials for turbine application. Long-term ageing tests have been performed at 550°C up to 10 000 h. The ductile brittle transition behaviour before and after ageing is compared in Fig. 7 for both alloys investigated. With alloy 917 a pronounced loss in impact toughness is observed. After 2809 h the upper shelf energy is about 40 J, compared to 80 J before ageing, and the DBTT increased from −50°C before ageing to about +60°C after ageing. At a longer ageing time of 10 000 h, no
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Table 3 Result of room temperature tensile tests and austenite content (VV (g)) before (quenched and tempered) and after long-term ageinga Quenched and tempered
Quenched and tempered +550°C/2809 h
Quenched and tempered +550°C/10 000 h
917 Rp0.2/Rm (MPa) A5/Z (%) VV (g) (%)
916/1118 24.2/66.0 26.3 9 3.3
944/1122 22.0/52.8 26.19 5.3
864/1099 21.3/49.0 29.7 9 0.8
919 Rp0.2/Rm (MPa) A5/Z (%) VV (g) (%)
756/1101 22.0/59.8 37.2 9 4.0
a
– – –
719/992 24.3/56.8 34.6 93.6
Rp0.2, 0.2% proof stress; Rm, ultimate tensile strength; A5, fracture elongation; Z, reduction of area.
further embrittlement takes place. Alloy 919 instead, shows nearly no loss in impact energy after ageing at 550°C for 10 000 h. The DBTT after ageing is − 10°C, compared to −20°C before ageing, and the upper shelf energy is 60–80 J in both cases. The effect of long-term ageing on tensile properties is small, compared to the effect on impact properties. The result of tensile tests before and after long-term ageing is shown in Table 3. After ageing for 2809 h at 550°C, alloy 917 shows a small increase in yield strength Rp0.2, coupled with a small decrease in fracture elongation A5. With longer ageing time (10 000 h), the yield strength decreases again, at nearly constant fracture elongation. Alloy 919 shows a decrease in yield strength and a small increase in fracture elongation after ageing at 550°C for 10 000 h, compared to the quenched and tempered condition. The austenite content before and after longterm ageing is also given in Table 3. For alloy 917 the austenite content decreases in the first 2809 h of ageing, but remains constant at further ageing up to 10 000 h. No data is available for 2809 h of ageing for alloy 919, however, after 10 000 h of ageing a smaller austenite content was also measured, compared to the quenched and tempered condition. Therefore, it is expected that alloy 919 shows a similar behaviour compared to alloy 917.
4.1. Quenched and tempered condition Investigations by means of TEM in the quenched and tempered condition showed a different chromium distribution in alloy 917 and alloy 919. In alloy 919 the formation of chromium-rich precipitates at the phase boundaries of austenite and tempered martensite during tempering takes place, as can be seen in the chromium distribution revealed by TEM-EDX analysis (Fig. 8). Those precipitates were shown to be chromium–vanadium carbonitrides. In alloy 917 instead, only fine precipitates of vanadium nitride are present in the tempered martensite (Fig. 2), while, as the TEM studies showed, chromium is evenly distributed. The formation of chromium–vanadium carbonitrides in alloy 919 is assumed to be responsible for the lower austenite stability after tempering (Fig. 4), because the amount of carbon and nitrogen in solid solution in the austenite is reduced. As shown in [13–15] the marten-
4. Discussion The experimental results revealed a great difference between steel No. 917 alloyed with nitrogen only, and steel No. 919 alloyed with carbon plus nitrogen. This was most obvious in the impact energy after long-term ageing, and in the stability of the austenite against martensitic transformation. In the ‘short-term’ properties instead, like yield strength or impact energy in the quenched and tempered condition, the effect of carbon addition is much lower.
Fig. 8. Chromium-distribution in alloy 919 in the quenched and tempered condition after tempering at 600°/4 h. Bright regions (arrows) represent high chromium content.
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4.2. Properties after long-term ageing
Fig. 9. Microstructure of alloy 917 after long-term ageing at 550°C/ 2809 h revealed by TEM. (a) Electron image with precipitates of molybdenum-rich s-phase at the phase boundaries (dark). (b) Element distribution of nickel and chromium revealed by X-ray diffraction in the STEM-mode. Grey areas have high nickel-content and represent austenite; bright areas are low in nickel and represent tempered martensite. The chromium-intensity is plotted in black.
site start temperature increases with decreasing nitrogen and/or carbon content. The chromium – vanadium carbonitrides are further assumed to be responsible for weaker tensile and impact properties of alloy 919. Because of their coarse and film-like nature at the phase boundaries, they reduce strength and impact toughness. The fine and evenly distributed vanadium nitrides in alloy 917 instead, are more suitable to achieve high strength and toughness.
TEM-EDX analysis of alloy 917 after long-term ageing (550°C/2809 h) revealed the formation of a molybdenum-rich phase and chromium–vanadium rich phase (Fig. 9) at the phase boundaries of austenite and tempered martensite. The molybdenum-rich phase was identified by X-ray diffraction as s-phase, the chromium–vanadium-rich phase as chromium–vanadium nitride. For two reasons the molybdenum-rich s-phase is supposed to play a minor role in embrittlement. First, because alloy 919 with the same molybdenum-content as alloy 917, showed no embrittlement, and second, because of the fine and homogeneous dispersion of the s-phase. The main difference between both alloys is the precipitation of nitrides in the case of alloy 917 or carbonitrides in the case of alloy 919. The precipitation differs by means of the amount of precipitates, the moment at which precipitation takes place, and the place where precipitates are formed. Therefore, as will be shown, this is assumed to be the main factor causing embrittlement. Due to the precipitation of nitrides in alloy 917 during long-term ageing, the amount of nitrogen and chromium in solid solution in austenite and tempered martensite is lowered, influencing the austenite composition. However, this influence is localised to the austenite lamellae close to the phase boundary where precipitates are formed. The MS temperature of austenite increases with decreasing nitrogen content in the austenite [14,15]. A higher MS temperature, i.e. a reduced stability of the austenite, enhances its susceptibility to martensitic transformation during plastic deformation, for example in an impact test. The formation of virgin martensite, however, may lead to embrittlement. In alloy 919 chromium–vanadium carbonitrides are formed during tempering (Fig. 8) together with the austenite. This is different from alloy 917, which showed no coarse nitrides at the phase boundaries of austenite and tempered martensite after tempering, but only very fine vanadium nitrides within the tempered martensite (Fig. 2a). The formation of carbonitrides during tempering lowers impact energy and strength in the quenched and tempered condition because of their coarse and slim-like nature at the phase boundaries. However, the formation during tempering has also consequences on the stability of the austenite after tempering as well as after long-term ageing. After tempering the austenite will show a higher MS temperature, because of its lower carbon and nitrogen contents, which is clearly shown in Fig. 4 for tempering temperatures below 675°C. During long-term ageing the tendency to further precipitation of carbonitrides is much lower, and, therefore, practically no embrittle-
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ment occurs in alloy 919 after ageing at 550°C for 10 000 h.
5. Conclusions The newly developed 9 – 12% chromium-steels with microduplex structure are characterised by a good combination of strength, and toughness even at elevated temperatures up to 550°C. The two steels investigated showed differences in their short-term and longterm properties. While the steel alloyed with nitrogen only is characterised by high strength and high toughness, it suffers from embrittlement after long-term ageing. The carbon – nitrogen-alloyed steel shows lower strength and toughness. However, it shows much better long-term properties, because no embrittlement took place. The reason for embrittlement is assumed to be the reduced stability of the austenite due to the formation of chromium – vanadium nitrides during ageing and the coarse and slim-like nature at the phase boundaries. The reduced nitrogen and chromium content in the austenite increases its proneness to martensitic transformation. Therefore, due to plastic deformation during mechanical testing, virgin martensite is formed, causing brittle fracture in impact tests.
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Acknowledgements The authors would like to acknowledge ABB Corporate Research for financial support in this work and carrying-out the ageing experiments. Drs Reiner Steins and Alkan Go¨cmen are acknowledged for helpful discussions. References [1] K.J. Irvine, D.J. Crowe, F.B. Pickering, J. Iron Steel Inst. 195 (1960) 386. [2] V. Foldyna, Z. Kubon, M. Filip, K.-H. Mayer, C. Berger, Steel Res. 67 (1996) 375. [3] W. Bendick, M. Ring, Steel Res. 67 (1996) 382. [4] J. Hald, Steel Res. 67 (1996) 369. [5] A. Goecmen, R. Steins, C. Solenthaler, P.J. Uggowitzer, M.O. Speidel, ISIJ Int. 36 (1996) 768. [6] M. Shiga, K. Hidaka, S. Nakamura, Y. Fukui, ISIJ Int. 35 (1995) 1400. [7] B. Anthamatten, PhD-thesis, ETH Zurich, No. 9047, 1990. [8] F. Krafft, Fortschritt-Berichte VDI Reihe 5 Nr. 222, VDI, Dusseldorf, 1991. [9] R. Steins, PhD-thesis, ETH Zurich, No. 11997, 1996. [10] P. Greenfield, J.B. Marriott, K. Pithan (Eds.), A Review of the Properties of 9 – 12% Chromium Steels for Use as HP/IP Rotors in Advanced Steam Turbines, 1988, p.61. [11] D.H. Jack, Acta Metall. 24 (1976) 137. [12] M. Pope, P. Grieveson, K.H. Jack, Scand. J. Metall. 3 (1973) 29. [13] L.D. Jaffe, J.H. Hollomann, Trans. AIME 167 (1946) 617. [14] J. Kunze, B. Beyer, Z. Metallk. 88 (1997) 722. [15] U.R. Lenel, B.R. Knott, Metall. Trans 18A (1987) 767.