Creep and microstructural processes in a low-alloy 2.25%Cr1.6%W steel (ASTM Grade 23)

Materials Characterization 109 (2015) 1–8

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Creep and microstructural processes in a low-alloy 2.25%Cr1.6%W steel (ASTM Grade 23) K. Kucharova a, V. Sklenicka a,b,⁎, M. Kvapilova a,b, M. Svoboda a,b a b

Institute of Physics of Materials, Academy of Sciences of the Czech Republic, CZ-616 62 Brno, Czech Republic CEITEC — IPM, Institute of Physics of Materials, Academy of Sciences of the Czech Republic, CZ-616 62 Brno, Czech Republic

a r t i c l e

i n f o

Article history: Received 13 July 2015 Received in revised form 10 August 2015 Accepted 12 August 2015 Available online 14 August 2015 Keywords: Bainitic steel Low-alloy steel Creep strength Microstructural changes Carbide precipitation

a b s t r a c t A low-alloy 2.25%Cr1%Mo steel (ASTM Grade 22) has been greatly improved by the substitution of almost all of the 1%Mo by 1.6%W. The improved material has been standardized as P/T23 steel (Fe–2.25Cr–1.6W–0.25V– 0.05Nb–0.07C). The present investigation was conducted on T23 steel in an effort to obtain a more complete description and understanding of the role of the microstructural evolution and deformation processes in hightemperature creep. Constant load tensile creep tests were carried out in an argon atmosphere in the temperature range 500–650 °C at stresses ranging from 50 to 400 MPa. It was found that the diffusion in the matrix lattice is the creep-rate controlling process. The results of an extensive transmission electron microscopy (TEM) analysis programme to investigate microstructure evolution as a function of temperature are described and compared with the thermodynamic calculations using the software package Thermo-Calc. The significant creep-strength drop of T23 steel after long-term creep exposures can be explained by the decrease in dislocation hardening, precipitation hardening and solid solution hardening due to the instability of the microstructure at high temperature. © 2015 Elsevier Inc. All rights reserved.

1. Introduction The continuing effort to reduce the costs of advanced power plants while striving to improve efficiency has focused attention on the need for high strength steels at moderate costs for critical high-temperature components. This target can only be reached by increasing plant efficiency through raising the steam pressure and temperature. The most recent advancement in bainitic low-alloy Cr steels is the improvement in the creep rupture strength of 2.25 and 3Cr steels [1]. The advanced low-chromium steel grade T/P23 (ASTM Grade 23) [1–3] is a candidate material for the components of ultra-supercritical power plants, and as a potential replacement material for conventional low-alloy ferritic steels such as ASTM Grade 22 (T/P22) in older plants. Thus, a 2.25Cr– 1.6W–0.1Mo–0.25V–0.05Nb–0.07C (ASTM Grade 23) steel was derived from conventional low-alloy 2.25%Cr1%Mo steel (ASTM Grade 22) by the addition of W, V, Nb and B and the optimization of the C and Mo content to enhance the creep resistance and ensure easier weldability. Substituting W for Mo retards the evolution of the microstructure and remarkably improves the creep rupture strength. To further enhance its creep rupture strength, the elements V, Nb and B are added to form stable carbonitrides of the MX type. In addition, the weldability of T23 is greatly improved by the optimization of the C content and it can be ⁎ Corresponding author at: Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Zizkova 22, 616 62 Brno, Czech Republic. E-mail address: [email protected] (V. Sklenicka).

http://dx.doi.org/10.1016/j.matchar.2015.08.008 1044-5803/© 2015 Elsevier Inc. All rights reserved.

used in the as-welded condition without preheating and post-weld heat treatment. The improved steel has been standardized in ASTM as Grade 23 (T/ P23) with the nominal composition of Fe–2.25Cr–1.6W–0.1Mo–0.25V– 0.05Nb–0.07C. This new bainitic steel T/P23 is well suited for manufacturing water wall panels for ultra-super critical boilers (USCB), it is also used for the superheaters and reheaters of conventional boilers and the heat recovery steam generators (HRSG) of combined cycle plants and in other potential fields of high-temperature application. During the last two decades, extensive investigations have been carried out in an attempt to understand the improved properties of T23 steel, to relate microstructures to observed creep properties, especially creep strength, and to gain insight to facilitate the development of improved low-alloy Cr bainitic creep-resistant steels [1–14]. However, only very limited reports are available describing stress dependences of the minimum creep rate and, prevailingly, the time to fracture of T23 steel [3,4,7,9–14]. Therefore, we are still some distance from a full understanding of the processes occurring during the long-term creep exposure of T23 steel and a number of problems remain unsolved. Thus, there is still disagreement concerning the dominating creepstrength degradation mechanism(s). In the present paper we will try to further clarify the creep behaviour and properties of the advanced low-alloy bainitic T23 steel. The standard isothermal creep tests of this steel and measured minimum creep strain rates at five testing temperatures ranging from 500 to 650 °C are analysed with the aim of ascertaining whether the minimum

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Table 1 Summary of creep data. Temperature Stress

Time to fracture Strain to fracture Minimum creep rate

T [°C]

σ [MPa] tf [h]

εf

500 500 500 500 500 500 550 550 550 550 550 550 550 600 600 600 600 600 600 600 600 600 600 600 625 625 625 650 650 650 650 650 650 650

250 275 300 350 375 400 175 200 200 225 250 300 350 120 125 125 140 150 150 160 175 200 250 300 125 150 200 50 75 100 125 150 175 200

12.1 9.5 7.3 7.5 11.2 10.4 13.3 12.7 12.2 11.1 9.8 8.7 9.8 3.0 3.5 3.2 2.9 11.7 7.3 20.4 11.2 13.5 10.2 10.0 12.4 10.1 10.0 10.4 9.9 17.3 16.4 18.5 15.4 15.4

6500.6 2100.2 867.8 99.5 20.0 1.8 12,246.0 1309.5 1538.8 289.7 156.9 32.2 1.8 11,456.8 10,263.4 12,269.8 12,547.9 2898.8 2582.5 652.7 277.2 63.3 5.8 0.4 1901.4 270.9 8.1 2561.5 3632.3 1571.3 284.4 65.1 9.2 1.7

[%]

εm [s−1]


8.0 × 10−11 5.5 × 10−10 1.4 × 10−9 1.7 × 10−8 1.6 × 10−7 4.1 × 10−6 6.8 × 10−10 2.1 × 10−9 2.9 × 10−9 5.4 × 10−9 1.0 × 10−8 4.6 × 10−8 2.4 × 10−6 4.7 × 10−10 8.3 × 10−11 8.9 × 10−11 2.6 × 10−9 2.8 × 10−9 1.5 × 10−8 2.9 × 10−8 7.0 × 10−8 5.8 × 10−7 8.7 × 10−6 5.1 × 10−9 2.4 × 10−8 6.6 × 10−7 3.0 × 10−10 5.4 × 10−10 5.5 × 10−9 2.6 × 10−8 1.2 × 10−7 7.8 × 10−7 3.8 × 10−6

strain rate of the steel under investigation is lattice-diffusion controlled. Further objectives of these creep tests are to determine the lifetimes over a range of applied stresses at selected testing temperatures. Finally, an attempt is made to identify the structural processes to obtain a more complete description of the role of microstructural stability in the creep of this steel. 2. Steel T23 and experimental procedures The low-alloy bainitic steel T23 was produced by Wakayama Steel Kainan Works, Japan, with the following chemical composition (in

Fig. 2. Standard creep curves of T23 steel with creep at different applied stresses and temperatures.

wt.%): 0.06C, 0.28Si, 0.28Mn, 0.02P, 0.003S, 2.24Cr, 0.09Mo, 0.24V, 0.0055B, 0.006N, 1.49W, and 0.04Nb. The steel was received in the form of a tube OD × WT 38 × 8 mm with the following heat treatment: 1045 °C/10 min/air + 770 °C/60 min/air [15]. Uniaxial constant load tensile creep tests were carried out using the flat creep specimens at temperatures from 500 to 650 °C with the testing temperatures continuously monitored and maintained constant to within ±0.5 °C of the desired value. The initial applied tensile stresses ranged from 50 to 400 MPa. The creep specimens having a gauge length of 50 mm and a cross-section area 5 × 3.2 mm were machined directly from the tube. The creep elongations were measured using a linear variable differential transducer (the strain was measured with a sensitivity of 5 × 10−6) and they were continuously recorded digitally and computer processed. All of the creep specimens were run to the final fracture. The creep tests were performed in a purified argon atmosphere to avoid the detrimental effect of oxidation [9]. Following creep testing, samples were prepared for microstructural examination by means of both transmission and scanning electron microscopy. Transmission electron microscopy (TEM) studies were carried out on carbon extraction replicas and thin foils prepared from both the head and gauge parts of the specimens subjected to creep using a Jeol 2100F microscope operating at 200 kV. Particles of the secondary phases

Fig. 1. Standard creep curves of T23 steel with creep at 600 °C and different applied stresses.

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Fig. 3. Modified creep curves of the standard creep curves shown in Fig. 2: (a) strain rate vs. time and (b) creep rate vs. strain.

extracted into carbon replicas were identified by means of selected area diffraction (SAED) and their local chemical composition was measured by energy dispersive X-ray spectroscopy (EDS) using the 10-nm probe size and thin sample approximation. The measured metallic element content was recalculated assuming stoichiometric carbon content for the carbides. The fracture surfaces were investigated using a scanning electron microscope Tescan Lyra 3. The experimental results of the carbide analyses were compared with the results of equilibrium thermodynamic calculations using the software package Thermo-Calc [16] and own thermodynamic database STEEL 16 [17–19].

as shown in Fig. 3. It is clear that neither curve exhibits a well-defined steady state. In fact, this stage is reduced to an inflection point (Fig. 3b). Further, the curves in Fig. 3 show that the very short primary stage is terminated by attainment of the minimum creep rate. It is important to note that the strain at which the minimum creep rate is reached (ε ~ 0.015, Fig. 3b) is very small. By contrast, the tertiary stage is fairly extensive and covers practically a dominant part of the creep tests (Fig. 3a). A variation of the strain to fracture εf, with stress and temperature is shown in Fig. 4. The values of strain to fracture are on an average of only 7–15% slightly depending on temperature and stress. In Fig. 5, the minimum creep rates εm are plotted against the applied stress σ on a bilogarithmic scale. Inspection of Fig. 5 leads to an observation that the stress dependences of the minimum creep rates for T23 steel at various temperatures show a different trend. The slopes and, therefore, the apparent stress exponents of the minimum creep rate n ¼


3. Experimental results 3.1. Creep results The results of the tensile creep tests are summarized in Table 1. Figs. 1 and 2 show selected standard creep curves for T23 steel for uniaxial tensile creep tests conducted at different applied stresses and duration of creep exposures. It is important to note that the standard creep curves shown in Figs. 1 and 2 do not clearly indicate the individual creep stages. However, these standards ε vs. t curves can be easily replotted in the form of the instantaneous strain rate ε vs. time t and/or ε vs. strain ε

ð∂ ln ε m =∂ lnσ ÞT depend strongly on the applied stress. In general, the values of the stress exponent decrease with increasing test temperature and decreasing applied stress. The double logarithmic plots of the time to fracture tf as a function of applied stress are shown in Fig. 6. Similarly, the stress exponents of the time to fracture m = −(∂ ln tf/∂ ln σ)T depend strongly on the value of the applied stress.

Fig. 4. Stress dependences of the strain to fracture εf.

Fig. 5. Stress dependences of the minimum creep rate ε m .













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energy values were determined: 208 ± 230 kJ/mol (150 MPa), 446 ± 151 kJ/mol (200 MPa) and 499 ± 100 kJ/mol (250 MPa). 3.2. Microstructural investigations

Fig. 6. Stress dependences of the time to fracture tf.

A very important characteristic of creep is the strong dependence of the minimum creep rate on temperature which can be described by the activation energy for creep QC. The activation energy QC can be defined as Q C ¼ ½∂ logε m =ð∂−1=kTÞσ and be obtained as a k-multiple of





the slope of log εm vs. 1/T plots which can be re-plotted from Fig. 5. It was found that the values of QC are dependent on stress. Due to an experimental scatter of the data the following apparent activation

In the normalized and tempered condition (i.e., in the as-received state) the microstructure is 100% tempered bainite (Fig. 7). Creep strength in bainitic steels is determined primarily by the precipitates, and to identify their types in the steel, they were examined by TEM after creep testing using both carbon replicas and thin foils. During tempering and creep, the prior austenitic grain boundaries were decorated by a discontinuous network of coarse particles (Fig. 8). These large precipitates were primarily chromium-rich M23C6; the presence of some iron- and tungsten-rich M6C was also detected. Further, after creep exposure many fine particles which decorated the bainite lath boundaries and were also precipitated within the matrix were found to be MX particles (Fig. 9). In most fine particles, EDX analyses revealed the simultaneous occurrence of vanadium and niobium. For the particle type determination (precipitated particles were extracted in carbon replicas) the chemical composition evaluated from X-ray spectra (EDAX software for thin specimens without corrections for absorption or fluorescence) and selected area diffraction (SAD) were employed. The experimental results of particle analyses together with the phase equilibrium thermodynamic calculations are given in Tables 2 and 3. The main difference between the theoretical and experimental results is the experimental evidence of the existence of the M23C6 carbide at 600 and 650 °C which was not predicted by the calculation. On the other hand, the more detailed modelling proved that small increases in C or Cr content will lead to the appearance of M23C6 carbides in the calculations. The discrepancy is probably caused by the fact, that the thermodynamic database was built upon the assessments of important Fe-based

Fig. 7. TEM micrographs of T23 steel in a normalized and tempered state.

Fig. 8. TEM micrographs of T23 steel after creep at 500 °C and 250 MPa (time to fracture 6500 h).

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Fig. 9. TEM micrographs of T23 steel after creep at 650 °C and 75 MPa (time to fracture 3632 h).

binary and ternary systems, using experimental data obtained at higher temperatures, where the structures are closer to the thermodynamic equilibrium. Consequent extrapolations to lower temperature can introduce this inaccuracy. It has to be noted that the decreasing number of M23C6 particles and their increasing size could lead to lower creep strength. Anyway, the agreement between the experimental and calculated chemical compositions of the precipitates under investigation seems to be satisfactory. Typical fracture surfaces after T23 steel creep are shown in Fig. 10. The SEM fractographs indicate that the creep fractures are of the ductile dimple fracture mode. 4. Discussion In principle, there may be two basic processes influencing creep behaviour and the creep-strength degradation in P23 steel [10]. One is oxidation causing the reduction in the cross section of the loading component and thus the reduction in creep life. The creep rupture behaviour and the effect of oxidation on the creep strength of T23 steel have been investigated by Sawada et al. [9]. They found that the detrimental effect of oxidation is not due to an intrinsic material property. The other possible degradation process occurs due to the acting deformation processes and microstructural changes. As mentioned earlier, in this work all the creep testing of T23 steel was carried out in a purified argon atmosphere to avoid oxidation of the creep specimens making it possible to omit this process from the following consideration. Thus, further analysis of the creep behaviour will focus on the creep deformation mechanism(s) and microstructural changes ignoring the oxidation process. Table 2 Experimentally determined and calculated chemical composition of the phases in T23 steel at 600 °C and 650 °C (in wt.% calculated composition normalized to 100% of metallic elements.). Phase

Element

V

Cr

Mn

Fe

Nb

Mo

W










Experimental M6C MX (V, W) MX (Nb) M23C6

600 °C 2.2 2.7 36.1 5.0 12.8 1.5 4.8 57.1

0.4 0.5 0.8 2.2

16.5 1.4 1.5 31.3

2.2 14.4 69.2 1.1

4.5 3.2 11.3 1.2

71.5 39.4 2.9 2.3

Calculated M6C MX

600 °C 1.7 74.3

0.0 0.0

25.3 0.0

0.0 19.6

2.1 1.0

70.2 3.6

0.7 1.6

In Fig. 11 the dependences of the minimum creep rates normalized to the coefficient of lattice diffusion DL(m2 s−1) = 1.8 × 10−4 exp(−237/ RT) [20] are plotted against the applied stress normalized to the shear modulus G (MPa) = 97,400 − 39T [20]. Considering a natural scatter of the creep data it can be seen that nearly all the experimental points could fit to a single curve independently of which of the laws describing the stress dependence of the minimum creep rate is chosen. This implies that the creep rate is controlled by lattice diffusion in the region of temperatures and applied stresses under consideration [21,22]. However, such a conclusion would imply that the temperature dependence of the stress exponent n could be ignored. From Fig. 12, in which the stress exponent of the minimum creep rate n is plotted against stress, it can be seen that n is a linear function of stress and that the exponent n is more or less independent of temperature (however, this would be hardly justifiable for a temperature of 550 °C at higher stresses). Sawada et al. [10] reported that the stress exponents of the minimum creep rate n at 625 and 650 °C were 7.8–13 for higher stresses and 3.9–5.3 for lower stresses in the P23/T23 steels. These values are in very good agreement with the values of the exponent n presented in Fig. 12. The highest values of n in Fig. 12 at higher stresses could correspond to the power-low breakdown region. Note that, at higher normalized stresses, ε m =DL increases with σ/G considerably quickly. It was found that the activation energy for creep QC is stress dependent. A similar trend and values of QC were reported by Sawada et al. [10] for P23/T23 steels at temperatures 500–650 °C. The activation energy for the lattice self-diffusion of ferritic steels is about 350 kJ/mol [23]. The proximity of this value with the experimentally determined values of QC once again suggests that the creep rate at higher stresses is latticediffusion controlled and the controlling creep deformation mechanism is the climbing of mobile dislocations. A downward tendency of QC at lower stresses may be explained by the increasing contribution of grain-boundary diffusion-controlled processes like grain boundary sliding. Experimentally determined values of the time to fracture tf and strain to fracture εf correlate well with the minimum creep rates ε m

Experimental M6C MX

650 °C 1.4 40.0

3.1 4.7

0.4 0.8

17.3 1.4

1.8 11.2

4.9 4.1

71.1 37.8

Calculated M6C MX

650 °C 1.4 72.2

0.9 2.7

0.0 0.0

26.0 0.0

0.0 18.2

1.9 1.0

69.8 6.0

μ

through the Monkman–Grant relationship ðε m Þ . tf = const M [24]; tf

Table 3 Calculated equilibrium and experimentally determined chemical composition of carbides in T23 after creep at 650 °C for 3632 h. Content of metallic elements normalized to 100%. Carbide

Element

Cr

M6C

Calculated Experimental Calculated Experimental Experimental Calculated Experimental

0.9 1.9 3 4 2.6 0.9 6 3 1 2 Not predicted 28 2

MC

M23C6

Mo

V

W

Nb

Fe

Mn

1.5 2 72.2 35 6

69.8 72 5.7 45 12

0 2 18.6 8 76

25.9 17 0 2 2

0 0 0 0 1

2

21

1

46

1

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K. Kucharova et al. / Materials Characterization 109 (2015) 1–8

Fig. 10. SEM micrographs of creep fracture of T23 steel with creep at 500 °C and 250 MPa (time to fracture is 6500 h).




is approximately inversely proportional to εm (Fig. 13). For μ = 1 the Monkman–Grant relationship predicts the temperature and stress dependence of the time to fracture to be inverse to that of the minimum creep rate, unless the strain to fracture εf is temperature and stress dependent. It can be assumed that the Monkman–Grant constant M corresponds to creep ductility εf. The validity of the Monkman–Grant relationship was previously confirmed by Sawada et al. [9] for P23 steel subjected to creep testing at 625 and 650 °C. These authors found that the constant M should be higher in the short-term tests since the creep ductility is usually high. By contrast, the constant M would be lower in the long-term tests due to the low creep ductility. Anyway, although the productεm  t f could vary with stress and temper





ature, the approximately inverse dependence of tf on ε m means that the creep life is determined by the overall creep rate, i.e., the fracture is strain controlled. Since the minimum creep strain rate is obviously lattice-diffusion controlled, most probably, the recovery involving dislocation climb and annihilation of mobile dislocations account for the creep behaviour of the T23 steel. The same applies to dislocation climb past the secondary (carbide) particles, provided the dislocations do not exhibit an attractive interaction with these particles [25]. However, the particles

Fig. 11. The normalized minimum creep rate plotted against the normalized applied stress.

identified in this investigation are certainly not coherent with the matrix. Hence, an attractive interaction of dislocations with these particles may be expected [26]. Then, the dislocation climb around a particle can be relatively easy as compared with the detachment of the dislocation from the particle after the climb around it is finished. The estimation of the shear stress, necessary to detach dislocations from interacting particles, needs the results of quantitative microstructural analysis of the secondary phase particles in P/T23 steel. However, at present, such an analysis has not been completed. The strengthening mechanism in the T23 steel is the combination of solid solution hardening due to W and precipitation hardening mostly due to fine (V,Nb)C particles in the fully bainitic matrix. The substitution of W for Mo also constrains the precipitation of M6C during creep which retards the evolution of the bainitic microstructure [27]. The precipitation of W-rich M6C carbides lowers W in solution (see Tables 2 and 3), which is the key factor for the solid solution strengthening in this steel. It was found [1] that the MC carbides in W steel kept coherence even after long-term exposure. Moreover, thermodynamic calculations show that the equilibrium volume fraction of MC carbides slightly increases with increasing temperature up to 650 °C, see Table 4. Therefore, the population of MC precipitates should be more resistant to the overall particle number decrease due to particle dissolution during the coarsening process taking place at ageing and/or in creep. This fact should contribute to a strong increase of strength of P/T23 steel and strength and ductility stability during long-term ageing compared to T22 as was reported by Bendick et al. [4]. This also suggests that W is superior to Mo in terms of precipitation hardening. A further way to

Fig. 12. Stress dependences of the stress minimum creep rate exponent n.

K. Kucharova et al. / Materials Characterization 109 (2015) 1–8

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of 2.5Cr–1Mo steel does not degrade at all when tested in a vacuum or argon atmosphere but does degrade more with decreasing specimen diameter when tested in air. Endo et al. [31] also confirmed that in 2.25Cr– 1Mo steel the effect of oxidation on creep life is much stronger than that of structural degradation on long-term creep life. Nevertheless, the microstructure of conventional annealed 2.25Cr–1Mo steel is ferritepearlite, while the normalized and tempered P23/T23 steels have a bainitic structure. Sawada et al. [9] reported that the creep rupture life of T23 steel at 625 °C was almost the same both in helium and air, indicating that factors other than oxidation are responsible for the reduction in long-term creep strength. Based on the present results, it can be concluded that although significant oxidation of the T23 steel specimens may reduce creep rupture life to some extent, a large drop in longterm creep rupture strength is mainly caused by degradation due to microstructural change during creep exposure [9,32].

5. Conclusions

Fig. 13. Relationship between the minimum creep rate and the time to fracture.

improve the creep rupture strength of T23 steel is to increase the overall vanadium content to the stoichiometric ratio V:C. Foldyna et al. [28] reported the effect of the V:C ratio on the time to creep fracture in lowalloy Cr steels. The theoretical modelling has shown that, in model alloys with 0.5%V, the amount of vanadium bounded in VC particles is increased. Due to the higher amount of VC particles, the inter-particle spacing decreases and the creep rupture strength increases. Foldyna et al. [28] suggested that the overall V content should be increased to the stoichiometric ratio V:C ~ 4 to 5 to increase the effect of precipitation strengthening in low-alloy Cr steels. The evolution of the microstructure in P23 steel after long-term high-temperature exposure has experimentally been assessed by Mariani et al. [3]. For investigation, two selected samples were used; the first was exposed at 550 °C for around 20,000 h and the second was exposed at 600 °C for about 52,000 h. Both aged microstructures consisted of tempered bainite and the grain boundaries were decorated mainly by M6C precipitates. TEM and XRD investigations of the aged samples show that M7C3 and M23C6 precipitates are not thermodynamically stable in the temperature range of creep service and that a new population of fine M6C particles precipitates with increasing exposure duration, while the M6C particles formed during the initial tempering tend to coarsen with increasing exposure time and/or temperature. The existing MC particles are stable against coarsening in both the analysed conditions. Finally, it should be mentioned that Igarashi et al. [11] studied the long-term creep properties of T23/P23 steels between 500 and 600 °C. They reassessed the creep results using the oxidation correction method proposed by Nakashiro et al. [29] and concluded that in P23/T23 steels no creep-strength degradation occurs, apart from that due to oxidation at high temperatures. Viswanathan et al. [30] studied creep in 2.25Cr– 1Mo steel using various sizes of specimens and atmospheres of air and argon, and in a vacuum. They concluded that the creep rupture strength

Table 4 Calculated equilibrium volume fraction of phases for different temperatures in T23 steel. Temperature (°C)

Ferrite (%)

M6C (%)

MC (%)

500 550 600 625 650

98.62 98.66 98.72 98.75 98.79

1.05 0.99 0.92 0.87 0.82

0.33 0.35 0.37 0.38 0.39

The isothermal constant load tensile creep tests of T23 steel were carried out in an argon atmosphere at five testing temperatures ranging from 500 to 650 °C and at various applied stress levels from 50 to 400 MPa. The creep data were analysed and it was found that the minimum creep rate is obviously lattice-diffusion controlled. The experimental results of the microstructural analysis of the particles of the secondary phases, which are responsible for precipitation hardening in the steel under investigation, were in good agreement with the thermodynamic calculations using the modified software package ThermoCalc. The significant creep-strength drop of T23 steel after long-term creep exposure can be explained by the decrease in dislocation hardening, precipitation hardening and solid solution hardening due to the instability of the microstructure at high temperature. It is suggested that, under restricted oxidation due to argon atmosphere, the microstructure instability is the main detrimental process in the degradation of the creep rupture strength of T23 steel under investigation. Acknowledgments The authors wish to thank Dr. Ales Kroupa of the Institute of Physics of Materials, Academy of Sciences of the Czech Republic for the thermodynamic calculations of carbide particles using the modified software package Thermo-Calc. Financial support for this work was provided by the Technology Agency of the Czech Republic under Project No. TA02010260. These experiments were realized at CEITEC — Central European Institute of Technology with the research infrastructure supported by Project CZ.1.05/1.1.00/02.0068 financed by the European Regional Development Fund.

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