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Influence of Thermal Ageing on Microstructure and Tensile Properties of P92 Steel

  • T. Sakthivel EMAIL logo , S. Panneer Selvi , P. Parameswaran and K. Laha
Published/Copyright: September 5, 2017
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High Temperature Materials and Processes
From the journal High Temperature Materials and Processes Volume 37 Issue 5

Abstract

Microstructure and tensile properties of P92 steel in the normalized and tempered, and thermal aged at 923 K for 5000 h and 10,000 h conditions have been investigated. Laves phase precipitate was observed in the thermal-aged steels. The size of Laves phase precipitate increased with increase in thermal exposure. This was also confirmed from the observation that the area fraction of Laves phase precipitate was higher in the 5000 h aged condition which decreased with further increase in thermal exposure. On the other hand, the size and area fraction of M23C6 precipitate were found increased in the 5000 h aged steel, further continued to enhanced precipitation of fine M23C6 in the 10,000 h aged steel. This resulted in significant increase in area fraction and comparable size with the steel aged for 5000 h. Hardness of the steel was decreased with increase in the duration of ageing. Thermal-aged steels exhibited lower yield stress, ultimate tensile strength and relatively higher ductility in comparison with steel in the normalized and tempered condition. The increase in lath width and recovery of dislocation structure under thermal-aged condition resulted in lower tensile strength and hardness. An extensive Laves phase formation and coarsening by loss of tungsten in the matrix led to decrease in the tensile strength predominantly in the 5000 h aged steel. The tensile strength of 10,000 h aged steel was comparable with that of 5000 h aged steel due to enhanced precipitation of fine M23C6 in the steel due to enhanced mobility of carbon in the absence of tungsten in the matrix.

Keywords: P92 steel; thermal ageing; microstructure; tensile strengths

Introduction

ASME P91 heat-resistant steel has widely been used in the fossil-fired and nuclear power plants. The increase in environmental concern and commercial demands on energy production have led to an emphasis on the development of steels operable at supercritical conditions [1, 2]. The addition of 1.8–2.0 wt.% tungsten and 0.002–0.005 wt.% boron with reduction of molybdenum to around 0.5 wt.% in P91 steel led to the development of P92 (9Cr-1.8W-0.5Mo-VNb) steel [2, 3, 4, 5, 6, 7, 8, 9, 10, 11, 12]. The 9Cr ferritic-martensitic steel has been considered for the liquid metal cooled fast reactors [13], which offers about 30 % higher creep strength over the P91 steel at higher temperature [4]. P92 steel is a promising material for the next generation nuclear and fossil power plants due to its excellent mechanical properties at high temperature, high thermal conductivity, low thermal expansion coefficient, good weldability, adequate corrosion and stress-corrosion cracking resistance, excellent resistance to irradiation in comparison with austenitic steels. The steel derives its high temperature strength from tempered martensite lath structure, stabilization of martensitic lath structure by M23C6 type of carbide, martensite phase transformation induced high dislocation density, intra-lath MX (M-V, Nb; X-C, N) type of carbide and nitride, and solid solution strengthening from tungsten [10, 11, 12, 13, 14, 15, 16, 17, 18, 19, 20, 21]. P92 steel is considered as the key material for thick section components in the ultra supercritical boilers at temperature up to 898 K [22]. In the service conditions at high temperature, themicrostructural evolution of the steel occurs such as recovery of martensitic lath structure, coarsening of precipitate and formation of newer precipitate which is likely to affect the mechanical properties. Vyrostkova et al. [4] has observed the formation of various phases during ageing at 898 K up to 9000 h. The precipitation and coarsening behaviour of Laves phase under thermal ageing of P92 steel at 923 K for 8000 h were reported by Xue Wang et al. [22]. The Laves phase precipitate size above 0.5 µm is considered as deleterious to mechanical properties [4]. The static mechanical properties of P92 steel aged at 923 K for different duration up to 5000 h were studied at 300 K by Xiaofeng Guo et al. [21]. On this pretext, it is essential to study in detail on the evolution of microstructure of the steel subjected to thermal ageing under longer duration and its influence on tensile properties.

In the present investigation, tensile properties of P92 steel in the normalized and tempered, and thermal aged at 923 K for 5000 h and 10,000 h conditions were studied. Microstructural examinations were performed to understand the distinct tensile behaviour of the steel in the three different microstructural state, namely- normalized and tempered, and thermal aged conditions.

Experimental details

The chemical composition of P92 steel is given in Table 1. The P92 steel plate of 12 mm thickness was normalized at 1323 K for 30 minutes and tempered at 1053 K for 2 h followed by air cooling. The steel blanks having dimensions of 12 mm (T) x 12 mm (W) x 120 mm (L) were used for thermal ageing at 923 K for 5000 h and 10,000 h in air. Button head cylindrical tensile specimen having 4.0 mm gauge diameter and 28.6 mm gauge length was used to carry out the tensile test. The tensile test was performed at various test temperature ranging from 300 to 923 K and at a nominal strain rate of 3 x 10‒4 s‒1 in air as per the ASTM E21. Prior to tensile testing, specimen was heated to test temperature using electrical resistance split type furnace having three heating zone. Separate temperature controller was used to maintain constant temperature across the different zones of the furnace. K-type thermocouple was used to monitor the temperature of the testing specimen. The specimen was held for 15 minutes at test temperature to ensure uniform temperature along the specimen within ±2 K. Single specimen was tested for each condition of the test. However, tensile test at 523 K and 573 K were repeated in order to confirm the variation in the strength values in the normalized and tempered steel. Vickers hardness measurement on normalized and tempered and thermal-aged steels have been carried out using 10 Kp load and with a 15 seconds dwell time. Optical, scanning electron (SEM) and transmission electron microscopic (TEM) investigations were carried out on the normalized and tempered, and thermal-aged steels. The chemical composition of M23C6 and Laves phase precipitate were obtained from Energy dispersive spectroscopy (EDS) attached with SEM/TEM. Immersion etching for 15 seconds using Villella’s reagent (picric acid 5 g, HCl acid 25 ml, ethyl alcohol 500 ml) has been employed to reveal the microstructure. Specimen for TEM investigation was thinned down to 50 µm by mechanical polishing under flowing water followed by electrolytic double jet thinning using 20 % perchloric acid and ethanol solution, at 238 K with an applied voltage of 20 V. The size and area fraction of the precipitate have been measured using SEM secondary electron (SE) and backscattered electron (BSE) images by image analysis (Image J) software. Analysis of Laves phase precipitation using field emission gun (FEG) SEM image has been reported in the literature to avoid the discrepancy due to the insufficient number of particle per TEM image [10].

Table 1:

Chemical composition (wt. %) of P92 steel.

ElementsCCrWMoMnSiVNNb
P92 steel0.109.21.90.510.360.270.220.0500.07
ElementsSPNiAlBTiN/AlFe
P92 steel0.0020.0100.0600.0100.0010.0025Bal

Results and discussion

Microstructure

The SEM micrographs of the steel in the normalized and tempered (NT), and thermal aged (TA) for 5000 h and 10,000 h conditions are shown in Figure 1. The grain and sub-grain boundaries of the NT steel were decorated by M23C6 precipitates. MX precipitate was observed in the intra-lath region of the steel (Figure 1(g)). TEM micrographs of the steel in the NT, and thermal aged for 5000 h and 10,000 h conditions are shown in Figure 2. The martensitic lath width of the steel has increased with increase in duration of thermal ageing. The growth rate of lath width was significantly higher from NT to 5000 h TA condition, and relatively lower growth rate of lath width was observed on further increase in the ageing duration from 5000 h to 10,000 h (Figures 2 and 3). Laves phase precipitate was not observed in the NT steel. Precipitation of Laves phase was observed in the TA steels. EDS spectrum of M23C6 precipitate observed in the NT steel and Laves phase precipitate in the TA steel is shown in Figure 4. Tungsten content about 7–9 wt. % has been observed in the M23C6 precipitate. The presence of tungsten in M23C6 precipitate led to delay in the coarsening rate of precipitate under service conditions.

Figure 1: SEM-SE/BSE micrographs of P92 steel in the normalized and tempered (a) SEM-SE and (b) SEM-BSE mode, and thermal aged at 923 K for 5000 h (c) SEM-SE and (d) SEM-BSE mode, and 10,000 h (e) SEM-SE and (f) SEM-BSE mode, and (g) FEG-SEM micrograph of the normalized and tempered P92 steel.
Figure 1:

SEM-SE/BSE micrographs of P92 steel in the normalized and tempered (a) SEM-SE and (b) SEM-BSE mode, and thermal aged at 923 K for 5000 h (c) SEM-SE and (d) SEM-BSE mode, and 10,000 h (e) SEM-SE and (f) SEM-BSE mode, and (g) FEG-SEM micrograph of the normalized and tempered P92 steel.

Figure 2: TEM micrographs of P92 steel in the (a) normalized and tempered, and thermal-aged condition at 923 K for (b) 5000 h and ((c) and (d)) 10,000 h.
Figure 2:

TEM micrographs of P92 steel in the (a) normalized and tempered, and thermal-aged condition at 923 K for (b) 5000 h and ((c) and (d)) 10,000 h.

Figure 3: Variations of lath width of P92 steel in the normalized and tempered, and thermal-aged conditions at 923 K for 5000 h and 10,000 h.
Figure 3:

Variations of lath width of P92 steel in the normalized and tempered, and thermal-aged conditions at 923 K for 5000 h and 10,000 h.

Figure 4: EDS spectra obtained from the (a) M23C6 precipitates in the normalized and tempered steel, and (b) Laves phase (Fe2W) precipitates in the thermal aged at 923 K for 5000 h steel.
Figure 4:

EDS spectra obtained from the (a) M23C6 precipitates in the normalized and tempered steel, and (b) Laves phase (Fe2W) precipitates in the thermal aged at 923 K for 5000 h steel.

SEM micrographs of the steels observed under SE and BSE modes are given in Figure 1. The increase in precipitate size has been noticed under TA condition (Figure 1 (a), (c) and (e)). Laves phase has been observed as bright phase in BSE mode SEM image. Laves phase appears brighter due to its higher average atomic number [10, 20]. The Laves phase precipitates predominantly nucleated on the prior-austenite and subgrain boundaries in the steel, and coarsen towards the sub-grain interior. The extensive formation of Laves phase precipitate was found in TA steels (Figure 1 (d) and (f)). Agglomeration of Laves phase precipitate at high angle boundaries of the steel were reported in the literature [23]. The lath boundaries in the steel were not clearly evident at the location where the chunky Laves phase precipitate was present (Figure 2(b) and (c)). Laves phase precipitate having small rod shape was noticed at the lath boundaries (Figure 2). The Laves phase precipitation along the boundaries and adjacent to the M23C6 precipitate due to segregation of Si and P has been reported by Isik et al. [14]. The formation of Laves phase precipitate adjacent to M23C6 precipitate led to consumption of the M23C6 precipitate which resulted in segregation of Cr and C in the boundaries and around the Laves phase. The tungsten content in the Laves phase precipitate was about 36 wt.% and 42 wt.% in the 5000 h and 10,000 h TA steels respectively. The formation of Laves phase led to decrease in the tungsten content of the matrix, and decreases the solid solution strengthening contribution by tungsten to the steel. The loss of solid solution strengthening effect due to tungsten in the matrix is predominantly seen up to 5000 h ageing than in the case of 10,000 h aged steel. Tungsten in P92 steel does not only act as the solute solution strengthener but also impedes the carbon diffusion and hence retards the further precipitation and coarsening of M23C6 precipitates. Hence, the presence of tungsten in P92 steel led to stabilize the microstructure including matrix and M23C6. The fine Fe2Nb precipitate (V:5.54 wt.%, Cr:6.11 wt.%, Fe:47.99 wt.%, Nb:40.35 wt.%) was also noticed in the 10,000 h aged steel.

The variation of precipitate size and area fraction with different conditions of the steel are shown in Figure 5. The Laves phase having average size of 0.160 µm in the 5000 h and 0.600 µm in the 10,000 h thermal-aged steels were observed. The significant increase in the size of the Laves phase has been observed in the steel aged for 10,000 h. However, the area fraction of the Laves phase decreased predominantly in the 10,000 h aged steel (Figure 5(b)). Laves phase having maximum of 1 µm size has been reported in the creep exposed steel for 16,000 h [24]. Laves phase precipitate having round or irregular shape was observed in the prior-austenite, packet and block boundaries in the steel. The growth of the Laves phase precipitate is controlled by the grain boundary or dislocation sub-structure diffusion of tungsten in the steel [23, 24, 25]. The extensive formation and coarsening of Laves phase led to destabilize the dislocation structure and lath boundaries by formation of dislocation cells and lower the pinning pressure on the lath boundary due to its coarser size. It is significant to note that the average size of the M23C6 precipitate increased from 0.060 µm in the NT steel to 0.187 µm in the 5000 h aged steel (Figure 5(a)). M23C6 precipitates with an average size about 0.2 µm has been observed in the 10,000 h aged steel. The formation of more number of fine M23C6 precipitate was significantly increased in the 10,000 h aged steel, this results in comparable average size of the precipitate with 5000 h aged steel (Figures 2 and 5). In addition, the area fraction of the M23C6 precipitate increased noticeably in the 10,000 h aged steel (Figure 5(b)). The increased precipitation of M23C6 precipitate in the 10,000 h aged steel resulted in the loss of solution strengthening contribution from carbon in the steel.

Figure 5: Variations of (a) precipitate size and (b) area fraction (%) of precipitates in the steel in normalized and tempered condition, thermal aged at 923 K for 5000 h and 10,000 h steels.
Figure 5:

Variations of (a) precipitate size and (b) area fraction (%) of precipitates in the steel in normalized and tempered condition, thermal aged at 923 K for 5000 h and 10,000 h steels.

Hardness

The variation of hardness with different conditions of the steel is shown in Figure 6. The decrease in hardness was noticed in the thermal-aged steels (220 HV10 in the 5000 h and 205 HV10 in the 10,000 h) in comparison with the steel in NT condition (247 HV10). The rate of decrement in hardness was relatively higher up to 5000 h TA steel than the 10,000 h TA steel. The increase in hardness of the steel in the initial stages of ageing at 873 K has been reported by Bumjoon Kim et al. [19]. The decrease in hardness in the 5000 h aged steel has been attributed to the extensive formation of Laves phase precipitate predominantly, coarsening of M23C6 precipitate, recovery of dislocation structure and coarsening of martensite lath width. The extensive precipitation of Laves phase in the 5000 h aged steel led to loss of tungsten in the matrix, the presence of tungsten in the matrix reduces the mobility of carbon atoms due to its higher atomic weight/mass. The loss of tungsten in the matrix resulted in enhanced mobility of carbon atoms in the steel which in turn resulted in extensive precipitation and further coarsening of M23C6 in the 10,000 h aged steel. This predominantly attributed to the lower hardness in the 10,000 h aged steel.

Figure 6: Variations of hardness with different conditions of P92 steel. The decrement shows different rate after 5000 h. Insert micrograph depicts the dominant precipitation which influences the hardness decrement.
Figure 6:

Variations of hardness with different conditions of P92 steel. The decrement shows different rate after 5000 h. Insert micrograph depicts the dominant precipitation which influences the hardness decrement.

Tensile properties

Tensile stress–strain curves of P92 steel in the NT, and TA for 5000 h and 10,000 h conditions at a strain rate of 3 x 10‒4 s‒1 at 300 K and 873 K are shown in Figure 7. The tensile strengths of the thermal-aged steels were significantly decreased in comparison with the steel in the NT condition. The minor increase in ductility was noticed in the thermal-aged steels. The steel aged for 10,000 h has shown comparable strengths and marginal increase in ductility with reference to the steel aged for 5000 h which tested at 300 K. The variation of yield stress (YS) at 0.2 % strain and ultimate tensile strength (UTS) with test temperature in the NT, and TA for 5000 h and 10,000 h conditions are shown in Figure 8. In all the conditions, the tensile strengths of the steel were found to decrease with increase in temperature. The decrease in strengths was relatively lower up to 473 K and it exhibits a plateau in the variation of strengths in the intermediate temperature regime (473 K to 723 K). The rapid decrease in strengths was observed with further increase in temperature. The lower extent of variation in strengths in the intermediate temperatures regime has been attributed to the dynamic strain ageing (DSA) behaviour. DSA occurs due to the interaction of mobile dislocations with solute atoms. This is manifested as the occurrence of serrations, plateau in the variation of properties, trough in the variation of ductility, and negative strain rate sensitivity of strengths [26]. The serrated flow has clearly been observed at the intermediate temperatures and in various microstructures of 9Cr ferritic steel [27]. However, in the present steel the appearance of serrated flow has been noticed at 573 K in the NT steel [28]. Serrated flow was not observed in the TA steels. Even though no extensive serrated flow in the tensile curve was observed in the intermediate temperature region which is generally a profound manifestation of DSA in this investigation, the lower extend of variations in the strengths and ductility of the steel with intermediate temperature supports the inference. The occurrence of DSA in the intermediate temperature regime causes an increased dislocation multiplication, uniform distribution of dislocation, and delay in recovery of dislocation structure rather than well-defined cell structures in many metals and alloys [28]. However, the dominance of DSA behaviour in the intermediate temperature was relatively reduced in the thermal-aged steels than the steel in the normalized and tempered condition might be due to decrease in dislocation density below the required critical dislocation density for the occurrence of noticeable DSA behaviour in the steel. The rapid decrease in strengths at high temperature has been attributed to dynamic recovery process which might have enabled by thermally activated process of climb of dislocations and cross slip. In the present steel, the recovery of dislocation lath structure and subgrain formation in the NT steel tensile tested at 873 K have been confirmed from the TEM investigation [28]. This clearly indicates that the rapid decrease in strengths and increase in ductility at high temperatures have been attributed to dominance of dynamic recovery.

Figure 7: Tensile stress–strain curves of P92 steel in the normalized and tempered, and thermal aged at 923 K conditions.
Figure 7:

Tensile stress–strain curves of P92 steel in the normalized and tempered, and thermal aged at 923 K conditions.

Figure 8: The variation of (a) yield stress and (b) ultimate tensile strength with test temperature of P92 steel in the normalized and tempered, and thermal-aged conditions, (c) YS and UTS with ageing time.
Figure 8:

The variation of (a) yield stress and (b) ultimate tensile strength with test temperature of P92 steel in the normalized and tempered, and thermal-aged conditions, (c) YS and UTS with ageing time.

The YS and UTS of TA steels significantly decreased in comparison with steel in the NT condition (Figure 8(c)). The rapid decrease in strengths in the 5000 h TA steel and lower rate of decrease in strengths in the 10,000 h aged steel have been observed. The extensive formation of Laves phase precipitate, recovery of dislocation structure, increase in martensite lath width and coarsening of M23C6 precipitate have been attributed to decrease in tensile strengths of 5000 h TA steel. The tensile strengths of 10,000 h TA steel has been attributed to extensive formation of fine M23C6 precipitate and recovery of martensitic lath structure. The presence of fine M23C6 precipitate might have effectively pinned the boundaries and dislocation structure in the 10,000 h TA steel resulted in comparable strengths with 5000 h TA steel (Figure 2(c) and (d), and Figure 8).

The uniform plastic strain (%), elongation (%) and reduction in area (%) of the steel in the NT, and thermal aged for 5000 h and 10,000 h conditions are shown in Figure 9. The uniform plastic strain was measured from the strain at the yield point at 0.2 % strain to the strain up to maximum stress (UTS) or before the initiation of necking in the specimen. The uniform plastic strain decreased with increase in temperature. Higher uniform plastic strain was observed in the thermal-aged steels in comparison with the steel in the NT condition. The uniform plastic strain has increased with increase in the duration of thermal ageing. The increase in elongation (%) was observed in the TA steels than the NT steel. The elongation (%) increases with increase in TA duration (Figure 9(b)). The reduction in area (%) increased with increase in temperature with a lesser extent of variation in the intermediate temperatures regime (473 K to 723 K). The lower reduction in area (RA) was observed in the NT condition as compared to the steel in the TA condition. The increase in ductility in TA condition has been attributed to softening of the matrix by recovery of dislocation structure, and the formation of Laves phase and M23C6 precipitate (Figure 2), which led to decrease in the tungsten and in the carbon content in the matrix. Tungsten in the matrix is responsible for the increase in the solid solution strengthening and impedes the carbon diffusion which retards the further precipitation and coarsening of M23C6 at the expense of ductility [4, 20, 21].

Figure 9: The variation of (a) uniform plastic strain, (b) elongation, and (c) reduction in area with test temperatures of P92 steel in the normalized and tempered, and thermal-aged conditions.
Figure 9:

The variation of (a) uniform plastic strain, (b) elongation, and (c) reduction in area with test temperatures of P92 steel in the normalized and tempered, and thermal-aged conditions.

Fractography analysis

Fracture surface of the tensile tested at 300 K and 923 K specimens in the NT, and TA conditions are shown in Figures 10 and 11. The chisel tip appearance in the radial direction of the specimen was observed in the steels tested at 300 K (Figure 10). The length of the chisel tip feature was higher in the NT steel in comparison with the TA steels. The presence of silicon and phosphorous in the carbide/matrix interfaces led to promote the chisel tip fracture in the steel [29, 30]. The tearing of lath structure was observed in all conditions of the steel, and more evidently in the NT steel. The segregation of Si and P around the carbide precipitates and various boundaries in the NT steel led to higher chisel tip fracture appearance. The presence of Si and P in the solid solution of Laves phase and reduced dislocation density led to lower chisel tip fracture appearance in the thermal-aged steels. Ductile mode of failure has been observed in both the conditions of the steels (Figure 11). The predominantly dimple mode of failure with bigger cell size (Figures 10 and 11) has been observed in the TA condition than the steel in NT condition. The steels tested at 923 K exhibited the cup and cone fracture in both the steels.

Figure 10: Fracture surface of the tensile tested P92 steel at 300 K in ((a) and (b)) normalized and tempered, and thermal aged ((c) and (d)) for 5000 h, and ((e) and (f)) for 10,000 h steels. Arrow mark indicates the chisel tip feature in the fracture surface which determines the extent of brittleness.
Figure 10:

Fracture surface of the tensile tested P92 steel at 300 K in ((a) and (b)) normalized and tempered, and thermal aged ((c) and (d)) for 5000 h, and ((e) and (f)) for 10,000 h steels. Arrow mark indicates the chisel tip feature in the fracture surface which determines the extent of brittleness.

Figure 11: Fracture surface of the tensile tested P92 steel at 923 K in ((a) and (b)) normalized and tempered, and thermal aged ((c) and (d)) for 5000 h, and ((e) and (f)) for 10,000 h steels.
Figure 11:

Fracture surface of the tensile tested P92 steel at 923 K in ((a) and (b)) normalized and tempered, and thermal aged ((c) and (d)) for 5000 h, and ((e) and (f)) for 10,000 h steels.

The decrease in hardness, tensile strengths and increase in tensile ductility in thermal-aged P92 steels in comparison with normalized and tempered steel have been attributed to (i) The decrease in solid solution strengthening by removal of tungsten and carbon from matrix by the formation and coarsening of Laves phase and M23C6 precipitates respectively, (ii) decrease in dislocation density by sub-grain formation and (iii) coarsening of martensitic lath. The formation and coarsening of Laves phase in the 5000 h and precipitation of fine M23C6 in the 10,000 h TA steels influenced the tensile strengths predominantly.

Conclusions

Microstructure and tensile properties of P92 steel have been studied under normalized and tempered, and thermal aged at 923 K for 5000 h and 10,000 h conditions.

  1. Laves phase precipitate was observed in the thermal-aged steels. Size of the Laves phase precipitate increased with increase in thermal exposure.

  2. Size of the M23C6 precipitate has been found to increase in the thermal-aged steels as compared to the steel in the normalized and tempered condition. However, the size of the M23C6 precipitate in the steel with thermal exposure (5000 h and 10,000 h) has not changed significantly due to enhanced fine M23C6 precipitate in the 10,000 h aged steel. This is confirmed by the significant increase in area fraction of the M23C6 precipitate in the 10,000 h aged steel.

  3. Hardness of the P92 steel decreased with increase in thermal exposure. Hardness decrement is found more dominant in 5000 h aged steel than the 10,000 h aged steel.

  4. The decrease in hardness of the 5000 h aged steel is due to (i) loss of solid solution strengthener tungsten in the matrix, (ii) increase in lath width and (iii) recovery of dislocation structure of the steel. The loss of carbon in the matrix due to (i) enhanced precipitation of M23C6 precipitate, (ii) increase in lath width and (iii) recovery of dislocation structure have been attributed to decrease in hardness in the 10,000 h aged steel.

  5. The 5000 h aged steel exhibited lower tensile strength in comparison with the normalized and tempered steel due to the loss of solid solution strengthener tungsten from the matrix by formation of Laves phase predominantly.

  6. However, the tensile strength of 10,000 h aged steel was comparable with 5000 h aged steel due to the presence of predominant fine and fresh M23C6 precipitate in the steel which pinned the laths, although the loss of carbon from the matrix which soften the steel.

  7. In conclusion, the competitive nature of precipitation of Laves and M23C6 in the long periods of ageing has a strong role to play in determining the strength.

Acknowledgements

The authors wish to thank Dr. Arun Kumar Bhaduri, Director, Indira Gandhi Centre for Atomic Research, and Dr. G. Amarendra, Director, Metallurgy and Materials Group, for their keen interest in the work and encouragement.

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Received: 2016-2-25
Accepted: 2017-4-22
Published Online: 2017-9-5
Published in Print: 2018-4-25

© 2018 Walter de Gruyter GmbH, Berlin/Boston

This article is distributed under the terms of the Creative Commons Attribution Non-Commercial License, which permits unrestricted non-commercial use, distribution, and reproduction in any medium, provided the original work is properly cited.

Articles in the same Issue

  1. Frontmatter
  2. Research Article
  3. Effect of Trailing Intensive Cooling on Residual Stress and Welding Distortion of Friction Stir Welded 2060 Al-Li Alloy
  4. Short Communication
  5. Study on the Growth Mechanism of K2Ti4O9 Crystal
  6. Research Articles
  7. Artificial Neural Network-Based Three-dimensional Continuous Response Relationship Construction of 3Cr20Ni10W2 Heat-Resisting Alloy and Its Application in Finite Element Simulation
  8. Influence of Thermal Ageing on Microstructure and Tensile Properties of P92 Steel
  9. A Novel Process for Joining Ti Alloy and Al Alloy using Two-Stage Sintering Powder Metallurgy
  10. Modeling and Finite Element Analysis for the Dynamic Recrystallization Behavior of Ti-5Al-5Mo-5V-3Cr-1Zr Near β Titanium Alloy During Hot Deformation
  11. Study on Dynamic Development of Three-dimensional Weld Pool Surface in Stationary GTAW
  12. Influence of Heat Treatment on Fracture Toughness and Wear Resistance of Nicral-Zro2 Multilayered Thermal Barrier Coating
  13. Kinetic Study on Phosphate Enrichment Behavior in CaO–SiO2–FeO–Fe2O3–P2O5 Steelmaking Slags
  14. Effect of Prestrain on Precipitation Behaviors of Ti-2.5Cu Alloy
  15. Study on Gamma Prime and Carbides of Alloy A286 by Traditional Thermodynamic Calculation
https://doi.org/10.1515/htmp-2016-0040
Keywords for this article
P92 steel; thermal ageing; microstructure; tensile strengths
Creative Commons
BY-NC-ND 3.0

Articles in the same Issue

  1. Frontmatter
  2. Research Article
  3. Effect of Trailing Intensive Cooling on Residual Stress and Welding Distortion of Friction Stir Welded 2060 Al-Li Alloy
  4. Short Communication
  5. Study on the Growth Mechanism of K2Ti4O9 Crystal
  6. Research Articles
  7. Artificial Neural Network-Based Three-dimensional Continuous Response Relationship Construction of 3Cr20Ni10W2 Heat-Resisting Alloy and Its Application in Finite Element Simulation
  8. Influence of Thermal Ageing on Microstructure and Tensile Properties of P92 Steel
  9. A Novel Process for Joining Ti Alloy and Al Alloy using Two-Stage Sintering Powder Metallurgy
  10. Modeling and Finite Element Analysis for the Dynamic Recrystallization Behavior of Ti-5Al-5Mo-5V-3Cr-1Zr Near β Titanium Alloy During Hot Deformation
  11. Study on Dynamic Development of Three-dimensional Weld Pool Surface in Stationary GTAW
  12. Influence of Heat Treatment on Fracture Toughness and Wear Resistance of Nicral-Zro2 Multilayered Thermal Barrier Coating
  13. Kinetic Study on Phosphate Enrichment Behavior in CaO–SiO2–FeO–Fe2O3–P2O5 Steelmaking Slags
  14. Effect of Prestrain on Precipitation Behaviors of Ti-2.5Cu Alloy
  15. Study on Gamma Prime and Carbides of Alloy A286 by Traditional Thermodynamic Calculation
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