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Effect of normalizing and tempering temperatures on microstructure and mechanical properties of P92 steel
International Journal of Pressure Vessels and Piping 132-133 (2015) 97e105
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International Journal of Pressure Vessels and Piping journal homepage: www.elsevier.com/locate/ijpvp
Effect of normalizing and tempering temperatures on microstructure and mechanical properties of P92 steel Dipika R. Barbadikar a, G.S. Deshmukh a, L. Maddi a, K. Laha b, P. Parameswaran c, A.R. Ballal a, D.R. Peshwe a, R.K. Paretkar a, M. Nandagopal b, M.D. Mathew b, * a b c
Department of Metallurgical and Materials Engineering, VNIT, Nagpur, India Mechanical Metallurgy Division, IGCAR, Kalpakkam, India Physical Metallurgy Group, IGCAR, Kalpakkam, India
a r t i c l e i n f o
a b s t r a c t
Article history: Received 5 March 2014 Received in revised form 26 June 2015 Accepted 1 July 2015 Available online 2 July 2015
In the present investigation, systematic studies on microstructure and mechanical properties of P92 steel subjected to various normalizing (1313e1353 K) and tempering (1013e1053 K) temperatures were carried out. The effect of heat treatment on microstructural parameters revealed an increase in grain size, lath width and decrease in the area fraction of the precipitates with an increase in normalizing temperature. The precipitate size has not changed significantly with increase in the normalizing temperature; rather it increased with increase in tempering temperature. Activation energy calculations confirmed the two fold mechanisms that dominate the tempering behavior. As a consequence, yield stress (YS) and ultimate tensile strength (UTS) were found to change with normalizing and tempering temperatures. P92 steel normalized at 1353 K and tempered at 1013 K was found to have the best combination of strength and ductility. © 2015 Elsevier Ltd. All rights reserved.
Keywords: P92 steel Normalizing Tempering Microstructure Mechanical properties
1. Introduction Ferritic/martensitic steel has been considered as a candidate material for power plant applications over austenitic stainless steel because of its excellent thermal conductivity, low coefficient of thermal expansion, good weldability accompanied with resistance to stress corrosion cracking and oxidation. They are widely used in the fabrication of high temperature components of fossil fired and steam generators of nuclear power plants. In order to increase the efficiency of the power plants operating at temperatures above 873 K and pressure of 250e300 bar [1], there is a need to develop materials in accordance with increased strength at particular service conditions. The development of 9Cr steels started few decades before, starting from P9, P91 and P92. P92 steel [2,3] is the next version of modified P91 steels, where Mo content is brought down to 0.5 wt.% from 1 wt.% and 1.8 wt.% tungsten is added. These steels derive their strength from tempered martensite lath structure which is stabilized by M23C6 type of carbide, intra-lath MX type carbide/nitride and martensite phase transformation induced high
* Corresponding author. E-mail address: [email protected] (M.D. Mathew). http://dx.doi.org/10.1016/j.ijpvp.2015.07.001 0308-0161/© 2015 Elsevier Ltd. All rights reserved.
dislocation density. P92 posses improved creep strength than P91 steel due to the solid solution strengthening and increased hardenability offered by tungsten addition. Enhancement in strength of P92 steel has been well explained by Ennis et al. [1] and has shown that the high degree of transient hardening in 9Cr steels is due to the presence of very high dislocation density obtained during normalization of the steel. The mechanical properties of these steels are found to be sensitive to normalizing and tempering temperatures as well as time, which alter the microstructural constituents [4]. The strengthening mechanisms of this steel have been discussed by many researchers. Kouichi et al. [5] have discussed the strengthening mechanism of creep resistant tempered martensitic steel. It has been reported that there is a large driving force for recovery due to the presence of high dislocation density and the pinning particles like M23C6, MX and solute atoms maintains the stable structure which is necessary for creep strength. The decrease in creep strength of 9Cr steels with service time has been attributed to decrease in dislocation density due to the formation of subgrain structure and to the coarsening of precipitates. Giroux et al. [6] have studied the softening behavior at high temperature in the uniaxial tensile test and found that the annihilation of subgrain boundaries and mobile dislocations leads to softening of the steel. The detailed
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Table 1 Chemical composition of P92 steel plate in wt. %.
Table 3 Hardness values of as-received and normalized conditions.
Element Wt. %
Cr 9.38
C 0.114
Mo 0.506
W 1.94
Nb 0.075
Mn 0.388
V 0.215
Element Wt. %
B 0.0018
N 0.0417
Al 0.0076
Cu 0.0162
P 0.0120
S 0.0035
Si 0.0239
TEM investigations of MN nitrides phases present in 9% Cr steel has been carried out by Yin Zhong Shen et al. [7] They have explained the role of precipitate in improving the creep strength of the steel. Investigation on precipitates present in P92 steel has been carried out by many researchers [8e11]. It has been reported that the ability of precipitates to act as dislocation barrier decreases as they coarsen. Thus finer precipitates lead to enhanced strength of the material. Ennis et al. [12] have studied the microstructural stability and creep strength of the P92 steel and have reported that the microstructural variations play a very important role in deciding the mechanical properties of the steel. However, the properties of steel deteriorate during service exposure. In particular, larger components with weldments have heat affected zones which are prone to type IV cracking, which depend on the initial microstructure of the material [13]. Further, weldments of the present steel consist of different regions of HAZ which are normalized for short times and are tempered. These zones develop heterogeneity in properties across the HAZ, which result in varied toughness and hence early creep rupture. Therefore achieving uniform properties have been a concern and it is as well a technological challenge to heat treat each individual zone of the HAZ at different tempering
Heat treatment
Hardness (Hv5)
Grain size (mm)
Initial heat treatment 1323 K/1053 K 1313 K 1333 K 1353 K
246 421 401 390
25 28 37 55
temperature. Considering this necessity, efforts are being made to systematically study the different combinations of normalizing and tempering temperatures which may provide an optimized combination of microstructure and mechanical properties. 2. Experimental details The chemical composition of the P92 steel analyzed using spectro-chemical analysis (ASTM-E-415-2008) is given in Table 1. This plate has been subjected to an initial heat treatment of normalizing and tempering at 1323 K for 30 min and 1053 K for 2 h respectively in order to homogenize the microstructure. This heat treated material has been considered as the as-received material in order to study the various normalizing and tempering treatments. The steel bars of 55 mm 12 mm 2 mm dimensions have been fabricated from the as-received material. The normalization and tempering conditions of these steel bars have been given in Table 2. Specimens for microstructure and hardness investigations have been prepared using standard metallographic methods. The specimens were etched with Villella's reagent solution (1 g picric acid þ 5 ml HCl þ 100 ml ethanol) to investigate the
Table 2 Heat treatment matrix. Initial heat treatment
1323 K (30 min)/1053 K (2 h)
Normalizing temperature (K) (for 30 min) Tempering temperature (K) (for 1 h)
1313 1013
1033
1053
1333 1013
1033
1053
135 1013
1033
Fig. 1. Optical micrographs of (a) as-received steel and steel normalized at (b) 1313 K, (c) 1333 K, (d) 1353 K. (Corresponding SEM micrograph is shown in inset).
1053
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Fig. 2. Optical micrographs of normalized and tempered conditions (a) 1313 K/1013 K, (b) 1313 K/1033 K, (c) 1313 K/1053 K, (d) 1333 K/1013 K, (e) 1333 K/1033 K, (f) 1333 K/1053 K, (g) 1353 K/1013 K, (h) 1353 K/1033 K, (i) 1353 K/1053 K. (Corresponding SEM micrograph is shown in inset).
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Table 4 Grain size of various normalized and tempered conditions. Heat treated condition
Grain size (mm) optical Grain size microscopy (mm) SEM
1313 K/1013 K 29 1313 K/1033 K 21 1313 K/1053 K 20
27 20 18
Heat treated condition
Grain size (mm) optical Grain size microscopy (mm) SEM
1333 K/1013 K 40 1333 K/1033 K 33 1333 K/1053 K 32
36 30 28
Heat treated condition
Grain size (mm) optical Grain size microscopy (mm) SEM
1353 K/1013 K 55 1353 K/1033 K 51 1353 K/1053 K 50
55 51 49
3. Results and discussion 3.1. Microstructure
Fig. 3. Variation of grain size with tempering temperature.
microstructures using optical (OM) and scanning electron microscope (SEM). Transmission electron microscopy (TEM) of heat treated samples has been carried out. TEM samples have been prepared by mechanically thinning up to 70 mm thickness using silicon carbide paper under flowing water. Discs of 3 mm diameter have been extracted followed by electrolytic double jet thinning using 10% perchloric acid and methanol solution as an electrolyte. Jet thinning has been carried out at 238 K and applied voltage of 20 V. For addition, extraction carbon replica was prepared out of etched sample to study the precipitates. Hardness measurement of as-received and heat treated steel samples have been carried using Vickers' hardness under the load of 5 kgf and dwell time of 15 s. The button head cylindrical tensile specimens having 4.0 mm gauge diameter and 28.6 mm gauge length of various heat treated conditions have been fabricated for the heat treated blocks. Tensile tests at room temperature have been carried out at constant cross-head speed with nominal strain rate of 3 103 sec1 in air. Load-elongation curves were recorded using data logger.
Microstructures of as-received steel and the steel normalized at various conditions are shown in Fig. 1. The steel exhibits a martensitic structure in 9 Cr steels with distinct prior-austenite grain boundaries. The tendency to form martensite is quite high in 9 Cr steel even if it is cooled very slowly as in the case of furnace cooling [14]. The prior-austenite grain size of the as-received steel was ~25 mm. The prior-austenite grain sizes of the steel normalized at 1313 K, 1333 K and 1353 K were about 28, 37 and 55 mm respectively (Table 3). In addition, the dissolution of precipitates has increased with higher normalizing temperature (Fig. 1). Microstructures of the various normalized and tempered conditions are shown in Fig. 2. The grain size measured by optical and scanning electron microscopy for all the heat treated conditions was in good agreement with each other, which is given in Table 4. It has been found that the grain size increased with increase in normalizing temperature [15]. However, the subgrain structure formation increased with increase in tempering temperature up to 1033 K; further increase in tempering temperature did not make any significant increase in the subgrains formation. This is due to the recovery taking place which resulted in decreased dislocation density [16] to its maximum up to 1033 K. SEM micrographs of the normalized and tempered conditions are shown in Fig. 2. The grain size increased with increase in normalizing temperature and decreased with increase in tempering temperature as discussed in previous section (OM) (Fig. 3). FEGeSEM micrographs of samples normalized at 1333 K and tempered at 1033 K are shown in Fig. 4. The M23C6 precipitates were decorated on the prior-austenitic and sub-boundaries (Fig. 4a), whereas intra-lath regions were uniformly decorated with MX precipitates (Fig. 4b). The martensitic laths were oriented in one direction within the packet boundary which is inside the prior austenite grains (Fig. 4). The area fraction of precipitates has been analyzed at constant magnification using image J software. It has been found that the area fraction of the precipitates decreased with increase in normalizing temperature and increased with increase in tempering temperature (Fig. 5). The decrease in area
Fig. 4. FEG-SEM micrographs of (a) M23C6 precipitates on the boundaries and (b) MX precipitates in the intra-lath region.
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Fig. 5. Variation of area fraction of precipitates with tempering temperature.
fraction of precipitates with increase in normalizing temperature might be due to the availability of lesser grain boundaries for the precipitation to occur. This can also be confirmed from the bulged grain boundaries observed in the case of the samples normalized at higher normalizing temperature and lower tempering temperature. Since the precipitates are less in these cases, there is less restriction to the grain boundary movement and hence the boundaries exhibits bulging (Fig. 2(a) and (g)). On the other hand, lower normalizing temperature and higher tempering temperature resulted in grain boundaries which are relatively straight (Fig. 2(c) and (i)), due to resistance to grain boundary movement offered by the more number of precipitates. TEM micrograph of sample normalized at 1333 K and tempered at 1033 K is shown in Fig. 6(a). M23C6 precipitates have been observed along the lath boundaries. The energy dispersive spectroscopy (EDS) spectrum confirming the chemical composition of the M23C6 precipitate is shown in Fig. 6(b). The diffraction pattern
101
of M23C6 is shown in Fig. 6(b) which matches with the diffraction pattern observed from JCPDS file 89-2724. The needle shaped MX precipitates have been observed inside the lath (Fig. 6(c)) and its electron energy loss spectrum (EELS) is shown in Fig. 6(d) which shows the presence of Vanadium in it. This confirms the presence of needle shaped vanadium nitride precipitate in the P92 steel and similar observation has been reported in the literature [7]. TEM micrographs of the samples normalized at 1313 K, 1333 K, 1353 K and subsequently tempered at 1013 K are shown in Fig. 7(a), (b) and (c) respectively. It has been observed that the lath size increased with increase in normalizing temperature (Fig. 8(a)). The variation of M23C6 precipitate size with normalizing temperature is shown in Fig. 8(b). The precipitate size has not changed significantly with change in the normalizing temperature. The samples normalized at 1353 K and tempered at 1013 K, 1033 K and 1053 K respectively are shown in Fig. 7 (c), (d) and (e). The lath width and precipitate size were found to increase with increase in tempering temperature. However the extent of increase in precipitate size was relatively lower up to 1033 K and higher from 1033 K to 1053 K. The higher extent of increase in precipitate size from 1033 K to 1053 K has been found to influence extensively on mechanical properties than the contribution from boundary strengthening, which is discussed in the subsequent sections. The activation energy for the growth of precipitates between the two temperature regimes 1013 K 1033 K and 1033 Ke1053 K has been calculated using the Equation (1) [17].
3R ln Q¼
d1 d2
ðT1 T2 Þ
T1 T2
(1)
where, Q e molar diffusion activation energy, (kJ mol1) R e universal gas constant, (J mol1 K1), d1 and d2 are the average particle sizes (nm) at respective temperatures of T1 (K) and T2 (K),
Fig. 6. (a) TEM micrograph depicts number of inter-lath M23C6 carbides, (b) EDS spectrum confirms Cr-rich nature of M23C6 carbide and diffraction pattern corresponding to [123] zonal axis of M23C6, (c) TEM micrograph of MX precipitate, (d) EELS spectrum of Vanadium rich MX precipitate.
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Fig. 7. TEM micrographs of (a) 1313 K/1013 K, (b) 1333 K/1033 K, (c) 1353 K/1013 K, (d) 1353 K/1033 K, (e) 1353 K/1053 K.
The above equation is obtained by substituting the precipitate size data for respective tempering temperatures and time in Equations (2) and (3), where the value of d0 is assumed as very small and negligible [17,18].
dnt dn0 ¼ kt
Q where k ¼ k0 exp RT
(2) (3)
where d0 initial size of the particle (nm) at t ¼ 0, dt is size of the particle after heating time in seconds (t) at a given temperature (T),
n is the growth exponent (3 for volume diffusion control) The activation energy for the coarsening of precipitates in the temperature regimes of 1013e1033 K and 1033e1053 K was obtained as 113 kJ mol1 and 272 kJ mol1 respectively. It can be observed that the activation energy for the growth of the precipitate increases with increase in the tempering temperature and corresponding increase in precipitate size [17]. In general, grain boundary energy increases with increase in tempering temperature [19]. This is due to the grain boundary energy being lowered by the segregation of precipitates along the grain boundary; segregation of precipitates retards the mobility of chromium atoms along the grain boundary. This increases the activation energy required for grain boundary diffusion [20]. In the present investigation, more
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Fig. 8. Variation of (a) lath width and (b) M23C6 precipitate size with tempering temperature.
precipitates have been observed at a tempering temperature of 1033 K in comparison with 1013 K. Hence grain boundary energy is less in the former case than the latter, which results in increased activation energy in the temperature regime of 1033e1053 K as compared to 1013e1033 K. Based on the values of activation energy in the temperature regime of 1013e1033 K and 1033e1053 K, it can be concluded that the rate controlling mechanism of the growth of the precipitate is the diffusion of interstitial solute carbon in the former case (97e133 kJ mol1) [21] and it is the diffusion of substitutional solute chromium in the later case (~210e~272 kJ mol1) [21]. Similar activation energy values have been reported by Pruthi et al [22] for volume diffusion of chromium (278 kJ/mol) in Inconel 600. 3.2. Hardness Hardness of the steel in the as-received and various normalized (untempered) conditions are given in Table 3. It has been observed that the hardness decreased with increase in normalizing temperature. This decrease in hardness with increase in normalizing temperature is due to increase in grain size. When these normalized samples were subjected to tempering treatment, for a given tempering temperature, the hardness decreased with decrease in the normalizing temperatures (Fig. 9). However, the extent of decrease in hardness between the samples normalized at 1353 Ke1333 K was found to be relatively higher, whereas the extent of decrease in hardness between the samples normalized at 1333 Ke1313 K was found to be relatively lower. This may be due to the presence of fine grain size (more grain boundaries) in the case of samples normalized at 1313 K, precipitation of large number of carbides (Fig. 5) and coarsening of undissolved precipitates [12] that occurs during tempering, which lead to relatively higher extent of decrease in hardness. Larger grain size present in the case of higher normalizing temperature led to lower precipitation of carbides (Fig. 5) during tempering and higher availability of the carbon in the matrix which led to increase the hardness of the material in comparison with other normalizing temperatures; similar observation has been reported by Ennis et al. [12]. This is reflected in the lower area fraction of the precipitate in the higher normalizing temperature and higher area fraction of the precipitates in the lower normalizing temperature (Fig. 5). Effect of decrease in area fraction of the precipitate dominate over the increase in the grain size and lath width with increase in normalizing temperature, that the competing influencing effects led to increase in hardness. The hardness was found to change with tempering temperatures for a given normalizing temperature (Fig. 9). The decrease in hardness with increase in tempering temperature
attributed to the increase in lath size, precipitate size, area fraction of precipitates and subgrain formation with decrease in dislocation density. It has been reported that the increase in dislocation density with the increase in normalizing temperature and decrease with increase in the tempering temperature, this also might have contributed to increase and decrease in hardness with normalizing and tempering conditions respectively [1,12]. 3.3. Mechanical properties The variation of yield stress (YS) and ultimate tensile strength (UTS) with normalizing and tempering temperatures have been shown in Fig. 10. For a given tempering temperature, the yield stress was found to increase with increase in normalizing temperature. The extent of increase in yield stress with normalizing temperature was higher up to the tempering temperature of 1033 K. However, yield stress with normalizing temperature had not significantly changed at tempering temperature of 1053 K. Similar behavior has been noticed in the variation of ultimate tensile strength with normalizing conditions for a given tempering temperature (Fig. 10(a) and (b)). The larger grain size and lath width in the steel normalized at 1353 K led to less precipitation of carbides due to the availability of lower boundary areas, which is reflected in the lower area fraction of the precipitates (Fig. 5). This resulted in the increase in hardness and strength of the steel normalized at 1353 K in comparison with the steel normalized at
Fig. 9. Variation of hardness with normalized and tempered conditions.
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Fig. 10. Variation of (a) YS, (b) UTS, (c) % elongation and (d) % uniform elongation with tempering temperature.
1333 K and 1313 K. The extent of increase in strength was significant for the steel tempered at 1013 K and 1033 K (Fig. 10). However, in the steel tempered at 1053 K, no significant change in the properties observed with various normalizing conditions due to relatively insignificant change in the formation of precipitates (Fig. 10(a), (b) and Fig. 6). Decrease in yield stress and tensile strength has been observed with increase in tempering temperature. The extent of decrease in strength was less up to the tempering temperature of 1033 K and further increase in tempering temperature led to decrease the strength (YS and UTS). This significant decrease in strength was strongly influenced by precipitate size than lath and grain size (Fig. 8(b), (a) and Fig. 3). This might be due to the fact that coarser the precipitates lesser is their tendency to act as dislocation movement barrier thereby decreasing the strength at the higher tempering temperature. Further, with increase in tempering temperature, carbon is consumed from the matrix for the precipitation and hence the solid solution strengthening effect would decrease. Although the decrease in strength was not significant up to 1033 K tempering temperature, the occurrence of recovery with decrease in dislocation density has been found to be dominant mechanism up to 1033 K [1]. The hardness and tensile results have been correlated and it was found that the correlation factor for change in yield stress and hardness is nearly 2.4 whereas for ultimate tensile strength and hardness is nearly 2.7. Busby et.al. have studied the relationship between hardness and yield stress in irradiated austenitic and ferritic steels they found that In ferritic steels, the correlation factor between change in yield stress and hardness was found to be 3.06 [23]. The variation of % elongation of the tensile tested samples with various normalizing and tempering conditions is shown in Fig. 10(c). The % elongation increases with increase in tempering temperature. However, the elongation increased significantly at tempering temperature of 1053 K. Similarly, uniform elongation
has increased with increase in tempering temperature up to 1053 K with a plateau up to 1033 K (Fig. 10(d)). On the other hand, it exhibited insignificant change with normalizing temperature. The steel when normalized at 1353 K exhibited higher ductility along with better strength and hardness (Fig. 10). The tempering temperature at 1013 K has been found to be suitable based on the good combination of strength and ductility in comparison with the tempering temperature at 1053 K, although the maximum ductility has been observed at 1053 K, the strengths of the steel were inadequate in comparison with other temperatures. Based on above discussions, it can be concluded that the influence of tempering temperature on the variation of microstructure and mechanical properties exhibited two distinct mechanisms in the temperature regimes (1313 Ke1333 K and 1333 Ke1353 K) as brought out in Fig. 11. In the temperature regime of 1313 Ke1333 K, the recovery of microstructure influenced the mechanical properties predominantly. On the other hand, coarsening of precipitates was found to be more dominant on mechanical properties in the temperature regime of 1333 Ke1353 K. When it is well known that the normalizing (1313e1353 K) and tempering temperature ranges (1013e1053 K) have been recommended in ASME A335 for the heat treatment of ASME P92 steel. Based on the present investigation, it can be suggested that normalizing temperature and tempering temperature for the steel can be 1353 K and 1013 K respectively. However, the efficacy of the choice of these temperatures needs to be considered with creep behavior as well. Therefore, creep studies are being planned to find out the long term microstructural stability of the steel consequent to the recommended heat treatment. 4. Conclusions P92 steel was subjected to varying normalizing and tempering temperatures in the temperature range of 1313 Ke1353 K and 1013 Ke1053 K respectively.
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mechanism of the growth of the precipitate in the former regime is the diffusion of interstitial solute carbon and it is the diffusion of substitutional solute chromium in the later regime Considering the observation of optimum combination of strength and ductility, normalizing at 1353 K and tempering at 1013 K has been suggested for ASME P92 steel. Acknowledgments The authors are grateful to UGC-DAE-CSR for funding this project. The first author is thankful to Dr G. Amarendra, ScientistIncharge UGC-DAE center Kalpakkam node for granting permission to use the facilities for conducting experiments, Mr. T. Sakthivel, Mr. Kudipudi Kiran Kumar, Mr. D. Vijayanand, Dr. Gopal Bhalerao and Dr. Shamima Hussain for technical discussion and support in the present investigation. References
Fig. 11. Correlation between microstructure and mechanical properties.
YS and UTS and hardness exhibited an increase with increase in normalizing temperature which was due to the decrease in area fraction of the precipitate The increase in precipitate size, lath width resulted in a decrease in the YS, UTS and hardness with increase in the tempering temperature The ductility increased with increase in normalizing and tempering temperature and a plateau was observed at the tempering temperature of 1033 K Activation energy calculations confirmed that two mechanisms dominate the tempering behavior. Thus based on the values of activation energy in the temperature regime of 1013e1033 K and 1033e1053 K, it can be concluded that the rate controlling
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International Journal of Pressure Vessels and Piping journal homepage: www.elsevier.com/locate/ijpvp
Effect of normalizing and tempering temperatures on microstructure and mechanical properties of P92 steel Dipika R. Barbadikar a, G.S. Deshmukh a, L. Maddi a, K. Laha b, P. Parameswaran c, A.R. Ballal a, D.R. Peshwe a, R.K. Paretkar a, M. Nandagopal b, M.D. Mathew b, * a b c
Department of Metallurgical and Materials Engineering, VNIT, Nagpur, India Mechanical Metallurgy Division, IGCAR, Kalpakkam, India Physical Metallurgy Group, IGCAR, Kalpakkam, India
a r t i c l e i n f o
a b s t r a c t
Article history: Received 5 March 2014 Received in revised form 26 June 2015 Accepted 1 July 2015 Available online 2 July 2015
In the present investigation, systematic studies on microstructure and mechanical properties of P92 steel subjected to various normalizing (1313e1353 K) and tempering (1013e1053 K) temperatures were carried out. The effect of heat treatment on microstructural parameters revealed an increase in grain size, lath width and decrease in the area fraction of the precipitates with an increase in normalizing temperature. The precipitate size has not changed significantly with increase in the normalizing temperature; rather it increased with increase in tempering temperature. Activation energy calculations confirmed the two fold mechanisms that dominate the tempering behavior. As a consequence, yield stress (YS) and ultimate tensile strength (UTS) were found to change with normalizing and tempering temperatures. P92 steel normalized at 1353 K and tempered at 1013 K was found to have the best combination of strength and ductility. © 2015 Elsevier Ltd. All rights reserved.
Keywords: P92 steel Normalizing Tempering Microstructure Mechanical properties
1. Introduction Ferritic/martensitic steel has been considered as a candidate material for power plant applications over austenitic stainless steel because of its excellent thermal conductivity, low coefficient of thermal expansion, good weldability accompanied with resistance to stress corrosion cracking and oxidation. They are widely used in the fabrication of high temperature components of fossil fired and steam generators of nuclear power plants. In order to increase the efficiency of the power plants operating at temperatures above 873 K and pressure of 250e300 bar [1], there is a need to develop materials in accordance with increased strength at particular service conditions. The development of 9Cr steels started few decades before, starting from P9, P91 and P92. P92 steel [2,3] is the next version of modified P91 steels, where Mo content is brought down to 0.5 wt.% from 1 wt.% and 1.8 wt.% tungsten is added. These steels derive their strength from tempered martensite lath structure which is stabilized by M23C6 type of carbide, intra-lath MX type carbide/nitride and martensite phase transformation induced high
* Corresponding author. E-mail address: [email protected] (M.D. Mathew). http://dx.doi.org/10.1016/j.ijpvp.2015.07.001 0308-0161/© 2015 Elsevier Ltd. All rights reserved.
dislocation density. P92 posses improved creep strength than P91 steel due to the solid solution strengthening and increased hardenability offered by tungsten addition. Enhancement in strength of P92 steel has been well explained by Ennis et al. [1] and has shown that the high degree of transient hardening in 9Cr steels is due to the presence of very high dislocation density obtained during normalization of the steel. The mechanical properties of these steels are found to be sensitive to normalizing and tempering temperatures as well as time, which alter the microstructural constituents [4]. The strengthening mechanisms of this steel have been discussed by many researchers. Kouichi et al. [5] have discussed the strengthening mechanism of creep resistant tempered martensitic steel. It has been reported that there is a large driving force for recovery due to the presence of high dislocation density and the pinning particles like M23C6, MX and solute atoms maintains the stable structure which is necessary for creep strength. The decrease in creep strength of 9Cr steels with service time has been attributed to decrease in dislocation density due to the formation of subgrain structure and to the coarsening of precipitates. Giroux et al. [6] have studied the softening behavior at high temperature in the uniaxial tensile test and found that the annihilation of subgrain boundaries and mobile dislocations leads to softening of the steel. The detailed
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Table 1 Chemical composition of P92 steel plate in wt. %.
Table 3 Hardness values of as-received and normalized conditions.
Element Wt. %
Cr 9.38
C 0.114
Mo 0.506
W 1.94
Nb 0.075
Mn 0.388
V 0.215
Element Wt. %
B 0.0018
N 0.0417
Al 0.0076
Cu 0.0162
P 0.0120
S 0.0035
Si 0.0239
TEM investigations of MN nitrides phases present in 9% Cr steel has been carried out by Yin Zhong Shen et al. [7] They have explained the role of precipitate in improving the creep strength of the steel. Investigation on precipitates present in P92 steel has been carried out by many researchers [8e11]. It has been reported that the ability of precipitates to act as dislocation barrier decreases as they coarsen. Thus finer precipitates lead to enhanced strength of the material. Ennis et al. [12] have studied the microstructural stability and creep strength of the P92 steel and have reported that the microstructural variations play a very important role in deciding the mechanical properties of the steel. However, the properties of steel deteriorate during service exposure. In particular, larger components with weldments have heat affected zones which are prone to type IV cracking, which depend on the initial microstructure of the material [13]. Further, weldments of the present steel consist of different regions of HAZ which are normalized for short times and are tempered. These zones develop heterogeneity in properties across the HAZ, which result in varied toughness and hence early creep rupture. Therefore achieving uniform properties have been a concern and it is as well a technological challenge to heat treat each individual zone of the HAZ at different tempering
Heat treatment
Hardness (Hv5)
Grain size (mm)
Initial heat treatment 1323 K/1053 K 1313 K 1333 K 1353 K
246 421 401 390
25 28 37 55
temperature. Considering this necessity, efforts are being made to systematically study the different combinations of normalizing and tempering temperatures which may provide an optimized combination of microstructure and mechanical properties. 2. Experimental details The chemical composition of the P92 steel analyzed using spectro-chemical analysis (ASTM-E-415-2008) is given in Table 1. This plate has been subjected to an initial heat treatment of normalizing and tempering at 1323 K for 30 min and 1053 K for 2 h respectively in order to homogenize the microstructure. This heat treated material has been considered as the as-received material in order to study the various normalizing and tempering treatments. The steel bars of 55 mm 12 mm 2 mm dimensions have been fabricated from the as-received material. The normalization and tempering conditions of these steel bars have been given in Table 2. Specimens for microstructure and hardness investigations have been prepared using standard metallographic methods. The specimens were etched with Villella's reagent solution (1 g picric acid þ 5 ml HCl þ 100 ml ethanol) to investigate the
Table 2 Heat treatment matrix. Initial heat treatment
1323 K (30 min)/1053 K (2 h)
Normalizing temperature (K) (for 30 min) Tempering temperature (K) (for 1 h)
1313 1013
1033
1053
1333 1013
1033
1053
135 1013
1033
Fig. 1. Optical micrographs of (a) as-received steel and steel normalized at (b) 1313 K, (c) 1333 K, (d) 1353 K. (Corresponding SEM micrograph is shown in inset).
1053
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Fig. 2. Optical micrographs of normalized and tempered conditions (a) 1313 K/1013 K, (b) 1313 K/1033 K, (c) 1313 K/1053 K, (d) 1333 K/1013 K, (e) 1333 K/1033 K, (f) 1333 K/1053 K, (g) 1353 K/1013 K, (h) 1353 K/1033 K, (i) 1353 K/1053 K. (Corresponding SEM micrograph is shown in inset).
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Table 4 Grain size of various normalized and tempered conditions. Heat treated condition
Grain size (mm) optical Grain size microscopy (mm) SEM
1313 K/1013 K 29 1313 K/1033 K 21 1313 K/1053 K 20
27 20 18
Heat treated condition
Grain size (mm) optical Grain size microscopy (mm) SEM
1333 K/1013 K 40 1333 K/1033 K 33 1333 K/1053 K 32
36 30 28
Heat treated condition
Grain size (mm) optical Grain size microscopy (mm) SEM
1353 K/1013 K 55 1353 K/1033 K 51 1353 K/1053 K 50
55 51 49
3. Results and discussion 3.1. Microstructure
Fig. 3. Variation of grain size with tempering temperature.
microstructures using optical (OM) and scanning electron microscope (SEM). Transmission electron microscopy (TEM) of heat treated samples has been carried out. TEM samples have been prepared by mechanically thinning up to 70 mm thickness using silicon carbide paper under flowing water. Discs of 3 mm diameter have been extracted followed by electrolytic double jet thinning using 10% perchloric acid and methanol solution as an electrolyte. Jet thinning has been carried out at 238 K and applied voltage of 20 V. For addition, extraction carbon replica was prepared out of etched sample to study the precipitates. Hardness measurement of as-received and heat treated steel samples have been carried using Vickers' hardness under the load of 5 kgf and dwell time of 15 s. The button head cylindrical tensile specimens having 4.0 mm gauge diameter and 28.6 mm gauge length of various heat treated conditions have been fabricated for the heat treated blocks. Tensile tests at room temperature have been carried out at constant cross-head speed with nominal strain rate of 3 103 sec1 in air. Load-elongation curves were recorded using data logger.
Microstructures of as-received steel and the steel normalized at various conditions are shown in Fig. 1. The steel exhibits a martensitic structure in 9 Cr steels with distinct prior-austenite grain boundaries. The tendency to form martensite is quite high in 9 Cr steel even if it is cooled very slowly as in the case of furnace cooling [14]. The prior-austenite grain size of the as-received steel was ~25 mm. The prior-austenite grain sizes of the steel normalized at 1313 K, 1333 K and 1353 K were about 28, 37 and 55 mm respectively (Table 3). In addition, the dissolution of precipitates has increased with higher normalizing temperature (Fig. 1). Microstructures of the various normalized and tempered conditions are shown in Fig. 2. The grain size measured by optical and scanning electron microscopy for all the heat treated conditions was in good agreement with each other, which is given in Table 4. It has been found that the grain size increased with increase in normalizing temperature [15]. However, the subgrain structure formation increased with increase in tempering temperature up to 1033 K; further increase in tempering temperature did not make any significant increase in the subgrains formation. This is due to the recovery taking place which resulted in decreased dislocation density [16] to its maximum up to 1033 K. SEM micrographs of the normalized and tempered conditions are shown in Fig. 2. The grain size increased with increase in normalizing temperature and decreased with increase in tempering temperature as discussed in previous section (OM) (Fig. 3). FEGeSEM micrographs of samples normalized at 1333 K and tempered at 1033 K are shown in Fig. 4. The M23C6 precipitates were decorated on the prior-austenitic and sub-boundaries (Fig. 4a), whereas intra-lath regions were uniformly decorated with MX precipitates (Fig. 4b). The martensitic laths were oriented in one direction within the packet boundary which is inside the prior austenite grains (Fig. 4). The area fraction of precipitates has been analyzed at constant magnification using image J software. It has been found that the area fraction of the precipitates decreased with increase in normalizing temperature and increased with increase in tempering temperature (Fig. 5). The decrease in area
Fig. 4. FEG-SEM micrographs of (a) M23C6 precipitates on the boundaries and (b) MX precipitates in the intra-lath region.
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Fig. 5. Variation of area fraction of precipitates with tempering temperature.
fraction of precipitates with increase in normalizing temperature might be due to the availability of lesser grain boundaries for the precipitation to occur. This can also be confirmed from the bulged grain boundaries observed in the case of the samples normalized at higher normalizing temperature and lower tempering temperature. Since the precipitates are less in these cases, there is less restriction to the grain boundary movement and hence the boundaries exhibits bulging (Fig. 2(a) and (g)). On the other hand, lower normalizing temperature and higher tempering temperature resulted in grain boundaries which are relatively straight (Fig. 2(c) and (i)), due to resistance to grain boundary movement offered by the more number of precipitates. TEM micrograph of sample normalized at 1333 K and tempered at 1033 K is shown in Fig. 6(a). M23C6 precipitates have been observed along the lath boundaries. The energy dispersive spectroscopy (EDS) spectrum confirming the chemical composition of the M23C6 precipitate is shown in Fig. 6(b). The diffraction pattern
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of M23C6 is shown in Fig. 6(b) which matches with the diffraction pattern observed from JCPDS file 89-2724. The needle shaped MX precipitates have been observed inside the lath (Fig. 6(c)) and its electron energy loss spectrum (EELS) is shown in Fig. 6(d) which shows the presence of Vanadium in it. This confirms the presence of needle shaped vanadium nitride precipitate in the P92 steel and similar observation has been reported in the literature [7]. TEM micrographs of the samples normalized at 1313 K, 1333 K, 1353 K and subsequently tempered at 1013 K are shown in Fig. 7(a), (b) and (c) respectively. It has been observed that the lath size increased with increase in normalizing temperature (Fig. 8(a)). The variation of M23C6 precipitate size with normalizing temperature is shown in Fig. 8(b). The precipitate size has not changed significantly with change in the normalizing temperature. The samples normalized at 1353 K and tempered at 1013 K, 1033 K and 1053 K respectively are shown in Fig. 7 (c), (d) and (e). The lath width and precipitate size were found to increase with increase in tempering temperature. However the extent of increase in precipitate size was relatively lower up to 1033 K and higher from 1033 K to 1053 K. The higher extent of increase in precipitate size from 1033 K to 1053 K has been found to influence extensively on mechanical properties than the contribution from boundary strengthening, which is discussed in the subsequent sections. The activation energy for the growth of precipitates between the two temperature regimes 1013 K 1033 K and 1033 Ke1053 K has been calculated using the Equation (1) [17].
3R ln Q¼
d1 d2
ðT1 T2 Þ
T1 T2
(1)
where, Q e molar diffusion activation energy, (kJ mol1) R e universal gas constant, (J mol1 K1), d1 and d2 are the average particle sizes (nm) at respective temperatures of T1 (K) and T2 (K),
Fig. 6. (a) TEM micrograph depicts number of inter-lath M23C6 carbides, (b) EDS spectrum confirms Cr-rich nature of M23C6 carbide and diffraction pattern corresponding to [123] zonal axis of M23C6, (c) TEM micrograph of MX precipitate, (d) EELS spectrum of Vanadium rich MX precipitate.
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Fig. 7. TEM micrographs of (a) 1313 K/1013 K, (b) 1333 K/1033 K, (c) 1353 K/1013 K, (d) 1353 K/1033 K, (e) 1353 K/1053 K.
The above equation is obtained by substituting the precipitate size data for respective tempering temperatures and time in Equations (2) and (3), where the value of d0 is assumed as very small and negligible [17,18].
dnt dn0 ¼ kt
Q where k ¼ k0 exp RT
(2) (3)
where d0 initial size of the particle (nm) at t ¼ 0, dt is size of the particle after heating time in seconds (t) at a given temperature (T),
n is the growth exponent (3 for volume diffusion control) The activation energy for the coarsening of precipitates in the temperature regimes of 1013e1033 K and 1033e1053 K was obtained as 113 kJ mol1 and 272 kJ mol1 respectively. It can be observed that the activation energy for the growth of the precipitate increases with increase in the tempering temperature and corresponding increase in precipitate size [17]. In general, grain boundary energy increases with increase in tempering temperature [19]. This is due to the grain boundary energy being lowered by the segregation of precipitates along the grain boundary; segregation of precipitates retards the mobility of chromium atoms along the grain boundary. This increases the activation energy required for grain boundary diffusion [20]. In the present investigation, more
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Fig. 8. Variation of (a) lath width and (b) M23C6 precipitate size with tempering temperature.
precipitates have been observed at a tempering temperature of 1033 K in comparison with 1013 K. Hence grain boundary energy is less in the former case than the latter, which results in increased activation energy in the temperature regime of 1033e1053 K as compared to 1013e1033 K. Based on the values of activation energy in the temperature regime of 1013e1033 K and 1033e1053 K, it can be concluded that the rate controlling mechanism of the growth of the precipitate is the diffusion of interstitial solute carbon in the former case (97e133 kJ mol1) [21] and it is the diffusion of substitutional solute chromium in the later case (~210e~272 kJ mol1) [21]. Similar activation energy values have been reported by Pruthi et al [22] for volume diffusion of chromium (278 kJ/mol) in Inconel 600. 3.2. Hardness Hardness of the steel in the as-received and various normalized (untempered) conditions are given in Table 3. It has been observed that the hardness decreased with increase in normalizing temperature. This decrease in hardness with increase in normalizing temperature is due to increase in grain size. When these normalized samples were subjected to tempering treatment, for a given tempering temperature, the hardness decreased with decrease in the normalizing temperatures (Fig. 9). However, the extent of decrease in hardness between the samples normalized at 1353 Ke1333 K was found to be relatively higher, whereas the extent of decrease in hardness between the samples normalized at 1333 Ke1313 K was found to be relatively lower. This may be due to the presence of fine grain size (more grain boundaries) in the case of samples normalized at 1313 K, precipitation of large number of carbides (Fig. 5) and coarsening of undissolved precipitates [12] that occurs during tempering, which lead to relatively higher extent of decrease in hardness. Larger grain size present in the case of higher normalizing temperature led to lower precipitation of carbides (Fig. 5) during tempering and higher availability of the carbon in the matrix which led to increase the hardness of the material in comparison with other normalizing temperatures; similar observation has been reported by Ennis et al. [12]. This is reflected in the lower area fraction of the precipitate in the higher normalizing temperature and higher area fraction of the precipitates in the lower normalizing temperature (Fig. 5). Effect of decrease in area fraction of the precipitate dominate over the increase in the grain size and lath width with increase in normalizing temperature, that the competing influencing effects led to increase in hardness. The hardness was found to change with tempering temperatures for a given normalizing temperature (Fig. 9). The decrease in hardness with increase in tempering temperature
attributed to the increase in lath size, precipitate size, area fraction of precipitates and subgrain formation with decrease in dislocation density. It has been reported that the increase in dislocation density with the increase in normalizing temperature and decrease with increase in the tempering temperature, this also might have contributed to increase and decrease in hardness with normalizing and tempering conditions respectively [1,12]. 3.3. Mechanical properties The variation of yield stress (YS) and ultimate tensile strength (UTS) with normalizing and tempering temperatures have been shown in Fig. 10. For a given tempering temperature, the yield stress was found to increase with increase in normalizing temperature. The extent of increase in yield stress with normalizing temperature was higher up to the tempering temperature of 1033 K. However, yield stress with normalizing temperature had not significantly changed at tempering temperature of 1053 K. Similar behavior has been noticed in the variation of ultimate tensile strength with normalizing conditions for a given tempering temperature (Fig. 10(a) and (b)). The larger grain size and lath width in the steel normalized at 1353 K led to less precipitation of carbides due to the availability of lower boundary areas, which is reflected in the lower area fraction of the precipitates (Fig. 5). This resulted in the increase in hardness and strength of the steel normalized at 1353 K in comparison with the steel normalized at
Fig. 9. Variation of hardness with normalized and tempered conditions.
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Fig. 10. Variation of (a) YS, (b) UTS, (c) % elongation and (d) % uniform elongation with tempering temperature.
1333 K and 1313 K. The extent of increase in strength was significant for the steel tempered at 1013 K and 1033 K (Fig. 10). However, in the steel tempered at 1053 K, no significant change in the properties observed with various normalizing conditions due to relatively insignificant change in the formation of precipitates (Fig. 10(a), (b) and Fig. 6). Decrease in yield stress and tensile strength has been observed with increase in tempering temperature. The extent of decrease in strength was less up to the tempering temperature of 1033 K and further increase in tempering temperature led to decrease the strength (YS and UTS). This significant decrease in strength was strongly influenced by precipitate size than lath and grain size (Fig. 8(b), (a) and Fig. 3). This might be due to the fact that coarser the precipitates lesser is their tendency to act as dislocation movement barrier thereby decreasing the strength at the higher tempering temperature. Further, with increase in tempering temperature, carbon is consumed from the matrix for the precipitation and hence the solid solution strengthening effect would decrease. Although the decrease in strength was not significant up to 1033 K tempering temperature, the occurrence of recovery with decrease in dislocation density has been found to be dominant mechanism up to 1033 K [1]. The hardness and tensile results have been correlated and it was found that the correlation factor for change in yield stress and hardness is nearly 2.4 whereas for ultimate tensile strength and hardness is nearly 2.7. Busby et.al. have studied the relationship between hardness and yield stress in irradiated austenitic and ferritic steels they found that In ferritic steels, the correlation factor between change in yield stress and hardness was found to be 3.06 [23]. The variation of % elongation of the tensile tested samples with various normalizing and tempering conditions is shown in Fig. 10(c). The % elongation increases with increase in tempering temperature. However, the elongation increased significantly at tempering temperature of 1053 K. Similarly, uniform elongation
has increased with increase in tempering temperature up to 1053 K with a plateau up to 1033 K (Fig. 10(d)). On the other hand, it exhibited insignificant change with normalizing temperature. The steel when normalized at 1353 K exhibited higher ductility along with better strength and hardness (Fig. 10). The tempering temperature at 1013 K has been found to be suitable based on the good combination of strength and ductility in comparison with the tempering temperature at 1053 K, although the maximum ductility has been observed at 1053 K, the strengths of the steel were inadequate in comparison with other temperatures. Based on above discussions, it can be concluded that the influence of tempering temperature on the variation of microstructure and mechanical properties exhibited two distinct mechanisms in the temperature regimes (1313 Ke1333 K and 1333 Ke1353 K) as brought out in Fig. 11. In the temperature regime of 1313 Ke1333 K, the recovery of microstructure influenced the mechanical properties predominantly. On the other hand, coarsening of precipitates was found to be more dominant on mechanical properties in the temperature regime of 1333 Ke1353 K. When it is well known that the normalizing (1313e1353 K) and tempering temperature ranges (1013e1053 K) have been recommended in ASME A335 for the heat treatment of ASME P92 steel. Based on the present investigation, it can be suggested that normalizing temperature and tempering temperature for the steel can be 1353 K and 1013 K respectively. However, the efficacy of the choice of these temperatures needs to be considered with creep behavior as well. Therefore, creep studies are being planned to find out the long term microstructural stability of the steel consequent to the recommended heat treatment. 4. Conclusions P92 steel was subjected to varying normalizing and tempering temperatures in the temperature range of 1313 Ke1353 K and 1013 Ke1053 K respectively.
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mechanism of the growth of the precipitate in the former regime is the diffusion of interstitial solute carbon and it is the diffusion of substitutional solute chromium in the later regime Considering the observation of optimum combination of strength and ductility, normalizing at 1353 K and tempering at 1013 K has been suggested for ASME P92 steel. Acknowledgments The authors are grateful to UGC-DAE-CSR for funding this project. The first author is thankful to Dr G. Amarendra, ScientistIncharge UGC-DAE center Kalpakkam node for granting permission to use the facilities for conducting experiments, Mr. T. Sakthivel, Mr. Kudipudi Kiran Kumar, Mr. D. Vijayanand, Dr. Gopal Bhalerao and Dr. Shamima Hussain for technical discussion and support in the present investigation. References
Fig. 11. Correlation between microstructure and mechanical properties.
YS and UTS and hardness exhibited an increase with increase in normalizing temperature which was due to the decrease in area fraction of the precipitate The increase in precipitate size, lath width resulted in a decrease in the YS, UTS and hardness with increase in the tempering temperature The ductility increased with increase in normalizing and tempering temperature and a plateau was observed at the tempering temperature of 1033 K Activation energy calculations confirmed that two mechanisms dominate the tempering behavior. Thus based on the values of activation energy in the temperature regime of 1013e1033 K and 1033e1053 K, it can be concluded that the rate controlling
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