- Email: [email protected]
Causes of heat-to-heat variation of creep strength in grade 91 steel
Author’s Accepted Manuscript Causes of heat-to-heat variation of creep strength in grade 91 steel K. Maruyama, J. Nakamura, N. Sekido, K. Yoshimi www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(17)30503-8 http://dx.doi.org/10.1016/j.msea.2017.04.050 MSA34951
To appear in: Materials Science & Engineering A Received date: 9 March 2017 Revised date: 12 April 2017 Accepted date: 12 April 2017 Cite this article as: K. Maruyama, J. Nakamura, N. Sekido and K. Yoshimi, Causes of heat-to-heat variation of creep strength in grade 91 steel, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2017.04.050 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Causes of heat-to-heat variation of creep strength in grade 91 steel K. Maruyama*, J. Nakamura, N. Sekido, K. Yoshimi Department of Materials Science, Tohoku University *
Corresponding author: 6-6-02 Aramaki-Aoba, Aoba-ku, Sendai 980-8579, Japan.
[email protected]
Abstract Heat-to-heat variation of creep strength is significant in grade 91 (Gr.91) steel, but the causes of the variation have not been well understood yet. In the present paper, creep rupture data of 14 heats of Gr.91 steel were analyzed paying attention to their chemical compositions and microstructures. The longest creep rupture lives analyzed are 2 105h at 500 and 550oC and 105h at 600oC. The causes of the heat-to-heat variation are different, depending on creep test conditions. At low temperature and high stress (creep rupture life of 104h at 500 and 550oC), creep rupture strength increases with increase of hardness after tempering. This suggests strengthening by a fine subgrain microstructure developed during normalizing and subsequent tempering. At higher temperature and intermediate time range (10 4h at 600oC), creep rupture strength depends on Cr concentration of the heats in addition to the hardness. This finding suggests an important contribution of recovery process of the subgrain microstructures to creep strength of the steel. In long-term creep (2 105h at 550oC and 105h at 600oC) creep rupture strength primarily increases with increasing grain size of the heats. This suggests that grain boundary sliding is an important deformation mode at low strain rate because of fine grain size usual with Gr.91 steel. Specifications on Ni concentration and N%/Al% ratio are newly introduced in the type II version of Gr.91 steel. They are not effective to eliminate a heat with low creep strength. Keywords: Grade 91 steel; Creep strength; Strengthening mechanism; Grain size effect; Creep life evaluation; Multi-region modelling
1. Introduction Creep strength enhanced ferritic (CSEF) steels are used for high temperature components in
1
ultra-super-critical steam boilers of electric power generation plants. Grade 91 (Gr.91) steel is most widely utilized among the CSEF steels. The components are designed to be used for 105h at 600oC, but creep defects are sometimes found in weldments of Gr.91 steel before the designed service duration [1]. Electric Power Research Institute (EPRI) has proposed a type II version of Gr.91 steel for preventing the premature creep failure [1,2]. ASME Boiler and Pressure Vessel Code has endorsed the type II version as a code case. In the type II version, allowable concentration of residual nickel is reduced from 0.4 to 0.2%, and upper limits of tramp elements and a lower limit of N%/Al% ratio are newly introduced. All the chemical concentrations in the present paper are in mass% otherwise indicated. The proposals are based mostly on effects of alloying elements on creep ductility and creep cracking. The EPRI report [1] mentions that creep rupture life and the consequent creep rupture strength increase with increasing creep ductility, but does not provide sufficient experimental evidence on the influence of ductility. Although normalizing temperature and duration have significant effects on creep rupture strength [2], the normalizing conditions of the type II version remain unchanged from the current ASME Boiler and Pressure Vessel Code. It is known that creep rupture strength of CSEF steels varies with Cr concentration [3-5], but the wide allowable range of Cr concentration, from 8 to 9.5%, is accepted in the type II version. Kimura et al. [6] have reported that 105h creep rupture strength of Gr.91 steel at 600oC increases monotonously with decrease in Ni concentration from 0.28 to 0.04%. The reduction of allowable Ni concentration in the type II version is partly based on this finding. They [6] took only four heats of Gr.91 steel in their examination, but six other heats are added to the original four heats in a recent examination [7]. From the recent examination, any correlation was not found between creep rupture strength and Ni concentration in the range less than 0.3%Ni. Kimura and Yaguchi [8] have confirmed absence of the Ni concentration dependence in their recent paper. Masuyama and Yamaguchi [9] and Komai et al. [10] have examined correlations between alloy concentrations and creep rupture strength of Gr.91 steel. They have suggested several alloying elements that may affect the creep rupture strength, but did not draw any quantitative conclusion on the correlations from their examinations. Abe [11] has studied heat-to-heat variation of creep rupture strength of several ferritic steels. He thought that formation of fine MX type precipitates is crucial for getting high creep rupture strength of the ferritic steels, and
2
paid attention to the free nitrogen concentration given by, (N concentration
Al concentration
Ti
concentration), since aluminum and titanium form thermally stable nitrides and consume N atoms necessary for forming MX precipitates. He has found that increase in the free nitrogen concentration improves creep rupture strength of 12Cr and 12Cr-1Mo-1W-0.3V steels. However, Gr.91 is not included in his study. The causes of heat-to-heat variation of creep strength in Gr.91 steel have not been well understood yet, though the variation is significant in the steel. In the present paper, creep strengths (creep rupture strength and creep deformation resistance) of 14 heats of Gr.91 steel are studied to clarify the causes of heat-to-heat variation of their creep strength and to assess the type II proposal. In creep of Gr.91 steel there are four regions with different creep behavior, in other words different stress exponents and activation energies for creep rupture life [7,12]. This fact suggests changes in dominant strengthening mechanism among the four regions. Therefore, causes of the variation are examined separately in each region. The examination in the multi-regions is different from the previous studies [1,2,6,9-11] on the heat-to-heat variation.
2. Materials studied The present paper studies creep strength of 6 heats of tube products and 8 heats of plate products of Gr.91 steel. Chemical concentrations of the major elements in each heat are listed in Table 1. Table 2 gives heat treatments and diffusion distances of Fe atoms during the normalizing and tempering of each heat along with its prior grain size and hardness measured after the treatments. The average distance X of diffusion is given by the following equation:
X 2Dt
(1)
where D is the diffusion coefficient, and t is the diffusion duration. The following values of the pre-exponential coefficient D0 and the activation energy QD for lattice self-diffusion of Fe atoms [13] are used for evaluating X: D0 = 2.8 10-4 m2/s, QD = 251 kJ/mol
(in phase)
D0 = 4.9 10-5 m2/s, QD = 284 kJ/mol
(in phase)
The heat treatment conditions given in Table 2 are used for the temperature and duration of diffusion.
3
Since diffusion of atoms occurs during heating before reaching the heat treatment temperatures, actual distances of diffusion may be slightly longer than the values listed in Table 2. The prior grain size is calculated from a grain size number assuming square shaped grains. Minimum creep rates and creep rupture lives of heats MGA, MGB, MGC, MGD, MGF, MGG, MgA, MgB, MgC and MgD have been reported in NIMS Creep Data Sheet 43A [14]. MGA to MGG are tube products, and MgA to MgD, plate products. Those data on plate products of heats JAA, JAB, JAC and JAD have been obtained by Japan Atomic Energy Agency [15]. Creep specimens of plate products were cut parallel to the rolling direction from center of the plates. The longest creep rupture lives of heats JAC and JAD at 500 and 550oC (close to 2 105h) were evaluated by analyzing their ongoing creep curves before fracture [12].
3. Analyses of creep data 3.1 Multi-region analysis Minimum Creep rate ̇ m and creep rupture life tr are represented by the following equations as functions of stress and temperature T: ̇ m = ̇ 0 n exp ( Q/R T)
(2)
tr = t0 -n exp (Q/R T)
(3)
where ̇ 0 and t0 are material constants, n is the stress exponent, Q is the activation energy, and R is the gas constant. Values of n and Q can change with creep test conditions. The changes in n and Q are often accompanied by a change in creep deformation or fracture mechanism. Stress dependence of creep rupture life in Gr.91 steel is shown in Fig.1. Heat MGC is taken as an example. The figure has originally been reported in Refs.[7,12]. The lines drawn in the figure are regression lines based on Eq.(3), and Refs.[7,12] explain how to determine the lines. As evident in the figure, there are four regions, H, M, L1 and L2 with different values of n. The appearance of the four regions is common to CSEF steels [7,16,17], and the four regions should appear in all the 14 heats of the present study. However, regions L1 an L2 are missing in some heats due to lack of high temperature or long-term data points. Stress exponents and activation energies in Fig.1 are as follows: n = 20 and Q =
4
718kJ/mol in region H, 8.9 and 606kJ/mol in M, 5.8 and 619kJ/mol in L1, and 2.6 and 232kJ/mol in L2. Figure 1 includes data points of heats JAC and JAD. The two heats are the same material, and were normalized at a higher temperature for a longer duration than those of heat MGC. This heat treatment makes their prior grain sizes of 17m, which is greater than 10.3m of heat MGC. Heat JAD was tempered for 60min at 760oC, and has Rockwell C hardness of HR = 22.5 (Vickers hardness of Hv = 252). Heat JAC was subjected to additional annealing for 500min at 740oC simulating post weld heat treatment (PWHT). This treatment gives its hardness of HR = 15.6 (Hv = 219) lower than HR = 18 of heat MGC. The creep rupture lives in region H (low temperature and high stress) increase with increasing the hardness of the heats. In long-term region L2, the creep rupture lives of heat JAC and JAD are longer than that of heat MGC due to their larger prior grain sizes. Contrary to these regions, the effects of hardness and prior grain size are indistinct in regions M and L1. The creep strengthening mechanisms are obviously different among regions H, M and L2. This finding points out that alloying elements and microstructures play different roles in strengthening among the three regions. Therefore, the causes of heat-to-heat variation in creep strength should be discussed separately in the three regions.
3.2 Evaluation of creep strength Creep data of each heat was subjected to the multi-region creep data analysis [7,16,17], like the example of heat MGC (see Fig.1), in order to specify the region each data point belongs to. Data points used in the present analyses are given in Fig.2: (a) creep rupture lives in regions H and M, (b) creep rupture lives in region L2, and (c) minimum creep rates in region L2. The multi-region creep data analyses give stress exponent n in each region of the heat. Their values are listed in Table3. All the data points of a heat in a region were used for determining the n values in addition to those plotted in Fig.2. n values are not listed in the table, when the heat does not have a sufficient number of data points to determine n in the region. As evident in the table, the scatter of n values within each region is not significant among the heats. A weighted average value of n is calculated in each region by taking a number of data points of each heat into account. The value is listed in the last line of the table. The lines drawn in Fig.2 are a regression line with the average n value and passing through the
5
center of the data points of each heat. The regression lines can represent stress dependence of creep rupture life and minimum creep rate of most of the heats, though the deviations of data points from the lines are not negligible in some heats. The regression lines are extrapolated or interpolated to the following conditions for getting creep rupture strength and creep deformation resistance under the conditions: 500 oC – 104 h, 550 oC – 104 h
(region H)
600 oC – 104 h
(region M)
550 oC – 2 x 105 h, 600 oC – 105 h, 600 oC – 10-7 h-1
(region L2)
The creep rupture lives vary by 40 times in region H, 5 times in region M and 3 times in region L2. The minimum value (59MPa) of 105h creep rupture strength at 600oC is two-thirds of its maximum value (91MPa) (see Fig.2(b)). The variation of the whole data points is wide, but the variation within a heat is narrow. The wide variation is primary caused by heat-to-heat variation in creep strength. The causes of the heat-to-heat variation is discussed in the next section on the basis of the creep strengths obtained in Fig.2.
4. Discussion 4.1 Factors not affecting creep strength Some heats of Gr.91 steel have substantially lower creep rupture strength than the values listed in ASME Code and ECCC Code. EPRI ascribes the low creep strength to excess amounts of harmful elements included in the weak heats [1]. To eliminate the low strength heats, EPRI has proposed the type II version of Gr.91 steel [1,2]. In this version, allowable Ni concentration is reduced from 0.4 to 0.2%, and N%/Al% ratio greater than 4 (in mass%) is newly introduced. Let us examine whether the revised specifications really affect creep strength of Gr.91 steel. Creep strengths of all the heats are plotted against residual Ni concentration in Fig.3. Open and solid symbols are tube and plate products, respectively. The error bar of each heat gives its largest upward and downward deviations of data points from its regression line in Fig.2. No error bar is attached to a data point when a heat has only one data point in a region at a temperature. At 550oC (see Fig.3(d)), for
6
example, every heat has only one data point in region L2 except for heat MGC. As evident in the figure, most error bars are substantially short when compared to the heat-to-heat variation in creep strength. Kimura at el. [6] have pointed out that creep rupture strength increases with decreasing Ni concentration, based on the creep rupture strength of heats MgC (0.04%Ni), MGA (0.12%Ni), MGB (0.20%Ni) and MGC (0.28%Ni). Figure 3 includes data points of 10 other heats in addition to the 4 heats. It is obvious that creep strength is independent of Ni concentration in all the regions H, M and L2. This finding accords with recent reports [7,8] on the effects of residual Ni concentration. Nitrogen forms MX type precipitates with vanadium and niobium. The precipitates are fine and thermally stable, and necessary to keep a fine subgrain microstructure that provides high creep strength of CSEF steels. Aluminum atoms have high affinity to N atoms, and form coarse AlN inclusions before the formation of MX precipitates. After the same service duration, weldments with N%/Al% ratio (in mass%) less than 4 have a larger number of creep cracks than those with a high N%/Al% ratio [1]. Based on this finding on Gr.91 steel, EPRI [1] recommends N%/Al% > 4 (N%/Al% > 7.7 in atomic ration) in the type II version. Brett [18] has also pointed out the importance of N%/Al% ratio for achieving high creep ductility of the steel. The creep strengths of the 14 heats are plotted against N%/Al% ratio (in atomic%) in Fig.4. The ratio varies from 4 (less than 7.7) to 100. Any systematic change of the creep strength with N%/Al% ratio cannot be seen despite the wide variation in the ratio. A recent report [8] supports this result. The result can be explained as follows on the basis of the literature [11]. The ratio is not necessary to be greater than 7.7 (in atomic ratio), if N%>Al% and free nitrogen concentration is large enough to form a sufficient number of fine MX precipitates. Correlations of the creep strength with Mo, Nb, V, Mn, Al and (C+N) concentrations are also studied. However, any clear correlations are not found on these variables since their concentrations are controlled within the ranges in which the elements do not cause variation of creep strength. It can be concluded that the variations in Ni, Mo, Nb, V, Mn, Al and (C+N) concentrations and N%/Al% ratio are not responsible for the variation in creep strength of Gr.91 steel. However, it should be recalled that 105h creep rupture strength at 600oC (Fig.4(e)) varies from 59 to 91 MPa. When the average creep rupture strength of Gr.91 steel is 90MPa [19], the creep rupture strength of the weakest heat is as low as an
7
allowable stress of Gr.91 steel in creep regime.
4.2 Effects of heat treatments on creep strength The creep rupture strengths of Gr.91 steel are plotted against diffusion distance during tempering (including the treatment simulating PWHT) in Fig.5. The creep rupture strength in region H decreases with increasing the diffusion distance, but the decrease is not evident in regions M (see Fig.5(c)) and L2. CSEF steels are strengthened by their fine subgrain microstructures [20,21]. Increase of diffusion distance during tempering results in coarsening (recovery) of subgrains, and their hardness after tempering is a measure of the recovery process. On the basis of these understandings, the creep strengths of the 14 heats are plotted in Fig.6 against their hardness. A better correlation is found in this figure than that in Fig.5, confirming the dominant role of subgrain microstructures in determining the creep strength in regions H and M. The slopes of the regression lines in Figs.6(a)-(c) decrease from 5.4MPa/HRC at 500oC to 3.8 at 550oC and further to 1.3 at 600oC. The effect of hardness becomes less significant with increasing creep testing temperature. Komai et al [22] have studied the effects of hardness after tempering on creep rupture strength of Gr.91 steel. Their result is very similar to the one obtained in the present study. The heat with the highest hardness in region L2 (see Figs.6(d)-(f)) has the high creep strength due to its large grain size. If this heat is omitted, the creep strength in region L2 is insensitive to the hardness. The creep strengths in region L2 are plotted in Fig.7 against average diffusion distance of Fe atoms during normalizing. The creep strength increases with increasing the diffusion distance contrary to the insensitiveness of creep strength in Figs.6(d)-(f). This increase is widely known, but the cause of strength increase has not been fully clarified yet [2]. It is to be noted that 104h creep rupture strength in regions H and M is independent of the diffusion distance. Coarsening of grains and dissolution of NbX precipitates proceed during the normalizing [2]. Dissolution of a larger amount of NbX gives a greater number of fine NbX precipitates formed during tempering and creep, and consequently can improve creep strength. Diffusion of Nb atoms is fast enough at temperatures above 1000oC, and Nb concentration around NbX reaches an equilibrium within a short
8
period of time. Therefore, the amount of NbX dissolved into matrix after normalizing is determined not by the diffusion distance in Fig.7 but by normalizing temperature [2,23]. This fact expects a close correlation between the creep strength and normalizing temperature, if the fine NbX precipitates determine the creep strength. The creep strengths of the 14 heats are plotted against normalizing temperature in Fig.8. As evident in Fig.8(a)-(c) the 104h creep rupture strength in regions H and M is independent of the normalizing temperature. As seen in Figs.8(d)-(f), the creep strength in region L2 increases with increasing normalizing temperature. This fact does not conflict with the strengthening due to the precipitation of NbX, but the scatter in data points is large in the figures, suggesting the presence of some other factor having a more significant effect on the creep strength. The other factor varied by the normalizing, namely grain size increases with increasing temperature and duration of normalizing [2]. The creep strengths in region L2 are plotted against prior grain size in Fig.9. Since a prior grain is divided into several packets after martensitic transformation, packet size should be used as a measure of grain size. It has been reported on low carbon steels that packet size is proportional to prior grain size and the proportional coefficient of 0.4 is independent of alloy addition [24]. Let us use the prior grain size instead of packet (grain) size in the present analysis, since packet sizes have not been reported on the 14 heats. The creep strength increases with increasing grain size in Fig.9. The scatter in data points in the figure is smaller than that in Fig.8. It follows from this fact that the increase in grain size is the major cause of the dependence of creep strength on the diffusion distance in Fig.7. It should be noted that the prior grain sizes are from 9 to 17m in Fig.9 and very small. Packet sizes are estimated to be 3.5 to 7m. Deformation modes related to grain boundary, such as grain boundary sliding and diffusion creep can contribute to creep strain in this range of grain size [25], and contributions of the deformation modes increase with decreasing grain size. Creep rupture life is related to packet size dp by the following equation in region L2 [25]: tr = to -2.5 dpP D-1
(4)
where dp is the packet size, and P is the packet size exponent. P = 2 and 3 in grain boundary sliding controlled by lattice diffusion and grain boundary diffusion, respectively [25]. The increase in packet size extends creep rupture life by four times when P = 2. This increase is similar to the heat-to-heat variation
9
in creep rupture life observed in Fig.2(b). It is to be noted that the 104h creep rupture strengths in regions H and M are independent of grain size.
4.3 Effects of Cr concentration on creep strength Among the elements listed in Table1, chromium is the only element that has some correlation with the creep strength. The creep strengths of the 14 heats are plotted against Cr concentration in Fig.10. The creep rupture strength (Figs.10(d) and (e)) and creep deformation resistance (Fig.10(f)) in region L2 increase with increasing Cr concentration. This is true for the 104h creep rupture strength in region M (Fig.10(c)). However, the scatter in data points is large in Figs.10(d) and (e) due to the variation in grain size. The same effect of Cr is not evident in region H (Figs.10(a) and (b)) in which diffusion distance is too short for recovery of subgrain microstructures to proceed. It can be concluded that creep strength of Gr.91 steel varies with Cr concentration in the range allowed in ASME Code (8 to 9.5%). Murata
and
his
co-workers
[3,4]
have
studied
creep
rupture
life
of
X%Cr-3.5%W-3%Co-VNbBN steels. The steels were normalized at 1070oC for 5h, and then double-tempered at 570oC for 20h and 680oC for 20h. They varied Cr concentration from X = 8.5 to 11.5% and carried out creep tests at 650oC. They have reported that the steels give the longest rupture life at 10%Cr in the rupture life range around 104h, at 9%Cr in the range around 3 104h and at 8.5%Cr in the range over 5 104h. Gasemi-Armaki et al [5] have studied creep of X%Cr-1.9%W-0.4%Mo-0.9%Cu-VNb steel (X = 5, 7, 9, 10.5 and 12%). The steels were normalized at 1050oC for 30min and then tempered at 770oC for 2h. Creep rupture lives of these steels increase with increasing Cr concentration in short-term creep, but the 9%Cr steel gives the longest rupture life in long-term creep. Creep rupture life of CSEF steels increases with decreasing subgrain width. Subgrain width after the tempering decreases with increasing Cr concentration, resulting in the longer creep rupture life of the higher Cr steel in short-term creep [5]. Ostwald ripening of M23C6 particles proceeds during creep, resulting in coarsening of subgrain width [21]. The ripening rate is faster at higher Cr concentration. Therefore, Cr concentration that gives the minimum subgrain width varies with the time ranges [5,26]. The 9%Cr steel has the smallest subgrain width after 104h annealing [5], resulting in its longest rupture life in the long-term creep. The Cr
10
concentration dependent creep behavior in Fig.10 is ascribed to the Cr concentration dependence of thermal stability of subgrain microstructures. This effect of chromium is especially important in Figs.10(c) and (f).
4.4 Quantitative analyses of creep strength The aforementioned analyses indicate that the hardness after tempering, the grain size and the Cr concentration are the major factors determining the creep strength. In more detail, the hardness has no effect in region L2, and the creep rupture strength is independent of grain size in regions H and M. Chromium concentration may not affect the creep rupture strength in region H. Let us assume that creep strength S is given by the following equations of the hardness HR, the Cr concentration Cr% and the prior grain size d: S = a (HR + b Cr%) + S1
(H and M)
(5)
S = c (d + e Cr%) +S2
(L2)
(6)
where a, b, c, e, S1 and S2 are constants. They are determined so that they give the best fit of the regression lines to the data points in Figs.11 and 12. Equations (5) and (6) include the two major variables only, but there may be some other variables that contribute to the variation in creep strength. The creep rupture strengths in regions H and M are plotted against (HR + b Cr%) in Fig.11, and those in region L2, against (d + e Cr%) in Fig.12. The values of b giving the best fit are slightly different in Figs.11(a) and (b). The average of these values is finally used in the two figures. A similar average value of e is used also in Figs.12(a) and (b). Good correlations hold in both figures, confirming the dominant roles of the hardness, the grain size and the Cr concentration in controlling creep strength of Gr.91 steel. In region H (Figs.11(a) and (b)), 99% of the change in abscissa is due to the hardness. The effect of Cr concentration on the creep rupture strength is negligible, since thermal recovery of subgrain microstructures does not proceeds in region H. In region M (Fig,11(c)), 60% of the change in abscissa comes from that in Cr concentration, and the other 40%, from the change in hardness. Thermal recovery of subgrain microstructures that are affected by Cr concentration contributes to the creep rupture strength in this region.
11
As evident in Figs.2(b) and (c), region L2 appears under the same creep test conditions both in creep rupture life and minimum creep rate. The rupture lives are 3 104 to 105h in region L2 at 600oC. The minimum creep rate appears at 10 to 20% of the rupture lives in region L2 [12]; namely it appears before 2 104h. Because of the early appearance of the minimum creep rate, the effect of Cr concentration is still substantial in Fig.12(c) as is the case in Fig.11(c). In Fig.12(c), 50% of the change in abscissa is due to Cr concentration and the other 50%, due to grain size. As shown in Figs.12(a) and (b), the 105h creep rupture strength at 600oC and the 2 105h creep rupture strength at 550oC increase with increasing the grain size and Cr concentration. The contribution of Cr concentration to the change in abscissa is 10%, and the grain coarsening is the dominant cause of the increase of creep rupture strength in region L2. The origin of the grain size dependence in creep strength is the contribution of grain boundary related deformation modes, probably grain boundary sliding to creep strain. The contribution increases with decreasing grain size. In region L2, stress exponent for creep rupture life is 2.5 (cf. 3 or greater in dislocation creep), and the activation energy for creep rupture life (230kJ/mol) is less than that for lattice self-diffusion (300kJ/mol [27]) rate-controlling dislocation creep. Stress exponents for creep rupture life are similar among CSEF steels in region L2, their welded joints creep-ruptured by type IV cracking in fine grain HAZ [28,29] and a material simulating fine grain HAZ [30]. These findings can be valid if grain boundary sliding is the dominant creep deformation mode in region L2. Heat MGG is the weakest in Fig.12(a); 0.12%Ni and N%/Al% = 23 (in mass%). In Fig.12(b) heat MgB is the weakest; 0.09%Ni and N%/Al% = 4.9 (in mass%). Both heats contain nickel less than 0.2% and have the ratios greater than 4. They fulfill the type II specifications on Ni concentration and N%/Al% ratio as well as on other elements, though concentrations of tramp elements have not been reported in the literature [14,15]. The type II specifications of Gr.91 steel cannot exclude the heats with low creep strength.
5. Summary 1) In creep of Gr.91 steel, there are four regions H, M, L1 and L2 with different values of stress exponent
12
n for creep rupture life and minimum creep rate. This fact is common to all the heats of the steel. The values of n in each region are similar in all the heats. 2) Concentrations of Mo, V, Nb, C+N, Ni, Mn and Al in Gr.91 steel are controlled within the ranges in which the elements do not cause variation of creep strength. Creep strength of the steel is independent of N%/Al% ratio when N%/Al% (in atomic%) is greater than 4. 3) Creep rupture strength of the steel depends primarily on its hardness after tempering in low-temperature and high-stress region H. It increases with increasing the hardness. 4) Creep rupture strength of the steel depends primarily on its grain size in long-term region L2, though increase in Cr concentration slightly contributes to strengthening. It increases with increasing grain size and Cr concentration. 5) Creep deformation resistance that gives minimum creep rate of 10-7h-1 at 600oC (region L2) increases with increasing Cr concentration and grain size of the steel. 104h creep rupture strength of the steel at 600oC (region M) increases with increasing its Cr concentration and hardness after tempering. The two variables equally contribute to the strengthening. 6) The type II specifications of Gr.91 steel cannot eliminate the weakest heats of the present study.
Acknowledgements A part of the present study was supported by an ALCA project, and the other part, by an NEDO project #1400070-0. Financial supports from Japan Science and Technology Agency (ALCA project) and New Energy and Industrial Technology Development Organization (NEDO project) are acknowledged. The author cannot accomplish the present analyses without the high quality data reported in NIMS Creep Data Sheet No.43A and JAEA-Data/Code 2008-030. The authors deeply acknowledge National Institute for Materials Science and Japan Atomic Energy Agency for their publication of the creep data sheets. Very useful suggestions on packet size were given by Prof. M. Mitsuhara, Department of Engineering Science for Electronics and Materials, Kyushu University.
13
References [1] EPRI Report No. 3002003472, The benefits of improved control of composition of creep-strength-enhanced ferritic steel grade 91, EPRI, Palo Alto, CA, 2014. [2] EPRI Report No. 3002004370, The influence of steel making and processing variables on the microstructure and properties of creep strength enhanced ferritic (CSEF) steel grade 91, EPRI, Palo Alto, CA, 2014. [3] Y. Murata, K. Yamashita, M. Morinaga, T. Hara, K. Miki, T. Azuma, T. Ishiguro, R. Hashizume, Dependence of precipitation behavior and creep strength on Cr content in high Cr ferritic heat resistant steels, J. Solid Mech. Mater. Eng. 3 (2009) 457-463. [4] K. Miki, T. Azuma, T. Ishiguro, O. Tamura, R. Hashizume, Y. Murata, M. Morinaga, Development of heat resistant ferritic steels containing 8.5–11.5% Cr based on long-term creep rupture strength, in: Proceedings of the 18th International Forgemasters Meeting, Pittsburg, PA, September 12-15, 2011, pp. 211-214. [5] H. Ghassemi-Armaki, R. Chen, K. Maruyama, M. Igarashi, Creep behavior and degradation of subgrain structures pinned by nanoscale precipitates in strength-enhanced 5 to 12 pct Cr ferritic steels, Metall. Mater. Trans. A 42 (2011) 3084-3094. [6] K. Kimura, K. Sawada, H. Kushima, Y. Toda, Influence of chemical composition and heat treatment on long-term creep strength of grade 91 steel, Procedia Eng. 55 (2013) 2-9. [7] K. Maruyama, J. Nakamura, K. Yoshimi, Assessment of long-term creep rupture strength of T91 steel by multiregion rupture data analysis, Trans. ASME, J. Pressure Vessel Technol. 138 (2016) 031407-1-9. [8] K. Kimura, M. Yaguchi, Re-evaluation of long-term creep strength of base metal of ASME grade 91 type steel, in: ASME 2016 Pressure Vessels and Piping Conference, 2016, Paper No. PVP2016-63355. [9] F. Masuyama, T. Yamaguchi, Chemical specification and design factor consideration for creep degradation in grade 91 steel, in: ASME 2016 Pressure Vessels and Piping Conference, 2016, Paper No. PVP2016-63307. [10] N. Komai, K. Arisue, N. Saito, K. Tominaga, M. Fujita, Effect of chemical composition on the creep
14
rupture strength of mod.9Cr-1Mo steel weldments, in: Proc. of the 1st International Conference on Advanced High-temperature Materials Technology for Sustainable and Reliable Power Engineering, Sapporo, Japan, June 29-July 3, 2015, pp. 57-60. [11] F. Abe, Heat-to heat variation in long-term creep strength of some ferritic steels, Intern. J. Press. Vess. Piping 87 (2010) 310-318. [12] K. Maruyama, J. Nakamura, K. Yoshimi, Y. Nagae, Evaluation of long-term creep rupture life of Gr.91 steel by analysis of on-going creep curves, in: J. Parker, J. Shingledecker, J. Siefert (Eds.), Advances in Materials Technology for Fossil Power Plants, ASM Intern., Materials Park, OH, 2016, pp. 467-478. [13] Metal Data Book, fourth ed., Edited by Japan Inst. Metals, Maruzen, Tokyo, 2004, pp. 20-25. [14] National Institute for Materials Science, Data sheets of the elevated-temperature properties of 9Cr-1Mo-V-Nb steel tubes for boilers and heat exchangers, 9Cr-1Mo-V-Nb steel plates for boilers and pressure vessels and 9Cr-1Mo-V-Nb seamless pipe for high temperature service, NIMS Creep Data Sheet No. 43A, NIMS, Tsukuba, Japan, 2014,. [15] S. Kato, T. Furukawa, E. Yoshida, Material test data of mod. 9Cr-1Mo steel (1), JAEA-Data/Code 2008-030, Japan Atomic Energy Agency, Tokai-mura, Japan, 2009. [16] K. Maruyama, K. Yoshimi, Influence of data analysis method and allowable stress criterion on allowable stress of Gr.122 heat resistant steel, Trans. ASME, J. Press. Vess. Technol. 129 (2007) 449-453. [17] K. Maruyama, J. Nakamura, K. Yoshimi, Prediction of long-term creep rupture life of grade 122 steel by multiregion analysis, Trans. ASME, J. Pressure Vessel Technol. 137 (2015) 021403-1-5. [18] S.J. Brett, UK experience with modified 9%Cr (grade 91) steel, Energy Mater., 2 (2007) 117-121. [19] W. Bendick, L. Cipolla, J. Gabrel, J. Hald, New ECCC assessment of creep rupture strength for steel grade X10CrMoVNb9-1 (Grade 91), in: I.A. Shibli, S.R. Holdsworth (Eds.), Creep and Fracture in High Temperature Components, DEStech Publications, Lancaster, PA, 2009, pp. 56-67. [20] K. Maruyama, H. Ghassemi Armaki, R.P. Chen, K. Yoshimi, M. Yoshizawa, M. Igarashi, Cr concentration dependence of overestimation of long term creep life in strength enhanced high Cr ferritic
15
steels, Int. J. Press. Vessels Piping 87 (2010) 276-281. [21] H. Ghassemi-Armaki, R.P. Chen, K. Maruyama, M. Igarashi, Contribution of recovery mechanisms of microstructure during long-term creep of Gr.91 steels, J. Nuclear Mater., 433 (2013) 23-29. [22] N. Komai, K. Arisue, K. Tominaga, M. Fujita, Influence of holding time during tempering on the long-term creep rupture strength of mod.9Cr-1Mo steel, in: J. Parker, J. Shingledecker, J. Siefert (Eds.), Advances in Materials Technology for Fossil Power Plants, ASM Intern., Materials Park, OH, 2016, pp. 430-440. [23] Y. Tsuchida, K. Tokuno, K. Hashimoto, Development of BOF manufacture of modified 9Cr-1Mo steel plates with excellent strength and toughness, Nippon Steel Technical Report 58 (1993) 27-35. [24] S. Morito, H. Yoshida, T. Maki, X. Huang, Effect of block size on the strength of lath martensite in low carbon steels, Mater. Sci. Eng. A 438-440 (2006) 237-240. [25] J. Cadek, Creep in Metallic Materials, Elsevier, Amsterdam, 1988. [26] H. Ghassemi-Armaki, R. Chen, K. Maruyama, M. Igarashi, Premature creep failure in strength enhanced high Cr ferritic steels caused by static recovery of tempered martensite lath structures, Mater. Sci. Eng. A 527 (2010) 6581-6588. [27] H. Oikawa, Y. Iijima, Diffusion behavior of creep-resistant steels, in: F. Abe, T. Kern, R. Viswanathan (Eds.), Creep-resistant Steels, Woodhead Publishing, Cambridge, 2008, pp. 241-264. [28] K. Maruyama, K. Yoshimi, Y. Hasegawa, H. Morimoto, F. Masuyama, Temperature and stress dependence of creep life of welded joints in strength enhanced high Cr ferritic steels, Energy Mater. 4 (2009) 70-75. [29] M. Yaguchi, S. Nagai, K. Sawada, K. Kimura, Microstructure and creep strength of grade 91 steel used in USC plants, in: J. Parker, J. Shingledecker, J. Siefert (Eds.), Advances in Materials Technology for Fossil Power Plants, ASM Intern., Materials Park, OH, 2016, pp. 447-458. [30] K. Yoshida, H. Tsuruta, M. Tabuchi, K.-I. Kobayashi, Creep damage evaluation method for welded joints of grade 91 steels, in: J. Parker, J. Shingledecker, J. Siefert (Eds.), Advances in Materials Technology for Fossil Power Plants, ASM Intern., Materials Park, OH, 2016, pp. 545-556.
16
300
Stress ( MPa )
200
100 80 60 40
20 1 10
550 oC
600 oC
650 o C
H
M
70
0 oC
L1
L2
JAD MGC JAC 10
2
10
3
10
4
10
5
Rupture Life ( h )
Figure 1 A comparison of creep rupture lives among heats MGC, JAC and JAD. Effects of their hardness and grain size are different among regions H, M and L2.
17
Figure 2 Creep rupture lives in regions (a) H (500 and 550oC), M (600oC), and (b) L2 (550 and 600oC) together with (c) minimum creep rates in region L2 (600oC).
18
Figure 2 (b) and (c)
19
360 o
(a) 500 C, 10kh 320 280 240
o
(b) 550 C, 10kh 220 200 180
Creep Strength ( MPa )
o
(c) 600 C, 10kh
140 120 100
o
(d) 550 C, 200kh
Open: Tube Solid: Plate
150 120 100 80 60 o
(e) 600 C, 100kh 100 80 60 o
−7 −1
(f) 600 C, 10 h 40 0
0.1
0.2
0.3
0.4
Ni Concentration ( mass% )
Figure 3 Correlation of Ni concentration (mass%) with creep rupture strength of Gr.91 steel in regions (a) (b) H, (c) M and (d) (e) L2, and (f) with creep deformation resistance in region L2.
20
360 o
(a) 500 C, 10kh 320 280 240
o
(b) 550 C, 10kh 220 200 180
Creep Strength ( MPa )
o
140
(c) 600 C, 10kh
120 100 o
(d) 550 C, 200kh
Open: Tube Solid: Plate
150 120 100 80 60 o
(e) 600 C, 100kh 100 80 60
−7 −1
o
(f) 600 C, 10 h 40 1
10
2
10
N% / Al% (Atomic Ratio)
Figure 4 Correlation of N%/Al% ratio with (a)-(e) creep rupture strength and (f) creep deformation resistance. (a) (b) Region H, (c) region M and (d)-(f) region L2.
21
360 o
Open: Tube Solid: Plate
(a) 500 C, 10kh
Creep Rupture Strength ( MPa )
320 280 240 o
(b) 550 C, 10kh 220 200 180 o
(c) 600 C, 10kh
140 120 100 0
1
2 1/2
(2 D t)
3
( m )
Figure 5 Variation of creep rupture strength with average lattice diffusion distance of Fe atoms during tempering and PWHT. (a) (b) Region H and (c) region M.
22
360 o
(a) 500 C, 10kh 320 280 240 o
(b) 550 C, 10kh 220 200 180
Creep Strength ( MPa )
o
(c) 600 C, 10kh
140 120
Open: Tube Solid: Plate
100 o
(d) 550 C, 200kh 150 120 100
o
(e) 600 C, 100kh 80 60 100 80 60
−7 −1
o
(f) 600 C, 10 h 40 12
16
20
24
Hardness (H RC)
Figure 6 Effect of hardness after tempering on (a)-(e) creep rupture strength and on (f) creep deformation resistance. (a) (b) Region H, (c) region M, and (d)-(f) region L2.
23
180 o
(a) 550 C, 200kh 150
Creep Strength ( MPa )
120 100 o
(b) 600 C, 100kh 80 60 100 o
−7 −1
(c) 600 C, 10 h 80 60 40 0.5
Open: Tube Solid: Plate
1
1.5
(2 D t)
1/2
2
2.5
( m )
Figure 7 Variation of creep strength in region L2 with average lattice diffusion distance of Fe atoms during normalizing. (a) (b) Creep rupture strength and (c) creep deformation resistance.
24
360 o
(a) 500 C, 10kh 320 280 240 o
(b) 550 C, 10kh 220 200 180
Creep Strength ( MPa )
o
140
(c) 600 C, 10kh
120 100 o
(d) 550 C, 200kh
Open: Tube Solid: Plate
150 120 100 o
(e) 600 C, 100kh 80 60 100 80 60 o
−7 −1
(f) 600 C, 10 h 40 1040
1050
1060 o
Normalizing Temperature ( C )
Figure 8 Correlation between creep strength and normalizing temperature. Creep rupture strength in regions (a) (b) H, (c) M, and (d) (e) L2, and (f) creep deformation resistance in region L2.
25
180 o
(a) 550 C, 200kh 150
Creep Strength ( MPa )
120 100 o
(b) 600 C, 100kh 80 60
Open: Tube Solid: Plate
100 o
−7 −1
(c) 600 C, 10 h 80 60 40 8
10
12
14
16
18
Grain Size ( m )
Figure 9 Effect of grain size on creep strength in region L2. (a) (b) Creep rupture strength and (c) creep deformation resistance.
26
360 o
(a) 500 C, 10kh 320 280 240 o
(b) 550 C, 10kh 220 200 180 o
Creep Strength ( MPa )
140
(c) 600 C, 10kh
120 100 o
(d) 550 C, 200kh
Open: Tube Solid: Plate
150 120 100 o
(e) 600 C, 100kh 80 60 100 80 60
−7 −1
o
(f) 600 C, 10 h 40 8.2
8.4
8.6
8.8
Cr Concentration ( mass% )
Figure 10 Effect of Cr concentration on creep strength in regions (a) (b) H, (c) M, and (d)-(f) L2. (a)-(e) Creep rupture strength and (f) creep deformation resistance.
27
HR + 0.19(Cr% − 8) 12 360
16
20
24
o
(a) 500 C, 10kh
Creep Rupture Strength ( MPa )
320 280
Open: Tube Solid: Plate
240 o
(b) 550 C, 10kh 220 200 180 o
140
(c) 600 C, 10kh
120 100 20
25 30 HR + 18(Cr% − 8)
Figure 11 Correlation of creep rupture strength with hardness HR after tempering and Cr concentration Cr% in (a) (b) region H and (c) region M.
28
(d/m) + 2.4 (Cr% − 8) 8
12
16
20
180 o
(a) 550 C, 200kh 150
Creep Strength ( MPa )
120 100 o
(b) 600 C, 100kh 80 60
Open: Tube Solid: Plate
100 −7 −1
o
(c) 600 C, 10 h 80 60 40 15
20
25
30
35
(d/m) + 22 (Cr% − 8)
Figure 12 Correlation of creep strength with grain size d and Cr concentration Cr% in region L2. (a) (b) Creep rupture strength and (c) creep deformation resistance.
Table 1
Chemical concentrations (mass%) of major elements in each heat studied. T and P stand for
tube and plate products, respectively. Heat
Cr
Mo
V
Nb
C
N
Al
Ni
Mn
MGA(T)
8.53
0.96
0.21
0.076
0.10
0.050
0.014
0.12
0.40
29
MGB(T)
8.51
0.90
0.205
0.076
0.09
0.042
0.016
0.20
0.45
MGC(T)
8.70
0.90
0.22
0.072
0.09
0.044
0.001
0.28
0.35
MGD(T)
8.41
0.90
0.185
0.07
0.10
0.048
0.016
0.10
0.41
MGF(T)
8.41
0.91
0.20
0.08
0.11
0.053
0.001
0.06
0.42
MGG(T)
8.60
0.95
0.190
0.08
0.10
0.0458
0.002
0.12
0.37
MgA(P)
8.34
0.89
0.23
0.070
0.08
0.059
0.012
0.09
0.49
MgB(P)
8.34
0.89
0.23
0.070
0.08
0.059
0.012
0.09
0.49
MgC(P)
8.74
0.94
0.21
0.076
0.10
0.0582
0.014
0.04
0.44
MgD(P)
8.44
0.99
0.21
0.09
0.11
0.050
0.023
0.08
0.45
JAA(P)
8.73
0.96
0.22
0.09
0.10
0.051
0.013
0.07
0.43
JAB(P)
8.75
0.97
0.21
0.089
0.10
0.0505
0.011
0.06
0.42
JAC(P)
8.76
0.94
0.21
0.08
0.09
0.0536
0.011
0.04
0.44
JAD(P)
8.76
0.94
0.21
0.08
0.09
0.0536
0.011
0.04
0.44
Table 2
Normalizing, tempering and post weld heat treatment (PWHT) conditions of each heat studied.
Diffusion distances X of Fe atoms during normalizing (in phase) and during tempering + PWHT (in phase) are listed together with its prior grain size (d) and Rockwell C hardness (HR) measured after the treatments. Normalizing
X in
d
Tempering
PWHT
o
min
m
m
o
C
min
o
MGA
1045
10
0.57
10.7
780
MGB
1050
60
1.5
14.
MGC
1050
10
0.60
MGD
1050
10
MGF
1045
MGG
Heat
X in
HR
C
min
m
HRC
60
-
-
0.84
16.
760
60
-
-
0.64
16.
10.3
765
30
-
-
0.49
18.
0.60
10.3
780
40
-
-
0.69
18.
60
1.4
9.0
780
60
-
-
0.84
18.
1050
15
0.74
8.7
790
60
-
-
0.97
18.
MgA
1050
10
0.60
9.6
770
60
740
500
2.1
13.
MgB
1050
10
0.60
9.6
770
60
740
60
1.2
14.
MgC
1060
90
2.0
14.
760
60
730
500
1.8
17.
MgD
1050
30
1.0
17.
780
30
-
-
0.60
16.
JAA
1050
50
1.3
21.
780
60
740
500
2.2
14.6
JAB
1050
65
1.5
39.
780
85
740
640
2.6
13.0
JAC
1060
60
1.6
17.
760
60
740
500
2.0
15.6
C
30
JAD
1060
60
1.6
17.
760
60
-
Table 3
Stress exponents for creep rupture life and minimum creep rate in each region. The values of
each heat are listed along with weighted average values of the whole heats. ̇ m (L2)
Heat
tr (H)
tr (M)
tr (L2)
MGA
16.5
9.06
3.01
2.26
MGB
16.7
8.89
2.18
2.54
MGC
19.7
8.88
2.58
4.38
MGD
19.4
7.53
-
5.05
MGF
19.0
8.23
-
4.94
MGG
14.3
7.99
2.94
4.49
MgA
18.5
8.47
2.00
5.83
MgB
15.5
10.1
2.13
5.96
MgC
20.2
9.74
2.09
6.44
MgD
16.0
9.03
-
-
JAA
14.5
9.23
-
-
JAB
13.6
9.53
-
-
JAC
-
-
-
-
JAD
-
-
-
-
17.1
8.81
2.48
4.64
Average
31
-
0.64
22.5
PII: DOI: Reference:
S0921-5093(17)30503-8 http://dx.doi.org/10.1016/j.msea.2017.04.050 MSA34951
To appear in: Materials Science & Engineering A Received date: 9 March 2017 Revised date: 12 April 2017 Accepted date: 12 April 2017 Cite this article as: K. Maruyama, J. Nakamura, N. Sekido and K. Yoshimi, Causes of heat-to-heat variation of creep strength in grade 91 steel, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2017.04.050 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Causes of heat-to-heat variation of creep strength in grade 91 steel K. Maruyama*, J. Nakamura, N. Sekido, K. Yoshimi Department of Materials Science, Tohoku University *
Corresponding author: 6-6-02 Aramaki-Aoba, Aoba-ku, Sendai 980-8579, Japan.
[email protected]
Abstract Heat-to-heat variation of creep strength is significant in grade 91 (Gr.91) steel, but the causes of the variation have not been well understood yet. In the present paper, creep rupture data of 14 heats of Gr.91 steel were analyzed paying attention to their chemical compositions and microstructures. The longest creep rupture lives analyzed are 2 105h at 500 and 550oC and 105h at 600oC. The causes of the heat-to-heat variation are different, depending on creep test conditions. At low temperature and high stress (creep rupture life of 104h at 500 and 550oC), creep rupture strength increases with increase of hardness after tempering. This suggests strengthening by a fine subgrain microstructure developed during normalizing and subsequent tempering. At higher temperature and intermediate time range (10 4h at 600oC), creep rupture strength depends on Cr concentration of the heats in addition to the hardness. This finding suggests an important contribution of recovery process of the subgrain microstructures to creep strength of the steel. In long-term creep (2 105h at 550oC and 105h at 600oC) creep rupture strength primarily increases with increasing grain size of the heats. This suggests that grain boundary sliding is an important deformation mode at low strain rate because of fine grain size usual with Gr.91 steel. Specifications on Ni concentration and N%/Al% ratio are newly introduced in the type II version of Gr.91 steel. They are not effective to eliminate a heat with low creep strength. Keywords: Grade 91 steel; Creep strength; Strengthening mechanism; Grain size effect; Creep life evaluation; Multi-region modelling
1. Introduction Creep strength enhanced ferritic (CSEF) steels are used for high temperature components in
1
ultra-super-critical steam boilers of electric power generation plants. Grade 91 (Gr.91) steel is most widely utilized among the CSEF steels. The components are designed to be used for 105h at 600oC, but creep defects are sometimes found in weldments of Gr.91 steel before the designed service duration [1]. Electric Power Research Institute (EPRI) has proposed a type II version of Gr.91 steel for preventing the premature creep failure [1,2]. ASME Boiler and Pressure Vessel Code has endorsed the type II version as a code case. In the type II version, allowable concentration of residual nickel is reduced from 0.4 to 0.2%, and upper limits of tramp elements and a lower limit of N%/Al% ratio are newly introduced. All the chemical concentrations in the present paper are in mass% otherwise indicated. The proposals are based mostly on effects of alloying elements on creep ductility and creep cracking. The EPRI report [1] mentions that creep rupture life and the consequent creep rupture strength increase with increasing creep ductility, but does not provide sufficient experimental evidence on the influence of ductility. Although normalizing temperature and duration have significant effects on creep rupture strength [2], the normalizing conditions of the type II version remain unchanged from the current ASME Boiler and Pressure Vessel Code. It is known that creep rupture strength of CSEF steels varies with Cr concentration [3-5], but the wide allowable range of Cr concentration, from 8 to 9.5%, is accepted in the type II version. Kimura et al. [6] have reported that 105h creep rupture strength of Gr.91 steel at 600oC increases monotonously with decrease in Ni concentration from 0.28 to 0.04%. The reduction of allowable Ni concentration in the type II version is partly based on this finding. They [6] took only four heats of Gr.91 steel in their examination, but six other heats are added to the original four heats in a recent examination [7]. From the recent examination, any correlation was not found between creep rupture strength and Ni concentration in the range less than 0.3%Ni. Kimura and Yaguchi [8] have confirmed absence of the Ni concentration dependence in their recent paper. Masuyama and Yamaguchi [9] and Komai et al. [10] have examined correlations between alloy concentrations and creep rupture strength of Gr.91 steel. They have suggested several alloying elements that may affect the creep rupture strength, but did not draw any quantitative conclusion on the correlations from their examinations. Abe [11] has studied heat-to-heat variation of creep rupture strength of several ferritic steels. He thought that formation of fine MX type precipitates is crucial for getting high creep rupture strength of the ferritic steels, and
2
paid attention to the free nitrogen concentration given by, (N concentration
Al concentration
Ti
concentration), since aluminum and titanium form thermally stable nitrides and consume N atoms necessary for forming MX precipitates. He has found that increase in the free nitrogen concentration improves creep rupture strength of 12Cr and 12Cr-1Mo-1W-0.3V steels. However, Gr.91 is not included in his study. The causes of heat-to-heat variation of creep strength in Gr.91 steel have not been well understood yet, though the variation is significant in the steel. In the present paper, creep strengths (creep rupture strength and creep deformation resistance) of 14 heats of Gr.91 steel are studied to clarify the causes of heat-to-heat variation of their creep strength and to assess the type II proposal. In creep of Gr.91 steel there are four regions with different creep behavior, in other words different stress exponents and activation energies for creep rupture life [7,12]. This fact suggests changes in dominant strengthening mechanism among the four regions. Therefore, causes of the variation are examined separately in each region. The examination in the multi-regions is different from the previous studies [1,2,6,9-11] on the heat-to-heat variation.
2. Materials studied The present paper studies creep strength of 6 heats of tube products and 8 heats of plate products of Gr.91 steel. Chemical concentrations of the major elements in each heat are listed in Table 1. Table 2 gives heat treatments and diffusion distances of Fe atoms during the normalizing and tempering of each heat along with its prior grain size and hardness measured after the treatments. The average distance X of diffusion is given by the following equation:
X 2Dt
(1)
where D is the diffusion coefficient, and t is the diffusion duration. The following values of the pre-exponential coefficient D0 and the activation energy QD for lattice self-diffusion of Fe atoms [13] are used for evaluating X: D0 = 2.8 10-4 m2/s, QD = 251 kJ/mol
(in phase)
D0 = 4.9 10-5 m2/s, QD = 284 kJ/mol
(in phase)
The heat treatment conditions given in Table 2 are used for the temperature and duration of diffusion.
3
Since diffusion of atoms occurs during heating before reaching the heat treatment temperatures, actual distances of diffusion may be slightly longer than the values listed in Table 2. The prior grain size is calculated from a grain size number assuming square shaped grains. Minimum creep rates and creep rupture lives of heats MGA, MGB, MGC, MGD, MGF, MGG, MgA, MgB, MgC and MgD have been reported in NIMS Creep Data Sheet 43A [14]. MGA to MGG are tube products, and MgA to MgD, plate products. Those data on plate products of heats JAA, JAB, JAC and JAD have been obtained by Japan Atomic Energy Agency [15]. Creep specimens of plate products were cut parallel to the rolling direction from center of the plates. The longest creep rupture lives of heats JAC and JAD at 500 and 550oC (close to 2 105h) were evaluated by analyzing their ongoing creep curves before fracture [12].
3. Analyses of creep data 3.1 Multi-region analysis Minimum Creep rate ̇ m and creep rupture life tr are represented by the following equations as functions of stress and temperature T: ̇ m = ̇ 0 n exp ( Q/R T)
(2)
tr = t0 -n exp (Q/R T)
(3)
where ̇ 0 and t0 are material constants, n is the stress exponent, Q is the activation energy, and R is the gas constant. Values of n and Q can change with creep test conditions. The changes in n and Q are often accompanied by a change in creep deformation or fracture mechanism. Stress dependence of creep rupture life in Gr.91 steel is shown in Fig.1. Heat MGC is taken as an example. The figure has originally been reported in Refs.[7,12]. The lines drawn in the figure are regression lines based on Eq.(3), and Refs.[7,12] explain how to determine the lines. As evident in the figure, there are four regions, H, M, L1 and L2 with different values of n. The appearance of the four regions is common to CSEF steels [7,16,17], and the four regions should appear in all the 14 heats of the present study. However, regions L1 an L2 are missing in some heats due to lack of high temperature or long-term data points. Stress exponents and activation energies in Fig.1 are as follows: n = 20 and Q =
4
718kJ/mol in region H, 8.9 and 606kJ/mol in M, 5.8 and 619kJ/mol in L1, and 2.6 and 232kJ/mol in L2. Figure 1 includes data points of heats JAC and JAD. The two heats are the same material, and were normalized at a higher temperature for a longer duration than those of heat MGC. This heat treatment makes their prior grain sizes of 17m, which is greater than 10.3m of heat MGC. Heat JAD was tempered for 60min at 760oC, and has Rockwell C hardness of HR = 22.5 (Vickers hardness of Hv = 252). Heat JAC was subjected to additional annealing for 500min at 740oC simulating post weld heat treatment (PWHT). This treatment gives its hardness of HR = 15.6 (Hv = 219) lower than HR = 18 of heat MGC. The creep rupture lives in region H (low temperature and high stress) increase with increasing the hardness of the heats. In long-term region L2, the creep rupture lives of heat JAC and JAD are longer than that of heat MGC due to their larger prior grain sizes. Contrary to these regions, the effects of hardness and prior grain size are indistinct in regions M and L1. The creep strengthening mechanisms are obviously different among regions H, M and L2. This finding points out that alloying elements and microstructures play different roles in strengthening among the three regions. Therefore, the causes of heat-to-heat variation in creep strength should be discussed separately in the three regions.
3.2 Evaluation of creep strength Creep data of each heat was subjected to the multi-region creep data analysis [7,16,17], like the example of heat MGC (see Fig.1), in order to specify the region each data point belongs to. Data points used in the present analyses are given in Fig.2: (a) creep rupture lives in regions H and M, (b) creep rupture lives in region L2, and (c) minimum creep rates in region L2. The multi-region creep data analyses give stress exponent n in each region of the heat. Their values are listed in Table3. All the data points of a heat in a region were used for determining the n values in addition to those plotted in Fig.2. n values are not listed in the table, when the heat does not have a sufficient number of data points to determine n in the region. As evident in the table, the scatter of n values within each region is not significant among the heats. A weighted average value of n is calculated in each region by taking a number of data points of each heat into account. The value is listed in the last line of the table. The lines drawn in Fig.2 are a regression line with the average n value and passing through the
5
center of the data points of each heat. The regression lines can represent stress dependence of creep rupture life and minimum creep rate of most of the heats, though the deviations of data points from the lines are not negligible in some heats. The regression lines are extrapolated or interpolated to the following conditions for getting creep rupture strength and creep deformation resistance under the conditions: 500 oC – 104 h, 550 oC – 104 h
(region H)
600 oC – 104 h
(region M)
550 oC – 2 x 105 h, 600 oC – 105 h, 600 oC – 10-7 h-1
(region L2)
The creep rupture lives vary by 40 times in region H, 5 times in region M and 3 times in region L2. The minimum value (59MPa) of 105h creep rupture strength at 600oC is two-thirds of its maximum value (91MPa) (see Fig.2(b)). The variation of the whole data points is wide, but the variation within a heat is narrow. The wide variation is primary caused by heat-to-heat variation in creep strength. The causes of the heat-to-heat variation is discussed in the next section on the basis of the creep strengths obtained in Fig.2.
4. Discussion 4.1 Factors not affecting creep strength Some heats of Gr.91 steel have substantially lower creep rupture strength than the values listed in ASME Code and ECCC Code. EPRI ascribes the low creep strength to excess amounts of harmful elements included in the weak heats [1]. To eliminate the low strength heats, EPRI has proposed the type II version of Gr.91 steel [1,2]. In this version, allowable Ni concentration is reduced from 0.4 to 0.2%, and N%/Al% ratio greater than 4 (in mass%) is newly introduced. Let us examine whether the revised specifications really affect creep strength of Gr.91 steel. Creep strengths of all the heats are plotted against residual Ni concentration in Fig.3. Open and solid symbols are tube and plate products, respectively. The error bar of each heat gives its largest upward and downward deviations of data points from its regression line in Fig.2. No error bar is attached to a data point when a heat has only one data point in a region at a temperature. At 550oC (see Fig.3(d)), for
6
example, every heat has only one data point in region L2 except for heat MGC. As evident in the figure, most error bars are substantially short when compared to the heat-to-heat variation in creep strength. Kimura at el. [6] have pointed out that creep rupture strength increases with decreasing Ni concentration, based on the creep rupture strength of heats MgC (0.04%Ni), MGA (0.12%Ni), MGB (0.20%Ni) and MGC (0.28%Ni). Figure 3 includes data points of 10 other heats in addition to the 4 heats. It is obvious that creep strength is independent of Ni concentration in all the regions H, M and L2. This finding accords with recent reports [7,8] on the effects of residual Ni concentration. Nitrogen forms MX type precipitates with vanadium and niobium. The precipitates are fine and thermally stable, and necessary to keep a fine subgrain microstructure that provides high creep strength of CSEF steels. Aluminum atoms have high affinity to N atoms, and form coarse AlN inclusions before the formation of MX precipitates. After the same service duration, weldments with N%/Al% ratio (in mass%) less than 4 have a larger number of creep cracks than those with a high N%/Al% ratio [1]. Based on this finding on Gr.91 steel, EPRI [1] recommends N%/Al% > 4 (N%/Al% > 7.7 in atomic ration) in the type II version. Brett [18] has also pointed out the importance of N%/Al% ratio for achieving high creep ductility of the steel. The creep strengths of the 14 heats are plotted against N%/Al% ratio (in atomic%) in Fig.4. The ratio varies from 4 (less than 7.7) to 100. Any systematic change of the creep strength with N%/Al% ratio cannot be seen despite the wide variation in the ratio. A recent report [8] supports this result. The result can be explained as follows on the basis of the literature [11]. The ratio is not necessary to be greater than 7.7 (in atomic ratio), if N%>Al% and free nitrogen concentration is large enough to form a sufficient number of fine MX precipitates. Correlations of the creep strength with Mo, Nb, V, Mn, Al and (C+N) concentrations are also studied. However, any clear correlations are not found on these variables since their concentrations are controlled within the ranges in which the elements do not cause variation of creep strength. It can be concluded that the variations in Ni, Mo, Nb, V, Mn, Al and (C+N) concentrations and N%/Al% ratio are not responsible for the variation in creep strength of Gr.91 steel. However, it should be recalled that 105h creep rupture strength at 600oC (Fig.4(e)) varies from 59 to 91 MPa. When the average creep rupture strength of Gr.91 steel is 90MPa [19], the creep rupture strength of the weakest heat is as low as an
7
allowable stress of Gr.91 steel in creep regime.
4.2 Effects of heat treatments on creep strength The creep rupture strengths of Gr.91 steel are plotted against diffusion distance during tempering (including the treatment simulating PWHT) in Fig.5. The creep rupture strength in region H decreases with increasing the diffusion distance, but the decrease is not evident in regions M (see Fig.5(c)) and L2. CSEF steels are strengthened by their fine subgrain microstructures [20,21]. Increase of diffusion distance during tempering results in coarsening (recovery) of subgrains, and their hardness after tempering is a measure of the recovery process. On the basis of these understandings, the creep strengths of the 14 heats are plotted in Fig.6 against their hardness. A better correlation is found in this figure than that in Fig.5, confirming the dominant role of subgrain microstructures in determining the creep strength in regions H and M. The slopes of the regression lines in Figs.6(a)-(c) decrease from 5.4MPa/HRC at 500oC to 3.8 at 550oC and further to 1.3 at 600oC. The effect of hardness becomes less significant with increasing creep testing temperature. Komai et al [22] have studied the effects of hardness after tempering on creep rupture strength of Gr.91 steel. Their result is very similar to the one obtained in the present study. The heat with the highest hardness in region L2 (see Figs.6(d)-(f)) has the high creep strength due to its large grain size. If this heat is omitted, the creep strength in region L2 is insensitive to the hardness. The creep strengths in region L2 are plotted in Fig.7 against average diffusion distance of Fe atoms during normalizing. The creep strength increases with increasing the diffusion distance contrary to the insensitiveness of creep strength in Figs.6(d)-(f). This increase is widely known, but the cause of strength increase has not been fully clarified yet [2]. It is to be noted that 104h creep rupture strength in regions H and M is independent of the diffusion distance. Coarsening of grains and dissolution of NbX precipitates proceed during the normalizing [2]. Dissolution of a larger amount of NbX gives a greater number of fine NbX precipitates formed during tempering and creep, and consequently can improve creep strength. Diffusion of Nb atoms is fast enough at temperatures above 1000oC, and Nb concentration around NbX reaches an equilibrium within a short
8
period of time. Therefore, the amount of NbX dissolved into matrix after normalizing is determined not by the diffusion distance in Fig.7 but by normalizing temperature [2,23]. This fact expects a close correlation between the creep strength and normalizing temperature, if the fine NbX precipitates determine the creep strength. The creep strengths of the 14 heats are plotted against normalizing temperature in Fig.8. As evident in Fig.8(a)-(c) the 104h creep rupture strength in regions H and M is independent of the normalizing temperature. As seen in Figs.8(d)-(f), the creep strength in region L2 increases with increasing normalizing temperature. This fact does not conflict with the strengthening due to the precipitation of NbX, but the scatter in data points is large in the figures, suggesting the presence of some other factor having a more significant effect on the creep strength. The other factor varied by the normalizing, namely grain size increases with increasing temperature and duration of normalizing [2]. The creep strengths in region L2 are plotted against prior grain size in Fig.9. Since a prior grain is divided into several packets after martensitic transformation, packet size should be used as a measure of grain size. It has been reported on low carbon steels that packet size is proportional to prior grain size and the proportional coefficient of 0.4 is independent of alloy addition [24]. Let us use the prior grain size instead of packet (grain) size in the present analysis, since packet sizes have not been reported on the 14 heats. The creep strength increases with increasing grain size in Fig.9. The scatter in data points in the figure is smaller than that in Fig.8. It follows from this fact that the increase in grain size is the major cause of the dependence of creep strength on the diffusion distance in Fig.7. It should be noted that the prior grain sizes are from 9 to 17m in Fig.9 and very small. Packet sizes are estimated to be 3.5 to 7m. Deformation modes related to grain boundary, such as grain boundary sliding and diffusion creep can contribute to creep strain in this range of grain size [25], and contributions of the deformation modes increase with decreasing grain size. Creep rupture life is related to packet size dp by the following equation in region L2 [25]: tr = to -2.5 dpP D-1
(4)
where dp is the packet size, and P is the packet size exponent. P = 2 and 3 in grain boundary sliding controlled by lattice diffusion and grain boundary diffusion, respectively [25]. The increase in packet size extends creep rupture life by four times when P = 2. This increase is similar to the heat-to-heat variation
9
in creep rupture life observed in Fig.2(b). It is to be noted that the 104h creep rupture strengths in regions H and M are independent of grain size.
4.3 Effects of Cr concentration on creep strength Among the elements listed in Table1, chromium is the only element that has some correlation with the creep strength. The creep strengths of the 14 heats are plotted against Cr concentration in Fig.10. The creep rupture strength (Figs.10(d) and (e)) and creep deformation resistance (Fig.10(f)) in region L2 increase with increasing Cr concentration. This is true for the 104h creep rupture strength in region M (Fig.10(c)). However, the scatter in data points is large in Figs.10(d) and (e) due to the variation in grain size. The same effect of Cr is not evident in region H (Figs.10(a) and (b)) in which diffusion distance is too short for recovery of subgrain microstructures to proceed. It can be concluded that creep strength of Gr.91 steel varies with Cr concentration in the range allowed in ASME Code (8 to 9.5%). Murata
and
his
co-workers
[3,4]
have
studied
creep
rupture
life
of
X%Cr-3.5%W-3%Co-VNbBN steels. The steels were normalized at 1070oC for 5h, and then double-tempered at 570oC for 20h and 680oC for 20h. They varied Cr concentration from X = 8.5 to 11.5% and carried out creep tests at 650oC. They have reported that the steels give the longest rupture life at 10%Cr in the rupture life range around 104h, at 9%Cr in the range around 3 104h and at 8.5%Cr in the range over 5 104h. Gasemi-Armaki et al [5] have studied creep of X%Cr-1.9%W-0.4%Mo-0.9%Cu-VNb steel (X = 5, 7, 9, 10.5 and 12%). The steels were normalized at 1050oC for 30min and then tempered at 770oC for 2h. Creep rupture lives of these steels increase with increasing Cr concentration in short-term creep, but the 9%Cr steel gives the longest rupture life in long-term creep. Creep rupture life of CSEF steels increases with decreasing subgrain width. Subgrain width after the tempering decreases with increasing Cr concentration, resulting in the longer creep rupture life of the higher Cr steel in short-term creep [5]. Ostwald ripening of M23C6 particles proceeds during creep, resulting in coarsening of subgrain width [21]. The ripening rate is faster at higher Cr concentration. Therefore, Cr concentration that gives the minimum subgrain width varies with the time ranges [5,26]. The 9%Cr steel has the smallest subgrain width after 104h annealing [5], resulting in its longest rupture life in the long-term creep. The Cr
10
concentration dependent creep behavior in Fig.10 is ascribed to the Cr concentration dependence of thermal stability of subgrain microstructures. This effect of chromium is especially important in Figs.10(c) and (f).
4.4 Quantitative analyses of creep strength The aforementioned analyses indicate that the hardness after tempering, the grain size and the Cr concentration are the major factors determining the creep strength. In more detail, the hardness has no effect in region L2, and the creep rupture strength is independent of grain size in regions H and M. Chromium concentration may not affect the creep rupture strength in region H. Let us assume that creep strength S is given by the following equations of the hardness HR, the Cr concentration Cr% and the prior grain size d: S = a (HR + b Cr%) + S1
(H and M)
(5)
S = c (d + e Cr%) +S2
(L2)
(6)
where a, b, c, e, S1 and S2 are constants. They are determined so that they give the best fit of the regression lines to the data points in Figs.11 and 12. Equations (5) and (6) include the two major variables only, but there may be some other variables that contribute to the variation in creep strength. The creep rupture strengths in regions H and M are plotted against (HR + b Cr%) in Fig.11, and those in region L2, against (d + e Cr%) in Fig.12. The values of b giving the best fit are slightly different in Figs.11(a) and (b). The average of these values is finally used in the two figures. A similar average value of e is used also in Figs.12(a) and (b). Good correlations hold in both figures, confirming the dominant roles of the hardness, the grain size and the Cr concentration in controlling creep strength of Gr.91 steel. In region H (Figs.11(a) and (b)), 99% of the change in abscissa is due to the hardness. The effect of Cr concentration on the creep rupture strength is negligible, since thermal recovery of subgrain microstructures does not proceeds in region H. In region M (Fig,11(c)), 60% of the change in abscissa comes from that in Cr concentration, and the other 40%, from the change in hardness. Thermal recovery of subgrain microstructures that are affected by Cr concentration contributes to the creep rupture strength in this region.
11
As evident in Figs.2(b) and (c), region L2 appears under the same creep test conditions both in creep rupture life and minimum creep rate. The rupture lives are 3 104 to 105h in region L2 at 600oC. The minimum creep rate appears at 10 to 20% of the rupture lives in region L2 [12]; namely it appears before 2 104h. Because of the early appearance of the minimum creep rate, the effect of Cr concentration is still substantial in Fig.12(c) as is the case in Fig.11(c). In Fig.12(c), 50% of the change in abscissa is due to Cr concentration and the other 50%, due to grain size. As shown in Figs.12(a) and (b), the 105h creep rupture strength at 600oC and the 2 105h creep rupture strength at 550oC increase with increasing the grain size and Cr concentration. The contribution of Cr concentration to the change in abscissa is 10%, and the grain coarsening is the dominant cause of the increase of creep rupture strength in region L2. The origin of the grain size dependence in creep strength is the contribution of grain boundary related deformation modes, probably grain boundary sliding to creep strain. The contribution increases with decreasing grain size. In region L2, stress exponent for creep rupture life is 2.5 (cf. 3 or greater in dislocation creep), and the activation energy for creep rupture life (230kJ/mol) is less than that for lattice self-diffusion (300kJ/mol [27]) rate-controlling dislocation creep. Stress exponents for creep rupture life are similar among CSEF steels in region L2, their welded joints creep-ruptured by type IV cracking in fine grain HAZ [28,29] and a material simulating fine grain HAZ [30]. These findings can be valid if grain boundary sliding is the dominant creep deformation mode in region L2. Heat MGG is the weakest in Fig.12(a); 0.12%Ni and N%/Al% = 23 (in mass%). In Fig.12(b) heat MgB is the weakest; 0.09%Ni and N%/Al% = 4.9 (in mass%). Both heats contain nickel less than 0.2% and have the ratios greater than 4. They fulfill the type II specifications on Ni concentration and N%/Al% ratio as well as on other elements, though concentrations of tramp elements have not been reported in the literature [14,15]. The type II specifications of Gr.91 steel cannot exclude the heats with low creep strength.
5. Summary 1) In creep of Gr.91 steel, there are four regions H, M, L1 and L2 with different values of stress exponent
12
n for creep rupture life and minimum creep rate. This fact is common to all the heats of the steel. The values of n in each region are similar in all the heats. 2) Concentrations of Mo, V, Nb, C+N, Ni, Mn and Al in Gr.91 steel are controlled within the ranges in which the elements do not cause variation of creep strength. Creep strength of the steel is independent of N%/Al% ratio when N%/Al% (in atomic%) is greater than 4. 3) Creep rupture strength of the steel depends primarily on its hardness after tempering in low-temperature and high-stress region H. It increases with increasing the hardness. 4) Creep rupture strength of the steel depends primarily on its grain size in long-term region L2, though increase in Cr concentration slightly contributes to strengthening. It increases with increasing grain size and Cr concentration. 5) Creep deformation resistance that gives minimum creep rate of 10-7h-1 at 600oC (region L2) increases with increasing Cr concentration and grain size of the steel. 104h creep rupture strength of the steel at 600oC (region M) increases with increasing its Cr concentration and hardness after tempering. The two variables equally contribute to the strengthening. 6) The type II specifications of Gr.91 steel cannot eliminate the weakest heats of the present study.
Acknowledgements A part of the present study was supported by an ALCA project, and the other part, by an NEDO project #1400070-0. Financial supports from Japan Science and Technology Agency (ALCA project) and New Energy and Industrial Technology Development Organization (NEDO project) are acknowledged. The author cannot accomplish the present analyses without the high quality data reported in NIMS Creep Data Sheet No.43A and JAEA-Data/Code 2008-030. The authors deeply acknowledge National Institute for Materials Science and Japan Atomic Energy Agency for their publication of the creep data sheets. Very useful suggestions on packet size were given by Prof. M. Mitsuhara, Department of Engineering Science for Electronics and Materials, Kyushu University.
13
References [1] EPRI Report No. 3002003472, The benefits of improved control of composition of creep-strength-enhanced ferritic steel grade 91, EPRI, Palo Alto, CA, 2014. [2] EPRI Report No. 3002004370, The influence of steel making and processing variables on the microstructure and properties of creep strength enhanced ferritic (CSEF) steel grade 91, EPRI, Palo Alto, CA, 2014. [3] Y. Murata, K. Yamashita, M. Morinaga, T. Hara, K. Miki, T. Azuma, T. Ishiguro, R. Hashizume, Dependence of precipitation behavior and creep strength on Cr content in high Cr ferritic heat resistant steels, J. Solid Mech. Mater. Eng. 3 (2009) 457-463. [4] K. Miki, T. Azuma, T. Ishiguro, O. Tamura, R. Hashizume, Y. Murata, M. Morinaga, Development of heat resistant ferritic steels containing 8.5–11.5% Cr based on long-term creep rupture strength, in: Proceedings of the 18th International Forgemasters Meeting, Pittsburg, PA, September 12-15, 2011, pp. 211-214. [5] H. Ghassemi-Armaki, R. Chen, K. Maruyama, M. Igarashi, Creep behavior and degradation of subgrain structures pinned by nanoscale precipitates in strength-enhanced 5 to 12 pct Cr ferritic steels, Metall. Mater. Trans. A 42 (2011) 3084-3094. [6] K. Kimura, K. Sawada, H. Kushima, Y. Toda, Influence of chemical composition and heat treatment on long-term creep strength of grade 91 steel, Procedia Eng. 55 (2013) 2-9. [7] K. Maruyama, J. Nakamura, K. Yoshimi, Assessment of long-term creep rupture strength of T91 steel by multiregion rupture data analysis, Trans. ASME, J. Pressure Vessel Technol. 138 (2016) 031407-1-9. [8] K. Kimura, M. Yaguchi, Re-evaluation of long-term creep strength of base metal of ASME grade 91 type steel, in: ASME 2016 Pressure Vessels and Piping Conference, 2016, Paper No. PVP2016-63355. [9] F. Masuyama, T. Yamaguchi, Chemical specification and design factor consideration for creep degradation in grade 91 steel, in: ASME 2016 Pressure Vessels and Piping Conference, 2016, Paper No. PVP2016-63307. [10] N. Komai, K. Arisue, N. Saito, K. Tominaga, M. Fujita, Effect of chemical composition on the creep
14
rupture strength of mod.9Cr-1Mo steel weldments, in: Proc. of the 1st International Conference on Advanced High-temperature Materials Technology for Sustainable and Reliable Power Engineering, Sapporo, Japan, June 29-July 3, 2015, pp. 57-60. [11] F. Abe, Heat-to heat variation in long-term creep strength of some ferritic steels, Intern. J. Press. Vess. Piping 87 (2010) 310-318. [12] K. Maruyama, J. Nakamura, K. Yoshimi, Y. Nagae, Evaluation of long-term creep rupture life of Gr.91 steel by analysis of on-going creep curves, in: J. Parker, J. Shingledecker, J. Siefert (Eds.), Advances in Materials Technology for Fossil Power Plants, ASM Intern., Materials Park, OH, 2016, pp. 467-478. [13] Metal Data Book, fourth ed., Edited by Japan Inst. Metals, Maruzen, Tokyo, 2004, pp. 20-25. [14] National Institute for Materials Science, Data sheets of the elevated-temperature properties of 9Cr-1Mo-V-Nb steel tubes for boilers and heat exchangers, 9Cr-1Mo-V-Nb steel plates for boilers and pressure vessels and 9Cr-1Mo-V-Nb seamless pipe for high temperature service, NIMS Creep Data Sheet No. 43A, NIMS, Tsukuba, Japan, 2014,. [15] S. Kato, T. Furukawa, E. Yoshida, Material test data of mod. 9Cr-1Mo steel (1), JAEA-Data/Code 2008-030, Japan Atomic Energy Agency, Tokai-mura, Japan, 2009. [16] K. Maruyama, K. Yoshimi, Influence of data analysis method and allowable stress criterion on allowable stress of Gr.122 heat resistant steel, Trans. ASME, J. Press. Vess. Technol. 129 (2007) 449-453. [17] K. Maruyama, J. Nakamura, K. Yoshimi, Prediction of long-term creep rupture life of grade 122 steel by multiregion analysis, Trans. ASME, J. Pressure Vessel Technol. 137 (2015) 021403-1-5. [18] S.J. Brett, UK experience with modified 9%Cr (grade 91) steel, Energy Mater., 2 (2007) 117-121. [19] W. Bendick, L. Cipolla, J. Gabrel, J. Hald, New ECCC assessment of creep rupture strength for steel grade X10CrMoVNb9-1 (Grade 91), in: I.A. Shibli, S.R. Holdsworth (Eds.), Creep and Fracture in High Temperature Components, DEStech Publications, Lancaster, PA, 2009, pp. 56-67. [20] K. Maruyama, H. Ghassemi Armaki, R.P. Chen, K. Yoshimi, M. Yoshizawa, M. Igarashi, Cr concentration dependence of overestimation of long term creep life in strength enhanced high Cr ferritic
15
steels, Int. J. Press. Vessels Piping 87 (2010) 276-281. [21] H. Ghassemi-Armaki, R.P. Chen, K. Maruyama, M. Igarashi, Contribution of recovery mechanisms of microstructure during long-term creep of Gr.91 steels, J. Nuclear Mater., 433 (2013) 23-29. [22] N. Komai, K. Arisue, K. Tominaga, M. Fujita, Influence of holding time during tempering on the long-term creep rupture strength of mod.9Cr-1Mo steel, in: J. Parker, J. Shingledecker, J. Siefert (Eds.), Advances in Materials Technology for Fossil Power Plants, ASM Intern., Materials Park, OH, 2016, pp. 430-440. [23] Y. Tsuchida, K. Tokuno, K. Hashimoto, Development of BOF manufacture of modified 9Cr-1Mo steel plates with excellent strength and toughness, Nippon Steel Technical Report 58 (1993) 27-35. [24] S. Morito, H. Yoshida, T. Maki, X. Huang, Effect of block size on the strength of lath martensite in low carbon steels, Mater. Sci. Eng. A 438-440 (2006) 237-240. [25] J. Cadek, Creep in Metallic Materials, Elsevier, Amsterdam, 1988. [26] H. Ghassemi-Armaki, R. Chen, K. Maruyama, M. Igarashi, Premature creep failure in strength enhanced high Cr ferritic steels caused by static recovery of tempered martensite lath structures, Mater. Sci. Eng. A 527 (2010) 6581-6588. [27] H. Oikawa, Y. Iijima, Diffusion behavior of creep-resistant steels, in: F. Abe, T. Kern, R. Viswanathan (Eds.), Creep-resistant Steels, Woodhead Publishing, Cambridge, 2008, pp. 241-264. [28] K. Maruyama, K. Yoshimi, Y. Hasegawa, H. Morimoto, F. Masuyama, Temperature and stress dependence of creep life of welded joints in strength enhanced high Cr ferritic steels, Energy Mater. 4 (2009) 70-75. [29] M. Yaguchi, S. Nagai, K. Sawada, K. Kimura, Microstructure and creep strength of grade 91 steel used in USC plants, in: J. Parker, J. Shingledecker, J. Siefert (Eds.), Advances in Materials Technology for Fossil Power Plants, ASM Intern., Materials Park, OH, 2016, pp. 447-458. [30] K. Yoshida, H. Tsuruta, M. Tabuchi, K.-I. Kobayashi, Creep damage evaluation method for welded joints of grade 91 steels, in: J. Parker, J. Shingledecker, J. Siefert (Eds.), Advances in Materials Technology for Fossil Power Plants, ASM Intern., Materials Park, OH, 2016, pp. 545-556.
16
300
Stress ( MPa )
200
100 80 60 40
20 1 10
550 oC
600 oC
650 o C
H
M
70
0 oC
L1
L2
JAD MGC JAC 10
2
10
3
10
4
10
5
Rupture Life ( h )
Figure 1 A comparison of creep rupture lives among heats MGC, JAC and JAD. Effects of their hardness and grain size are different among regions H, M and L2.
17
Figure 2 Creep rupture lives in regions (a) H (500 and 550oC), M (600oC), and (b) L2 (550 and 600oC) together with (c) minimum creep rates in region L2 (600oC).
18
Figure 2 (b) and (c)
19
360 o
(a) 500 C, 10kh 320 280 240
o
(b) 550 C, 10kh 220 200 180
Creep Strength ( MPa )
o
(c) 600 C, 10kh
140 120 100
o
(d) 550 C, 200kh
Open: Tube Solid: Plate
150 120 100 80 60 o
(e) 600 C, 100kh 100 80 60 o
−7 −1
(f) 600 C, 10 h 40 0
0.1
0.2
0.3
0.4
Ni Concentration ( mass% )
Figure 3 Correlation of Ni concentration (mass%) with creep rupture strength of Gr.91 steel in regions (a) (b) H, (c) M and (d) (e) L2, and (f) with creep deformation resistance in region L2.
20
360 o
(a) 500 C, 10kh 320 280 240
o
(b) 550 C, 10kh 220 200 180
Creep Strength ( MPa )
o
140
(c) 600 C, 10kh
120 100 o
(d) 550 C, 200kh
Open: Tube Solid: Plate
150 120 100 80 60 o
(e) 600 C, 100kh 100 80 60
−7 −1
o
(f) 600 C, 10 h 40 1
10
2
10
N% / Al% (Atomic Ratio)
Figure 4 Correlation of N%/Al% ratio with (a)-(e) creep rupture strength and (f) creep deformation resistance. (a) (b) Region H, (c) region M and (d)-(f) region L2.
21
360 o
Open: Tube Solid: Plate
(a) 500 C, 10kh
Creep Rupture Strength ( MPa )
320 280 240 o
(b) 550 C, 10kh 220 200 180 o
(c) 600 C, 10kh
140 120 100 0
1
2 1/2
(2 D t)
3
( m )
Figure 5 Variation of creep rupture strength with average lattice diffusion distance of Fe atoms during tempering and PWHT. (a) (b) Region H and (c) region M.
22
360 o
(a) 500 C, 10kh 320 280 240 o
(b) 550 C, 10kh 220 200 180
Creep Strength ( MPa )
o
(c) 600 C, 10kh
140 120
Open: Tube Solid: Plate
100 o
(d) 550 C, 200kh 150 120 100
o
(e) 600 C, 100kh 80 60 100 80 60
−7 −1
o
(f) 600 C, 10 h 40 12
16
20
24
Hardness (H RC)
Figure 6 Effect of hardness after tempering on (a)-(e) creep rupture strength and on (f) creep deformation resistance. (a) (b) Region H, (c) region M, and (d)-(f) region L2.
23
180 o
(a) 550 C, 200kh 150
Creep Strength ( MPa )
120 100 o
(b) 600 C, 100kh 80 60 100 o
−7 −1
(c) 600 C, 10 h 80 60 40 0.5
Open: Tube Solid: Plate
1
1.5
(2 D t)
1/2
2
2.5
( m )
Figure 7 Variation of creep strength in region L2 with average lattice diffusion distance of Fe atoms during normalizing. (a) (b) Creep rupture strength and (c) creep deformation resistance.
24
360 o
(a) 500 C, 10kh 320 280 240 o
(b) 550 C, 10kh 220 200 180
Creep Strength ( MPa )
o
140
(c) 600 C, 10kh
120 100 o
(d) 550 C, 200kh
Open: Tube Solid: Plate
150 120 100 o
(e) 600 C, 100kh 80 60 100 80 60 o
−7 −1
(f) 600 C, 10 h 40 1040
1050
1060 o
Normalizing Temperature ( C )
Figure 8 Correlation between creep strength and normalizing temperature. Creep rupture strength in regions (a) (b) H, (c) M, and (d) (e) L2, and (f) creep deformation resistance in region L2.
25
180 o
(a) 550 C, 200kh 150
Creep Strength ( MPa )
120 100 o
(b) 600 C, 100kh 80 60
Open: Tube Solid: Plate
100 o
−7 −1
(c) 600 C, 10 h 80 60 40 8
10
12
14
16
18
Grain Size ( m )
Figure 9 Effect of grain size on creep strength in region L2. (a) (b) Creep rupture strength and (c) creep deformation resistance.
26
360 o
(a) 500 C, 10kh 320 280 240 o
(b) 550 C, 10kh 220 200 180 o
Creep Strength ( MPa )
140
(c) 600 C, 10kh
120 100 o
(d) 550 C, 200kh
Open: Tube Solid: Plate
150 120 100 o
(e) 600 C, 100kh 80 60 100 80 60
−7 −1
o
(f) 600 C, 10 h 40 8.2
8.4
8.6
8.8
Cr Concentration ( mass% )
Figure 10 Effect of Cr concentration on creep strength in regions (a) (b) H, (c) M, and (d)-(f) L2. (a)-(e) Creep rupture strength and (f) creep deformation resistance.
27
HR + 0.19(Cr% − 8) 12 360
16
20
24
o
(a) 500 C, 10kh
Creep Rupture Strength ( MPa )
320 280
Open: Tube Solid: Plate
240 o
(b) 550 C, 10kh 220 200 180 o
140
(c) 600 C, 10kh
120 100 20
25 30 HR + 18(Cr% − 8)
Figure 11 Correlation of creep rupture strength with hardness HR after tempering and Cr concentration Cr% in (a) (b) region H and (c) region M.
28
(d/m) + 2.4 (Cr% − 8) 8
12
16
20
180 o
(a) 550 C, 200kh 150
Creep Strength ( MPa )
120 100 o
(b) 600 C, 100kh 80 60
Open: Tube Solid: Plate
100 −7 −1
o
(c) 600 C, 10 h 80 60 40 15
20
25
30
35
(d/m) + 22 (Cr% − 8)
Figure 12 Correlation of creep strength with grain size d and Cr concentration Cr% in region L2. (a) (b) Creep rupture strength and (c) creep deformation resistance.
Table 1
Chemical concentrations (mass%) of major elements in each heat studied. T and P stand for
tube and plate products, respectively. Heat
Cr
Mo
V
Nb
C
N
Al
Ni
Mn
MGA(T)
8.53
0.96
0.21
0.076
0.10
0.050
0.014
0.12
0.40
29
MGB(T)
8.51
0.90
0.205
0.076
0.09
0.042
0.016
0.20
0.45
MGC(T)
8.70
0.90
0.22
0.072
0.09
0.044
0.001
0.28
0.35
MGD(T)
8.41
0.90
0.185
0.07
0.10
0.048
0.016
0.10
0.41
MGF(T)
8.41
0.91
0.20
0.08
0.11
0.053
0.001
0.06
0.42
MGG(T)
8.60
0.95
0.190
0.08
0.10
0.0458
0.002
0.12
0.37
MgA(P)
8.34
0.89
0.23
0.070
0.08
0.059
0.012
0.09
0.49
MgB(P)
8.34
0.89
0.23
0.070
0.08
0.059
0.012
0.09
0.49
MgC(P)
8.74
0.94
0.21
0.076
0.10
0.0582
0.014
0.04
0.44
MgD(P)
8.44
0.99
0.21
0.09
0.11
0.050
0.023
0.08
0.45
JAA(P)
8.73
0.96
0.22
0.09
0.10
0.051
0.013
0.07
0.43
JAB(P)
8.75
0.97
0.21
0.089
0.10
0.0505
0.011
0.06
0.42
JAC(P)
8.76
0.94
0.21
0.08
0.09
0.0536
0.011
0.04
0.44
JAD(P)
8.76
0.94
0.21
0.08
0.09
0.0536
0.011
0.04
0.44
Table 2
Normalizing, tempering and post weld heat treatment (PWHT) conditions of each heat studied.
Diffusion distances X of Fe atoms during normalizing (in phase) and during tempering + PWHT (in phase) are listed together with its prior grain size (d) and Rockwell C hardness (HR) measured after the treatments. Normalizing
X in
d
Tempering
PWHT
o
min
m
m
o
C
min
o
MGA
1045
10
0.57
10.7
780
MGB
1050
60
1.5
14.
MGC
1050
10
0.60
MGD
1050
10
MGF
1045
MGG
Heat
X in
HR
C
min
m
HRC
60
-
-
0.84
16.
760
60
-
-
0.64
16.
10.3
765
30
-
-
0.49
18.
0.60
10.3
780
40
-
-
0.69
18.
60
1.4
9.0
780
60
-
-
0.84
18.
1050
15
0.74
8.7
790
60
-
-
0.97
18.
MgA
1050
10
0.60
9.6
770
60
740
500
2.1
13.
MgB
1050
10
0.60
9.6
770
60
740
60
1.2
14.
MgC
1060
90
2.0
14.
760
60
730
500
1.8
17.
MgD
1050
30
1.0
17.
780
30
-
-
0.60
16.
JAA
1050
50
1.3
21.
780
60
740
500
2.2
14.6
JAB
1050
65
1.5
39.
780
85
740
640
2.6
13.0
JAC
1060
60
1.6
17.
760
60
740
500
2.0
15.6
C
30
JAD
1060
60
1.6
17.
760
60
-
Table 3
Stress exponents for creep rupture life and minimum creep rate in each region. The values of
each heat are listed along with weighted average values of the whole heats. ̇ m (L2)
Heat
tr (H)
tr (M)
tr (L2)
MGA
16.5
9.06
3.01
2.26
MGB
16.7
8.89
2.18
2.54
MGC
19.7
8.88
2.58
4.38
MGD
19.4
7.53
-
5.05
MGF
19.0
8.23
-
4.94
MGG
14.3
7.99
2.94
4.49
MgA
18.5
8.47
2.00
5.83
MgB
15.5
10.1
2.13
5.96
MgC
20.2
9.74
2.09
6.44
MgD
16.0
9.03
-
-
JAA
14.5
9.23
-
-
JAB
13.6
9.53
-
-
JAC
-
-
-
-
JAD
-
-
-
-
17.1
8.81
2.48
4.64
Average
31
-
0.64
22.5