Alteration in nanohardness of matrix phase associated with precipitation during long-term aging of type 316 stainless steel

Materials Science and Engineering A 489 (2008) 85–92

Alteration in nanohardness of matrix phase associated with precipitation during long-term aging of type 316 stainless steel T. Ohmura a,∗ , K. Sawada b , K. Kimura c , K. Tsuzaki a a

Structural Metals Center, National Institute for Materials Science, Tsukuba 305-0047, Japan Materials Reliability Center, National Institute for Materials Science, Tsukuba 305-0047, Japan c Materials Data Sheet Station, National Institute for Materials Science, Tsukuba 305-0047, Japan b

Received 21 May 2007; received in revised form 30 November 2007; accepted 5 December 2007

Abstract Mechanical characterization for a long-term aged type 316 austenitic stainless steel was carried out using the nanoindentation technique. Samples were aged isothermally at 973 K for various times up to 39,332 h (4.5 years). The precipitation behavior was characterized with transmission electron microscope (TEM), and the hardness of the matrix phase was evaluated by the nanoindentation technique. The nanohardness of the matrix starts to decrease simultaneously with the initiation of M23 C6 precipitation and is reduced significantly after 100 h aging, while the macroscopic Vickers hardness shows a considerable increase. Since the reduction in the matrix hardness is synchronized with the M23 C6 precipitation, the softening results from a depression of the solid-solution hardening by solute elements such as carbon, chromium and molybdenum, which transform into M23 C6 . An accurate amount of precipitation hardening is evaluated by considering the reduction in the matrix hardness. The strengthening by the second phase can be understood quantitatively based on the Orowan mechanism for 183 h aging and the composite strengthening for 39,332 h, respectively. © 2007 Elsevier B.V. All rights reserved. Keywords: Type 316 stainless steel; Long-term aging; Precipitation; Nanoindentation; Hardness; Solid-solution hardening

1. Introduction Precipitation and dispersion of the second phase are major factors for the mechanical properties of various materials. The precipitation/dispersion hardening is basically understood as an interaction between dislocation and the second phase in the traditional models such as the coherency strain field, the cutting of a particle and the Orowan loop process [1]. Second phases also have the role of a component of the composite strengthening based on a continuum mechanics model. In conventional investigations, the strengthening by the second phase, for example precipitation hardening, is estimated as the difference of macroscopic strength before and after precipitation. This is not an accurate estimation because the strength of the matrix phase after precipitation must be lower than that before precipitation due to the reduction in solid-solution strengthening by solute elements brought into the precipitates. The analysis of the composite

∗

Corresponding author. Tel.: +81 298592164; fax: +81 298592101. E-mail address: [email protected] (T. Ohmura).

0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.12.003

strengthening requires the strength of both the second phase and the matrix phase to calculate the contribution of each phase. Therefore, an evaluation of the strength of the matrix phase in a dual-phase microstructure after precipitation is important for estimating the accurate contribution of the strengthening by the second phase. However, the evaluation of the matrix strength is not easy because the precipitation structure is very fine and the typical size of dispersion spacing is in a micron or sub-micron scale. Instrumented nanoindentation techniques [2,3] that have the capability of in situ scanning probe microscopy (SPM) can measure the nano-scale mechanical properties of a precise site in a microstructure with an accuracy of within 1 nm on a SPM image. Ohmura et al. have applied this technique to bulk steels including those with fine structures. They made indent marks on the matrix phase in the interior of a sub-micron grain, and then estimated the contribution of the matrix and the grain boundary separately to understand the strengthening mechanisms of the martensitic steel and ultra-fine grained steel more clearly [4–6]. These results demonstrate the capability of the nanoindentation technique applicable to the analysis of the grain boundary effect for the ultra-fine grained materials. The technique can also

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Table 1 Chemical composition of type 316 austenitic stainless steel C

Si

Mn

P

S

Ni

Cr

Mo

Cu

Ti

Al

B

N

Nb + Ta

0.05

0.74

1.74

0.035

0.006

11.19

17.6

2.32

0.28

0.014

0.005

0.0007

0.023

0.03

be applied to the analysis of the strengthening mechanisms by second phases for fine-structured dual-phase materials. In the present study, heat-resistant steels used in power plants are employed as the target material because the precipitation behavior during long-term aging significantly affects the creep deformation and the fracture property of the steel. In particular, void nucleation and growth occur in a very small scale such as at a grain boundary, a second phase and a hetero-phase interface [7]; hence an evaluation of the change in the local mechanical properties of the matrix and the other phases associated with the precipitation behavior is a key issue for interpreting the deformation and fracture mechanisms and obtaining a guide principle for improving the properties. The present authors evaluated the matrix and the ␴ phase of a type 316 stainless steel aged for about 40,000 h and found that the matrix phase of the aged sample softened significantly while the macroscopic strength increased [8]. The softening of the matrix hardness in the long-term aged sample is attributed to the reduction in the solid-solution strengthening by solute elements such as carbon, molybdenum and chromium, which form significant amount of precipitates [9–12]. For further understanding of the creep deformation and fracture behavior, an alteration in the nanohardness of the matrix phase during aging should be revealed especially in relation with the precipitation behavior. In this paper, nanohardness evaluation was carried out for the matrix phase of an 18Cr–12Ni–Mo austenitic stainless steel (type 316) that was aged isothermally for various times up to about 40,000 h. Then the alteration of the matrix hardness and the contribution of the strengthening by the second phase were considered.

2. Experimental 2.1. Sample The chemical composition of type 316 austenitic stainless steel (JIS SUS 316HP) is shown in Table 1. The sample was solid-solution treated at 1323 K for 80 min followed by water cooling (virgin state). The average grain size was about 100 ␮m. Some samples were subsequently aged for 0.1, 1.0, 10, 183, 3283 and 39,332 h at 973 K. In the three samples with 0.1, 1.0 and 10 h aging times, the specimens were heated in an argon gas atmosphere then quenched in room temperature to fix the aged microstructure. In the other three samples with longer aging times, the specimens were cut from the head portion of a crept specimen under the creep testing condition at 973 K under 137, 78 and 53 MPa [13]. A virgin sample was also provided for making comparisons with the aged samples. Samples for transmission electron microscopy (TEM) observations were finished by electropolishing in a solution of 10% perchloric acid and 90% ethanol at 258–273 K under a potential of 15–18 V. All the

specimen surfaces for nanoindentation testing were mechanically polished, and subsequently electropolished in a solution of 8% perchloric acid, 10% butylcellosolve, 60% ethanol, and 22% water at 273 K under a potential of 40 V to remove the damaged layer. 2.2. Microstructural and mechanical characterization TEM observations were carried out with an accelerating voltage of 200 kV using HITACHI HF2000. Scanning probe microscopy images were obtained by nanoindentation equipment with a SPM capability. In this paper, SPM images are shown not as a topographic image but as a gradient image because the indent marks can be seen more clearly in the latter. Nanoindentation experiments were carried out using a Hysitron, Inc. Triboindenter. A Berkovich indenter was employed, and the tip truncation was calibrated using a reference specimen of fused silica. Analyses for the tip calibration and the calculation of nanohardness, Hn, were conducted using the Oliver and Pharr method [3]. The probed sites and the shape of the indent marks on the specimen surface were confirmed before and after the indentation measurements with the SPM. Conventional Vickers hardness tests which were used for macroscopic strength evaluation were conducted with a load of 1.96 N. A typical size of an indent mark in the Vickers hardness tests had a diagonal length of around 40–50 ␮m. 3. Results TEM micrographs of microstructure evolution, particularly for precipitation behavior during the isothermal aging, are shown in Fig. 1. No precipitates are observed in (a) virgin and (b) 0.1 h aged samples even on the grain boundary and are confirmed by observations in higher magnifications up to 50k. Some precipitates that are smaller than 100 nm appear on the grain boundary in (c) the 1.0 h aged sample. Subsequently, the precipitates coarsen to become a couple of hundred nm as shown in (d) the 10 h aged condition. However, there are still very few precipitates in the grain interior of this sample. In (e) the 183 h aged sample, precipitates are observed within the grain interior and a coarsening of the precipitates also occurs on the grain boundary. The precipitates observed in (e) are considered to be M23 C6 type carbide according to the previous studies [9–12]. At a further aged state in (f) 3283 h, a Laves phase and/or a ␴ phase appear on the grain boundary and in the grain interior. It is not easy to estimate the volume fraction of the precipitates accurately because of their fine structures; however, the fraction of M23 C6 within the grain interior seems to be the same as that of (e) the 183 h aged sample based on the TEM observation. The ␴ phase on the grain boundary appears clearly in the final state

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Fig. 1. Transmission electron microscopy micrographs of type 316 stainless steels with isothermal aging at 973 K. Aging times are (a) virgin, (b) 0.1 h, (c) 1.0 h, (d) 10 h, (e) 183 h, (f) 3283 and (g) 39,332 h.

in (g) the 39,332 h aged sample, and the formation of a precipitation free zone near the ␴ phase is recognized. The ␴ phase also precipitates within the grain interior while the volume fraction of the M23 C6 type carbide seems to be constant. SPM images of the samples that were aged for more than 10 h are shown in Fig. 2. Each image includes a grain boundary and indent marks which are indicated by black and white arrows, respectively. Note that the magnitude of image (d) is lower than that of the other images. In Fig. 2(a) of the 10 h aged sample, few precipitates exist within the grain interior, and fine precipitates

are present on the grain boundaries. The size of these precipitates is around a few 100 nm. When the aging time is close to 200 h in Fig. 2(b), precipitates are observed within the grain interior, and the size of the precipitates on the grain boundary increases. In the much longer aged time of 3283 h in Fig. 2(c), the precipitates on the grain boundary have coarsened considerably with a width of 1 ␮m. For the longest time of the 39,332 h aged sample in Fig. 2(d), the precipitates on the grain boundary, which is considered to be a ␴ phase [8,14], have a much larger size with a few ␮m width. Additionally, the precipitation free

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Fig. 2. Scanning probe microscopy images of type 316 stainless steels with isothermal aging times of (a) 10 h, (b) 183 h, (c) 3283 h and (d) 39,332 h. Black and white arrows indicate grain boundaries and indent marks, respectively.

zone near the precipitates on the grain boundary is clearly recognized. All the features of the microstructure on the SPM images, therefore, coincide with those in the TEM micrographs in Fig. 1. Most of the precipitates on the SPM images have a triangular pyramid shape and their size is relatively larger than that on the TEM micrographs. This is caused by the so called “convolution effect” on SPM imaging [15]. When a probe scans a convexity with an aspect ratio that is lower than that of the probe, the image shows the geometry of the probe at a slightly larger size rather than the actual shape and size of the convexity itself. All the indented positions were determined on the SPM images to prevent the plastic zone from including the fine precipitates as demonstrated on the SPM images in Fig. 2. The other samples of virgin, 0.1 and 1.0 h aging times are not shown in Fig. 2 because no specific microstructure appeared on their SPM images. Typical load–displacement curves obtained by the nanoindentation measurements are shown in Fig. 3. The three curves correspond to the samples of virgin, 183 and 39,332 h aging times. The penetration depth increases significantly with aging time, indicating a remarkable softening by the isothermal aging. The penetration depth at the peak load is around 50 nm, and the size of the corresponding plastic deformation zone is estimated to be about 500 nm in diameter using the hemispherical approximation [16]. The nanohardness Hn obtained by an analysis of the unloading curves for each sample was plotted as a function of aging time with solid circles in Fig. 4. The error bars are the

standard deviation. The nanohardness starts to decrease after about one h aging, and subsequently softens monotonically up to about 40,000 h aging. The Vickers hardness Hv, which corresponds to the macroscopic strength, was also plotted with solid squares. The value of the standard deviation for the Hv is very

Fig. 3. Typical load–displacement curves of type 316 stainless steels with isothermal aging at 973 K. Open squares, solid squares and open circles represent aging times of virgin, 183 h and 39,332 h, respectively.

T. Ohmura et al. / Materials Science and Engineering A 489 (2008) 85–92

89

Fig. 5. Time–temperature-precipitation diagram of type 316 stainless steel [13].

Fig. 4. Nanohardness Hn and Vickers hardness Hv of type 316 stainless steels plotted as a function of aging time with solid circles and squares, respectively. Open squares and open triangles represent estimated Vickers hardness Hvmatrix and strengthening by precipitates Hvp obtained as a difference between Hv and Hvmatrix .

small and overlaps with the square marks. The Vickers hardness starts to increase after 10 h aging, which is the same time that the nanohardness starts to decrease. The marks of the open squares and triangles, which represent the estimated data from measured Hn and Hv, are explained in Section 4. Some important aging times for the beginning of precipitation are indicated on the top side of the time axis in Fig. 4 based on the time–temperature-precipitation (TTP) diagram of Fig. 5 reported in the previous study [14]. For the aging at 973 K, the precipitation of M23 C6 on the grain boundary and the grain interior starts after about 1 and 5 h aging, respectively. A Laves phase subsequently precipitates after around 1000 h, and then a ␴ phase forms after more than 10,000 h aging. 4. Discussion 4.1. Quantitative evaluation of the reduction of the matrix phase An accurate estimation of the matrix hardness is important for understanding the hardening by the second phase for the macroscopic strength. To do the estimation quantitatively, the matrix hardness should be evaluated in the unit of the Vickers hardness Hvmatrix so that a comparison can be made with the measured macroscopic Hv of the dual-phase. Once a reduction ratio in the Hn of the each aging time was normalized by the nanohardness of the virgin sample, the Hnvirgin was calculated as scaled on

the right-upper axis in Fig. 4. Then the Hvmatrix for each aging time was estimated using the ratio of Hn/Hnvirgin in the corresponding aging time from the measured Hv of the virgin sample as plotted with the open squares. For example, the nanohardness Hn of the 39,332 h aged sample is 72% smaller than that of the virgin sample. Using the value of the ratio Hn/Hnvirgin , the Hvmatrix of the matrix phase of the 39,332 h aged sample is calculated to be 1.1 GPa as 72% of 1.5 GPa of the virgin sample. This estimation is reasonable because the virgin sample has a single austenitic phase, and the strengthening factors are the same in both the macro and nano-scale. Therefore, the softening behavior in the Hvmatrix should have the same trend as that of the measured nanohardness Hn. The values of the open squares also coincide with the values estimated by the empirical relation of Hn = 3Hv [17]. Accordingly, the accurate amount of hardening by the second phase Hvp can be obtained by the difference between the Hv of the solid squares and the Hvmatrix of the open squares, as plotted by the open triangles in the bottom of Fig. 4. 4.2. Softening mechanism of the matrix phase during aging Since the softening of the matrix nanohardness during aging is thought to occur from the reduction of solid-solution hardening, contributions of each alloying element can be estimated along with the time difference in the precipitation of M23 C6 and the other phases. In the aging time of shorter than 1 h, there is no change in either the nanohardness Hn or the Vickers hardness Hv; hence remarkable precipitation does not occur. The reduction in the Hn for the samples of 10 and 183 h aging was associated with the precipitation of M23 C6 because the softening started at the same time as the initiation of the precipitation of M23 C6 , and the Laves phase did not precipitate even after about 200 h aging. The metal elements M included in the M23 C6 and the mass fractions fx of each element at the 100 h aging state that are shown in the previous study [14] are represented in Table 2. If we assume that the entire amount of the 0.05 mass% carbon in the sample transforms into M23 C6 , the mass fraction mx of the metal elements that balanced with the amount of carbon can be

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Table 2 Estimation of the reduction in nanohardness of the matrix phase associated with the precipitation of M23 C6 for 100 h aging. Element Cr Fe Mo Ni C Total

fx (mass%) 60 20 15 5 – –

mx (mass%)

−σfss (MPa)

−Hvss (MPa)

0.54 0.18 0.13 0.045 0.05

8 – 2 – 30

24 – 6 – 90

–

40

120

fx : Mass fraction of the alloying element M in M23 C6 . mx : Mass fraction of the alloying element M that is balanced with 0.05 mass% carbon in M23 C6 , namely the reduction of the solute elements in the matrix phase. −σfss : Flow stress reduction in solid-solution hardening by the alloying element M and carbon. −Hvss : Vickers hardness reduction in solid-solution hardening converted from −σfss through the empirical relation Hv ∼ 3σ f [18].

calculated as shown in the table. Solid-solution hardening in a flow stress σ f of Mo and Cr for the fcc austenitic phase is approximately 15 MPa/mass% while Ni does not make any significant hardening contribution [8]. Therefore, the flow stress reductions with an amount of solid-solution hardening −σfss corresponding to 0.54 mass% Cr and 0.13 mass% Mo are about 8 and 2 MPa, respectively. Additionally, the reduction by the 0.05 mass% C is about 30 MPa. Accordingly, the total reduction of −σfss associated with the precipitation of M23 C6 is estimated to be 40 MPa. Since σ f is about 1/3 of the Hv in metals [18], the corresponding reduction −Hvss is 120 MPa. On the other hand, the previous study [19] on microstructural characterization shows a change in the number density of M23 C6 in the grain interior during an isothermal aging at 973 K for a type 316 stainless steel as represented in Fig. 6. Since the density shows a peak at around 100 h for two different heats, the precipitation of M23 C6 in the present steel could be completed at around 100 h aging. In Fig. 4, the nanohardness Hn of the virgin state and the 183 h aging time

Table 3 Mass fractions of each element in the Laves and ␴ phases [14]

Laves ␴

Cr

Ni

Fe

45 10

15 35

5 5

35 50

are 4.5 and 4.0 GPa, respectively, and the reduction in measured Hn by 183 h aging is estimated to be around 500 MPa. Assuming that the Hn can be converted to the Hv by the empirical equation of Hv = 1/3Hn [18], the measured reduction is around 160 MPa in Hv, which coincides roughly with the amount of the reduction −Hvss estimated above by the M23 C6 precipitation. With further aging time, the softening of the matrix is associated with the reduction of in-solution elements of Mo and Cr by the precipitation of Laves and ␴ phases. Mass fractions of the main elements of Mo, Cr, Ni and Fe included in the precipitates are shown in Table 3. The values are picked out roughly from the plot in literature [14]. Since the ␴ phase starts to precipitate after about a 10,000 h aging, the softening of the matrix between 1000 and 10,000 h is affected by a relatively large contribution of Mo. Subsequently, the precipitation of the ␴ phase depresses an amount of in-solution Cr leading to a further softening of the matrix phase. However, the solid-solution hardening of Mo and Cr is almost the same, so that no significant change occurs in the slope of the softening curve. 4.3. Hardening of macroscopic hardness by the precipitates The major strengthening mechanisms by the precipitates are (1) particle cutting, (2) Orowan process and (3) composite strengthening. The former two are based on the dislocation theory and the other on a continuum mechanics theory. To understand the strengthening mechanisms by the precipitates responsible for the Hvp , we selected the two aging times of 183 and 39,332 h to be the characteristic precipitation stage. For the 183 h aging time, all the precipitates could be M23 C6 because the Laves phase precipitates mainly after 1000 h aging as shown in Fig. 4. The particle size shown in Fig. 1(e) is relatively large with a couple of hundred nm, and the M23 C6 precipitate is very hard. Therefore, the particles could never be cut, and the most probable mechanism is the Orowan process. A critical shear stress τ c for overcoming obstacles is given by [20] τc =

Fig. 6. Change in the number density of M23 C6 particles precipitated within a grain interior during isothermal aging of type 316 stainless steel [17] (courtesy of Shinya).

Mo

0.8μb L

(1)

where μ is the shear modulus, b a magnitude of the Burgers vector, L the spacing between particles for square lattice. The coefficient 0.8 is obtained empirically from simulation works [21,22]. The Hvp for the 183 aging are 0.4 GPa, which is equivalent to about 133 MPa in σ f based on the same relation between Hv and σ f as described above. A flow stress σ f of a polycrystal can be converted to a shear stress τ c through the relation σ f = Mτ c using the Taylor factor M of 3.06 for fcc metals. Substituting the values of 43.5 MPa for τ c , 80 GPa for μ and 0.258 nm for b into Eq. (1), the value L for 183 h aging of 375 nm is obtained. The

T. Ohmura et al. / Materials Science and Engineering A 489 (2008) 85–92

value coincides roughly with the TEM observation in Fig. 1, and is also in the same order with the value of 580 nm given as an inverse of the square root of the number density of about 3.0 ␮m−2 at the peak in Fig. 6. Therefore, the value L appears to be reasonable, and the main mechanism for the 183 h aging state could be the Orowan mechanism. Note that the average spacing is almost the same or slightly smaller than the plastic deformation zone size estimated in Fig. 3. However, the indenter was carefully positioned to avoid the plastic deformation zone from interacting with the precipitates as shown in Fig. 2; hence the measured nanohardness corresponds to the matrix hardness in principle. The same evaluation of L was carried out for the 39,332 h aging. The Hvp for the 39,332 h aging is 1.1 GPa as given in Fig. 4, and the calculated L of 138 nm was obtained through Eq. (1). However, the spacing between the particles shown in Fig. 1(g) is obviously larger than the calculated L; therefore, the Orowan mechanism is inappropriate for the strength of the 39,332 aging sample. Note that the microstructure of the precipitates of the 39,332 h aging is remarkably different from that of the 183 h aging. The volume fraction of the second phases in the 39,332 aging sample is much larger than that of the 183 aging sample, however, the size of the precipitate is very large and the microstructure is not fine as shown in Figs. 1 and 2. Therefore, one of the probable strengthening mechanisms is the composite strengthening. The strength of a composite Hcomp can be obtained based on the simple rule of mixture given as H comp = (1 − V␴ − Vp )Hm + V␴ H␴ + Vp Hp ,

(2)

where V and H stand for the volume fraction and the hardness of each constituent phase, respectively. The subscripts of m, ␴ and p refer to the matrix, the ␴ phase and the other phases such as the M23 C6 and the Laves phase. From Eq. (2), the composite hardening of the ␴ phase and the other precipitates is given as Hcomp = V␴ H␴ + Vp Hp . The volume fraction V␴ is evaluated to be 0.16 by point counting [23], and the hardness H␴ is measured directly as 17 GPa in a unit of nanohardness Hn [8]. Unfortunately, it is hard to measure the Vp and Hp because the size of the M23 C6 and the Laves phase is very small at less than 100 nm. However, we can presume the Vp as 0.03 from the TEM micrograph and the Hp as the same level of the ␴ phase [8]. Using these estimations for each term, the Hn-based Hncomp can be calculated as 3.23 GPa, which is converted to a Hv-based Hvcomp as 1.08 GPa using the empirical relation. The estimated Hvcomp of 1.08 GPa is very close to the Hvp of 1.1 GPa. The transition of the strengthening mechanisms could originate from the change of the contributing precipitates with different microstructures. In the 183 h aging, the M23 C6 is the major precipitate with the fine structure; hence the dispersion hardening based on the Orowan mechanism is understood as the main hardening factor. However, the number density of the M23 C6 has a peak around 100 h, which must be also the time for saturating the volume fraction, and becomes much lower for the 39,332 h aging than that of the 183 h aging as shown in Fig. 6. Therefore, the contribution of the M23 C6 must be much lower in the 39,332 h aging time. On the other hand, Hvp is higher

91

for the 39,332 h aging than that of the 183 h while the number density of the M23 C6 is much lower for the 39,332 h, suggesting another strengthening factor for the 39,332 h aging. In the aging time 39,332 h, the main precipitates in volume fraction are the Laves and the ␴ phases. Particularly, the volume fraction of the ␴ phase is over five times higher than that of the M23 C6 as described above, but the size of the precipitate is very large and microstructure is not fine as shown in the TEM and SPM images. Therefore, it is reasonable that the main hardening factor in the 39,332 h aging is the contribution of the Laves and the ␴ phases and the most probable mechanism is composite strengthening by the significant volume fraction. Since the long-term aging causes the remarkable change in the microstructures of the precipitate, the corresponding strengthening mechanism is totally different in the different aging stage. 5. Conclusions Nanomechanical characterization was carried out for an isothermally aged type 316 austenitic stainless steel. The nanohardness of the matrix phase of the aged sample was successfully evaluated by the nanoindentation technique. Then the alteration in the hardness of the matrix phase and the macroscopic hardness associated with the precipitation behavior were considered. The following conclusions were drawn. (1) The nanohardness of the matrix phase starts to soften at the same time as the initiation of the M23 C6 precipitation. The macroscopic hardness of the Vickers hardness also increases with the precipitation behavior of M23 C6 . (2) The amount of softening in the matrix nanohardness from the virgin state to the 183 h aging, which is around the peak of the number density of M23 C6 particles, coincides with the reduction of solid-solution strengthening by solute elements transforming into M23 C6 precipitates including the entire amount of the 0.05 mass% C of the sample. (3) The accurate contribution of precipitation hardening can be evaluated by considering the reduction in the matrix hardness. In the 183 h aging, the precipitation hardening based on the Orowan mechanism is effective. In the 39,332 h aging, the composite strengthening by the Laves and the ␴ phases as well as the M23 C6 is dominant for the age hardening behavior. Acknowledgments The authors would like to thank Dr. N. Shinya for the courtesy of providing the data in Fig. 6. A part of this study was financially supported by IKETANI Science and Technology Foundation and the Budget for Nuclear Research of the Ministry of Education, Culture, Sports, Science and Technology, based on the screening and counseling by the Atomic Energy Commission. References [1] P. Haasen, Physical Metallurgy, Cambridge University Press, Cambridge, 1978, pp. 370–375.

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