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Microstructural stability of an as-fabricated 12Cr-ODS steel under elevated-temperature annealing
Accepted Manuscript Microstructural stability of an as-fabricated 12Cr-ODS steel under elevatedtemperature annealing Jingjie Shen, Huilong Yang, Yanfen Li, Sho Kano, Yoshitaka Matsukawa, Yuhki Satoh, Hiroaki Abe PII:
S0925-8388(16)33487-9
DOI:
10.1016/j.jallcom.2016.11.029
Reference:
JALCOM 39524
To appear in:
Journal of Alloys and Compounds
Received Date: 24 July 2016 Revised Date:
15 October 2016
Accepted Date: 2 November 2016
Please cite this article as: J. Shen, H. Yang, Y. Li, S. Kano, Y. Matsukawa, Y. Satoh, H. Abe, Microstructural stability of an as-fabricated 12Cr-ODS steel under elevated-temperature annealing, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.11.029. 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 proof before it is published in its final 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.
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Microstructural stability of an as-fabricated 12Cr-ODS steel under elevated-
Jingjie Shen
a,*
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temperature annealing , Huilong Yang b, Yanfen Li b, Sho Kano c, Yoshitaka Matsukawa b, Yuhki
Satoh b, Hiroaki Abe b,c
Graduate School of Engineering, Tohoku University, 2-1-1 Katahira, Sendai 980-8577,
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a
Japan b
Institute for Materials Research, Tohoku University, 2-1-1 Katahira, Sendai 980-8577,
c
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Japan
School of Engineering, The University of Tokyo, 2-22 Shirakata Shirane, Tokai-mura,
Ibaraki 319-1188, Japan *
Corresponding author:
Abstract
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E-mail address: [email protected] (J.J. Shen)
The effects of elevated-temperature annealing in the range of 1473 K - 1673 K on the
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coarsening kinetics of nanoparticles and microstructural evolution in an as-fabricated 12CrODS steel was systematically investigated. Results show that the nanoparticles were
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extremely stable and continued precipitating during annealing at 1473 K, whereas, significant coarsening took place during exposure at 1673 K, increasing by 177% and 423% in diameter for 1 h and 24 h, respectively. Furthermore, heterogeneously distributed voids were developed, sizes of which increased distinctly at 1573 K and 1673 K for longer holding time. This is probably due to the growth and/or aggregation of Ar-entrapped and pre-existed voids in the as-fabricated condition. Additionally, no obvious changes were observed in microstructure and texture at 1473 K, nevertheless, full recrystallization were achieved at 1673 K for 24 h, and the major texture was not altered except reduction of the
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intensity. The apparent activation energy for recrystallization was experimentally calculated as 424 ± 22 kJ/mol, which was almost the same as that for nanoparticle coarsening, i.e., 412
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± 55 kJ/mol, suggesting that recrystallization proceeded accompanying with nanoparticle coarsening that would in turn stimulate the recrystallization process. It was deduced that recrystallization in the as-fabricated 12Cr-ODS steel was presumably not able to occur until the nanoparticles were dissolving and/or coarsening during annealing.
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Keywords: oxide dispersion strengthened steels; nanoparticles; annealing; coarsening;
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voids; recrystallization
1. Introduction
Development of oxide dispersion strengthened (ODS) steels for structural materials in future fission and fusion reactors has been of increasing interest during the past decades, because of their adequately high thermal stability and strength, and excellent irradiation
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damage resistance [1-5].
In order to produce reduced-activation ODS steels, alloying elements with a slow decaying rate, such as Mo, Nb in conventional steels, are substituted by W, V, Ti, which have a higher decay rate of radioactivity, because the strengthening function of Mo and Nb
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in steels could be compensated by W, V and Ti [6, 7]. Owing to the moderate tensile and creep properties at high temperatures, the utilization of conventional high-chromium
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ferritic/martensitic steels, like modified 9Cr-1Mo, or the reduced-activation steels, such as F82H, ORNL 9Cr-2WVTa, EUROFER, and JLF-1, was limited to ~823 K [3, 8]. One way to increase the operating temperature is to add extremely stable oxide particles in the matrix. The high number density of nanoscale oxide particles induced by mechanical alloying play an important role in the performances. Not only preventing the dislocation migration and providing obstacles to resist the motion of grain boundaries, thereby ensuring the adequate strength and thermal stability at high temperatures, but also it can serve as sinks for irradiation induced defects and nucleation sits for helium bubbles, making ODS steels more
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resistant to irradiation hardening and swelling [9-11]. Therefore, the stability of those nanoscale oxide particles during both the fabrication processing (e.g., hot isotropic pressing,
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hot extrusion, hot rolling, annealing) and the operating temperatures (up to ~1073 K) or even higher is of crucial importance, which is directly responsible for the deterioration of strength, creep properties, and irradiation resistance.
Several investigations about the effects of high-temperature annealing on the thermal
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stability of ODS steels have been done. Schneibel et al. [12] reported that the density of 14YWT (Fe-14Cr-3W-0.4Ti-0.25Y2O3) decreased due to pore formation during annealing at 1273 K up to 30 h. Besides, Williams et al. [13] suggested that Ti addition did not affect
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the coarsening kinetics of oxide particles during annealing at 1473 K for 100 h, which was in accordance with the diffusion-controlled mechanism. Furthermore, Miller et al. [14, 15] found that the size of (Y, Ti, O)-enriched oxide particles in MA957 (Fe-14Cr-1Ti-0.3Mo0.25Y2O3) increased from 2.4 nm to 9.2 nm in diameter and the number density decreased from
~2 × 1024 m-3 to ~8 × 1022 m-3 after aging at 1573 K for 24 h. Although the
coarsening behavior of the nanoscale oxide particles had been observed in the previous
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studies, most of which were focused on a certain annealing temperature, so that the coarsening kinetics is not able to be analyzed, leaving the dependence of the annealing temperatures still unclear.
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In this work, we are aiming to investigate the microstructural evolution in an asfabricated 12Cr-ODS steel, and discuss the coarsening kinetics of Y-Ti-O nanoparticles
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under annealing at temperatures ranging from 1473 K to 1673 K. The findings are of the essence to provide experimental dataset for thermodynamic and kinetic study on the nucleation-growth-coarsening model of nanoparticles, so as to optimize the fabrication processing, and predict the long-term aging stability of ODS steels.
2. Experimental procedure
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The 12Cr-ODS steel with a nominal composition of Fe-12Cr-2W-0.3Ti-0.25Y2O3 was utilized in this work. The fabrication procedures and chemical composition can be found in
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detail elsewhere [16, 17]. The as-fabricated specimen was cut into small pieces and then isothermally annealed in a vacuum furnace at 1473 K, 1573 K and 1673 K for various time intervals ranging from 1 h - 60 h. The heating rate was 300 K/h, and the cooling rate was ~4000 K/h. In the following, specimens were designated after the annealing parameters.
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For instance, the specimen annealed at 1473 K for 1 h is hereafter marked as 1473 K-1 h. Microstructure observations were performed with conventional scanning electron microscope (SEM), and a JEOL JXA 8530 field-emission SEM equipped with electron
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backscatter diffraction (EBSD) device. Texture intensity in an area of 400 × 400 µm2 was described by orientation distribution function (ODF), which was computed by the series expansion method using TSL OIM analysis software. Besides, the orientation was presented in the form of the Euler angles {φ1, Ф, φ2} (Bunge’s notation) [18], which were calculated under the assumption of orthonormal sample symmetry, i.e., 0° ≤ {φ1, Ф, φ2} ≤ 90° [19]. The nanoscale oxide particles were characterized by a JEM-2100 transmission
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electron microscope (TEM) operating at 200 kV. TEM foils were firstly mechanically thinned down to ~80 µm and then electrochemically polished in a solution of 5 vol% perchloric acid in acetic acid by a twin-jet unit (Struers TenuPol-5) at room temperature.
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Thickness of the observed area was determined under two beam condition. The diameter of nanoparticles was measured using ImageJ software from at least 600 particles based on the
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TEM micrographs.
For the hardness and EBSD testing, the annealed specimens were ground with silicon
carbide paper up to 2000 grit on the rolling plane, and then polished with alumina and colloidal silica suspension. Vickers hardness was carried out under a load of 1 kg for a period of 15 s, and randomly 12 measurements were averaged after eliminating their minimum and maximum.
3. Results and discussion 4
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3.1 Effect of annealing on nanoscale oxide particles 3.1.1 Size and number density of nanoscale oxide particles
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Fig. 1 shows bright-field TEM micrographs of nanoscale oxide particles in the annealed specimens. It is noted that some of the oxide particles were cuboidal at 1473 K and 1573 K, while it became spherical at 1473 K. This is probably due to variation of interfacial structures between oxide particles and matrix, changing from coherent to
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incoherent structure [20, 21]. The average diameter and number density of the nanoscale oxide particles in each specimen were plotted as a function of annealing time in Fig. 2. As
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for the as-fabricated condition, the average diameter and number density were 3.5 ± 0.7 nm and ~1.3 × 1023 m-3 [17]. As shown in Fig. 2, the size of the nanoscale oxide particles did not obviously vary, and the number density slightly increased after annealing at 1473 K, indicating that Y, Ti, O elements were still supersaturated in the as-fabricated matrix, and numerous oxide particles were precipitating during annealing. In contrast, it was significantly coarsened after exposure at higher temperatures, resulting in lower number
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density. For instance, the average diameter amounted to 9.7 ± 2.3 nm and 18.3 ± 4.5 nm after annealing at 1673 K for 1 h and 24 h, which increased by 177% and 423%, respectively. Dependence of coarsening rate on temperature and holding time was clearly observed, which is presumably because the solute atoms (i.e., Y, Ti and O) associated with
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the oxide particles diffused more rapidly at higher temperatures.
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3.1.2 Coarsening kinetics of nanoscale oxide particles In general, growth of precipitates in alloys controlled by solute diffusion, so-called
Ostwald ripening [22], can be expressed by the Lifshitz-Slyozov-Wagner (LSW) theory [23, 24], which was applied to analyze the coarsening kinetics of the oxide particles in ODS steels [21, 25]. dt3 - d03 = kt
(1)
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where dt and d0 are the average diameter of nanoparticles in the annealed and as-fabricated specimens, t is the annealing time, and k is the coarsening rate constant that is dependent on
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annealing temperature. Eq. (1) demonstrates that the increment of volume of the nanoscale oxide particles has a linear correlation with respect to annealing time t, and k can be derived from linear fitting, which is presented in Fig. 3(a). Moreover, the coarsening rate k can be determined by Arrhenius equation:
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k = k0 exp (-Q/RT) (2)
where k0 is a constant, T is the absolute temperature, Q is the apparent activation energy for coarsening, R is the universal gas constant (8.314 J/mol/K). Taking natural logarithm of Eq.
lnk = lnk0 - Q/RT
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(2), it yields: (3)
According to Eq. (3), it is easy to get the slope from correlation of lnk against 1/T, as plotted in Fig. 3(b). Consequently, the apparent activation energy for coarsening of the nanoscale oxide particles, Q, was experimentally estimated as 412 ± 55 kJ/mol. For
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coarsening process of the nanoparticles, the system tends to minimize the free energy by reducing the interface area between particles and matrix [22]. In terms of this fact, Y, Ti, O atoms need to dissolve from the surface of smaller particles, diffuse through the matrix and redeposit onto the larger particles, resulting in the growth of larger ones. In fact, the size of
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Ti, Y atoms is much larger than Fe, Cr atoms [15], and the diffusion of Y atoms in the FeCr matrix is rather difficult, thereby giving rise to a sluggish diffusion process. Williams et
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al. [13] presumed that diffusivity of Y atoms dominated the coarsening rate of Y-Ti-O enriched nanoparticles since the diffusion coefficient of Y is significantly lower than that of Ti and O atoms. However, the activation energy of Y for diffusion in matrix is 218.1 kJ/mol [26], which is clearly smaller than the experimental estimation in this work. Thus, the coarsening of Y-Ti-O nanoparticles may not be controlled only by the diffusion of Y atoms. Previous research [15, 27, 28] indicated that a large amount of TiO, YO, CrO molecular ions were detected in matrix, because of the strong solute-oxygen affinity. Those solute-oxygen nanoclusters may diffuse as an entity during annealing, which would retard
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the coarsening of the nanoscale oxide particles, resulting in much larger activation energy. Additionally, the high concentration of vacancies in the nanoclusters probably stimulate the
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diffusivity of solute-oxygen complex in matrix [29, 30]. As a result, the coarsening equation for nanoscale oxide particles in this 12Cr-ODS steel can be written by: (4)
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dt = {1.5 × 1015 · exp(-(412 ± 55)/(8.314·T))·t + d03}1/3
where T , t and d0 are in Kelvin, hour and nanometer, respectively.
Therefore, it is probably able to estimate the coarsening rate and predict corresponding
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average diameter of the nanoparticles under a certain annealing condition with the aid of Eq. (4). For instance, comparison of growth time of nanoparticles from the as-fabricated condition to 5.2 nm, the average diameter in 1473 K-24 h specimen, would correspond to ~ 7.7 × 106 h during exposure at 1073 K, suggesting the nanoparticles should be remarkably stable at this or lower service temperatures. Nevertheless, the long-term annealing
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experiments at lower temperatures are needed to validate this prediction. 3.2 Effect of annealing on microstructure 3.2.1 Formation of voids upon annealing
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Accompanying with the coarsening of nanoparticles, microstructure evolution of the matrix was also characterized. Fig. 4 shows SEM micrographs of the annealed specimens.
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Inhomogeneous microstructure development associated with voids was evident. The voids were observed in the annealed specimens. The number density and size of voids increased due to annealing at higher temperature and longer duration. In order to elucidate the formation mechanism of these voids, the microstructure of as-
fabricated specimens was also examined. Fig. 5 represents the micrographs of the asfabricated specimen in SEM and TEM observations. As shown in Fig. 5(a), it is clear that voids with a diameter ranging from tens of nanometers to several micrometers existed in the as-fabricated matrix. Moreover, in Fig. 5(b), voids trapped around the nanoparticles and
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embedded in matrix were also observed, which were firstly confirmed as Ar cavities by the electron energy loss spectrometer in Ref. [31], retaining from mechanical alloying in the
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argon atmosphere. Accordingly, two potential reasons for the voids formation are proposed: one is likely that there are pre-existed voids in the as-fabricated condition, and then growth and/or agglomeration took place in the subsequent heat treatment; the other one is probably ascribed to entrapment of argon gas at the surface of the oxide particles. Turker [32] reported that, in PM2000 (Fe-20Cr-5.5Al-0.5Ti-0.5Y), less porosity was observed after
3.2.2 Recrystallization behavior
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which was explained by the entrapped-gas mechanism.
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mechanical alloying in hydrogen atmosphere compared to that in an argon atmosphere,
Fig. 6 and Fig. 7 demonstrate the inverse pole figure (IPF) images and the corresponding φ2 = 45° sections of ODF maps of the annealed specimens. As can be seen, no significant changes in microstructure were observed at 1473 K, as well as texture shown
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in Fig. 7. It is obviously notable that the dominant texture was composed of intense {001}<110> component and relatively weaker γ-fiber with <111> axis parallel to normal direction. This texture was akin to that of as-fabricated condition [17], even after annealing up to 24 h, revealing that the microstructure of the as-fabricated 12Cr-ODS steel was stable.
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Two possible factors are responsible for the thermal stability: on the one hand, {001}<110> texture retards the recrystallization, due to containing the lowest stored energy that is the
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driving force for recrystallization [33, 34]; on the other hand, large number density of nanoscale oxide particles inhibit the migration of dislocations and grain boundaries during annealing. It is due to the strong pinning effect of nanoparticles and the lowest stored energy of {100}<110> texture [33], thereby preventing dislocations and grain boundaries from migration and retarding the recrystallization process. At 1573 K, it was found that only a few proportion of grains were growing during exposure up to 24 h (Fig. 6), and the intensity of texture gradually decreased (Fig. 6). In contrast, as for 1673 K-12 h specimen, evident growth of grains was observed, and full recrystallization was achieved with further
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annealing for 24 h. As shown in Fig. 7, the texture has two characteristic peaks, namely, {001}<110> and {111}<110>. The structure was principally kept throughout the
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experimental conditions. The evident alteration was observed in the intensity reduction in {001}<110> and {111}<110> components. This suggests that the nucleation sites for recrystallization were limited by the original crystallographic distribution.
Furthermore, as well known that recrystallization is a migration process of high-angle
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grain boundaries. Misorientation, which is defined as a minimum rotation angle required to bring the neighboring grain into coincidence, was also measured by EBSD for each specimen in Fig. 6. Fig. 8 shows the distributions of misorientation angles in the specimens.
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It is evident that the average misorientation angle slightly increased at the early stage of annealing by comparison with the as-fabricated specimen, which was composed of a majority of highly oriented {001}<110> grains [17], giving a lower misorientation (~ 17.5°). It seems to be very stabilized during exposure at 1473 K, indicating recovery may be still in progress. On the other hand, owing to partial recrystallization, the proportion of low misorientation reduced, and corresponding high misorientation increased, resulting in
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obvious increase of the average misorientation. As shown in Fig. 8, the proportion of low misorientation decreased a little larger in the 1573 K-24 h and 1673 K-12 h specimens, respectively. As a consequence of full recrystallization and/or grain growth, the
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misorientation angle increased up to ~ 38.7° after heating at 1673 K up to 24 h, in which the grain boundaries changed into high angle grain boundaries.
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In addition, it is found that the recrystallized grains grew along the rolling direction, forming elongated grains, which is a common feature in recrystallized ODS ferritic steels. This is presumably ascribed to the alignment of oxide particles along the extrusion direction (or cold rolling direction), resulting in an anisotropic pinning force that hinders grain boundary movements effectively in the transverse direction [35-37]. Furthermore, in particle embedded materials, it is widely accepted that diameter of the critical maximum grain (Dc) can be quantitatively estimated by Zener equation [38], whose general form is written as:
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Dc = Kd / f m (5) where K is a dimensionless constant, d and f are the diameter and volume fraction of
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pinning nanoparticles, and m is an index for f. According to Ref. [38], parameters K = 0.17 and m = 1.0 were reasonably chose for this work, where the volume fraction f was 0.67% for the fully recrystallized specimen (1673 K-24h). Substituting those parameters into Eq. (5), the calculated diameter of grains was ~0.46 µm, which was approximately two orders
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of magnitude lower than the experimental measurements (~46.8 µm). This deviation may probably stem from the artificial error of determination on parameters (k, m, f), or other
3.3 Effect of annealing on hardness 3.3.1 Variations of hardness
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underlying mechanisms tailoring the grain size, which is necessary to be verified in future.
Due to softening during annealing is associated with the recrystallization, the volume
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fraction of recrystallization (Xrex) was generally calculated from Vickers hardness testing [39, 40], which can be given by:
Xrex = (HV0 – HVt)/(HV0 – HVrex)
(6)
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where HV0 and HVt are the Vickers hardness for the as-fabricated and annealed specimens at a given condition, respectively. HVrex is the hardness of a fully recrystallized specimen.
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Here, the 1673 K-24 h specimen was considered as full recrystallization (~ 100%) based on EBSD measurements, as shown in Fig. 6. Fig. 9 presents Vickers hardness and volume fraction of recrystallization as a function
of annealing condition. As shown in Fig. 9(a), it can be noticed that the hardness slightly decreased at 1473 K up to 24 h, however, obvious reduction was observed after annealing at 1573 K and 1673 K for more than 6 h. The corresponding recrystallization fraction based on hardness variations is shown in Fig. 9(b). Evidently, the recrystallization fraction levels off at 1473 K from 6 h to 24 h, indicating that no obvious changes were emerged in
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microstructure, and the specimen remained in recovery stage. This can be confirmed from microstructure evolution at 1473 K, as shown in Fig. 6. On the contrary, the fraction
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increased directly during annealing at 1573 K for 6 h, and 1673 K for 1 h, due to occurrence of recrystallization, leading to more proportion of softening. Together with Fig. 6, one can conclude that the strain recovery occurred at 1473 K for 24 h, partial
3.3.2 Apparent activation energy of recrystallization
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recrystallization started at 1573 K, and grain growth took place at 1673 K for 24 h.
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To describe the kinetics of isothermal recrystallization, Johnson-Mehl-AvramiKolmogorov (JMAK) equation is commonly accepted as [41-43]: Xrex = 1 – exp (– ktn)
(7)
where k is a temperature dependent constant and n is referred to as the Avrami exponent, which is a function of both nucleation and growth rates. Eq. (7) can also be given as: (8)
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ln[ln(1/(1 – Xrex))] = nlnt + lnk
Eq. (8) yields a straight line whose slope is equal to n when the left-hand side is plotted against lnt, which is shown in Fig. 10(a). The values of Avrami exponent at 1473 K, 1573
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K and 1673 K ranged from 0.41 to 0.9, which is deviated from the theoretical values, i.e., n = 2 for 2D growth, and n = 3 for 3D growth under assumption of site-saturated nucleation
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during recrystallization [44]. Several researches on the static recrystallization kinetics of Fe-based alloys claimed the similar results. For instance, Xie et al. [45] reported that the value of Avrami exponent was estimated as 0.54 in a low-carbon steel, and Malekjani et al. [46] calculated Avrami exponent of austenitic and ferritic stainless steels, which was in the range of 0.75 - 1.1. For the deviation of the Avrami exponent in this study, it is probably an open question.
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The annealing time up to a certain level of recrystallization can be described by Arrhenius equation. It was defined the annealing time as the time that was required for 50%
t0.5 = t0 exp (Qrex / RT)
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recrystallization, t0.5, which is expressed by the following formula [47]. (9)
where t0 is a constant, and Qrex is the apparent activation energy of recrystallization.
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Taking natural logarithm of both sides of Eq. (9), the following expressions are given by:
k0 = Qrex/R
(10)
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lnt0.5 = k0/T + lnt0 (11)
As presented in Fig. 10(a), if the values of n and lnk were obtained by linear fitting, then, lnt0.5 at each temperature is able to be calculated by substituting Xrex = 0.5 into Eq. (8). Fig. 10(b) illustrates that values of lnt0.5 were plotted as a function of 1000/T. The slope, Qrex/R, was attained by linear fitting. Accordingly, the apparent activation energy of
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recrystallization for the as-fabricated 12Cr-ODS steel was estimated as 424 ± 22 kJ/mol, which is much larger than that of self-diffusion in α-iron and δ-iron with 251 kJ/mol [48] and 296 kJ/mol [49]. Nevertheless, similar results were reported in the previous research.
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For instance, Liu et al. [50] estimated the activation energy for recrystallization in the coldrolled V-doped steel, which was 474 kJ/mol, suggesting that the dispersed V(C,N) precipitates prevented the growth of recrystallized grains. The high value in this study was
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presumably resulted from the pinning effect of nanoscale oxide particles against the motion of dislocations and grain boundaries. Additionally, it is surprisingly noteworthy that the apparent activation energy for
recrystallization was nearly the same as that for nanoparticle coarsening, which was 412 ± 55 kJ/mol as estimated in Section 3.1.2. This reveals that coarsening of nanoparticles took place accompanying with the recrystallization process. In other words, the nanoparticles play a pivotal role in retarding the recrystallization and the following grain growth, acting
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as the significant barriers. Nevertheless, once the nanoparticles were coarsened during the high-temperature annealing, both the number of pinning sites and pinning force would be
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largely decreased, thereby in turn facilitating recovery and recrystallization process. As a consequence, it can be inferred that recrystallization in the as-fabricated 12Cr-ODS steel is presumably not able to occur until the nanoparticles are dissolving and/or coarsening during annealing, owing to the nanoscale and high concentration of oxide particles acting as strong
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barriers in the matrix.
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4. Conclusions
The microstructural stability of the as-fabricated 12Cr-ODS steel under elevatedtemperature annealing was investigated in the present work. The dependence of nanoparticle coarsening, microstructural evolution, and hardness variations upon the annealing conditions was systematically studied. The main conclusions can be drawn as
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follows:
(1) The nanoparticles exhibited high stability and continued precipitating during annealing at 1473 K. On the contrary, evident coarsening occurred at 1673 K, making the diameter of nanoparticles rise to 9.7 ± 2.3 nm and 18.3 ± 4.5 nm after annealing for 1 h and
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24 h, respectively. The apparent activation energy for coarsening was estimated as 412 ± 55 kJ/mol, and the coarsening function was derived as:
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dt = {1.5 × 1015 · exp [ - (412 ± 55)/(8.314·T)]·t + d03}1/3 where T , t and d0 are in Kelvin, hour and nanometer, respectively. (2) Development of porosity was observed in the annealed specimens, and the voids
distributed heterogeneously and increased distinctly at 1573 and 1673 K for longer annealing time. This is probably due to the growth and/or aggregation of Ar-entrapped and pre-existed voids formed during the manufacture processing.
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(3) No significant changes were observed in microstructure and texture at 1473 K. Nevertheless, during exposure at 1673 K, partial and full recrystallization were achieved,
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and the major texture was not altered except that its intensity reduced. The apparent activation energy for recrystallization was calculated as 424 ± 22 kJ/mol, which was almost the same as that for nanoparticle coarsening, indicating that recrystallization proceeded simultaneously with the coarsening of nanoparticles that would probably in turn stimulate
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the recrystallization process.
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Acknowledgements
This study was sponsored in part by a project “R&D of nuclear fuel cladding materials and their environmental degradations for the development of safety standards” entrusted to Tohoku University by the Ministry of Education, Culture, Sports, Science and Technology in Japan (MEXT). The authors would like to thank the support of China Scholarship
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Figures
Fig. 1 Bright-field TEM micrographs of nanoparticles in the specimens, which were annealed at 1473 K, 1573 K, and 1673 K for 1 h, 12 h, and 24 h, respectively.
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100 10
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Diameter (nm)
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Annealing time (h)
Fig. 2 The dependence of average diameter and number density of nanoparticles on the annealing time at 1473 K, 1573 K, and 1673 K. The solid dark and open blue symbols denote average
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diameter and number density, respectively.
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Fig. 3 (a) Plot of (dt3 - d03) as a function of annealing time, (b) Relationship of lnk against reciprocal absolute temperature. Some error shown in (a) is smaller than the symbol.
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Fig. 4 SEM micrographs of the specimens after annealing at 1473 K, 1573 K and 1673 K for 1 h, 12 h, and 24 h, respectively.
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Fig. 5 (a) The SEM and (b) bright-field TEM micrographs of voids in the as-fabricated specimens. The yellow triangles denote the voids in matrix.
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Fig. 6 IPF images of specimens after annealing at 1473 K, 1573 K and 1673 K for 1 h, 12 h and 24 h, respectively. The arrows present the rolling direction (RD) and transverse direction (TD), and dark lines denote the high angle grain boundaries (>15°).
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Fig. 7 φ2 = 45° sections of ODF maps of the specimens after annealing at 1473 K, 1573 K and 1673 K for 1 h, 12 h and 24 h. The symbols ■, ▲, ► denote {001}<110>, {111}<110>, and
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{111}<112> components, respectively.
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Fig. 8 Misorientation angle distribution of the specimens after annealing at 1473 K, 1573 K and
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1673 K for 1 h, 12 h and 24 h, respectively.
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Fig. 9 (a) Vickers hardness and (b) corresponding volume fraction of recrystallization as a function of annealing temperature and time.
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Fig. 10 (a) Plot of ln[ln(1/(1-Xrex))] as a function of lnt, (b) Correlation of lnt0.5 against reciprocal absolute temperature.
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Highlights The nanoparticles were stable at 1473 K, whereas, evident coarsening occurred
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at 1673 K.
Heterogeneous formation of voids was observed in the annealed specimens. Full recrystallization was achieved after annealing at 1673 K for 24 h.
The apparent activation energy for recrystallization was coincident with that for
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nanoparticle coarsening.
S0925-8388(16)33487-9
DOI:
10.1016/j.jallcom.2016.11.029
Reference:
JALCOM 39524
To appear in:
Journal of Alloys and Compounds
Received Date: 24 July 2016 Revised Date:
15 October 2016
Accepted Date: 2 November 2016
Please cite this article as: J. Shen, H. Yang, Y. Li, S. Kano, Y. Matsukawa, Y. Satoh, H. Abe, Microstructural stability of an as-fabricated 12Cr-ODS steel under elevated-temperature annealing, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.11.029. 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 proof before it is published in its final 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.
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Microstructural stability of an as-fabricated 12Cr-ODS steel under elevated-
Jingjie Shen
a,*
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temperature annealing , Huilong Yang b, Yanfen Li b, Sho Kano c, Yoshitaka Matsukawa b, Yuhki
Satoh b, Hiroaki Abe b,c
Graduate School of Engineering, Tohoku University, 2-1-1 Katahira, Sendai 980-8577,
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a
Japan b
Institute for Materials Research, Tohoku University, 2-1-1 Katahira, Sendai 980-8577,
c
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Japan
School of Engineering, The University of Tokyo, 2-22 Shirakata Shirane, Tokai-mura,
Ibaraki 319-1188, Japan *
Corresponding author:
Abstract
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E-mail address: [email protected] (J.J. Shen)
The effects of elevated-temperature annealing in the range of 1473 K - 1673 K on the
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coarsening kinetics of nanoparticles and microstructural evolution in an as-fabricated 12CrODS steel was systematically investigated. Results show that the nanoparticles were
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extremely stable and continued precipitating during annealing at 1473 K, whereas, significant coarsening took place during exposure at 1673 K, increasing by 177% and 423% in diameter for 1 h and 24 h, respectively. Furthermore, heterogeneously distributed voids were developed, sizes of which increased distinctly at 1573 K and 1673 K for longer holding time. This is probably due to the growth and/or aggregation of Ar-entrapped and pre-existed voids in the as-fabricated condition. Additionally, no obvious changes were observed in microstructure and texture at 1473 K, nevertheless, full recrystallization were achieved at 1673 K for 24 h, and the major texture was not altered except reduction of the
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intensity. The apparent activation energy for recrystallization was experimentally calculated as 424 ± 22 kJ/mol, which was almost the same as that for nanoparticle coarsening, i.e., 412
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± 55 kJ/mol, suggesting that recrystallization proceeded accompanying with nanoparticle coarsening that would in turn stimulate the recrystallization process. It was deduced that recrystallization in the as-fabricated 12Cr-ODS steel was presumably not able to occur until the nanoparticles were dissolving and/or coarsening during annealing.
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Keywords: oxide dispersion strengthened steels; nanoparticles; annealing; coarsening;
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voids; recrystallization
1. Introduction
Development of oxide dispersion strengthened (ODS) steels for structural materials in future fission and fusion reactors has been of increasing interest during the past decades, because of their adequately high thermal stability and strength, and excellent irradiation
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damage resistance [1-5].
In order to produce reduced-activation ODS steels, alloying elements with a slow decaying rate, such as Mo, Nb in conventional steels, are substituted by W, V, Ti, which have a higher decay rate of radioactivity, because the strengthening function of Mo and Nb
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in steels could be compensated by W, V and Ti [6, 7]. Owing to the moderate tensile and creep properties at high temperatures, the utilization of conventional high-chromium
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ferritic/martensitic steels, like modified 9Cr-1Mo, or the reduced-activation steels, such as F82H, ORNL 9Cr-2WVTa, EUROFER, and JLF-1, was limited to ~823 K [3, 8]. One way to increase the operating temperature is to add extremely stable oxide particles in the matrix. The high number density of nanoscale oxide particles induced by mechanical alloying play an important role in the performances. Not only preventing the dislocation migration and providing obstacles to resist the motion of grain boundaries, thereby ensuring the adequate strength and thermal stability at high temperatures, but also it can serve as sinks for irradiation induced defects and nucleation sits for helium bubbles, making ODS steels more
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resistant to irradiation hardening and swelling [9-11]. Therefore, the stability of those nanoscale oxide particles during both the fabrication processing (e.g., hot isotropic pressing,
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hot extrusion, hot rolling, annealing) and the operating temperatures (up to ~1073 K) or even higher is of crucial importance, which is directly responsible for the deterioration of strength, creep properties, and irradiation resistance.
Several investigations about the effects of high-temperature annealing on the thermal
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stability of ODS steels have been done. Schneibel et al. [12] reported that the density of 14YWT (Fe-14Cr-3W-0.4Ti-0.25Y2O3) decreased due to pore formation during annealing at 1273 K up to 30 h. Besides, Williams et al. [13] suggested that Ti addition did not affect
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the coarsening kinetics of oxide particles during annealing at 1473 K for 100 h, which was in accordance with the diffusion-controlled mechanism. Furthermore, Miller et al. [14, 15] found that the size of (Y, Ti, O)-enriched oxide particles in MA957 (Fe-14Cr-1Ti-0.3Mo0.25Y2O3) increased from 2.4 nm to 9.2 nm in diameter and the number density decreased from
~2 × 1024 m-3 to ~8 × 1022 m-3 after aging at 1573 K for 24 h. Although the
coarsening behavior of the nanoscale oxide particles had been observed in the previous
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studies, most of which were focused on a certain annealing temperature, so that the coarsening kinetics is not able to be analyzed, leaving the dependence of the annealing temperatures still unclear.
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In this work, we are aiming to investigate the microstructural evolution in an asfabricated 12Cr-ODS steel, and discuss the coarsening kinetics of Y-Ti-O nanoparticles
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under annealing at temperatures ranging from 1473 K to 1673 K. The findings are of the essence to provide experimental dataset for thermodynamic and kinetic study on the nucleation-growth-coarsening model of nanoparticles, so as to optimize the fabrication processing, and predict the long-term aging stability of ODS steels.
2. Experimental procedure
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The 12Cr-ODS steel with a nominal composition of Fe-12Cr-2W-0.3Ti-0.25Y2O3 was utilized in this work. The fabrication procedures and chemical composition can be found in
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detail elsewhere [16, 17]. The as-fabricated specimen was cut into small pieces and then isothermally annealed in a vacuum furnace at 1473 K, 1573 K and 1673 K for various time intervals ranging from 1 h - 60 h. The heating rate was 300 K/h, and the cooling rate was ~4000 K/h. In the following, specimens were designated after the annealing parameters.
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For instance, the specimen annealed at 1473 K for 1 h is hereafter marked as 1473 K-1 h. Microstructure observations were performed with conventional scanning electron microscope (SEM), and a JEOL JXA 8530 field-emission SEM equipped with electron
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backscatter diffraction (EBSD) device. Texture intensity in an area of 400 × 400 µm2 was described by orientation distribution function (ODF), which was computed by the series expansion method using TSL OIM analysis software. Besides, the orientation was presented in the form of the Euler angles {φ1, Ф, φ2} (Bunge’s notation) [18], which were calculated under the assumption of orthonormal sample symmetry, i.e., 0° ≤ {φ1, Ф, φ2} ≤ 90° [19]. The nanoscale oxide particles were characterized by a JEM-2100 transmission
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electron microscope (TEM) operating at 200 kV. TEM foils were firstly mechanically thinned down to ~80 µm and then electrochemically polished in a solution of 5 vol% perchloric acid in acetic acid by a twin-jet unit (Struers TenuPol-5) at room temperature.
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Thickness of the observed area was determined under two beam condition. The diameter of nanoparticles was measured using ImageJ software from at least 600 particles based on the
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TEM micrographs.
For the hardness and EBSD testing, the annealed specimens were ground with silicon
carbide paper up to 2000 grit on the rolling plane, and then polished with alumina and colloidal silica suspension. Vickers hardness was carried out under a load of 1 kg for a period of 15 s, and randomly 12 measurements were averaged after eliminating their minimum and maximum.
3. Results and discussion 4
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3.1 Effect of annealing on nanoscale oxide particles 3.1.1 Size and number density of nanoscale oxide particles
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Fig. 1 shows bright-field TEM micrographs of nanoscale oxide particles in the annealed specimens. It is noted that some of the oxide particles were cuboidal at 1473 K and 1573 K, while it became spherical at 1473 K. This is probably due to variation of interfacial structures between oxide particles and matrix, changing from coherent to
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incoherent structure [20, 21]. The average diameter and number density of the nanoscale oxide particles in each specimen were plotted as a function of annealing time in Fig. 2. As
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for the as-fabricated condition, the average diameter and number density were 3.5 ± 0.7 nm and ~1.3 × 1023 m-3 [17]. As shown in Fig. 2, the size of the nanoscale oxide particles did not obviously vary, and the number density slightly increased after annealing at 1473 K, indicating that Y, Ti, O elements were still supersaturated in the as-fabricated matrix, and numerous oxide particles were precipitating during annealing. In contrast, it was significantly coarsened after exposure at higher temperatures, resulting in lower number
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density. For instance, the average diameter amounted to 9.7 ± 2.3 nm and 18.3 ± 4.5 nm after annealing at 1673 K for 1 h and 24 h, which increased by 177% and 423%, respectively. Dependence of coarsening rate on temperature and holding time was clearly observed, which is presumably because the solute atoms (i.e., Y, Ti and O) associated with
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the oxide particles diffused more rapidly at higher temperatures.
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3.1.2 Coarsening kinetics of nanoscale oxide particles In general, growth of precipitates in alloys controlled by solute diffusion, so-called
Ostwald ripening [22], can be expressed by the Lifshitz-Slyozov-Wagner (LSW) theory [23, 24], which was applied to analyze the coarsening kinetics of the oxide particles in ODS steels [21, 25]. dt3 - d03 = kt
(1)
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where dt and d0 are the average diameter of nanoparticles in the annealed and as-fabricated specimens, t is the annealing time, and k is the coarsening rate constant that is dependent on
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annealing temperature. Eq. (1) demonstrates that the increment of volume of the nanoscale oxide particles has a linear correlation with respect to annealing time t, and k can be derived from linear fitting, which is presented in Fig. 3(a). Moreover, the coarsening rate k can be determined by Arrhenius equation:
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k = k0 exp (-Q/RT) (2)
where k0 is a constant, T is the absolute temperature, Q is the apparent activation energy for coarsening, R is the universal gas constant (8.314 J/mol/K). Taking natural logarithm of Eq.
lnk = lnk0 - Q/RT
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(2), it yields: (3)
According to Eq. (3), it is easy to get the slope from correlation of lnk against 1/T, as plotted in Fig. 3(b). Consequently, the apparent activation energy for coarsening of the nanoscale oxide particles, Q, was experimentally estimated as 412 ± 55 kJ/mol. For
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coarsening process of the nanoparticles, the system tends to minimize the free energy by reducing the interface area between particles and matrix [22]. In terms of this fact, Y, Ti, O atoms need to dissolve from the surface of smaller particles, diffuse through the matrix and redeposit onto the larger particles, resulting in the growth of larger ones. In fact, the size of
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Ti, Y atoms is much larger than Fe, Cr atoms [15], and the diffusion of Y atoms in the FeCr matrix is rather difficult, thereby giving rise to a sluggish diffusion process. Williams et
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al. [13] presumed that diffusivity of Y atoms dominated the coarsening rate of Y-Ti-O enriched nanoparticles since the diffusion coefficient of Y is significantly lower than that of Ti and O atoms. However, the activation energy of Y for diffusion in matrix is 218.1 kJ/mol [26], which is clearly smaller than the experimental estimation in this work. Thus, the coarsening of Y-Ti-O nanoparticles may not be controlled only by the diffusion of Y atoms. Previous research [15, 27, 28] indicated that a large amount of TiO, YO, CrO molecular ions were detected in matrix, because of the strong solute-oxygen affinity. Those solute-oxygen nanoclusters may diffuse as an entity during annealing, which would retard
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the coarsening of the nanoscale oxide particles, resulting in much larger activation energy. Additionally, the high concentration of vacancies in the nanoclusters probably stimulate the
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diffusivity of solute-oxygen complex in matrix [29, 30]. As a result, the coarsening equation for nanoscale oxide particles in this 12Cr-ODS steel can be written by: (4)
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dt = {1.5 × 1015 · exp(-(412 ± 55)/(8.314·T))·t + d03}1/3
where T , t and d0 are in Kelvin, hour and nanometer, respectively.
Therefore, it is probably able to estimate the coarsening rate and predict corresponding
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average diameter of the nanoparticles under a certain annealing condition with the aid of Eq. (4). For instance, comparison of growth time of nanoparticles from the as-fabricated condition to 5.2 nm, the average diameter in 1473 K-24 h specimen, would correspond to ~ 7.7 × 106 h during exposure at 1073 K, suggesting the nanoparticles should be remarkably stable at this or lower service temperatures. Nevertheless, the long-term annealing
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experiments at lower temperatures are needed to validate this prediction. 3.2 Effect of annealing on microstructure 3.2.1 Formation of voids upon annealing
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Accompanying with the coarsening of nanoparticles, microstructure evolution of the matrix was also characterized. Fig. 4 shows SEM micrographs of the annealed specimens.
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Inhomogeneous microstructure development associated with voids was evident. The voids were observed in the annealed specimens. The number density and size of voids increased due to annealing at higher temperature and longer duration. In order to elucidate the formation mechanism of these voids, the microstructure of as-
fabricated specimens was also examined. Fig. 5 represents the micrographs of the asfabricated specimen in SEM and TEM observations. As shown in Fig. 5(a), it is clear that voids with a diameter ranging from tens of nanometers to several micrometers existed in the as-fabricated matrix. Moreover, in Fig. 5(b), voids trapped around the nanoparticles and
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embedded in matrix were also observed, which were firstly confirmed as Ar cavities by the electron energy loss spectrometer in Ref. [31], retaining from mechanical alloying in the
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argon atmosphere. Accordingly, two potential reasons for the voids formation are proposed: one is likely that there are pre-existed voids in the as-fabricated condition, and then growth and/or agglomeration took place in the subsequent heat treatment; the other one is probably ascribed to entrapment of argon gas at the surface of the oxide particles. Turker [32] reported that, in PM2000 (Fe-20Cr-5.5Al-0.5Ti-0.5Y), less porosity was observed after
3.2.2 Recrystallization behavior
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which was explained by the entrapped-gas mechanism.
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mechanical alloying in hydrogen atmosphere compared to that in an argon atmosphere,
Fig. 6 and Fig. 7 demonstrate the inverse pole figure (IPF) images and the corresponding φ2 = 45° sections of ODF maps of the annealed specimens. As can be seen, no significant changes in microstructure were observed at 1473 K, as well as texture shown
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in Fig. 7. It is obviously notable that the dominant texture was composed of intense {001}<110> component and relatively weaker γ-fiber with <111> axis parallel to normal direction. This texture was akin to that of as-fabricated condition [17], even after annealing up to 24 h, revealing that the microstructure of the as-fabricated 12Cr-ODS steel was stable.
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Two possible factors are responsible for the thermal stability: on the one hand, {001}<110> texture retards the recrystallization, due to containing the lowest stored energy that is the
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driving force for recrystallization [33, 34]; on the other hand, large number density of nanoscale oxide particles inhibit the migration of dislocations and grain boundaries during annealing. It is due to the strong pinning effect of nanoparticles and the lowest stored energy of {100}<110> texture [33], thereby preventing dislocations and grain boundaries from migration and retarding the recrystallization process. At 1573 K, it was found that only a few proportion of grains were growing during exposure up to 24 h (Fig. 6), and the intensity of texture gradually decreased (Fig. 6). In contrast, as for 1673 K-12 h specimen, evident growth of grains was observed, and full recrystallization was achieved with further
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annealing for 24 h. As shown in Fig. 7, the texture has two characteristic peaks, namely, {001}<110> and {111}<110>. The structure was principally kept throughout the
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experimental conditions. The evident alteration was observed in the intensity reduction in {001}<110> and {111}<110> components. This suggests that the nucleation sites for recrystallization were limited by the original crystallographic distribution.
Furthermore, as well known that recrystallization is a migration process of high-angle
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grain boundaries. Misorientation, which is defined as a minimum rotation angle required to bring the neighboring grain into coincidence, was also measured by EBSD for each specimen in Fig. 6. Fig. 8 shows the distributions of misorientation angles in the specimens.
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It is evident that the average misorientation angle slightly increased at the early stage of annealing by comparison with the as-fabricated specimen, which was composed of a majority of highly oriented {001}<110> grains [17], giving a lower misorientation (~ 17.5°). It seems to be very stabilized during exposure at 1473 K, indicating recovery may be still in progress. On the other hand, owing to partial recrystallization, the proportion of low misorientation reduced, and corresponding high misorientation increased, resulting in
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obvious increase of the average misorientation. As shown in Fig. 8, the proportion of low misorientation decreased a little larger in the 1573 K-24 h and 1673 K-12 h specimens, respectively. As a consequence of full recrystallization and/or grain growth, the
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misorientation angle increased up to ~ 38.7° after heating at 1673 K up to 24 h, in which the grain boundaries changed into high angle grain boundaries.
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In addition, it is found that the recrystallized grains grew along the rolling direction, forming elongated grains, which is a common feature in recrystallized ODS ferritic steels. This is presumably ascribed to the alignment of oxide particles along the extrusion direction (or cold rolling direction), resulting in an anisotropic pinning force that hinders grain boundary movements effectively in the transverse direction [35-37]. Furthermore, in particle embedded materials, it is widely accepted that diameter of the critical maximum grain (Dc) can be quantitatively estimated by Zener equation [38], whose general form is written as:
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Dc = Kd / f m (5) where K is a dimensionless constant, d and f are the diameter and volume fraction of
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pinning nanoparticles, and m is an index for f. According to Ref. [38], parameters K = 0.17 and m = 1.0 were reasonably chose for this work, where the volume fraction f was 0.67% for the fully recrystallized specimen (1673 K-24h). Substituting those parameters into Eq. (5), the calculated diameter of grains was ~0.46 µm, which was approximately two orders
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of magnitude lower than the experimental measurements (~46.8 µm). This deviation may probably stem from the artificial error of determination on parameters (k, m, f), or other
3.3 Effect of annealing on hardness 3.3.1 Variations of hardness
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underlying mechanisms tailoring the grain size, which is necessary to be verified in future.
Due to softening during annealing is associated with the recrystallization, the volume
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fraction of recrystallization (Xrex) was generally calculated from Vickers hardness testing [39, 40], which can be given by:
Xrex = (HV0 – HVt)/(HV0 – HVrex)
(6)
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where HV0 and HVt are the Vickers hardness for the as-fabricated and annealed specimens at a given condition, respectively. HVrex is the hardness of a fully recrystallized specimen.
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Here, the 1673 K-24 h specimen was considered as full recrystallization (~ 100%) based on EBSD measurements, as shown in Fig. 6. Fig. 9 presents Vickers hardness and volume fraction of recrystallization as a function
of annealing condition. As shown in Fig. 9(a), it can be noticed that the hardness slightly decreased at 1473 K up to 24 h, however, obvious reduction was observed after annealing at 1573 K and 1673 K for more than 6 h. The corresponding recrystallization fraction based on hardness variations is shown in Fig. 9(b). Evidently, the recrystallization fraction levels off at 1473 K from 6 h to 24 h, indicating that no obvious changes were emerged in
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microstructure, and the specimen remained in recovery stage. This can be confirmed from microstructure evolution at 1473 K, as shown in Fig. 6. On the contrary, the fraction
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increased directly during annealing at 1573 K for 6 h, and 1673 K for 1 h, due to occurrence of recrystallization, leading to more proportion of softening. Together with Fig. 6, one can conclude that the strain recovery occurred at 1473 K for 24 h, partial
3.3.2 Apparent activation energy of recrystallization
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recrystallization started at 1573 K, and grain growth took place at 1673 K for 24 h.
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To describe the kinetics of isothermal recrystallization, Johnson-Mehl-AvramiKolmogorov (JMAK) equation is commonly accepted as [41-43]: Xrex = 1 – exp (– ktn)
(7)
where k is a temperature dependent constant and n is referred to as the Avrami exponent, which is a function of both nucleation and growth rates. Eq. (7) can also be given as: (8)
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ln[ln(1/(1 – Xrex))] = nlnt + lnk
Eq. (8) yields a straight line whose slope is equal to n when the left-hand side is plotted against lnt, which is shown in Fig. 10(a). The values of Avrami exponent at 1473 K, 1573
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K and 1673 K ranged from 0.41 to 0.9, which is deviated from the theoretical values, i.e., n = 2 for 2D growth, and n = 3 for 3D growth under assumption of site-saturated nucleation
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during recrystallization [44]. Several researches on the static recrystallization kinetics of Fe-based alloys claimed the similar results. For instance, Xie et al. [45] reported that the value of Avrami exponent was estimated as 0.54 in a low-carbon steel, and Malekjani et al. [46] calculated Avrami exponent of austenitic and ferritic stainless steels, which was in the range of 0.75 - 1.1. For the deviation of the Avrami exponent in this study, it is probably an open question.
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The annealing time up to a certain level of recrystallization can be described by Arrhenius equation. It was defined the annealing time as the time that was required for 50%
t0.5 = t0 exp (Qrex / RT)
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recrystallization, t0.5, which is expressed by the following formula [47]. (9)
where t0 is a constant, and Qrex is the apparent activation energy of recrystallization.
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Taking natural logarithm of both sides of Eq. (9), the following expressions are given by:
k0 = Qrex/R
(10)
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lnt0.5 = k0/T + lnt0 (11)
As presented in Fig. 10(a), if the values of n and lnk were obtained by linear fitting, then, lnt0.5 at each temperature is able to be calculated by substituting Xrex = 0.5 into Eq. (8). Fig. 10(b) illustrates that values of lnt0.5 were plotted as a function of 1000/T. The slope, Qrex/R, was attained by linear fitting. Accordingly, the apparent activation energy of
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recrystallization for the as-fabricated 12Cr-ODS steel was estimated as 424 ± 22 kJ/mol, which is much larger than that of self-diffusion in α-iron and δ-iron with 251 kJ/mol [48] and 296 kJ/mol [49]. Nevertheless, similar results were reported in the previous research.
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For instance, Liu et al. [50] estimated the activation energy for recrystallization in the coldrolled V-doped steel, which was 474 kJ/mol, suggesting that the dispersed V(C,N) precipitates prevented the growth of recrystallized grains. The high value in this study was
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presumably resulted from the pinning effect of nanoscale oxide particles against the motion of dislocations and grain boundaries. Additionally, it is surprisingly noteworthy that the apparent activation energy for
recrystallization was nearly the same as that for nanoparticle coarsening, which was 412 ± 55 kJ/mol as estimated in Section 3.1.2. This reveals that coarsening of nanoparticles took place accompanying with the recrystallization process. In other words, the nanoparticles play a pivotal role in retarding the recrystallization and the following grain growth, acting
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as the significant barriers. Nevertheless, once the nanoparticles were coarsened during the high-temperature annealing, both the number of pinning sites and pinning force would be
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largely decreased, thereby in turn facilitating recovery and recrystallization process. As a consequence, it can be inferred that recrystallization in the as-fabricated 12Cr-ODS steel is presumably not able to occur until the nanoparticles are dissolving and/or coarsening during annealing, owing to the nanoscale and high concentration of oxide particles acting as strong
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barriers in the matrix.
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4. Conclusions
The microstructural stability of the as-fabricated 12Cr-ODS steel under elevatedtemperature annealing was investigated in the present work. The dependence of nanoparticle coarsening, microstructural evolution, and hardness variations upon the annealing conditions was systematically studied. The main conclusions can be drawn as
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follows:
(1) The nanoparticles exhibited high stability and continued precipitating during annealing at 1473 K. On the contrary, evident coarsening occurred at 1673 K, making the diameter of nanoparticles rise to 9.7 ± 2.3 nm and 18.3 ± 4.5 nm after annealing for 1 h and
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24 h, respectively. The apparent activation energy for coarsening was estimated as 412 ± 55 kJ/mol, and the coarsening function was derived as:
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dt = {1.5 × 1015 · exp [ - (412 ± 55)/(8.314·T)]·t + d03}1/3 where T , t and d0 are in Kelvin, hour and nanometer, respectively. (2) Development of porosity was observed in the annealed specimens, and the voids
distributed heterogeneously and increased distinctly at 1573 and 1673 K for longer annealing time. This is probably due to the growth and/or aggregation of Ar-entrapped and pre-existed voids formed during the manufacture processing.
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(3) No significant changes were observed in microstructure and texture at 1473 K. Nevertheless, during exposure at 1673 K, partial and full recrystallization were achieved,
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and the major texture was not altered except that its intensity reduced. The apparent activation energy for recrystallization was calculated as 424 ± 22 kJ/mol, which was almost the same as that for nanoparticle coarsening, indicating that recrystallization proceeded simultaneously with the coarsening of nanoparticles that would probably in turn stimulate
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the recrystallization process.
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Acknowledgements
This study was sponsored in part by a project “R&D of nuclear fuel cladding materials and their environmental degradations for the development of safety standards” entrusted to Tohoku University by the Ministry of Education, Culture, Sports, Science and Technology in Japan (MEXT). The authors would like to thank the support of China Scholarship
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Figures
Fig. 1 Bright-field TEM micrographs of nanoparticles in the specimens, which were annealed at 1473 K, 1573 K, and 1673 K for 1 h, 12 h, and 24 h, respectively.
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100 10
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30
1473 K 1573 K 1673 K
10 0
0
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0.1
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Diameter (nm)
40
12
16
20
0.01
Number density (1022 m-3)
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0.001 24
1E-4
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Annealing time (h)
Fig. 2 The dependence of average diameter and number density of nanoparticles on the annealing time at 1473 K, 1573 K, and 1673 K. The solid dark and open blue symbols denote average
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diameter and number density, respectively.
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Fig. 3 (a) Plot of (dt3 - d03) as a function of annealing time, (b) Relationship of lnk against reciprocal absolute temperature. Some error shown in (a) is smaller than the symbol.
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Fig. 4 SEM micrographs of the specimens after annealing at 1473 K, 1573 K and 1673 K for 1 h, 12 h, and 24 h, respectively.
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Fig. 5 (a) The SEM and (b) bright-field TEM micrographs of voids in the as-fabricated specimens. The yellow triangles denote the voids in matrix.
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Fig. 6 IPF images of specimens after annealing at 1473 K, 1573 K and 1673 K for 1 h, 12 h and 24 h, respectively. The arrows present the rolling direction (RD) and transverse direction (TD), and dark lines denote the high angle grain boundaries (>15°).
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Fig. 7 φ2 = 45° sections of ODF maps of the specimens after annealing at 1473 K, 1573 K and 1673 K for 1 h, 12 h and 24 h. The symbols ■, ▲, ► denote {001}<110>, {111}<110>, and
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{111}<112> components, respectively.
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Fig. 8 Misorientation angle distribution of the specimens after annealing at 1473 K, 1573 K and
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1673 K for 1 h, 12 h and 24 h, respectively.
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Fig. 9 (a) Vickers hardness and (b) corresponding volume fraction of recrystallization as a function of annealing temperature and time.
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Fig. 10 (a) Plot of ln[ln(1/(1-Xrex))] as a function of lnt, (b) Correlation of lnt0.5 against reciprocal absolute temperature.
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Highlights The nanoparticles were stable at 1473 K, whereas, evident coarsening occurred
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at 1673 K.
Heterogeneous formation of voids was observed in the annealed specimens. Full recrystallization was achieved after annealing at 1673 K for 24 h.
The apparent activation energy for recrystallization was coincident with that for
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nanoparticle coarsening.