Tailoring the secondary phases and mechanical properties of ODS steel by heat treatment

1. Introduction

Nuclear energy has been rapidly developed for several decades in order to solve the problem of worldwide energy shortage [1,2]. Greater demands are emphasized on the nuclear reactors and structural materials due to the increasing development of nuclear energy [3,4]. Oxide dispersion strengthened (ODS) steels are promising candidates for the cladding materials of the Generation IV nuclear fission reactors and the blanket of the future fusion reactors. ODS steels possess excellent performances in harsh environment such as high temperature, neutron irradiation, corrosion and oxidation [5,6]. This can be attributed to the stable nanoparticles in the matrix, which effectively hinder the dislocation glide and grain boundary migration [7,8].

ODS steels are usually fabricated by mechanical alloying and subsequent consolidation [2]. As is well known, the addition of Ti is effective to refine and stabilize the nanoparticles [9], [10], [11], [12], [13], [14]. The correlation between chemical composition and size of the Y—Ti—O nanoparticles has been investigated by Sakasegawa et al. [9]. In addition, the Y—Ti—O nanoparticles such as Y₂Ti₂O₇, Y₂TiO₅ and YTiO₃ have been identified by some researchers [12,14]. The formation of non-stoichiometric Y—Ti—O nanoclusters and stoichiometric Y—Ti—O particles has been proved to be related to the increase of titanium content. More stable nanoclusters can be obtained by adjusting the composition of ODS steels [9,12]. The addition of Al in ODS steels can obviously improve the corrosion resistance in (lead-bismuth eutectic) LBE [15], while the coarser Y—Al—O compounds are detrimental to the strength of the ODS steels [16], [17], [18], [19], [20]. Zr and Hf are added into the Al-containing ODS steels to balance the corrosion resistance and the mechanical properties by forming finer nanoparticles [21], [22], [23], [24], [25], [26], [27]. Hf can not only improve the strength and ductility of the ODS steel, but also suppress the Cr depletion at grain boundaries by inhibiting the combination of Cr and self-interstitials [23,26,[22], [23], [24], [25], [26], [27].

Some secondary phases with large size in the ODS steel were also identified except for nanoparticles. Li et al. [28] found Cr-rich phases, AlN, TiN, YAl and Al₂O₃ in the 18Cr-4Al ODS ferritic steel. Song et al [29] investigated the M₂₃C₆ particles and Cr-rich oxides in a 12Cr ODS steel and refined these large secondary phases by equal channel angular extrusion. Li et al. [30] also found Ti-rich particles in the ODS steels. Auger et al. [31] found that the Cr-rich particles in the ODS steel are effective to inhibit the recovery and grain growth process. Olier et al. [32] mentioned that the formation of Ti(C,N) compounds and some FeCrW carbides (M₂₃C₆ type) was due to the high carbon contamination during mechanical alloying. Rouffié et al. [33] has mentioned the existence of M₂₃C₆ particles in the ODS steel but not carried on the further analysis about such secondary phases.

Among the large secondary phases mentioned above, MX and M₂₃C₆ have been widely investigated as the normal carbides in the non-ODS steels. Li et al. [34] concluded that the large M₂₃C₆ formed during normalization could be eliminated by austenitizing in the martensitic heat-resistant steels. Such heat treatment is beneficial to the toughness of the heat-resistant steels. Lots of researchers have studied the evolution of different carbides during heat treatment in various steels [35], [36], [37], [38], [39], [40], [41]. Taneike et al. [35] investigated the precipitate locations and distribution conditions of the MX and M₂₃C₆ in the steels with different C contents and analyzed the mechanisms. The investigations on secondary phases in the non-ODS steels are always related to the heat treatment [35,42]. Heat treatment is likewise an indispensable step during the production process of the ODS steels. The long-term thermal-stability of ODS steels has been investigated by heat treatment, as well as the Ti-rich precipitates and M₂₃C₆ particles [43,44]. Xu et al. [45] annealed the ODS powders at different temperatures, and concluded that the grain and nanoparticle size increased with the increasing annealing temperature. Oksiuta et al. [46,47] and Cunningham et al. [48] found the same phenomenon in the HIPed and hot extruded ODS steels. Besides, Shen et al. [49] obtained the activation energy of recrystallization and nanoparticle coarsening in ODS steels by heat treatment, and analyzed their voids evolution. Nanoparticle coarsening and some other large precipitates have also been identified in the HIPed ODS steel during annealing [50]. Most emphases have been put on the nanoparticles when the ODS steels were heat treated [51], [52], [53]. Both the nano size and large size secondary phases influence the microstructure and mechanical properties of the ODS steel. The effect of heat treatment on the secondary phases with different size has not been adequately developed.

In the present work, the evolution of density, microstructure, secondary phases and mechanical properties of a 14Cr ODS steel during annealing at different temperatures has been investigated. The evolution of secondary phases and its influence on the mechanical properties have been adequately discussed.

2. Experimental

The ODS steel was fabricated by powder metallurgy process. The atomization method was utilized to prepare the pre-alloyed powders with the nominal composition of Fe-14Cr-2W-0.2V-0.07Ta (wt%). The pre-alloyed powders were then mechanically alloyed with 0.3 wt% Ti and 0.3 wt% Y₂O₃ for 30 h at 400 rpm with a ball to powder ratio of 15:1 under high purity argon atmosphere. The milling process was performed in a planetary ball mill (QM-3SP4, Nanjing NanDa Instrument Plant, China). The consolidation of the mechanically alloyed powders was conducted by hot isostatic pressing (HIP) at 1150 °C for 3 h to obtain the ODS steel with the composition of Fe-14Cr-2W-0.2V-0.07Ta-0.3Ti-0.3Y₂O₃. The pressure was 150 MPa. The as-HIPed ODS steel was cut up for the subsequent heat treatment. Heat treatments of these ODS samples were performed at 800 °C, 1000 °C, and 1200 °C for 5 h with water cooling, respectively.

The overall microstructures and precipitates were observed by scanning electron microscopy (SEM, SU1510) and transmission electron microscopy (TEM, JEM-2100F). The ODS samples for SEM observation were prepared by mechanical grinding and polishing. The polished samples were then etched in a solution composed of 5 g copper chloride, 100 ml hydrochloric acid, 100 ml ethyl alcohol. Two kinds of samples were utilized for the TEM observation: thin foil samples and carbon extraction replicas. For the preparation of thin foils, the 300 μm thick slices were cut from the as-HIPed ODS steel, mechanically grinded to 50 μm, punched into 3 mm diameter discs, and then electropolished in a double jet electropolishing device at -20 °C. The electrolyte is a solution of 5% perchloric acid and 95% ethanol. The carbon replicas were obtained by coating a thin layer of C on the SEM samples, detaching the layers and collecting them on the Cu grids. The chemical extraction method was used to prepare the X-ray diffraction (XRD) samples. The matrix was dissolved in the 20% HCl solution. The remaining precipitates were centrifuged, washed and dried before XRD observation. The carbon content was measured by Electron Probe Micro-analyzer (EPMA, JEOL JXA-8100), and it refers to the content in matrix. The tensile tests were carried out at room temperature with a nominal strain rate of 10^-3 s^-1. Two measurements were performed on each sample. The gauge length of the samples was 4 mm. The fracture surfaces of the post-tensile samples were investigated by SEM. Densities of the ODS samples were measured by electronic density balance based on Archimedes principle. Ten measurements were performed on each sample. The average value of tested results is considered as the density of each sample.

3. Results and discussion

3.1. Microstructure

Fig. 1 shows the overall microstructures of the as-HIPed and annealed ODS steels. As can be seen in Fig. 1(a), the as-HIPed ODS steel possesses a heterogeneous microstructure consisting of recrystallized and recovered/deformed grains. The regions marked by dotted ellipses present the recrystallized grains. The size of the recrystallized grains approximately floats from 1 to 10 μm. The regions marked by rings are recovered/deformed grains. Such regions always possess a lamellar microstructure which is similar with the as-milled particles. All the ODS samples expect for the one annealed at 1200 °C possess such lamellar microstructure. After repeating welding and fracture during mechanical alloying process, the particles with folded, lamellar structure form [1,5,54]. During the sintering process, the deformed grains which fail to meet the driving force of recrystallization retain the recovered/deformed morphologies. The grains that can achieve the recrystallization driving force show recrystallized morphologies after consolidation. Fig. 1(b) presents the microstructure of the ODS steel annealed at 800 °C for 5 h. The microstructure also consists of recrystallization and recovered/deformed grains. But the amount of recrystallized grains slightly increases when compared with the as-HIPed ODS steel. Such phenomenon indicates that annealing promotes the recrystallization process to some extent. Whether the recrystallization occurs depends on the heat treatment temperature and the strain energy stored in grains. At a specific temperature, recrystallization is more likely to occur in the grains stored high strain energy [1,54]. When the annealing temperature is increased, recrystallization can be further promoted, which can be seen in Fig. 1(c). More recrystallized grains can be found in the ODS steel annealed at 1000 °C for 5 h. The nanoscale precipitates in ODS steel are capable of impeding the dislocation glide and grain boundary motion, which effectively inhibits the recrystallization process. Therefore, the ODS steels annealed at 800 °C and 1000 °C for 5 h failed to obtain a fully recrystallized microstructure. Still, there is an obvious difference between the ODS steel annealed at 800 °C and the one annealed at 1000 °C. Minimal voids exist in the ODS steel annealed at 1000 °C. The size and number density of the voids in the ODS steel annealed at 1000 °C cannot be accurately measured because of the small amount. Micropores cannot be avoided in the as-HIPed ODS steel, and they are difficult to be observed in the as-fabricated samples as their sizes are small. One of the reasons for the micropore formation is related to the mechanical alloying process which results in large amounts of dislocations and vacancies in the grains. In addition, the powder metallurgy method is utilized in the fabrication of ODS steel [1,3,5,6,54]. It is impossible to achieve a theory density during the consolidation process, and the micropores are inherent defects in the as-HIPed ODS steels. In general, the micropores are trapped by the nanoscale precipitates or dislocations in the microstructure. However, the micropores are temperature-sensitive defects and may diffuse in the microstructure at an appropriate temperature. The diffuse micropores tend to accumulate and form voids on the grain boundaries [30]. More voids can be seen in the microstructure of the ODS steel annealed at 1200 °C for 5 h (Fig. 1(d)). The ODS steel possesses a fully recrystallized microstructure with large amount of voids distributed along grain boundaries. The size and number density of the voids in Fig. 1(d) are 0.56 ± 0.18 μm and 3.9 × 10¹⁶/m³, respectively, which are obviously increased when compared with that in Fig. 1(c). Shen et al. [49] have found that the number density and size of voids increased with the increasing annealing temperature when the ODS steel was annealed in the temperature range of 1200 °C-1400 °C. When the annealing temperature is high enough, the nanoscale precipitates are not effective to trap the micropores and inhibite the grain boundaries migration. Moreover, the dislocation glide toward grain boundaries is accelerated, thus facilitating the motion of micropores pinned by the dislocations [30,55,56].

The TEM graphs of the as-HIPed and annealed ODS steels are presented in Fig. 2. The as-HIPed ODS steel in Fig. 2(a) presents the equiaxed grains with a few hundred nanometers in size. A large recrystallized grain surrounded by some small grains can be seen in Fig. 2(b), which proves the occurrence of recovery and recrystallization. The recrystallization mechanism maybe attributed to the theory of subgrains coalescence and growth based on the microstructure in TEM graphs [51]. Identical phenomenon is also depicted in the ODS steels annealed at 1000 and 1200 °C, which is shown in Fig. 2(c) and (d) respectively. The recovery and recrystallization phenomenon occurred during the annealing process can be inferred according to the SEM graphs in Fig. 1. The analysis based on the TEM graphs provides the direct evidence for such phenomenon.

The grain size distribution of ODS samples is shown in Fig. 3. Both large and small grains exist in the ODS samples. The grain size distribution results in Fig. 3(a) and (b) possess the bimodal characteristics, and the mean sizes of grains are approximately 6.03 μm and 5.78 μm. After annealing at 1000 °C and 1200 °C, the grain size distributions in the ODS samples exihibit unimodal characteristics. The amount of submicron grains decreases with the increasing annealing temperature. The mean grain sizes in Fig. 3(c) and (d) are 4.51 μm and 6.02 μm. The increase of annealing temperature improves the recovery and recrystallization process. Therefore, the average grain size of the ODS steel decreases after annealing at 1000 °C. The fully recrystallized microstructure is obtained after annealing at 1200 °C, and the average grain size increases when compared with the sample annealed at 800 °C and 1000 °C. This can be attributed to the grain growth phenomenon after recrystallization.

Fig. 4 displays the XRD patterns and the calculated dislocation densities of the ODS samples. The dislocation densities are obtained on the basis of XRD results by using the following equation [57]:

ρ= $\frac{\beta ^{2}}{(4.3 b \times 10^{-2})^{5}}$ (1)

where b is Burger vector (2.5 × 10^-8 m [58]), and β is the half-width of the XRD patterns. The dislocation densities gradually decrease with the increasing annealing temperature, which is related to the dislocation recovery and recrystallization process during annealing.

3.2. Density

Densities of the as-HIPed and annealed ODS steels are displayed in Fig. 5. The actual density of ODS steels decreases with the increase of heat treatment temperature. As annealing temperature increases from 800 °C to 1200 °C, the actual density decreases from (7.7353 ± 0.006 ± 0.006) g/cm³ to (7.7000 ± 0.0037 ± 0.0037) g/cm³. Densities of the ODS steels show a slight drop after annealing at 800 °C and 1000 °C. Nevertheless, an obvious decline of density occurs after annealing at 1200 °C. The results observed above are in good agreement with the microstructure depicted in Fig. 1. No voids are investigated in the as-HIPed ODS steel and the one annealed at 800 °C A handful of voids exists in the ODS steel annealed at 1000 °C. The ODS steel annealed at 1200 °C possesses large amounts of voids. The phenomenon depicted above is consistent with the density results. Density is one of the factors needed to be considered for choosing an appropriate heat treatment temperature.

3.3. Secondary phases

As the EDS results of the same kind of particles show little differences, we choose one of the results to represent the whole element contents of the particles. The M₂₃C₆ and TiC particles in the as-HIPed ODS steel are marked by white and red arrows respectively in Fig. 6(a). The size of M₂₃C₆ particles is larger than that of TiC. The size of the TiC particles is approximately several hundred nanometers with a flaky or spherical morphology. Most of the carbides tend to distribute along the grain boundaries. The defects located on the grain boundaries are beneficial for the nucleation of carbides. Moreover, Ti tends to accumulate on the grain boundaries, thus promoting the formation of the TiC particles [7,10]. In general, the existence of C is related to the fabrication process. Powder metallurgy is used for the fabrication of ODS steel. Therefore, C is unavoidable in the pre-alloyed powders. Moreover, the stainless steel jars would cause C contamination during mechanical alloying [28,32]. The M₂₃C₆ and TiC particles also exist in the ODS steel annealed at 800 °C. Similar with the as-HIPed ODS steel, the M₂₃C₆ and TiC particles are also marked by white and red arrows in Fig. 6(b). The distribution of carbides in the ODS steel annealed at 1000 °C is shown in Fig. 6(c). The quantity and size of the M₂₃C₆ particles in Fig. 6(c) seem to be smaller than that in the as-HIPed ODS steel and the ODS steel annealed at 800 °C. The distribution of carbides in the ODS steel annealed at 1200 °C is exhibited in Fig. 6(d). No M₂₃C₆ particles can be seen in the matrix except for large amounts of TiC particles. The above results have led us to speculate that the M₂₃C₆ particles may gradually dissolve in the matrix during annealing, thus promoting the formation of TiC particles.

The carbon extraction replicas were utilized to investigate the TiC particles. Fig. 7(a)-(d) depicts the TiC particles in the as-HIPed ODS steel and the ODS steels annealed at 800 °C, 1000 °C and 1200 °C, respectively. The inserted graphs are the EDS patterns of the TiC particles marked by Arabic numerals. The size and morphology of the TiC particles are identical as that described in Fig. 6. The carbon replicas are collected on the Cu grids, and the peaks of Cu and C in the EDS patterns are inherent. In the as-HIPed ODS steels, the EDS patterns also show the peaks of Ta except for the peaks of Ti, which means the particles marked in Fig. 7(a) are rich in Ti and Ta. As mentioned in the experimental procedure, 0.07 wt% Ta is inherent in the pre-alloyed powders, and 0.3 wt% Ti is added into the pre-alloyed powders during the mechanical alloying process. Both Ta and Ti are strong carbide forming elements. They can combine with C to form TaC and TiC, which efficiently trap the detrimental C. In addition, the formation of TaC and TiC particles can effectively inhibit the generation of harmful M₂₃C₆ particles [35]. The amount of TiC particles is probably more than the TaC particles because the content of Ti is more than that of Ta. Moreover, TaC is able to dissolve in TiC [35,59]. This may make it difficult to investigate the TaC particles. Therefore, it is assumed that the carbides with similar morphology and size to the particles depicted in Fig. 7 are mainly TiC.

Based on Fig. 6, Fig. 7, the average sizes of M₂₃C₆ and TaC in the ODS samples are calculated and shown in Table 1. The average size of M₂₃C₆ decreases with the increasing annealing temperature, indicating the dissolution phenomenon of M₂₃C₆ during annealing process. Almost all the M₂₃C₆ particles are dissolved in the matrix after annealing at 1200 °C. The average size of TiC in the as-HIPed ODS steel and the steel annealed at 800 °C and 1000 °C shows no obvious differences. When the annealing temperature reaches to 1200 °C, the average size of TiC increases, meaning the slight coarsening of TiC. In order to further support the above statement about the carbides evolution during annealing, the contents of C in the matrix of ODS steels are measured and shown in Table 2. C in the pre-alloyed powders is inevitable, and the content of C increases during the fabrication process. The content of C in the annealed ODS steels increases with the increasing annealing temperature. This phenomenon implies that the increase of annealing temperature promotes the dissolution of M₂₃C₆.

The distribution of nanoparticles in the grains was examined by using TEM foil samples in order to make a though and reliable investigation. As shown in Fig. 8, the distributions of nanoparticles in the as-HIPed and annealed ODS samples are uniform. The average size and number density of nanoparticles are listed in Table 3. The average size gradually increases with the increasing annealing temperature. But the increase in average size is small, and no obvious coarsening of nanoparticles can be observed. The number density of nanoparticles located inside the grains decreases as the annealing temperature increases.

No matter the carbides or the nanoparticles, the results obtained above are all qualitative. The chemical extraction method and XRD investigation were utilized aiming at performing a comprehensive analysis on all kinds of precipitations. The XRD patterns of the precipitates are displayed in Fig. 9. Fig. 9(a) shows the XRD pattern of the secondary phases in the as-HIPed ODS steel. Peaks of M₂₃C₆, TiC and Y₂Ti₂O₇ are observed in the XRD pattern. No peaks of the TaC particles are found because of the low content, which is consistent with the analysis about Fig. 7. The characteristic peaks of the M₂₃C₆ particles are higher in intensity than that of the TiC and Y₂Ti₂O₇ particles. The XRD pattern of the secondary phases in the ODS steel annealed at 800 °C is depicted in Fig. 9(b). The peaks of the M₂₃C₆, TiC, and Y₂Ti₂O₇ particles can also be seen in the XRD pattern. Similar peaks of secondary phases in the ODS steel annealed at 1000 °C are shown in Fig. 9(c). The relative peak intensity of these secondary phases varies with the annealing temperature. No peaks of M₂₃C₆ particles can be seen in the XRD pattern of the secondary phases in the ODS steel annealed at 1200 °C (Fig. 9(d)). The nanoparticles in the ODS samples are identified as Y₂Ti₂O₇, and the evolution of carbides can be inferred according to the XRD patterns.

In order to quantify the content of the secondary phases in the ODS steels, semi quantitative calculation based on the XRD patterns is performed, and the results are depicted in Fig. 10. The content of M₂₃C₆ decreases as the increase of the annealing temperature, while the content of Y₂Ti₂O₇ shows the contrary regularity. The content of TiC increases as the decline of the M₂₃C₆ when just concerning the transformation of carbides. With respect to the carbides in the ODS steels, the transformation between the M₂₃C₆ and TiC is concerned with the dissolution and precipitation process [35,60,61]. The calculation results shown in Fig. 10 correspond to the analysis of the SEM and XRD results. The M₂₃C₆ particles gradually dissolve in the matrix, and part of the released C can combine with Ti which is the strong carbide forming element. Such phenomenon continually improves the precipitation of TiC at appropriate temperature. When the annealing temperature increases to 1200 °C, the M₂₃C₆ particles totally dissolve in the matrix. The relative content of Y₂Ti₂O₇ nanoparticles is proportional to the increase of annealing temperature. The above analysis means that the annealing treatment is an effective method to regular the secondary phases.

Fig. 11 gives the schematic diagram of the secondary phases in the ODS samples. The relative volume fractions and types of the secondary phases in the as-HIPed and annealed ODS samples show various characterizations. In the schematic diagram, the phases in black, blue and red sections represent M₂₃C₆, TiC and Y₂Ti₂O₇, respectively. The variety regulation displayed in Fig. 11 corresponds with the analysis about Fig. 9, Fig. 10 depicted above. The annealing temperature has an important influence on the volume fraction of the secondary phases in the ODS samples, which will finally affect the mechanical properties of the ODS samples.

3.4. Tensile properties

The tensile strain-stress curves and the corresponding values of the ODS steels are shown in Fig. 12. Fig. 12(a) shows all the measurement results of the ODS samples, and the average values of strength and elongation are calculated and depicted in Fig. 12(b). The as-HIPed ODS steel and the ODS steels annealed at 800 °C exhibit the near yield and tensile strength according to the tensile stain-stress curves. However, the maximum strain of the ODS steels increases with the annealing temperature increasing. The ODS steel annealed at 1000 °C exhibits the highest maximum strain when compared with the as-HIPed ODS steel and the one annealed at 800 °C. The ODS steel annealed at 1200 °C shows the obviously different tensile properties with very high strength and low maximum strain. According to the statistical results shown in Fig. 12(b), the slight downward slope of yield strength (YS) and ultimate tensile strength (UTS) occurs when the annealing temperature increases to 1000 °C. The YS and UTS of the ODS steel annealed at 1200 °C even reach to approximately 1200 MPa and 1600 MPa, respectively. The total elongation (TE) and uniform elongation (UE) of the ODS steels annealed at 800 °C and 1000 °C increase with the increasing temperature. But the TE and UE of the annealed ODS steel show a drastic decline when the annealing temperature reaches to 1200 °C. The ODS steel annealed at 1000 °C possesses the highest TE and UE, implying the higher elastic and plastic deformation level than the other ODS samples. Better comprehensive mechanical properties of the ODS steel are obtained after annealing at 1000 °C.

The tensile results are directly related to the microstructure and secondary phases of the ODS steels. As mentioned above, the annealing treatment can efficiently control the secondary phases which results in the improvement of the final properties. As the increase of the annealing temperature, the harmful M₂₃C₆ particles gradually dissolve in the matrix leading to the formation of TiC particles to same extent, which is beneficial for the whole distribution of secondary phases. Two kinds of carbides (M₂₃C₆ and TiC) exist in the as-HIPed ODS steel and ODS steels annealed at 800 °C and 1000 °C. Among them, the ODS steel annealed at 1000 °C possesses the smallest amount of M₂₃C₆ particles, and the size of M₂₃C₆ are smaller than that in the as-HIPed ODS steel and the ODS steel annealed at 800 °C. In addition, the stable Y₂Ti₂O₇ nanoparticles plays an important role in improving the properties [10]. The lower average grain size and dislocation densities are also helpful to the improvement of mechanical properties. The ODS steel annealed at 1200 °C shows the highest YS, UTS, and the lowest TE and UE. Mcclintock et al. [62] and Klueh et al. [63] have obtained the similar values of YS and UTS in 14YWT and MA957, and the TE and UE are higher than the 1200 °C-annealed ODS steel because of the high densities of nanoclusters and small precipitates in the 1200 °C-annealed ODS steel. The lowest TE and UE of the ODS steel annealed at 1200 °C can be attributed to the low bulk density and high fraction of voids which are detrimental to the ductility. C contents listed in Table 2 shows the drastic increase of C in the ODS matrix after annealing at 1200 °C. Large amount of C released from M₂₃C₆ significantly improves the solid solution strength, which leads to the high UTS of the 1200 °C-annealed ODS steel.

The fracture surfaces of the ODS steels are depicted in Fig. 13. Figs. 13(a)-(d) shows the fractographs of the as-HIPed ODS steel and the ODS steels annealed at 800 °C, 1000 °C, 1200 °C, respectively. As can be seen in Fig. 13(a) and (b), cracks are found in the fracture surfaces of the as-HIPed ODS steel and the ODS steel annealed at 800 °C. The secondary cracks may originate from the large amount of harmful M₂₃C₆ particles. The initiation and propagation of cracks are prone to occur along the interface between the M₂₃C₆ particles and the matrix, which is detrimental to the ductility of the ODS steel [28,29,31]. The as-HIPed ODS steel and the one annealed at 800 °C possess a ductile-brittle fracture mechanism. After annealing at 1000 °C, the size of carbides have been optimized, leading to a better ductility when compared with the as-HIPed ODS steel and the one annealed at 800 °C. The fracture surface shown in Fig. 13(c) consists of homogeneous dimples and few secondary cracks. The fracture characterization indicates a ductile fracture mechanism of the ODS steel annealed at 1000 °C. Fig. 13(d) shows the fractograph of the ODS steel annealed at 1200 °C A large amount of cleavages and secondary cracks exist in the fracture surface, indicating a brittle fracture mode. This phenomenon is attributed to the formation of massive voids when annealed at a high temperature [30]. All the analysis about the tensile results corresponds to the discussions on the microstructure and secondary phases.

4. Conclusions

(1) The M₂₃C₆ particles gradually dissolve into the matrix with the increasing annealing temperature. The relative fraction of TiC increases with the dissolution of M₂₃C₆ particles.

(2) The average size of Y₂Ti₂O₇ nanoparticles shows slow growth character as the annealing temperature increases, and no obvious coarsening of nanoparticles can be found after annealing.

(3) The ODS steel annealed at 1000 °C possesses the best tensile properties among all the ODS samples due to the optimization of the secondary phases and stable microstructure.

(4) Heat treatment is an effective way to improve the mechanical properties of ODS steels by tailoring the secondary phases, as well as the grain size, bulk density, and dislocation densities.

Acknowledgments

The authors are grateful to the National Natural Science Foundation of China (Grant Nos. 51325401, 51474156 and U1660201) and the National Magnetic Confinement Fusion Energy Research Project (granted No. 2014GB125006) for grant and financial support.

The authors have declared that no competing interests exist.

---