The main design idea of ODS steels is to greatly reduce the C content for inhibiting the precipitation of M23C6 carbides and concurrently introduce numerous nanoscale oxide precipitates for offsetting the strength reduction caused by the reduction of M23C6 carbides. Oxides have better thermal stability because of their higher dissolution temperature than M23C6 and MX carbides . Thus, the induced nanoscale oxide particles improve the precipitation strengthening and thermal stability of steels simultaneously. In the 1960s, SCK CEN in Belgium developed the first ODS steel . At first, Y2O3 is the main oxide in ODS steel because Y2O3 particles have strong anti-radiation decomposition and anti-oxidation peeling abilities . However, the size of Y2O3 in ODS steels is usually large (~30–40 nm), leading to the limited strengthening effect. Klueh et al.  reported that adding Ti to the ODS steel results in the formation of fine Y–Ti–O particles with a size of 3–5 nm. By contrast, the size of MX carbides in the EUROFER97 steel is 20–30 nm . Subsequently, Al was introduced to improve the corrosion resistance of ODS steels. Given the strong oxygen affinity of Al, Y–Al–O particles are preferentially precipitated in Ti-containing ODS steels. However, the large Y–Al–O particles (~10–20 nm) cause the significant reduction of mechanical properties of the steel compared with the fine Y–Ti–O precipitates . To overcome this problem, Dou et al.  refined the oxide particles to 7 nm by adding Zr in an Al-containing ODS steel. Yan further refined the oxide particles to 4 nm by adding Hf in an Al-containing ODS steel . As a result, the high-temperature strength and radiation resistance of the ODS steel are much higher than those of the RAFM steel in the same period. Fig. 7 shows a comparison of the high-temperature creep performance and radiation resistance between ODS and RAFM steels in recent years [100–103]. The creep life of the 9Cr ODS steel is ~1000 times longer than that of the F82H and Eurofer 97 steels under 650°C/130 MPa. Thus, the ODS steel is a possible candidate to meet the service requirements of next-generation nuclear fusion reactors.
Figure 7. Comparison of high-temperature creep strength at 650°C (a) and irradiation resistance at 450–480°C (b) between ODS and RAFM steels in recent years.
Nevertheless, the preparation process of ODS steels is complex. Thus, the fabrication is difficult to be industrialized at present. ODS steels are mainly prepared via powder metallurgy , and a schematic of the preparation process is shown in Fig. 8. In brief, the metal and oxide particles are separately milled to nanometer powders, and then the metallic and oxide powders are mixed by mechanical alloying to promote the homogeneity of the material. The mechanical alloying powders are sintered via hot isostatic pressing to enhance the compactness of the material and promote the nucleation of nano-sized oxide particles in steels. Finally, the microstructure of the steel is further controlled by deformation processing and heat treatment. Ball milling and mechanical alloying are difficult to control. Thus, they may lead to the discrepancy of chemical compositions of ODS steels. Moreover, the preparation process is complex, time consuming, and inefficient, leading to a great increase in the production cost of ODS steels . At present, the development of ODS steels is mainly focused on how to improve production capacity and efficiency to achieve mass production of bulk ODS steels.
Figure 8. Schematic of the preparation processes of ODS steels . Reprinted from J. Nucl. Mater., 518, T. Jaumier, S. Vincent, L. Vincent, and R. Desmorat, Creep and damage anisotropies of 9%Cr and 14% Cr ODS steel cladding, 274, Copyright 2019, with permission from Elsevier.
Machining strengthening can be divided into cold deformation and thermal deformation. The main idea of the cold deformation strengthening of RAFM steels is to greatly refine subgrains and carbides by introducing severe plastic deformation (SPD) [106–107] at low temperatures in order to increase subgrain boundary strengthening and precipitation strengthening. In 2011, Wang et al.  first prepared ultrafine subgrained P91 steel through surface mechanical attrition treatment (SMAT). They successfully reduced the average diameter of subgrains and M23C6 carbides to 35 and 20 nm, respectively, from 320 nm and 125 nm in the traditional RAFM steel, respectively. However, SMAT is a SPD method to synthesize the nanostructured surface layer on metallic materials; thus, it cannot fabricate samples with large dimensions. Later, Chen et al.  prepared an ultrafine subgrained RAFM steel by cold forging and post-annealing, which reduced the subgrain and M23C6 carbide size to 240 and 60 nm, respectively. Compared with the traditional quenching plus tempering process, cold forging and post-annealing successfully increase the yield strength of the RAFM steel from 550 to 650 MPa and elongation from 18% to 25%. Recently, Jin et al.  have prepared an ultrafine subgrained 9Cr2WVTa steel by multi-pass rotary-swaging and post-annealing, and reduced the subgrain and M23C6 carbide sizes to 330 and 50 nm, respectively. The creep life of ultrafine grained 9Cr2WVTa steel at 550°C has been greatly improved, i.e., 550 h at an applied stress of 270 MPa and exceeding 1300 h at an applied stress of 250 MPa. By contrast, the creep lives of the Eurofer 97 and F82H steels are merely 17, 50, and 150 h under identical creep conditions. The creep life of the cold deformed 9Cr2WVTa steel is 30 times higher than that of the F82H and Eurofer 97 steels under 550°C/270 MPa. The comparison of creep lives among ultrafine subgrained 9Cr2WVTa, typical Eurofer 97, and F82H steels is shown in Fig. 9.
Figure 9. Creep properties of cold deformation and thermal deformation-strengthened steels compared with typical Eurofer 97and F82H steels.
The thermal deformation of RAFM steels is mainly associated with thermo-mechanical treatment (TMT) . The main idea of TMT is to reduce the M23C6 carbide precipitation strengthening and replace it with MX carbide precipitation strengthening, grain boundary strengthening, and subgrain boundary strengthening. The goal is to increase thermal stability and improve the high-temperature creep properties without losing the strength of steels. Fig. 10 shows a schematic of the structural change in a steel under TMT processing. First, the steel is heated above the austenitizing temperature (1050–1300°C) to dissolve the precipitate and then cooled to 700–1000°C for hot rolling. These procedures introduce high-density dislocations and reduce subgrains. Meanwhile, the precipitation of MX carbides instead of M23C6 carbides is promoted. Finally, the steel is annealed at 650–800°C to obtain completely recrystallized grains and completely precipitate M23C6 carbides . In 2016, Hoffmann et al.  treated the Eurofer 97 steel by using TMT at a thermal deformation temperature of 700°C and annealing temperature of 750°C. TMT significantly improves the high-temperature creep properties of the Eurofer 97 steel to >2300 h at 650°C/100 MPa and >4200 h at 600°C/140 MPa. By contrast, the creep lives of the typical Eurofer 97 and F82H steels at identical creep conditions are about 1000 and 200 h, respectively. So far, no systematic analysis is conducted on the microstructural change before and after TMT. Prakash et al.  prepared a 9Cr–1W–0.06Ta RAFM steel via TMT at a thermal deformation temperature of 700°C and annealing temperature of 760°C. Compared with traditional quenching plus tempering, prior austenitic grain size significantly increases from (10 ± 2) μm to (175 ± 12) μm, the M23C6 carbide size decreases from 50–150 nm to 30–100 nm, the MX carbide size decreases from 20–50 nm to 5–30 nm, and the subgrain size decreases from 272 nm to 179 nm . As expected, the high-temperature creep life of the steel significantly increases to 1200 h at 550°C/260 MPa  compared with ~50 h under the same creep conditions for typical Eurofer 97 and F82H steels . The creep life of the TMT-processed 9Cr–1W–0.06Ta is about 20 times higher than that of the F82H and Eurofer 97 steels under 550°C/260 MPa. Comparison of the creep lives between two types of TMT steels and the typical Eurofer 97 and F82H steels is summarized in Fig. 9.
Microalloying elements, such as Ti, Mo, V, and Ta, are strong carbide-forming elements. In steels, these elements easily form stable MX carbides . Theoretically, the addition of microalloying elements can significantly promote the precipitation of MX carbides and inhibit the precipitation of M23C6 carbides because the precipitation temperature of the former is considerably higher than that of the latter [77–78]. For example, Xiao et al. [115–116] suggested that adding V and Ta in RAFM steels not only leads to the precipitation of MX carbides but also refines M23C6 carbides. Basing from this idea, we prepared a CLAM steel containing 0.10wt% Ti by controlling the tempering and other preparation processes. Strength reduction caused by the decreasing M23C6 is compensated by increasing the subgrain strengthening . Table 1 shows the microstructure information of 0.10wt% Ti CLAM steels treated with different tempering processes , and the conventional CLAM steel without Ti is included for comparison. Due to the same rolling and quenching process, the grain size and MX carbide size of the steels show insignificant variations, i.e., 7 and 17 nm, respectively. With the decrease in tempering temperature and time, the sizes of the subgrains and M23C6 carbides decrease significantly. To be explicit, they are 650 and 90 nm, 450 and 60 nm, and 370 and 40 nm, after tempering processing at 760°C/90 min, 760°C/30 min, and 730°C/30 min, respectively. The corresponding densities of the M23C6 carbides are 2.85 × 1019, 4.30 × 1019, and 1.10 × 1020 m−3, respectively. Compared with those in the conventional CLAM steels, the sizes of grains and densities of M23C6 and MX carbides decrease, whereas the sizes of subgrains and densities of M23C6 and MX increase in 0.10wt% Ti CLAM steels after tempering at 760°C/90 min. With decreasing tempering temperature and time, the sizes of subgrains and M23C6 carbides continuously decrease and the densities of M23C6 carbides increases. The above-mentioned strengthening mechanism (Section 2) indicates that Ti addition promotes MX precipitate strengthening but restrains the strengthening from subgrain boundaries, grain boundaries, and M23C6 precipitates.
Steel Grain size / nm Subgrain / nm M23C6 MX Size / nm Density / m−3 Size / nm Density / m−3 0.1 Ti CLAM (760°C/90 min) 7 650 90 2.85 × 1019 17 8.20 × 1024 0.1 Ti CLAM (760°C/30 min) 7 450 60 4.30 × 1019 17 8.20 × 1024 0.1 Ti CLAM (730°C/30 min) 7 370 40 1.10 × 1020 17 8.20 × 1024 CLAM  20 320 125 6.05 × 1019 30 3.20 × 1024
Table 1. Microstructure information of 0.10wt% Ti CLAM steel fabricated via different tempering processes compared with their counterparts without Ti addition
The microstructural change related to Ti addition and tempering is due to the fact that Ti effectively promotes the nucleation of MX carbides, whereas the dispersed fine MX particles refine grains by hindering grain growth at the solid solution stage in martensitic steels. In consideration that the melting point of M23C6 is remarkably lower than that of MX carbides, the formation of M23C6 carbides is inhibited. The small M23C6 carbides induce the increment of martensite lath size by hindering the growth of martensite lath during tempering. The tempering processes are 760°C/90 min, 760°C/30 min, and 730°C/30 min, respectively, and those temperatures are close to the precipitation temperature of M23C6 carbides. Therefore, the size of M23C6 carbides decreases significantly with decreasing tempering temperature and time. The different tempering processes show insignificant effect on MX carbides because the precipitation temperature of the carbides is high. In addition, the austenitizing temperature of 805°C  is above the tempering temperature. Hence, the size of the primary austenitic grains remains constant despite the slightly decreasing subgrain size with reduced temperature and time for tempering.
Fig. 11 shows the engineering stress–strain plots of 0.10wt% Ti CLAM steels with different tempering processes at 600°C compared with the standard CLAM steel . The tensile strength (σTS) and yield strength (σy) of the 0.10wt% Ti CLAM steel treated by 760°C/90 min tempering are 285 and 260 MPa, respectively, and the maximum elongation is 47%. With the decrease in tempering temperature and time, σTS and σy increase while elongation decreases, which are respectively 340 MPa, 325 MPa, and 33% at 760°C/30 min and 370 MPa, 345 MPa, and 29% at 730°C/30 min. For the conventional CLAM steel, the σTS, σy, and elongation are 305 MPa, 325 MPa, and 27%, respectively. Except for the 0.10wt% Ti CLAM steel tempered at 760°C/90 min, the σTS, σy, and elongation of the 0.10wt% Ti CLAM steel are higher than those of the conventional CLAM steel. Fig. 12 shows the creep curves of the 0.10wt% Ti CLAM steels fabricated by different tempering processes. Creep experiments were conducted at 600°C/170 MPa. At this condition, the creep life and elongation of the 0.10wt% Ti CLAM steels tempered at 760°C for 90 min are 20 h and 50%, respectively. The creep life of the 0.10wt% Ti CLAM steel increases and its creep elongation decreases with decreasing tempering temperature and time, reaching 145 h and 30% at 760°C/30 min and 190 h and 25% at 730°C/30 min. The creep life and ductility are significantly higher than 120 h and 20% for the conventional CLAM steel, respectively . Obviously, adding Ti and adjusting the tempering process increases the high-temperature creep life of the CLAM steel successfully without sacrificing the ductility. The processing of the Ti CLAM steel is simple compared with that of the traditional CLAM steel, which is beneficial to commercial production.
Figure 11. Engineering stress–strain plots for the 0.10wt% Ti CLAM steels via different tempering processing at 600°C. The stress–strain curve for the steel without Ti addition is also shown for comparison.
Strengthening mechanisms of reduced activation ferritic/martensitic steels: a review
20 April 2020
Revised: 11 June 2020
Accepted: 17 June 2020
Available online: 21 June 2020
Abstract: This review summarizes the strengthening mechanisms of reduced activation ferritic/martensitic (RAFM) steels. High-angle grain boundaries, subgrain boundaries, nano-sized M23C6, and MX carbide precipitates effectively hinder dislocation motion and increase high-temperature strength. M23C6 carbides are easily coarsened under high temperatures, thereby weakening their ability to block dislocations. Creep properties are improved through the reduction of M23C6 carbides. Thus, the loss of strength must be compensated by other strengthening mechanisms. This review also outlines the recent progress in the development of RAFM steels. Oxide dispersion-strengthened steels prevent M23C6 precipitation by reducing C content to increase creep life and introduce a high density of nano-sized oxide precipitates to offset the reduced strength. Severe plastic deformation methods can substantially refine subgrains and MX carbides in the steel. The thermal deformation strengthening of RAFM steels mainly relies on thermo-mechanical treatment to increase the MX carbide and subgrain boundaries. This procedure increases the creep life of TMT(thermo-mechanical treatment) 9Cr–1W–0.06Ta steel by ~20 times compared with those of F82H and Eurofer 97 steels under 550°C/260 MPa.