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Page 2 of 20 Li et al. J Mater Inf 2024;4:27 https://dx.doi.org/10.20517/jmi.2024.44

INTRODUCTION
Reduced activation ferritic-martensitic (RAFM) steels are derived from conventional 8-12 wt.% CrMoVNb
ferritic-martensitic (FM) steels. To obtain low activation capability, the long-lived transmutation elements
such as Mo, Nb and Ni in FM steels are replaced with short-lived transmutation elements such as W and
[1]
Ta . This modification is to simplify the waste management of highly radioactive components after their
service in fusion reactors. The blanket system in fusion reactors requires materials that can withstand
extreme conditions including high heat loads and exposure to high-energy (14 MeV) and high-fluence
[2]
neutron irradiation . With decades of development, RAFM steels have become promising candidate
structural materials for blanket components of fusion reactors, because of their advantages in thermal
[3-5]
mechanical properties, irradiation resistance, etc. . The outstanding properties of RAFM steels are closely
associated with their tempered ferritic/martensitic microstructure, comprising mainly martensitic laths and
small precipitates such as M C (M = Cr, W, Fe) and MX (M = V, Ta; X = C, N) . However, the upper
[6,7]
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application temperature limit of conventional RAFM steels is ~550 °C [8-10] , which restricts the operation of
fusion reactors at higher temperatures and consequently limits the potential enhancement of power
generation.
Conventional RAFM steels such as F82H and China low activation martensitic (CLAM) steel exhibit
comparable tensile properties but poorer high-temperature creep resistance, compared with similar
engineering FM steels such as Grade 91 [11-14] . During the creep process of RAFM steels, the coarsening of
[15]
martensitic laths, packets and blocks occurs, leading to material softening . The microstructural evolution
is mainly influenced by the coarsening of M C and MX precipitates. The M C precipitates at lath, block,
6
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packets and prior austenite grain boundaries (PAGBs) provide limited resistance to grain boundary
migration during creep because of their easy coarsening to > 200 nm from initial ~100 nm [16,17] . In contrast,
the nanoscale MX precipitates randomly distributed in the matrix exhibit good thermal stability, which acts
as obstacles to pin dislocations during creep and decelerates the microstructure degradation [18,19] .
Unfortunately, conventional RAFM steels contain only trace amounts of nanoscale MX precipitates. For
example, the calculated volume fraction of MX precipitates (V ) in F82H is around 0.007% , significantly
[20]
MX
lower than the approximately 0.35% in Grade 91 steel . Consequently, RAFM steels have lower creep
[21]
resistance compared to engineering FM steels, which is mainly attributed to the significantly lower amount
of MX precipitates.
In order to optimize V in RAFM steels, researchers have mainly made great efforts to regulate the
MX
compositions and heat treatment conditions through two methods: trial-and-error experiment and
calculation of phase diagrams (CALPHAD) modeling. For the former method, the focus is primarily on
adjusting the content of elements (such as Ti, Ta, and Zr [22-25] ). For example, the addition of 0.015 wt.% Ti
increased the proportion of nano-sized TiC particles, resulting in a twofold improvement in creep-fatigue
life of base RAFM steel . Mao et al. enhanced 0.1C-RAFM by adding Ta and Zr elements, which promoted
[22]
the precipitation of the MX phase . This modification resulted in a nearly 100 MPa increase in ultimate
[23]
tensile strength (UTS) at 600 °C and an approximately 1,900 h increase in creep rupture time at 600 °C/
180 MPa. The study by Jun et al. indicated that a new RAFM steel with added Ta and Ti had a creep rupture
time of 1,823 h under 550 °C/200 MPa condition, which is approximately three times that of EUROFER97
steel (592 h) . This improvement is primarily attributed to the higher V (0.12% vs. 0.0642%). However,
[24]
MX
the trial-and-error experiment faces challenges in optimizing multiple elements simultaneously, limiting
further improvement of V . For the latter method, it can compensate for the shortcomings of trial-and-
MX
error experiments in simultaneously optimizing multiple variables, thereby improving design efficiency.
[26]
Klueh et al. proposed a CALPHAD-based outline for optimizing MX precipitates . Following this
approach, Tan et al. developed a variety of castable nanostructure alloys (CNAs) with high V (> 0.4%) by
MX