**4. Discussion**

Previous work on the thermal creep deformation of V–Ti alloy indicated that 3% Ti addition to V resulted in the lowest activation energy for creep strain behavior in the V–xTi alloys (x = 0% to 20%) [17–19]. The low activation energy of creep deformation reduces the creep strain rate at a given temperature in the material and leads to a high creep strength. The creep strength is strengthened by Ti addition to V in the V–Ti system alloy from the increase in yield and tensile strengths because of a solid–solution hardening mechanism and/or a dispersed particle-strengthening mechanism. The main reason why the V–3Ti alloy exhibits the strongest creep strength in the V–Ti system alloy (0% to 20% Ti addition) cannot be explained by the solid–solution hardening mechanism because solid–solution hardening due to Cr and Ti addition to V–Cr–Ti alloys was not apparent in this work as determined from the nano-indentation test of unirradiated V–Cr–Ti alloys as shown in Figure 2. In the thermal creep deformation in V–Ti alloy at an elevated temperature, the formation of a dislocation-network as well as formation of a titanium-oxycarbonitride precipitate, Ti(CON) have been observed as thermal vacancy migration and dislocation slip motion. Figure 5 provides an example of precipitate formation in V–4Cr–4Ti alloys in creep deformation at 600 ◦C with a stress of 200 MPa for 2800 h in a liquid Na environment [20]. Small precipitates formed near the grain boundary or large bulk precipitates and the precipitate was Ti(CON) titanium-oxycarbonitride precipitate on a habit pane of {100}. The nature of the Ti(CON) precipitate has been reported by Impagnatiello [21–23]. In the thermal creep test, it is likely that Ti(CON) precipitates are formed along with defect sinks such as grain boundaries and bulk precipitates where the thermal vacancy is absorbed preferentially. It is expected that the Ti(CON) precipitate is more likely to form in an environment where that migration of Ti is enhanced due to vacancy flux to sinks such as grain boundaries through the Kirkendall effect for a long period and the nucleation of Ti(CON) precipitate occurs around sinks such as grain boundaries and large bulk precipitates due to the thermal heat treatment. Since He irradiation at high temperature also provides a large defect flux of not only excess vacancies but also displaced Ti atoms produced by He ion bombardment in the matrix, this kind of irradiation may lead to the formation of Ti(CON) precipitates in the matrix. Previous transmission electron microscopy (TEM) work for the microstructural observation of neutron- and ion-irradiated V–Cr–Ti and V–Ti alloys reported that Ti(CON) precipitate formation occurred above 350 ◦C with more

than 0.1 dpa of damage level [24–28]. It is deduced that the significant irradiation hardening in V–Cr–Ti alloys in this work was caused by irradiation-induced Ti(CON) precipitation during 500 ◦C and 700◦C irradiation, even though microstructural observation by TEM work remained unexamined in this work. The high density of small Ti(CON) precipitate formed in V–4Cr–1Ti alloys and V–4Cr–1Ti shows the highest obstacle resistance against dislocation slip based on the Orowan equation termed as dispersed barrier hardening in [29];

$$
\Delta \sigma\_{-} \mathbf{y} = \mathbf{M} \mathbf{a} \mu \mathbf{b} \sqrt{\mathbf{N} \mathbf{D}} \tag{1}
$$

where M is the Taylor factor (3.06 for fcc polycrystals), µ the shear modulus of the matrix, b the magnitude of the Burgers vector of the moving dislocation, and ∆σ<sup>y</sup> represents the increment in yield strength because of the obstacles of size D, number density, N and barrier strength α. It is assumed that the density of small precipitate in V–4Cr–1Ti will be highest among all V-4Cr–xTi and shows a significant irradiation hardening based on the dispersed precipitate hardening produced by He-ion irradiation.

**Figure 5.** TEM micrographs of thermal-creep-deformed V–4Cr–4Ti alloy at 600 ◦C and 150 MPa for 2800 h. The left side of the low-magnification image shows that precipitates gathered along a grain boundary and the right side of the high magnification image shows that precipitates on {100} habit plane were formed. [L] and [H] indicate the direction of the longitudinal direction and horizontal direction of the creep tube, respectively. (after Fukumoto et al. [20]).

The nano-indentation test results from the V–Cr–Ti alloys show that the unirradiated V–6Cr–1Ti and V–8Cr–1Ti alloy hardness did not change much compared with that of the V–4Cr–1Ti alloys. Therefore, Cr addition may be ineffective for solid–solution hardening in V–Cr–Ti alloys. The He-irradiated V–yCr–1Ti alloy results show that the effect of Cr addition for irradiation hardening is independent of Cr addition in the V–(4–8)Cr–Ti alloys.

The effect of gas impurity level for irradiation hardening appears in irradiation hardening of V–4Cr–1Ti alloys irradiated at 500 ◦C and 700 ◦C, and the highly purified V–Cr–Ti alloys reduce the irradiation hardening as shown in Figures 2 and 4. This reduction of irradiation hardening in highly purified V–Cr–Ti alloys may be caused by the formation of Ti(CON) precipitates during He-ion irradiation. The nucleation and growth of Ti(CON) precipitate should be rate-limited to the concentration of gas impurities of C, N and O, and Ti atoms and the nucleation rate is proportional to the product Cimp·CTi of the gas impurity concentration Cimp and Ti concentration CTi from the kinetics of the point defect reaction [30], when the nuclei of Ti(CON) are assumed to be TiO- or TiC-type [31]. The reduction in gas impurities in the V–Cr–Ti alloy matrix is connected with the nucleation rate of Ti(CON) precipitate and a decrease of irradiation hardening in the V–Cr–Ti alloys. Microstructural observation by TEM is required to clarify the correlation between irradiation hardening and microstructural evolution, and especially Ti(CON) formation in the future.

Because V–4Cr–1Ti alloys show significant irradiation hardening and more than 2% Ti addition to V–4Cr–xTi alloys results in a lower irradiation hardening than the V–4Cr– 1Ti alloy, the optimum amount of Ti addition to candidate alloys for structural materials of fusion reactor application should be 2% to avoid surplus irradiation hardening at low-temperature irradiation. In terms of the swelling behavior of V–Cr–Ti alloys, 1% Ti addition to V–Cr–Ti and V–Fe–Ti alloys is enough to suppress void swelling from 400 to 600 ◦C with heavy damage levels to 30 dpa [32]. Hence, 2% Ti addition to V– Cr–Ti alloys helps to suppress void swelling at a high temperature with heavy damage. Consequently, from the viewpoints of the suppression of surplus irradiation hardening and void swelling, an optimum composition of V–Cr–Ti alloys for structural materials of fusion reactor engineering is proposed to be a highly purified V–(6–8)Cr–2Ti.
