In a nuclear collision, if the energy transmitted to the target atom exceeds a certain threshold value, called atomic displacement threshold energy Ed, the atom is ejected from its
site and becomes a projectile; a Frenkel pair is hence created. This atom is then called primary atom and can eject other atoms. All the primary, secondary, etc. displacements form a cascade of displacements, which leads to regions containing defects [55]. Depending on the target material, this defect accumulation may cause the amorphization of the structure. This amorphization occurs beyond a dose, called critical amorphization dose, depending on the material, but also on both the mass and the energy of the incident particle [56-61].
In a nuclear reactor, the nuclear interactions are mainly induced by neutron elastic collisions. As previously explained (Section 2), the simulation of such interactions is usually performed with low energy ions. However, since the electronic interactions are not harmful to Ti3SiC2 (Section 3.1), irradiations with higher energy ions (e.g. with ions of about 1
MeV amu-1) would induce damage only attributable to nuclear shocks; such a damage increases along the ion path (Figure 1b) [24]. Also, in this Section, to compare the nuclear damage induced by irradiations with different energy ions, the results will be as often as possible presented in terms of number of displacements per atom (dpa) induced by these irradiations.
3.2.1. Defects and Disorder
Figure 4 shows cross-sectional transmission electron micrographs of a Ti3SiC2 sample
irradiated at room temperature with 92 MeV Xe [24].
Figure 4. Cross-sectional transmission electron micrographs of Ti3SiC2 grains irradiated at room
temperature with 92 MeV Xe ions to 1015 cm-2, (a) at the beginning and (b) at the end of the ion range.
In a material containing no defects before irradiation, similar to that shown in Figure 2a, the nuclear shocks create many defects that appear as black dots (Figure 4a). These defects are too small to be identified by transmission electron microscopy; they may be clusters of Frenkel pairs or dislocation loops [62]. Along the 92 MeV Xe ion path, an increase of the concentration of irradiation induced defects has been observed, confirming first that they are induced by nuclear shocks, and second that electronic interactions do not damage Ti3SiC2
(Figure 1b). Figure 4b shows the end of the ion course, where the damage is ten times higher than at the beginning of the course. A Ti3SiC2 partly irradiated grain can be observed:
compared to the virgin area, the irradiated area of the grain seems amorphous. To verify this observation, high-resolution transmission electron microscopy has been carried out (Figure 5).
This Figure shows the evolution of the Ti3SiC2 structure as a function of the number of
dpa [24]. These high-resolution micrographs were obtained after orientation of the Ti3SiC2
basal planes parallel to the electron beam, in order to appreciate the nanolamellar structure of this ternary compound (Figure 5a); this structure can also be observed in the diffraction pattern, through three diffraction spots of low intensity between two more intense ones.
Through Figures 5b-d, it may be noted that when the number of dpa increases, the nanolamellar structure disappears, as well in the micrographs as in the diffraction patterns [24, 62]. However, the close-packed hexagonal stack is still observable, indicating that Ti3SiC2 does not become amorphous even for 7 dpa; for comparison, the critical
amorphization dose of silicon carbide at room temperature ranges between 0.1 and 0.5 dpa [56, 63-70]. The cause of this nanolamellar structure disappearance has apparently not been determined yet. However, considering the crystal structure of Ti3SiC2 [71-73], it seems
that the most likely hypothesis to explain this nuclear-shock caused damage would be the substitution of titanium atoms by silicon atoms (and/or vice versa), inducing the creation of defects such as SiTi (and/or TiSi) [24, 62].
Figure 5. High resolution transmission electron micrographs of (a) an unirradiated Ti3SiC2 grain, and of
grains irradiated at room temperature to (b) 1 dpa, (c) 3 dpa and (d) 7 dpa.
When the irradiation temperature increases, it has been observed that both the concentration of defects decreases, and the Ti3SiC2 nanolamellar structure is less damaged
(Figure 6a), to appear as intact for an irradiation performed at 950 °C generating 7 dpa (Figure 6b). The creation of irradiation defects being an athermal phenomenon, this damage decrease is attributed to a defect annealing by temperature [24]. Indeed, the higher the temperature, the higher the diffusion of the species constituting the material, and consequently the higher the probability of recombination of Frenkel pairs.
Figure 6. High resolution transmission electron micrographs of Ti3SiC2 irradiated to 7 dpa at (a) 500 °C
3.2.2. Change in Lattice Parameters
The creation of defects by nuclear shocks usually leads to other structural damage, the most often reported of which is the variation of lattice parameters [17, 74-77]. Figure 7 shows different diffraction patterns obtained by low incidence X-ray diffraction on samples irradiated with both low energy (Figure 7a) and medium energy (Figure 7b) ions; incidence angle utilized are 1° for low energy irradiations and 3° for medium energy ones, corresponding to an estimated analyzed thickness of 230 and 760 nm respectively [24]. Despite the low incidence used, it has been shown that the diffractograms of samples irradiated with 4 MeV Au ions present a contribution of both the irradiated area and the virgin one located at about 1 micron [24].
Figure 7. Low-incidence X-ray diffraction patterns of a commercial Ti3SiC2 irradiated at room
temperature with (a) 4 MeV Au ions and (b) 92 MeV Xe ions; TSC, TC and TS stand for Ti3SiC2,
TiC0.92 and TiSi2 phases respectively.
From these diffractograms, one can see a shift of the Ti3SiC2 (008) line towards lower
diffraction angles as a function of the dose applied in the samples. This reflects a change in lattice parameters of the ternary compound. Figure 8 shows the evolution of both a and c parameters of Ti3SiC2 as a function of the number of dpa, as well as the variation in unit
volume [24].
Thus, it has been shown that the Ti3SiC2 lattice both expands along the c axis (Figure 8b),
and shrinks slightly along the a axis (Figure 8a). These leads to an increase in the unit volume up to a critical dose beyond which the unit volume decreases (Figure 8c); this critical dose has been estimated at 0.1 dpa for irradiations performed at room temperature. It is interesting to note that these results obtained on a sample of Ti3SiC2 differ slightly from that obtained by X-
ray diffraction in Bragg-Brentano configuration on Ti3Si0.90Al0.10C2 samples irradiated under
similar conditions [78]: indeed, if the expansion of the lattice along c axis was confirmed on these latter samples (same order of magnitude), the shrinkage along a has not been observed, resulting in an increase of the unit volume with the dose.
Eventually, an increase in irradiation temperature reducing the damage caused by nuclear shocks (Section 3.2.1), the changes in lattice parameters are not large enough to observe variations as a function of dpa such as those noticeable in Figure 8. Nevertheless, it has been shown that the defect annealing allows both the parameters to be less modified, and the critical dose (beyond which the unit volume decreases) to increase with increasing temperature [24].
Figure 8. Change in the lattice parameters as a function of dpa for irradiations performed at room temperature; the fitting curve of the unit cell volume data was obtained from those of a and c parameters.
3.2.3. Change in Microstrain Yield
The previously observed Ti3SiC2 structural changes may induce microdistorsions in the
crystal network. To verify this, a study of X-ray diffraction line broadening by Williamson and Hall method has been performed [24, 78]. Thus, Figure 9 shows the evolution of the microstrain as a function of both 74 MeV Kr and 92 MeV Xe ion fluence [24], considering that the crystallite size of the ternary compound (a few microns before irradiation [79]) is not affected by irradiations, and therefore does not induce any line broadening. For this Figure,
the microstrain yields were evaluated from the broadening of the Ti3SiC2 (104) line alone, but
the results seems in agreement with others obtained considering many lines [78].
Figure 9. Variation of the microstrain yield as a function of the ion fluence for some medium energy irradiations.
In Figure 9, it can be observed an increase of the microstrain yield as a function of the fluence. This increase is smaller for 74 MeV Kr ion irradiations than for 92 MeV Xe ions, which induce more damage (Figure 1b). Moreover, since no microstrain could be measured for some samples irradiated with high energy ions (930 MeV Xe [80]), it can be conclude that the microstrains are induced by nuclear shocks. Finally, an increase of the irradiation temperature is still beneficial for the Ti3SiC2 damage: the higher the irradiation temperature,
the fewer the microdistorsions in the crystal network (Figure 9).