*3.8. Defects in 3C-SiC: SFs and APBs*

In 3C-SiC, the most important defects that hinder its use in the microelectronic industry are related to SFs and dislocations. Stacking faults (SFs) are the most important ones dominating over the entire 3C-SiC layer thickness. In literature, three types of SFs are observed depending on the number of atomic planes with the wrong orientation: SFs can have 1, 2, or 3 errors in the stacking sequence and they are called intrinsic (or SF<1>), extrinsic (or SF<2>), or conservative (or SF<3>) [40,55,56].

In 3C-SiC, mechanisms for SFs' self-annihilation exist but there is also the possibility for SFs to be generated [6,40,57,58]. The concomitant presence of these mechanisms leads to the fact that SFs in 3C-SiC can be hardly reduced. The minimum SFs' densities achieved so far amount to about 10<sup>4</sup> cm−<sup>1</sup> in thin films [58]. In Figure 11a, a TEM image in inplane view shows four stacking faults that are generated from a grain boundary, while in Figure 11b, a TEM image in cross-view shows the annihilation of the SF. In Figure 11a, a vertical grain boundary generates three clearly visible SFs lying in (111), (111), and (111) planes. Interestingly, two SFs in the (111) planes limit the SF in the (111) plane; they are limited on the other side by the grain boundary. The intersection between the SFs (111) and (111) is the Lomer−Cotrell partial dislocation and has a Burgers vector of a/6[011 The place in which the grain boundary intersects the SF (111) is the place in which the SF (111) is generated. In Figure 11b, the crossing of several SFs is shown. This image proves that there are two possible intersections of the SF lying in the (111) and (111) planes. The intersection indicated as "1" has an inverted V-shape typical for the formation of a Lomer–Cottrel dislocation. This dislocation, as already reported, has a Burgers vector (a/6[110]) lower than the usual partial dislocation Burger vector (a/6[112]) that borders the stacking fault. The intersection called "2" has a lambda shape; it forms for kinetical reasons. These two configurations can decrease the amount of SF approaching the surface and improve the quality of the epitaxial film.

The annihilation mechanism is considered in more detail in Reference [59]. It is found that the key parameter for the formation of a lambda-shape or an inverted V-shape is the distance between the PDs and the mutual orientation of their Burgers vectors. In the case in which the PDs have Burgers vectors that sum in such a way that the resulting Burger vector is shorter than the initial ones, the partial dislocation attracts to each other and if they are closer by less than about 15 nm, the propagation of both SFs is suppressed with the formation of a Lomer–Cottrell lock. In the case in which the two PDs are far more than 15 nm, they do not interact with each other and can form a Lambda-shape structure. The Lambda-shape can form even if the partial dislocations are close enough, but they repulse. In Figure 12 on the left, a sequence of MD simulation snapshots of the formation of

"inverted V"-shaped intersection of stacking faults have been shown. In Figure 12 on the right, MD simulation snapshots of the formation of "λ"-shaped intersections of stacking faults in the case of a large distance between partial dislocations (a–c) and repulsing dislocations (d–f) have been reported. Blue atoms correspond to the Si and C atoms in the cubic diamond lattice, while orange atoms belong to the stacking faults.

**Figure 11.** (**a**) TEM image in in-plane view shows four stacking faults that are generated from a grain boundary. (**b**) TEM image in cross-view showing the annihilation of the SF with two different structures (adapted from References [40,60]).

**Figure 12.** (**A**) Molecular dynamics simulation snapshots of the inverted V-shape configuration. (a–c) The simulation time: (a)—0, (b)—120 ps, and (c)—180 ps. Blue atoms correspond to the Si and C atoms in the cubic diamond lattice, orange atoms belong to the stacking faults. Inset in panel (c) shows the atomic configuration of the formed Lomer–Cottrell lock dislocation. (**B**) Molecular dynamics simulation snapshots of the lambda-shape configuration. in the case of the large distance between the 30◦ leading dislocations (a–c) and as a result of the interaction of closely spaced 30◦ dislocations with equal screw components of Burgers vectors (d– f). Simulation time: (a)—0, (b)—360 ps, (c)—540 ps, (d)—0, (e)—60 ps, (f)—200 ps. Inset in panel (c) shows the atomic configuration of the intersection of 30◦ partial dislocation with crossing stacking fault, also corresponding to the intersection in panel (f). (adapted from Reference [59]).

The SFs can interact also with other extended defects, such as the inverted domain boundary (IDB) (sometimes called the anti-phase boundary, APB). In 3C-SiC grown on (100) "on-axis" silicon, due to the symmetries of the Si lattice, two equivalent dispositions of the SiC crystal are possible. The two possible orientations are rotated 90◦ around [001] and, due to the SiC symmetries, a rotation of 90◦ is equivalent to flip the crystal upside-down. The boundary between two such domains is called IDB (or APB). The SFs can interact strongly with this kind of extended defect of SiC. In Figure 13, a sequence of STEM images showing an IDB interacting with SFs is shown. The image is the projection of the TEM lamellae in the (110) plane. The SFs are observed as straight lines, while the IDB has different lying planes and appears as a ribbon. A close inspection of these images shows that SFs can be generated and annihilated by the IDB: several SFs can be recognized in the figure and some of these are apparent only in the crystal below the boundary, while some others are apparent only in the crystal above the boundary. SFs that are in the lower crystal are not allowed to propagate in the top crystal and in this case, we observe an annihilation of the SF due to the presence of IDB. On the contrary, SFs that belong to the top crystal and are not present in the lower crystal are generated in the IDB. The SFs can be generated during the growth due to interface instability that creates seeds for nucleation; after the nucleation, it expands following the growth of the surface. Eventually, it can collide on an IDB and be annihilated [40]. SFs can be also generated during the cooling down of the temperature after the growth; indeed, temperature gradients can induce stress in the layer. Above critical shear stress, it is known that the formation of dislocations and, in 3C-SiC, the formation of partial dislocation is a thermodynamically favored process.

**Figure 13.** Sequence of STEM (110) cross-view images showing an IDB and its interaction with SFs. The lying planes of IDB and SFs are indicated (adapted from Reference [40]).

As previously discussed at the beginning of Section 3.7, different kinds of SFs are present in 3C-SiC. These different types of SFs can be seen as inclusions of different hexagonal polytypes in the cubic structure. In particular, the intrinsic SF can be called a 2H-like SF, the extrinsic one can be seen as a 4H-like SF, and the conservative one can be seen as a 6H-like SF. These different SFs have different energies [60] and different behaviors of these defects should be expected. The room temperature µ-PL map at 540 nm, taken on a 3C-SiC sample in cross-section, is shown in Figure 14a [61]. Moving from the Si-SiC interface towards the top (from 0 µm to 35 µm), the band-edge peak intensity rises, showing a considerable improvement of the crystalline quality, increasing the growth thickness. Figure 14b,c exhibits µ-Raman maps obtained in the same location and indicate certain areas as well as the 3C-SiC/Si interface with greater signal magnitude at 778 cm−<sup>1</sup> and 784 cm−<sup>1</sup> , respectively. Figure 14d–f displays the mean Raman spectra obtained in areas (1), (2), and (3), which reveal the 3C-SiC TO mode centered at about 796.5 <sup>±</sup> 0.2 cm−<sup>1</sup> . Conversely, the mean Raman spectrum obtained in points (2) and (3) reveal an extra peak at 778.3 cm−<sup>1</sup> , as well as two more peaks correspondingly at 778.0 cm−<sup>1</sup> and 784.0 cm−<sup>1</sup> . These additional peaks can be related to the presence of extrinsic (4H-like and 6H-like)

stacking faults. From these data, we can observe that, despite the low energy of the 6H-like SF, it appears that this kind of stacking fault can be observed in larger regions and closer to the surface with respect to the 4H-like SF. More investigations should be done concerning this aspect but we suspect that this large presence of 6H-like SFs could be due to kinetic reasons more than energetic ones. In fact, from the energetic point of view, this SF has the lowest formation energy.

**Figure 14.** Micro-PL mapping (**a**) at 540 nm and micro-Raman mapping of a 3C-SiC cross-section located at (**b**) 778 cm−<sup>1</sup> and (**c**) 784 cm−<sup>1</sup> .The interface with the removed Si substrate is shown by point 0 on the Y-axis. Average Raman spectra achieved in the (**d**) area (1) of the map (**b**), (**e**) zone (2) of the map (**b**), and (**f**) zone (3) of the map (**b**,**c**). The laser probe created the peak located at 828.37 cm−<sup>1</sup> (\*) (see Reference [61]).
