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Figure 4 shows the surface relief within one grain of fatigued ferritic steel.
The number of parallel PSMs denoted A up to H is present.
In Fig. 5 the height h in austenitic steel and the mean height h in ferritic steel are plotted vs. the number of loading cycles.
Plastic replicas of the austenitic grain after different number of loading cycles with εap = 1x10 -3.
Figure 6. shows plastic replica corresponding to one grain at different loading cycles.

Research [9,10] show that the appearance of third brittle zone is related not only to the film-like ferrite formed along the original austenite grain boundaries, but also to the embrittlement of austenite grain boundary due to the precipitation of Nb, Al, V carbonitride compound.
In addition, a large number of dislocations are found, as recovery or recrystallization which can eliminate the dislocations is difficult to occur at 750 oC (Fig. 6d).
Generally, the embrittlement mechanism of second brittle zone is that S, O soluted in the high temperature austenite precipitate at the austenite grain boundaries in the form of (Fe, Mn)O or (Fe, Mn)S, and then these precipitates grow up along the grain boundaries to reduce the grain boundary strength, causing the decrease of hot ductility.
The test steel in this study has good hot ductility at 950oC and above, and the RA at 900oC is slightly lower than 60%, which is owing to the embrittlement of grain boundary contributed by precipitates of (Fe, Mn)O, (Fe, Mn) S distributing along austenite grain boundary.
The decrease of plasticity is caused by the precipitation of (Fe, Mn)S weakening the grain boundary.

Therefore, the number of heat-conducting particles in the bond can be increased without the risk of loss of strength.
The combination of components so completely different in hardness as corundum and graphite (corundum hardness is 9 on the Mohs scale, graphite hardness - 1) requires an additional operation of welding the samples in the so-called "canvas". [20] To produce polished sections of corundum, a method is used for the stepwise polishing of samples on boron carbide with grain numbers: M-28, M-14, M-10, M-7, M-5, M-3, M-1.
The sample is polished for at least 15 minutes using each of the above grain numbers.
Therefore, the number of heat-conducting particles in the bond can be increased without the risk of loss of strength.
For example, in the absence of a grinding machine, the plane can be manually ground on ceramic circles with the grain size of 16N, 25N.

The technique was developed initially for aluminum alloys, however, since then the FSW was found suitable for joining a large number of materials (Sato et al. 1999).
The SEM micrographs of Nugget (b), TMAZ (c) regions show finer grain size diameters than those of HAZ (d), and BM (e) regions.
The grain sizes in the stir welded zone present diameters between 25-30 μm, while those in unaffected base metal are in the range of 60-70 μm.
The microstructure of the stir welded zone presents fine stirred grains.
On the other hand, the unaffected base material and the heat affected zone present coarse grains.

Meanwhile, Mo layer can not only induce the c-axis orientation of NdFeB, but also improve the crystallization of NdFeB grains.
A number of preparationparameters, including buffer layer material and its thickness, deposition temperature and/or annealing temperature, target composition, and even sputter media[8], can affect the structure of the films and their final magnetic properties.
The above results indicate that introducing Mo can improve the c-axis orientation and the crystallization of NdFeB grains.
It indicates that introducing Mo can induce the crystallization of NdFeB grains.
Good c-axis orientation and crystallization of NdFeB grains was obtained at 40W NdFeB sputter power.

The starting materials were α-Si3N4 powders with average grain size of approximately 0.5 µm, purity 99.5%, nano α-Si3N4 particles with average grain size of approximately 140 µm, purity 99.5% and nano TiC particles with average grain size of 130 nm, purity 99.8%.
For the specimen quenched at ∆T=800�, micro cracks initiated on the surface and the number of micro voids increased as shown in Fig.2(b), leading to a drastic drop in strength.
The thermal shock at ∆T=900� increased the number of micro voids accompanied with oxidation, crack interlinking which resulted in surface chapping as shown in Fig.2(c), with very low strength retained.
The matrix grains of the material were refined by adding nano Si3N4 and nano TiC particles as shown in Fig.3(a).
The strengthening mechanisms by grain boundary strengthening and grain refinement, and the toughening mechanisms by rod-like grain bridging, crack deflection and transgranular fracture are the main cause for the improved thermal shock resistance.

The maximum is decreased only at the grain boundaries. d) PL integrated intensity mapping of the same area, step size 30μm.
Although the first assumption breaks down in long lifetime layers produced, for example, by oxidations, it holds in a large number of as-grown layers.
Assumption number two holds reasonably well except in regions of structural defects, where the decreased initial signal would lead to on overestimation of carrier lifetime in this region.
Fig. 3: PL intensity mappings of grain boundaries in different samples. a) Mapping of layer C with step size 30μm.
The etched grain boundaries (marked G), a carrot (marked C) and downfalls (marked D) are visible.

After annealing, however, recrystallized grains along shear bands were mainly Goss grains regardless of reduction.
Table 1 The cold rolling condition Specimens Reduction (%) Initial thickness (mm) Final thickness (mm) Number of passes a 28 2.85 2.05 12 b 40 2.85 1.71 16 c 50 2.85 1.43 24 d 68 2.85 0.92 29 Results Formation Angles of Shear Bands.
In 28% of reduction (Fig.3(a)), elongated recrystallized grains along SBs were observed and almost half of them were Goss grains.
In 40% (Fig.3(b)), almost all SBs were occupied by recrystallized grains, and most of which were Goss grains.
After annealing, however, recrystallized grains along shear bands were mainly Goss grains regardless of reduction.

In the 4XH sample a very fine and equiaxed grain structure with d ≈ 0.3 µm is observed.
On the other hand, the furnace cooled samples are characterized by coarse precipitates, mostly located along grain or sub-grain boundaries rather than within the grain.
The other parameters are the level of ECAP deformation and the concurrent grain size reduction, both opposing dislocation accumulation.
As for grain size, a definitive comment regarding its effect is not possible at this stage due to the poor visibility of the boundaries (or sub-boundaries).
A smaller grain size means a shorter distance for intragranular dislocation annihilation at grain boundaries [ ], and again, comparison of 1XS – 4XS confirms this assertion.

Grain boundary spacing was determined using the linear intercept method with at least 65 intercepts recorded.
Overall, the level of the grain refinement is very low.
The optical microscopy revealed the formation of shear / microbands within the number of ferrite grains for all deformation routes (these grains have darker contrast in the optical images), Fig. 1a-c.
The distribution of the darker grains is quite heterogeneous in the ND direction for all samples.
The number of the darker grains in the symmetrically rolled specimens is higher close to the sheet surface and is lower in the middle of the sheet thickness (Fig. 1a).