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BLSFs have the general formula (Bi2O2) 2+- (Am-1BmO3m+1) 2, in which pseudoperovskite (Am-1BmO3m+1) 2 layers are interleaved with (Bi2O2) 2+ layers, where m is the number of BO6 octahedra in the pseudo-perovskite layers (m = 1 to 5).
There are several possible ways, such as Hot-Forging (HF) method and Templated Grain Growth (TGG) method, to obtain grain oriented samples.
Grain-oriented samples were prepared by the hot-forging (HF) method � [5].
The grain-oriented factor, F, was calculated using the Lotgering method [7].
This is due to the grain orientation.

Refining grains are important to obtain sound cast billets suitable for further processing.
With increasing of applying LFEF treated time, the fine grain occupied the whole section of ingot, and the ratio of fine grain zone to the whole section was proportional to the treated time.
The maximum change rate (the largest nucleation rate) was observed from 30s to 70s, which means that solidification front is fragile, and a large number of nuclei can be torn off from the solidification front by the effect of LFEF treatment in this period.
With increasing of LFEF treated time, the number of new nuclei has increased, and the ratio of fine grain zone to the whole section has increased gradually with the increase of treated time.
Qian, A Brief History of the Grain Refinement of Cast Light Alloys, Mater.

To date, the number of research work and publications concerning the superplasticity of CP alloys are very limited [11], since the majority of research are concentrated on Ti-6Al-4V alloys.
In the present work, fine grains are defined as grains having a diameter of ≤10 µm.
The volume fraction of fine grain ratio was calculated by dividing the fine grain area by the total image area.
The large grain size is not suitable for grain boundary sliding (GBS), in view of the fact that only grains with an average size of less than 10µm can deform by GBS [12, 13].
However, many coarse grains with grain size larger than 10 µm can still be found in Fig. 7 (d).

Among the procedures devised for grain refinement, severe plastic deformation (SPD) has been proven capable of producing ultrafine grained (UFG) materials with sub micron grain size.
Increasing the number of ECAP passes an increase in hardness across the normal direction was observed.
The sample deformed by one ECAP pass at room temperature presented a large number of deformation twinning (Fig. 3a).
After HPT (Fig. 5d) the microstructure is characterized by both elongated and equiaxed grains which are separated by sharp grain boundaries as shown in the micrograph, and the diffraction pattern indicate a great misorientation between individual grains.
At 480 oC only a small number of crystals exhibit deformation twins; grains are equiaxed and with a amean diameter of 75 nm (Fig. 5e).

In this study, the effects of HIP treatment, WC grain size and content of βt phase on bending fatigue characteristics were investigated for fine and coarse-grained WC-6.7mass%βt-10.4mass%Co alloys （6.7βt(F) and 6.7βt(C) alloys） and fine-grained WC-20mass%βt-11.3mass%Co alloy (20βt(F) alloy) by comparing with the case of fine-grained WC-10mass%Co alloy（10Co(F) alloy） without βt phase.
In fatigue test, the stress in the range from 0 to maximum stress and with 10Hz frequency was applied up to fatigue failure and then S-N curves showing the relation between applied stress and number of cycles up to failure were obtained.
The fatigue strength of HIP-ed 6.7βt(F) alloy at a definite number of cycle was remarkably low in high stress range and nearly equal in low stress range, compared with that of HIP-ed 10Co(F) alloy.
(a), agglomerate of βt grains; (b), Co pool with a step-shaped form.
Conclusion The effects of HIP treatment, WC grain size and content of βt phase on bending fatigue characteristics were investigated for fine and coarse-grained WC-6.7mass%βt-10.4mass%Co alloys （6.7βt(F) and 6.7βt(C) alloys） and fine-grained WC-20mass%βt-11.3mass%Co alloy (20βt(F) alloy) by comparing with the case of fine-grained WC-10mass%Co alloy（10Co(F) alloy） without βt phase.

The initial microstructure of the sample B is shown in Fig.1b, in which a large number of fine precipitated particles are smaller than 1mm.
In Fig.2a, a larger number of uniform distribution second-phase particles can be observed.
In addition, some of original grain boundaries become bent or waved, which suggests that sub-grains may be formed along grain boundaries.
It can be seen that original grains elongated and a few recrystallized grains can be observed along grain boundaries when the height reduction is 60%(Fig.4a).
From Fig.5c, it can be seen that a large number of sub-grains retained in the elongated grains.

Most the grains present rosette-like and globular-like, there hardly being the grains with dendritic-like.
As the addition of lanthanum increases to 0.8%, more primary grain presents the rosette-like, and the grain with particle-like is decreased.
It is seen from the Fig.3 that there are strip-like or needle-like bright areas at the grain boundaries, which should be the enriching area of La theoretically, because the atomic number of La is biggest during the elements (such as Al, Si, Mg, La) contained in the alloy used in this test.
Moreover, lanthanum is a surface active element to reduce the interface tension of liquid alloy and to decrease the nucleation power of the grain, in which the critical nucleus radius decreases and the formation of nucleus is easy so that a great number of nuclei is produced in the melt to fine the microstructure.
Therefore, the latent heat released from the growth of dendrite raises the temperature at the linking section of arm of the dendritie to impel the dendrite locally to be neck shrunk and fused, thus the more crystallizing nuclei can be produced in the melt. 2) the eutectic reaction in Al-La alloy preferentially happens at the temperature higher than the liquidus temperature in A356 alloy, so that the α-Al grains crystallized from the eutectic reaction can provides the effective nuclei for formation of primary grain in A356 alloy. 3) La is a surface active element to reduce the interface tension of liquid alloy and to decrease the nucleation power of the grain, in which the critical nucleus radius decreases and the formation of nucleus is easy so that a great number of nuclei is produced in the melt to fine the as-cast microstructure.

The change of the damping properties of graphitic steel caused by thermomagnetic treatment is relatively small, due to a large number of diamagnetic inclusions of graphite and quite fine-grained ferrite matrix.
This is explained by high density of diamagnetic graphite inclusions in a relatively fine-grained ferrite matrix (the average grain diameter equals to about 20 μm).
The sources of enrichment can be, firstly, segregation of Cr atoms at the grain boundaries, which are diffused throughout the grain body under thermomagnetic treatment, which is most likely in the fine-grained alloys; and, secondly, inclusions dissolve under thermomagnetic treatment.
The numbers on the distribution function of ultra-fine magnetic fields (Fig. 3) in the order 1–9 correspond to atomic interactions: Fe-Fe, Fe-Cr2, Fe-Cr1, Fe-(Cr1,Cr2), Fe-(Cr1,2Cr2), Fe-(2Cr1,Cr2), Fe-(2Cr1,2Cr2), Fe-(2Cr1,3Cr2), Fe-(3Cr1,2Cr2), where the index 1 or 2 is the number of the coordination sphere in which Cr atom is located, and the numerical coefficients are the quantity of Cr atoms in these coordination spheres.
In graphitized high carbon steels with ferrite-graphite structure, the change in damping properties as a result of thermomagnetic treatment is quite insignificant, which is explained by a large number of diamagnetic graphite inclusions in a relatively fine-grained ferrite matrix. 3.

The results show that the as-cast microstructure of the Mg-Zr alloys was composed of non-dendritic, equiaxed Mg grains, with a few Zr particles within the Mg grains and along grain boundaries.
A few zirconium particles are distributed within the Mg grains and along grain boundaries.
The number of large particles(e.g.>1μm)also increased with increasing in Mn concentration.
Summary Solidification of dilute Mg-Zr alloys produced equiaxed grains, with a small number of Zr particles within the Mg grains and along grain boundaries.
The Mg- Mn alloys solidified as columnar grains.

In the AA1100 alloy sheets, the grains were elongated along RD and the thickness of the elongated grains gradually decreased with increasing the number of ARB cycles.
Specimen Annealed 2 Passes 4 Passes 7 Passes 10 Passes 13 Passes Grain Thickness (µm) 35 4.5 1.8 1.1 0.5 0.45 Grain Length (µm) 33 14 6.1 3.4 1.3 1.1 In the sheet ARB processed by 2 cycles, a number of dislocations were generated in the original grains and it made subgrain structures.
The fraction of the UFG regions increased with increasing the number of ARB cycles, i.e. strain.
With increasing ARB cycles the grains were elongated along RD and the thickness of the grains gradually decreased
(2) The reduction of the grains thickness and length was carried out with increasing the number of ARB cycles, i.e. strain