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After a certain number of operating cycles of making and breaking operations, contacts are often destroyed catastrophically although their total commutation wear is not very severe at the moment.
Microstructure transformation investigations of the contact material were carried out successively after different numbers of making and breaking operation cycles by Hitachi S-570 scanning electron microscope and Neophot II optical microscope.
If there are oxides exist along grain boundaries, the binding force of recrystallized columnar grains is weakened.
Based on vast number of electrical tests simulating the servicing conditions on equipment similar to that described in reference [6] (U=380V, I=30A, cosϕ=0.35, with no arc quenching), the relation between the matrix grain size and duration times of the Cu-Cd-Nb-CP composite was concluded and results are given in fig. 5.
It shows that the number of cycles N increases 4-5 times when the particles size of Nb decreases from ∼40 µm to ∼20 µm and then changes slightly.

ECAP has a number of advantages compared with traditional metal processes, such as deformation without changing the shape of objective material [5,6,7,8,9].
Apparently, with the increase in the pressing temperature, the equiaxed α phase becomes smaller in number and the secondary lamellar α phase begins to be separated out.
Additionally, the number of acicular α phase at boundaries was much larger than that in the coarse β grains from Fig. 4b.
Because of no annealing softening, a large number of tangled dislocations which are produced by ECAP deformation have been seen clearly (Fig. 6).
With the increase in pressing temperature, equiaxed grains have become coarser and the number of α phase has decreased.

Cracks on a broken section of DWTT samples can (a) penetrate the coarse grains directly, (b) propagate in Zig-Zaga way in the fine grains, and (c) be around the boundary of original austenite grains.
In order to establish the steel processing for stable and high quality steel plates, especially to guarantee the DWTT property, a number of industrial trials were carried out for developing OHTP technology in Shagang.
A lot of continuous coarse austenite grains present on the middle of 1#sample, in comparison of small austenite grains in 2# sample.
In the case of X80 pipeline steel with acicular ferrite, the original austenite grain plays a key role for the fracture toughness because these grains are the effective grains when a fracture is taken place [5,6].
Fig.13 shows some crack propagations between coarse and fine austenite grains of 1# DWTT sample, in which (a) cracks penetrate the coarse grains directly, (b) cracks propagate in Zig-Zaga way in the fine grains, and (c) cracks are around the boundary of original austenite grains.

Grain sliding or other types of grain boundary accommodation mechanisms are thought to predominate in this regime.
The differences in the grain growth observed can therefore be attributed to differences in grain boundary mobility.
Grain boundary sliding, or motion of the grains relative to each other in a direction parallel to the grain boundary does not require mass transport to maintain the preferred equilibrium partition of the impurities between grain boundary and bulk.
On the other hand, grain boundary migration and grain growth imply motion of the grain boundaries perpendicular to the grain boundary plane and require mass transport of the impurity together with GB motion in order to maintain the equilibrium segregation ratio.
Our technique will allow the engineering of interfaces for desired grain boundary mobility and mechanical response, something that is crucial for a number of applications such as building better targets for the National Ignition Facility which Figure 9: Results of the shock wave simulation in pure and micro alloyed samples (a) as prepared) (b) pure Cu sample (c) Sample containing 3% Fe.

Fig. 1 shows the general algorithm for Markov Chain - Monte Carlo based model of intergranular crack propagation where N is the number of Monte Carlo steps, n is the number of different GBCDs, M is the maximum crack length and Lavg is the average crack length.
[Rh] = [V(A1), V(A2), V(A3), … … … …, V(An)] [Ri1] = [V(B1), V(B2), V(B3), … … … …, V(Bm)] [Ri2] = [V(C1), V(C2), V(C3), … … … …, V(Cm)] n = number of type 1 TJ m = number of type 2 TJ TJ type 1: T11(i) = 1 - V(Ai) T12(i) = V(Ai) i = 1, 2, 3, … … …, n TJ type 2: T22(j) = 1 - {V(Bj)+V(Cj)-V(Bj)*V(Cj)} T23(j) = V(Bj) T24(j) = V(Cj) j = 1, 2, 3, … … … …, m C1 1 3 4 A1 B1 2 Table 1: Transition Matrix for Fig. 2 Table 2: Grain Boundary Ranking for HM The initial crack location can be expressed in matrix form of the same column size of Tr: p0 = [1 0 0 0]
The susceptibilties (hypothetical) of grain boundaries are given in Table 2.
Contrary to the equiaxed hexagonal grain structure, the voronoi microstructure (see, e.g.
It will enable one to incorporate different grain shapes, grain size distributions etc. in the model microstructure and, determine their effects on the crack propagation behaviour.

Larger tensile thermal strains develop when the (0001) pole of adjacent grains lies closer to the grain boundary normal.
This is a complex problem, and there have been a number of different 2D and 3D approximations (e.g. [2], [3]), using abstract representations of the microstructure.
The numbers of identifiable crystallographic orientations, grains and grain boundaries were 886, 1047 and 5725 respectively (isolated regions of the same orientation were treated as separate grains).
a) b)c)d) Figure 4: The variation of average stress normal to grain boundaries as a function of a) relative misorientation between the (0001) poles in the adjacent grains, b) grain boundary area c) boundaries within a range of misorientation of the (0001) pole and the grain boundary pole for both grains, and d) boundaries within a range of misorientation of the (0001) pole and the grain boundary pole for one or both grains.
The cumulative numbers of grain boundary facets in each range of misorientation or size are also shown.

The results showed that with decreasing grain size the number of twin annealing added And four types of annealing twin in the microstructure, in the end they all become one twin and then turn into grain.
Twin boundaries act like grain boundaries.
On grain boundaries there are {111} plates.
Some important factors affect the rate of twins formation during grain growth including: rate of deformation before annealing, temperature and time cycle of annealing heat treatment, grain size, energy of grain boundaries, migration velocity of grain boundaries, SFE, context of grains and impurities [4].
A stands for corner twins, B stands for a given type of twins connecting both ends of a grain to each other, C stands for a given type of twins restricted within grains and D stands for twins restricted within grains but inclined towards grain boundaries.

The number that comes after "E" is the number of ECAP passes as a pre-state for the plane strain compression.
Fig. 5 (a) shows the evolution of cell structure size with number of passes.
Fig. 5 (b) gives the variation of average misorientation angle with number of passes.
Fig. 5 (c) gives the area fraction with grains less than 1 mm based on a 2o grain tolerance angle.
Fig. 5 (d) represent the evolution of number fraction of high angle grain boundaries (HAGBs), that is boundaries with misorientation across them greater than 15o.

The sequence of phase transformations of zirconium-doped ultrafine-grained alloy Al-Mg-Li in heating is revealed.
Temperature dependence of the lattice parameter for ultrafine-grained alloy at in-situ heating in the diffractometer chamber (1).
After annealing at 380 °C and subsequent quenching metastable S-phase particles of type I are observed to dissolve with increasing number and volume fraction of more stable S-phase particles of type II that are precipitated at high-angle boundaries of general-type grains (Fig. 2b).
With increasing temperature the coagulation of grain-boundary S-phase particles leads to recrystallization.
Microstructure of the ultrafine-grained alloy after 20-minute annealing at the temperatures 300°С (a) and 380°С (b).

These bands normally initiated at grain boundaries and were restricted to the area of one grain.
In the medium-SFE Cu-10%Zn alloy, the grain morphology became blurry and SBs were readily observed in most of the deformed grains.
After two (Fig. 2(b)) and four (Fig. 2(c)) passes, the microstructures in the three materials were further refined and the density of SBs increased with the pass number.
The incomplete central symmetry can be attributed to the coarse grain size at this stage, such that only a small sampling of the grains were measured in the XRD measurements.
For a given pass number, the texture strength was always higher in the pure Cu than in the Cu-10%Zn and then Cu-30%Zn alloys, indicating a weakening of texture with the decrease of SFE.