Hypothetical alloys

Compounds discussed in some detail in the text hypothetical alloy type... 

These principles are illustrated with reference to the hypothetical alloy system shown in Fig. 12-3. This system contains two substitutional terminal solid solutions a and p, both assumed to be face-centered cubic, and an intermediate phase y, which is body-centered cubic. The solubility of either A or B in y is assumed to be negligibly small the lattice parameter of y is therefore constant in all alloys in which this phase appears. On the other hand, the parameters of a and P vary with composition in the manner shown by the lower part of Fig. 12-3. Since the B atom is assumed to be larger than the A atom, the addition of B expands the A lattice, and the parameter of a increases from a for pure A to 03 for a solution of composition x, which represents the limit of solubility of B in A at room temperature. In two-phase (a -t- y) alloys containing more than x percent B, the parameter of a remains constant at its saturated value a. Similarly, the addition of A to B causes the parameter of P to decrease from 02 to 04 at the solubility limit, and then remain constant in the two-phase (y -1- P) field. 

Fig. 12-3 Phase diagram and lattice constants of a hypothetical alloy system.

Hypothetical alloys A, B, and C described in Fig. 4.17 are used to illustrate the alloy evaluation under reducing and oxidizing conditions. For reducing (or active state of the alloy) conditions (a), the alloy C is superior because of the lowest corrosion rates compared with the other two alloys in the active region. 

The main part of Figure 4 refers to data for pure elements. It is so regular that one expects it to be relevant also for, e.g., binary (thermodynamically stable or hypothetical) alloys connecting two adjacent elements. There are few extensive sets of data for transition-metal alloys with components from the same row in the Periodic Table, apart from systems where magnetism may strongly affect the composition-dependence of the elastic constants. However, for the d -row Zr-Nb and Nb-Mo bcc alloys there are detailed experimental data for all three elastic constants. The results for C, marked with plus-signs in Fig. 4, fit well the general trend for the Sd elements. 

The atomic weight, density, and atomic radius O for three hypothetical alloys are listed in the following table. For each, determine whether its crystal structure is FCC, BCC, or simple cubic and then justify your determination. 

Some hypothetical alloy is composed of 25 wt% of metal A and 75 wt% of metal B. If the densities of metals A and B are 6.17 and 8.00 g/cm, respectively, and their respective atomic weights are 171.3 and 162.0 g/mol, determine whether the crystal structure for this alloy is simple cubic, face-centered cubic, or body-centered cubic. Assume a unit cell edge length of 0.332 nm. 

Metals A and B form an alloy or solid solution. To take a hypothetical case, suppose that the structure is simple cubic, so that each interior atom has six nearest neighbors and each surface atom has five. A particular alloy has a bulk mole fraction XA = 0.50, the side of the unit cell is 4.0 A, and the energies of vaporization Ea and Eb are 30 and 35 kcal/mol for the respective pure metals. The A—A bond energy is aa and the B—B bond energy is bb assume that ab = j( aa + bb)- Calculate the surface energy as a function of surface composition. What should the surface composition be at 0 K In what direction should it change on heaf)pg, and why ... 

Reaction 5.45 is at least partly hypothetical. Evidence that the Cl does react with the Na component of the alanate to form NaCl was found by means of X-ray diffraction (XRD), but the final form of the Ti catalyst is not clear [68]. Ti is probably metallic in the form of an alloy or intermetallic compound (e.g. with Al) rather than elemental. Another possibility is that the transition metal dopant (e.g. Ti) actually does not act as a classic surface catalyst on NaAlH4, but rather enters the entire Na sublattice as a variable valence species to produce vacancies and lattice distortions, thus aiding the necessary short-range diffusion of Na and Al atoms [69]. Ti, derived from the decomposition of TiCU during ball-milling, seems to also promote the decomposition of LiAlH4 and the release of H2 [70]. In order to understand the role of the catalyst, Sandrock et al. performed detailed desorption kinetics studies (forward reactions, both steps, of the reaction) as a function of temperature and catalyst level [71] (Figure 5.39). 

In February 1909, the results of the experiments on nitride formation had led to the outline of a patent application which covered the preparation of metal nitrides in the presence of auxiliary substances. Following a hypothetical concept of the action of these additions, they were defined as flux promoters." This draft of an application ended with the following sentence Finally, it is also advantageous to add a flux promoter to metals or alloys which serve as catalysts for the ammonia synthesis. This statement was made in view of the early catalytic experiments in which we had observed the synthesis of traces of ammonia in the presence of catalysts similar to those which acted favorably for the nitride formations. 

As evident from Fig. 4, the Curie temperature of the austenitic phase decreases upon substitution of Mn for Ni. This is due to the fact that upon such a substitution a number of Ni atoms with smaller, as compared to Mn, magnetic moment increases. A similar tendency presumably takes place for a hypothetical Curie temperature of the martensitic phase. This follows from the fact that, as is seen from Fig. 3, in the low-temperature martensitic phase the magnetization of the Ni2+a Mni a Ga alloys with high Ni excess decreases more rapidly with increasing temperature. 

The collected data allow us a rough estimation of the hypothetical Curie temperature of the martensitic phase. Figure 10 shows temperature dependencies of the reduced spontaneous magnetization m = Ms(T)/Ms(0) of the alloys as a function of reduced temperature t = T/Tc. It is seen that the magnetization of the austenitic phase and... 

How the hypothetical reaction pathway represented by Eqs. (1) to (3) may be accomphshed in a real bimetallic alloy nanoparticle Recently, Bard and co-workers discussed the possibihty to completely remove Pt from the alloy systems and proposed thermodynamic guidehnes for the design of bimetallic catalysts for dioxygen elecfroieduction. Furthermore, Wang and Balbuena... 

In Figure 22(a), an approximately 2-mm catalyst fragment shows an evidently similar peripheral penetration of alloy at X40 magnification. In Fig. 22(b), the penetration is increased toward the interior. At the highest pressure, 50 atm, there is substantial penetration throughout the whole interior, and Fig. 22(c) shows behavior exactly similar to the virtual section in Fig. 21(c). Also clearly evident in Fig. 22(c) are dark spaces identifiable with unpenetrated pore voids. These have an apparently similar appearance and spatial distribution to the dark unpenetrated hypothetical voids shown black in Fig. 21(c). 

A very clear distinction between the total DOS D s) and the LDOS D e,t) is shown by calculations on binary alloys [51]. Look, e.g., at the (partly hypothetical) series of isoelectronic 1 1 alloys TcTc, MoRu, NbRh, ZrPd, YAg, where the alloying partners have nominal valence differences between 0 (for pure Tc) and 8 (for YAg). As shown in Figure 5, the overall density of states curves look more or less alike for all five alloys, but the partial densities on the sites of the individual partners are very different. Such curves also show why the so-called collective electron model does not work for catalytic activity [52, p. 458], or even for alloy properties in general. 

Fig. 3.20. (b) Free-enthalpy curves of a hypothetical A-B alloy (lower) and (a) the resultant... 

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