PHASE DIAGRAMS AS TOOLS FOR ADVANCED MATERIALS DESIGN : APPLICATIONS TO NON-FERROUS ALLOYS

It is recognized that phase diagrams and a thermodynamic database constructed by the CALPHAD (Calculation of Phase Diagrams) approach play an important role in alloy design and materials development. In this paper, recent progress on the thermodynamic database for micro-solders and Cu-base alloys, which is useful for the development of Pb-free solders and predictions of interfacial phenomena between solders and the Cu substrate in electronic packaging technology, is presented. In addition, typical examples of the application of phase diagrams are presented to facilitate the development of shape memory alloys, Co-Cr base magnetic recording media, ferromagnetic Heusler alloys and Co base superalloys.


INTRODUCTION
Since the pioneering work in the computer calculation of phase diagrams by Kaufman and Bernstein [1] in 1970, the CALPHAD (computer coupling of phase diagrams and thermochemistry) method [2,3] has been established due to progress in computer software such as Thermo-Calc [4], Fact-Sage [5,6], Pandat [7], etc., which are widely used for materials design.Thermodynamic databases are now available for several alloy systems, such as Fe, Ni.Ti, Al, Mg-based alloys, ceramic systems and alloy semiconductors, as shown in Table 1.
These thermodynamic databases have been developed based on experimental data of thermodynamic properties and phase equilibria in the literature, as shown in Figure 1.
Recently, calculated results obtained by first principle calculation have also been applied to thermodynamic assessment in the same way with the experimental data.Once the thermodynamic databases have been constructed, many types of properties indispensable for alloy design, computer simulation, etc., can be predicted by the above-mentioned software for calculation.Tecnol.Metal.Mater.Miner., São Paulo, v. 13, n. 1, p. 46-63, jan./mar.2016 of Pb [13].For the development of suitable alternatives to Pb-Sn solders, the melting temperature is one of the most important properties to be optimized.A thermodynamic database of micro-soldering materials named ADAMIS (Alloy Database for Micro-Solders) listed in Table 1 has been developed [8][9][10], by which phase equilibria can be calculated in the whole composition range for the elements Ag, Bi, In, Cu, Sn, Sb, Zn and Pb and in a limited composition range for Al and Au. Figure 2 shows the liquidus surface projection of six Sn-base ternary systems calculated using the ADAMIS database [14], in which the isothermal liquidus lines and reaction types are indicated.Melting temperatures of Sn-base ternary alloys can be predicted at a glance using these calculated diagrams.
The solidification process of solders is also an important factor for the alloy design of Pb-free solders.Although the Scheil model assumes that local equilibrium occurs at the liquid/solid interface and that diffusion is absent in the solidified phases, such calculations can still provide useful information for the prediction of reality.Solidification simulation of a promising candidate as an alternative to Pb-Sn solders, Sn-2.0Ag-0.5Cu-7.5Bi(mass%) alloy, has been carried out.Figure 3 shows the variation of the calculated mass fraction of the solid phases with the temperature of this alloy under equilibrium and Scheil model solidification conditions.
In both cases, the solidification starts with the primary crystals of η-Cu 6 Sn 5 , and the liquid phase disappears at 177.9°C under the equilibrium solidification condition.According to the Scheil model, however, Bi is concentrated in the liquid phase during solidification, which causes an extensive fall of the terminating temperature of solidification (139.9°C),corresponding to the eutectic reaction of the Sn-Bi binary system.Many other properties related to phase diagrams can be calculated and can be used for the development of Pb-free solders not only for conventional temperature ranges around 200°C but also for high temperature applications above 300°C.

Egg-type Microstructures
Recently, we have shown that an egg-like core/shell microstructure can be obtained in powder particles prepared by conventional gas atomization [15], in which the minor liquid phase always forms the core of the egg.Since this finding, extensive studies have been carried out on the mechanism [16][17][18] of the formation of the egg-type microstructure, the role of interfacial energy [19] and bulk ingots with a core-type microstructure [20].Figure 4 shows a schematic illustration of the relation between the microstructure of powder particles and a liquid phase miscibility gap in a phase diagram [21].An egg-type core/shell structure is formed when the chemical composition of an alloy is located in the middle portion of the miscibility gap, while the in situ composite which shows that a fine second phase dispersed in the matrix is formed just inside both ends of the miscibility gap.Therefore, thermodynamic calculation of the liquid miscibility gap plays a key role in the study of the microstructural evolution of powder.
Bi-based alloys are candidates for high-temperature Pb-free solders.Figure 5d shows a calculated phase diagram of the Bi-Cu-Sn ternary system at 650°C, in which a metastable liquid phase miscibility gap in the Bi-Cu binary system is stabilized by the addition of Sn.
Thus, the formation of the egg-type microstructure can be expected in this ternary system.Since the details of Bi-based solder properties such as phase relation, microstructure, interfacial reaction, reliability, etc., have already been reported [21][22][23], only typical microstructures of an egg-type core/shell and in situ composite powder particles are shown in Figures 5a and 5c, respectively.The shell of egg-type particles and the matrix of the in situ composite consist of a Bi-rich phase which plays the role of high-temperature solder.On the other hand, a Cu-Sn intermetallic compound forms the core and dispersion particles, which are expected to improve the mechanical properties.The melting temperatures have been found to be about 270°C and 730°C, which correspond to those of Bi-rich phase and Cu 3 Sn-compound [21].
On the other hand, three kinds of Ag-based powders with high electrical conductivity, i.e., Ag-Cu-Ni, Ag-Cu-Co and Ag-Cu-Fe alloys, were prepared making use of the egg-type microstructure caused by immiscible liquid [24].In these ternary systems, the morphology of the powder obtained by the gas-atomizing method is schematically summarized in Figure 6.
The in situ composite and egg-type micro-powders can be produced from immiscible alloys whose chemical composition is located near both ends and around the middle of the miscibility gap, respectively.In between these two composition regions shown by the shaded areas in Figure 6, mixed powder of in situ composite and egg-type microstructures can be obtained.

High-strength and High-conductive Cu-Ni-Al Based Alloys
In the field of materials science, combinatorial techniques, for instance making use of a diffusion couple (DC) or triple (DT) samples, has been recognized as a powerful and effective method for the determination of phase diagrams, including not only equilibrium compositions but also transition boundaries of martensite/austenite, ferro-/para-magnetic, order/disorder phases, etc., simultaneously in a wide range of chemical composition [25].Here we would like to show how the combinatorial technique was applied to develop high-strength and high-conductive Cu-Ni-Al based alloys.Figure 7 shows the isothermal section phase diagram of the Cu-Ni-Al ternary system at 500°C calculated from the thermodynamic database, where the two-phase region of fcc (Cu) and L1 2 (Ni 3 Al) phases are predicted.
This two-phase region is expected to have high strength and high electrical conductivity strengthened by Ni 3 Al precipitation.To find the optimum compositions, the combinatorial method was used.Ingots of pure Cu, Cu-30Ni and Cu-20Al (mass%) were prepared by induction melting.Each ingot was cut into small cubic pieces, the surfaces of which were mechanically polished to remove contaminated layers.Cubic samples of Cu and Cu-30Ni were first diffusion bonded at 900°C for 1 hour.The Cu/Cu-30Ni diffusion couples and Cu-Al alloys were diffusion bonded as the Cu/Cu-30Ni interface was aligned perpendicular to the bonded surface of the Cu-Al samples.The prepared DT samples were heat-treated at 1000°C for 60 days to obtain the solid solution fcc-(Cu) phase, in which considerable amounts of Ni and Al were solved to form a wide range of chemical compositions in one DT sample.After that, the solution-treated DT samples were aged at 500°C for 6 to 18 hours to obtain a precipitated microstructures, which can be predicted from the calculated phase diagram shown in Figure 7.Chemical compositions of the solution-treated DT samples were determined by an Election Probe Micro Analyser (EPMA) (JEOL, JXA-8100).Vickers hardness and conductivity were measured by a micro Vickers hardness tester (Akashi, MVK-H1) and an electrical resistance tester for tiny areas equipped with a 0.1 mm-interstices four-point probe, respectively.Precipitated microstructures of aged DT samples were examined by a TEM (JEOL, JEM-2000EXII).
Figure 8 shows a dark field image taken from a (100) L1 2 super-lattice reflection of Cu-12.3Ni-1Al-0.3Si(mass%) alloy aged at 500°C for 12 h, the chemical composition of which was optimized based on results of combinatorial experiments.
Fine particles of the L1 2 -Ni 3 Al phase were dispersed presumably due to the spinodal decomposition, which results in a considerable increase of both hardness and electrical conductivity.The relationship between the tensile strength and electrical conductivity is summarized in Figure 9. Optimized Cu-Ni-Al-Si alloys, which show especially high-strength properties, can be substituted for conventional Cu-Be alloys.This alloy system has been commercialized as suspension wire for optical pickup in DVD drives.

Cu-Mn-Al Shape Memory Alloys
Cu-based shape memory alloys (SMAs) such as Cu-Al-and Cu-Zn-based alloys are commercially attractive for practical applications of the shape memory effect (SME) and pseudo elasticity (PE) because of their low cost and their advantages with regard to electrical and thermal conductivities.However, the polycrystalline Cu-based SMAs are generally brittle and tend to cause intergranular fractures due to their high degree of order in the parent β phase with B2, D0 3 or L2 1 structure and the extremely high elastic anisotropy ratio in the β phase [26].Kainuma et al. [27] have reported that excellent ductility can be obtained by controlling the degree of order in the β phase of the Cu-Al-Mn SMAs.Figures 10a and 10b show the isothermal section phase diagram at 850°C [28] and vertical ones of Cu-Al-10at.%Mn alloys [29], respectively.This system is well-known for the  Heusler phase Cu 2 AlMn, which is shown by the asterisk in Figure 10a, but this phase is no longer stable at 850°C, where a large solubility range of the β phase can be seen.It is also seen from Figure 10b that both the transition temperatures of order-disorder transitions, β (A2) → β 2 (B2) and β 2 → β 1 (L2 1 ), drastically decrease with decreasing Al content.The Cu-Al-Mn SMAs with an Al content below 18 at.%having a low degree of order in the β phase show excellent cold-workability of over 60% in cold-rolling reduction and also exhibit SME and PE based on cubic β 1 (L2 1 ) to monoclinic β' 1 (18R), as shown in Figure 11.The large grain size of this alloy system shows excellent SME and PE [30,31], which can be achieved by special heat treatment for obtaining the abnormal grain growth [32].This alloy has been practically used for an ingrown nail correcting device [33] and also applied as a seismic damping material [34].

Ni-base Metamagnetic Shape Memory Alloys
Large magnetic-field-induced strains have been observed in Heusler alloys and have been explained as resulting from the rearrangement of martensite structural variants due to an external magnetic field [35][36][37].We have reported some new ferromagnetic shape memory alloys in the NiMnIn, NiMnSn and NiMnSb alloy systems [38].Figure 12 shows the magnetic phase diagram, including the M s and A f temperatures of Ni 50 Mn 50-y X y (X=In, Sn, Sb) alloy systems [38].
The M s and A f temperatures decrease with increasing In, Sn and Sb contents and the slopes in the M s and A f temperatures as a function of X content in the ferromagnetic state are larger than those in the paramagnetic state.This is because the T 0 temperature (=(M s + A f )/2) at which the parent and martensite phases are in equilibrium vary with the magnetic contribution to the Gibbs energy.Especially in the NiMnIn system, the martensitic transformation temperatures in the ferromagnetic state drastically decrease with increasing In content.It is noted that the compositional dependence of the Curie temperatures T C of the austenite phase is small, while the T C of the martensite phase strongly depends on the alloy compositions and drastically decreases with decreasing In, Sn and Sb contents in all the systems.In particular, in the NiMnIn alloy system, the compositional dependence of the T C temperature of the martensite phase is the largest of those in the three alloy systems and the T C of the martensite phase in the Ni 50 Mn 35 In 15 alloy is estimated to  be below -200°C Figure 13 shows the thermo-magnetization curves for the 13.4In alloy at magnetic field strengths of H = 0.5, 20 and 70 kOe [39].
During the initial heating in a magnetic field of H = 0.5 kOe, the thermo-magnetization curve drastically decreased between 375 K and 390 K owing to the transformation from ferromagnetic (or ferrimagnetic) to paramagnetic in the parent phase, where the Curie temperature T C was calculated to be T C = 382 K.The magnetism of the parent phase below T C is considered to be ferromagnetic rather than ferrimagnetic because the temperature dependence of the reciprocal susceptibility in the paramagnetic state follows the normal Curie-Weiss law.After heating to 473 K followed by furnace cooling, the magnetization was cyclically measured at temperatures between 200 K and 330 K in magnetic fields of H = 0.5, 20 and 70 kOe.The magnetization curve of 0.5 kOe declined suddenly at about 292 K as a result of the martensitic transformation and entered a steady region with a very low magnetization on cooling.Furthermore, the magnetization drastically increased at about 300 K as a result of the reverse transformation with a small thermal hysteresis of less than 10 K on heating.Figure 14 shows the magnetization curves at several temperatures.While the curves at 200 K and 320 K exhibit typical non-magnetic and ferromagnetic M-H behaviors, respectively, the curves at 270 K and 290 K show a metamagnetic behavior with a large hysteresis with the magnetic field inducing reverse transformation from non-magnetic martensite to the ferromagnetic parent phase [39,40].These results are in accordance with the thermo-magnetization behavior shown in Figure 13.

Co-Cr-based Magnetic Recording Media
Co-Cr-based sputtered thin films have a high coercivity (Hc) with a high magnetization (M) and have been widely used as magnetic recording media of hard disks [41].The thin films have a hexagonal close-packed (hcp) structure and a fine microstructure with the cobalt-rich ferromagnetic phase precipitated in the chromium-rich paramagnetic matrix [42].The distribution of chromium in each phase and the microstructure have a significant effect on the recording properties of this alloy.Ishida and Nishizawa [43] have suggested that a likely origin of the chromium segregation in a Co-Cr alloy thin film is the miscibility gap in the hcp phase because the thermodynamic calculations of Co-Cr binary system by Hasebe et al. [44] indicate a magnetically  The calculated metastable magnetically induced phase separation in the hcp phase agrees well with the experimental results, strongly implying that the magnetically induced phase separation is responsible for the compositional heterogeneity in Co-Cr-based films.
From the thermodynamic calculation of the Co-Cr-X ternary system, the shape of the magnetically induced phase separation in the cobalt-rich portion can be classified into four types as shown in Figure 16 [46].
In Type I, the chromium contents in both the ferromagnetic and paramagnetic phases decrease with an increasing amount of X.In Type II, the chromium content decreases in the ferromagnetic phase and increases in the paramagnetic phase with an increasing amount of X.In Type III, the chromium content of the paramagnetic phase is increased by the addition of X, whereas that in the ferromagnetic phase is not effectively reduced.Finally, in Type IV, the three-phase equilibrium of one ferromagnetic and two paramagnetic phases is formed.Figure 17 shows the classification of the effect of the alloying elements on the phase equilibria of the Co-Cr system in the periodic table based on the calculated results.
Pd, Rh and Ir show a type of miscibility gap similar to that of the Co-Cr-Pt phase diagram.The magnetic properties of the Co-Cr-Pd thin-film alloy would be expected to be similar to those of the Co-Cr-Pt system.On the other hand, the addition of iridium to cobalt strongly decreases the magnetic anisotropy energy (MAE), which in turn would cancel out the beneficial effect of decreasing chromium content in the ferromagnetic phase by the alloying of iridium [47].This means that the MAE of each binary system, as well as the phase diagram, should be considered for the material design of Co-Cr-based high-density magnetic recording media.(Y=Ti, V, Cr, Mn, Fe and Z=Al, Ga) Heusler alloys [49][50][51].T C increases with the increase of N 3d and the values are not so different between the Al and Ga series.On the other hand, T t B2/L21 tends to indicate behavior opposite to that of T C , i.e., T t B2/L21 decreases with the increase of N 3d except for Y = Cr.Although the tendency of T t B2/L21 as a function of N 3d is similar with respect to one another, Z = Al and Ga, actual values of T t B2/L21 for Z = Ga are higher than that of the Z = Al series in whole.The Co-based Heusler alloys with higher values of T t B2/L21 are easier to order to the L2 1 -type structure.This means that the confirmation of T t B2/L21 is very valuable to efficiently heighten the degree of order by thermal annealing in industrial processing and is a very important component in the technology of spintronics [52].

Ni-base Superalloys
Ni-base superalloys are widely used as high-temperature materials in aircraft engines, gas turbines, reactors, chemical industry, etc., where the γ' phase of Ni 3 Al with the L1 2

Ferromagnetic Heusler Alloys
Since the ferromagnetic Cu 2 MnAl phase was discovered by Fritz Heusler in 1903, more than 1000 compounds have become known as Heusler compounds [48].These compounds are very attractive for use as multi-functional materials in magnetic recording, superconductors, shape memory, thermoelectrics, etc.The phase stability of Heusler compounds is the basic information for the development of functional materials.The present authors' group has investigated the phase stability of Co-based and Ni-based Heusler compounds.Figure 18 shows the ordering and phase separation between the B2 and L2 1 (Heusler) phases in the X 2 AlTi (X=Fe, Co, Ni, Cu) compounds [49].
Surprisingly, the phase stability, i.e., the maximum temperature of the B2/L2 1 ordering ), the tricritical temperature (T t

B2/L21
) and the phase boundaries of miscibility gap between B2 and L2 1 phases, can be predicted by the valence electron concentration of the X site atoms.Figure 19 shows the relation between the critical temperatures and 3d + 4s electron concentration in the X site.The β (NiAl, CoAl:B2) phase has been used for the surface coating of Ni-base superalloys because of its high resistance to oxidation [55].Therefore, data on phase equilibria involving the γ (A1), γ' (L1 2 ) and β (B2) phases in the Ni-Al-base system are fundamentally important for the practical design of Ni-base superalloys.Figure 22 shows the isothermal phase diagrams relating to the γ, γ' and β phases in the Ni-Al-X ternary systems [56], where the phase relations can be classified into six categories depending on the tie line and the existence of the γ' phase.The upper part of the diagrams in Figure 22 shows that the γ' phase disappears with alloying elements such as Mn, Cr, Fe, Co and Cu, while lower part of the diagrams shows that Ti, V, Ta, Mo, W and Si stabilize the γ' phase and that no direct γ/β equilibrium is obtained.
The partition behavior of alloying elements among the γ, γ' and β phases is shown in Figure 23.
The tendency for portioning among the γ, γ' and β phases is strongly related to the stability of Ni 3 X, NiX, and XAl compounds in the binary systems.For example, the elements X forming stable Ni 3 X, such as Ti, Zr, Hf, V, Nb, Ta, Si, Ge, and Sn, are strong γ' formers.Such a concept about the expected tendency of alloying elements in stabilizing the γ, γ' and β phases is very useful in designing and improving the ductility of γ' (Ni 3 Al) and β (NiAl) base alloys.

Co-base Superalloys
Although no γ' Co 3 Al compound is formed in the Co-Al binary system as shown in Figure 21, from the viewpoint of enthalpy of formation of Ni 3 Al and Co 3 Al estimated by Miedema et al. [57], the difference in ΔH for both compounds is small [58], which suggests that a metastable compound with the L1 2 structure of the Co 3 X would be expected to appear.In fact, the metastable γ' phase of Co 3 Al [59] and of Co 3 Ta [60] is confirmed.In 2006, the authors' group found the γ' phase of Co 3 (Al,W) with the L1 2 structure [61].The SEM micrographs of this γ + γ' structure of Co-base ternaries are shown in Figure 24, where the cuboidal γ' phase homogeneously precipitates in the γ (A1) matrix, which is very similar to the morphology of conventional Ni-base superalloys [61].
The phase diagram of the Co-Al-W ternary system at 900°C is shown in Figure 25, and the composition of cuboidal precipitates observed in Figure 24 is the ternary compound of Co 3 (Al,W), where the composition of Al and W has an almost equiatomic ratio.Although the γ' phase at 900°C in Figure 25 is metastable [62], the γ' phase is stabilized by alloying, Ni [63], Ta, Ti, etc. Figure 26 shows the effect of alloying elements on the partitioning behavior between the γ and γ' phases in Co-8.8Al-9.8W-2X(at.%) alloys [64] as a function of temperature, which is compared with that of the Ni-Al-X ternary systems [56].structure plays a key role in maintaining strength at higher temperatures [53], while the strengthening mechanism utilized Co-base high temperature alloys is principally due to the solid solution hardening with refractory element and carbide precipitation [54].This difference in the strengthening phase of Ni-base and Co-base alloy is due to the existence of the γ' phase as shown in Figure 21, where the stable γ'   The partitioning behavior of alloying elements in the Co-Al-W-X system is very similar to that of the Ni-Al-X system, where Ta, Nb and Ti are strong γ' stabilizing elements, while Cr, Mn and Fe are γ-forming elements.
High temperature strength and deformation of the γ + γ' two-phase structure of Co-Al-W base alloys have been studied [65][66][67], results showing that the positive temperature dependence of flow stress is observed above 600°C in Co-Al-W-base alloys similar to that of Ni-base superalloys.The creep properties of Co-Al-W-base alloys have been reported [68], the strength of these alloys being comparable to that of the Ni-base superalloys of the first generation.
Cast Co-base superalloys have been developed for application in friction stir welding (FSW) tools [69].FSW is a solid state joining technology that can be used to weld most Al alloys, where the tool steel is used as a joining tool.However, FSW tools that enable welding of high-softeningtemperature materials such as steels and titanium alloys are very limited.W-Re and polycrystalline cubic boron nitride (PCBN) tools have been developed, but they are too expensive to be used.Cast Co-Ni-Al-W-base alloys have been applied as a new tool material, which enables FSW for Ti-base alloys, mild steel, Cu-and Zr-base alloys [69].The Co-base tools for FSW are produced by a lost-wax process and then heat-treated to obtain the γ' phase in the matrix.The weld appearance of Ti-6Al-4V is shown in Figure 27, where a new Co-base alloy tool produces a defect-free weld with a smooth, shiny surface.
Although, development of Co-base superalloys using the γ' Co 3 (Al,W) phase has just been started, significant results have already been obtained.This is because phase equilibria, alloying behavior, microstructural control, mechanical properties, etc., of Co-base superalloys are very

SUMMARY
A thermodynamic tool for design of materials in electronic packaging has been developed on the basis of the CALPHAD method.This tool can provide much valuable information on thermodynamics such as stable and metastable phase equilibria, phase fraction, etc. Examples were herein given to demonstrate the great utility of these databases for designing Pb-free solders and high-strength Cu-base alloys.In addition, examples for development of shape memory alloys, Co-Cr base magnetic recording media and high-temperature alloys using the γ' phase with the aid of phase diagrams were also introduced.

Figure 2 .
Figure 2. Liquidus surface projection diagrams in the Sn-base ternary systems.

Figure 4 .
Figure 4. Schematic illustration of microstructural evolution of powder particles related to the phase diagram with liquid phase miscibility gap.

Figure 7 .
Figure 7. Calculated phase diagram at 500°C in the Cu-Ni-Al ternary system.Shaded area represents a miscibility gap between the fcc (Cu) and L1 2 Ni 3 Al phases.

Figure 9 .
Figure 9. Relationship between tensile strength and electrical conductivity of developed Cu-Ni-Al-Si alloy in comparison with conventional Cu-base alloys.

Figure 10 .
Figure 10.(a) Isothermal section phase diagram of the Cu-Al-Mn system at 850°C; (b) Vertical section phase diagram of Cu-Al-10at.%Mn alloys.

Figure 12 .
Figure 12.Martensitic and magnetic transition temperatures of the (a) Ni 50 Mn 50-y In y ; (b) Ni 50 Mn 50-y Sn y and (c) Ni 50 Mn 50-y Sb y alloys, where Para and Ferro mean paramagnetic and ferromagnetic, respectively, and A and M indicate the austenite and martensite, respectively.See Sutou et al. [38] for additional information.

Figure 13 .
Figure 13.Thermo-magnetization curves of the Ni 45 Co 5 Mn 36.6In 13.4 alloy measured in several magnetic fields.

Figure 15 .Figure 16 .
Figure 15.Phase diagram of the Co-Cr system in the Co-rich portion from Oikawa et al. [45].See Oikawa et al. [45] for further information.

Figure 17 .
Figure 17.Classification of elements in the Periodic Table for the miscibility gap of the hcp phase in the Co-Cr-X system.

Figure 19 .
Figure 19.The critical temperatures of max T C B2/L21 and T t B2/L21 evaluated by extrapolation from the experimental data vs 3d+4s electron number of Fe, Co, Ni and Cu.See Ishikawa et al. [49] for details.

Figure 20 .Figure 21 .
Figure 20.Order-disorder phase transition temperatures from L2 1 to B2 phase, T t B2/L21 and Curie temperature, T C , as a function of the number of valence electrons for 3d in Y elements for Co 2 YZ (Y=Ti, V, Cr, Mn, Fe and Z=Al, Ga) alloys.

Figure 25 .
Figure 25.Isothermal section phase diagrams of the Co-Al-W ternary system in the Co-rich portion at 1173K.

Figure 26 .
Figure 26.Effect of alloying elements on distribution coefficient between the γ and γ' phases in the Co-Al-W-X systems.