SUMMARY INTRODUCTIONA historical perspective on the development of AI-Li alloys was given by Balmuth and Schmidt at the recent International Aluminum-Lithium Conference and is included in the proceedings published by The Metallurgical Society of AIME.l The development of AI-base alloys containing lithium began in Germany in the 1920s and was primarily concerned with additions of small amounts of lithium to age-hardening alloys to increase their strength. In the 1950s metallurgists at Alcoa recognized that lithium increased the elastic modulus of aluminum and they developed the high strength AI-Cu-Li alloy 2020. However, this alloy had low ductility and fracture toughness in the maximum strength temper. These limitations, as well as production problems, led to its withdrawal as a commercial alloy in 1969.Since 1973, the rapid escalation of fuel costs has accelerated research on developing more fuel-efficient aircraft and, consequently, on developing new Al-Li alloys, for one way to achieve more fuel-efficient aircraft is to reduce weight. Recent studies by Lewis and coworkers 2 have shown that this is most effectively accomplished by reducing the density and increasing the specific modulus and strength of the materials used for airframe components. Since aluminum alloys comprise 80 wt. % of the airframe, metallurgists are actively pursuing ways to effect such improvements in these critical materials. Besides beryllium, which has associated manufacturing and health related problems, lithium is the only known metal which improves both modulus and density when alloyed with aluminum. Each weight percent lithium added to an aluminum alloy reduces the density approximately 3% and increases the elastic modulus approximately 6%3 for lithium additions up to 4%.

Zirconium diboride and a zirconium diboride/tantalum diboride mixture were synthesized by solution‐based processing. Zirconium n‐propoxide was refluxed with 2,4‐pentanedione to form zirconium diketonate. This compound hydrolyzed in a controllable fashion to form a zirconia precursor. Boria and carbon precursors were formed via solution additions of phenol–formaldehyde and boric acid, respectively. Tantalum oxide precursors were formed similarly as zirconia precursors, in which tantalum ethoxide was used. Solutions were concentrated, dried, pyrolyzed (800°–1100°C, 2 h, flowing argon), and exposed to carbothermal reduction heat treatments (1150°–1800°C, 2 h, flowing argon). Spherical particles of 200–600 nm for pure ZrB2 and ZrB2–TaB2 mixtures were formed.

Minor additions of Sc are effective in controlling the recrystallization resistance of 5xxx, 2xxx, and 7xxx aluminum. The addition of Sc to aluminum results in the rapid precipitation of homogeneously distributed Al 3 Sc dispersoids, which are coherent with the matrix and have the L1 2 structure. The presence of Al 3 Sc dispersoids increases the recrystallization resistance of wrought alloys. The higher coarsening rate of Al 3 Sc compared to that of Al 3 Zr may limit its applications as a single ancillary addition. When both scandium and zirconium are used in the same alloy, Al 3 (Sc 1-x , Zr x ) dispersoids form. These dispersoids are more effective recrystallization inhibitors than either Al 3 Sc or Al 3 Zr. The Al 3 (Sc 1-x , Zr x ) dispersoids precipitate more rapidly than Al 3 Zr but have a slower coarsening rate than Al 3 Sc. Furthermore, the distribution of Al 3 (Sc 1-x , Zr x ) is significantly more homogeneous than Al 3 Zr. It was also established that alloys containing up to 3.5Mg showed improvement in recrystallization resistance when both Sc and Zr were present. Several morphologies of Al 3 Sc and Al 3 (Sc 1-x , Zr x ) were also observed.

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Fatigue crack growth (FCG) experiments were conducted on controlled variations of Type 7075 and 7050 aluminum alloys. Alloy FCG resistance was ranked under constant amplitude and simple variable amplitude load spectra. Fracture mechanics and fractographic approaches were used to interpret causes for variation in ranking of 7XXX aluminum alloy FCG resistance with loading conditions. The interpretation is built around clarification of a controlling FCG mechanism that is dependent upon interaction of microstructure and load history. This clarification represents a necessary first step toward knowing which microstructure or which design (test) procedure is optimum for a particular class of application, for example, fighter as opposed to bomber or transport aircraft.

A thermodynamic modeling of phase diagrams in binary alkali silicate systems is described by using the Gibbs free energy of the phases. Simple models employing regular, quasi-regular, and subregular solution models were applied to describe the liquid phase. The interaction parameters of the liquid phases were obtained by using the liquidus curves of silica (and subliquidus data) through a multiple linear regression method. The present calculated liquidus curves agree very well with CszO-SiOz and RbzO-SiOz systems. In the KzO-, NazO-, and LizO-SiOz systems, the calculations permitted discrimination between sets of data and, based upon the calculation choices, could be correctly made to select the most appropriate phase diagram. Also, the spinodals were calculated where phase separation can occur by a spinodal decomposition process. [

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A range of compositions and temperatures below the monotectic temperature exists where there are thermodynamic restrictions that prevent the equilibrium solid from forming directly from the undercooled homogeneous liquid. In this region, the solid can form only after liquid-liquid phase separation has occurred. As suggested by earlier research, the thermodynamic restrictions on the crystallization process may be useful to control the crystallized grain structure in glass-ceramic systems. Thus, understanding the thermodynamic limitations on the formation of the solid in monotectic systems could have commercial significance. In the present paper, the metastable liquidus boundaries, liquid miscibility gaps, and spinodal curves in binary MgO-SiO 2 , CaO-SiO 2 , and SrO-SiO 2 systems are calculated by using analytical expressions for the Gibbs free energies of the liquid phases. Calculating the metastable liquidus rather than using a simple extrapolation as originally proposed in the aforementioned previous research provides greater control of the heat-treatment processes and, thus, greater control over the resulting microstructure.