STRUCTURE AND PROPERTIES OF RAPIDLY SOLIDIFIED ALUMINUM-LITHIUM ALLOYS

Transcription

1 STRUCTURE AND PROPERTIES OF RAPIDLY SOLIDIFIED ALUMINUM-LITHIUM ALLOYS N. Kim, R. Bye, S. Das To cite this version: N. Kim, R. Bye, S. Das. STRUCTURE AND PROPERTIES OF RAPIDLY SOLIDIFIED ALUMINUM-LITHIUM ALLOYS. Journal de Physique Colloques, 1987, 48 (C3), pp.c3-309-c < /jphyscol: >. <jpa > HAL Id: jpa Submitted on 1 Jan 1987 HAL is a multi-disciplinary open access archive for the deposit and dissemination of scientific research documents, whether they are published or not. The documents may come from teaching and research institutions in France or abroad, or from public or private research centers. L archive ouverte pluridisciplinaire HAL, est destinée au dépôt et à la diffusion de documents scientifiques de niveau recherche, publiés ou non, émanant des établissements d enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

2 JOURNAL DE PHYSIQUE Colloque C3, supplgment au n09, Tome 48, septembre 1987 STRUCTURE AND PROPERTIES OF RAPIDLY SOLIDIFIED ALUMINUM-LITHIUM ALLOYS N.J. KIM, R.L. BYE and S.K. DAS Metals and Ceramics Laboratory, Allied-Signal Incorporation, P.O. Box 1021 R, Morristown, NJ 07960, U.S.A. Considerable improvements have been made during the last decade in the properties of A1-Li alloys for aerospace applications. However, the development of A1-Li alloys for applications that require higher stiffness and lower density than those achieved with ingot casting technology necessitates the use of rapid solidification processing to minimize microsegregation related property deficiencies. The present paper discusses the current status of rapidly solidified A1-Li alloy development at Allied-Signal. The main emphasis of the program has been placed on the achievement of optimum properties by simple heat treatment process, which appears to be essential for the development of high-performance near net-shape A1-Li alloy forgings. INTRODUCTION Aluminum-lithium alloys, because of their low density and high elastic modulus, have received great attention for aerospace applications (see e.g., 1-3). The addition of one wt% lithium (- 3.5 at%) to aluminum decreases the density by - 3% and increases the.elastic modulus by - 6%, hence giving a substantial increase in specific modulus. Their tendency to brittleness, however, has been a major factor which hinders the development of high performance aluminum-lithium alloys. Although there has been much progress in the development of ingot metallurgy (I/M) aluminum-lithium alloys in recent years, the problem of segregation encountered in ingot casting limits the maximum lithium content to wt%, giving weight saving potential to approximately 8%. Increasing the lithium contents to further improve the density and elastic modulus usually results in a reduction in ductility and fracture toughness. Moreover, some of the approaches utilized in ingot cast aluminum-lithium alloys to improve toughness, such as stretching, cannot be utilized for forging applications. In aluminum-lithium alloys prepared by rapid solidification, however, possible weight savings of 10-15% can be achieved and elastic modulus can be increased further by using higher lithium contents (> 3 wt%). In addition, rapidly solidified aluminum-lithium alloys do not require a stretching operation, thereby making them particularly well suited for forging applications. The present paper discusses the structure and properties of the recently developed rapidly solidified aluminum-lithium alloys with high lithium (> 3 wt%) and zirconium (0.5 wt%) contents. The main principles of alloy design and rapid solidification processing utlilized in our development program have been discussed in previous papers (4,5) MATERIALS AND EXPERIMENTAL PROCEDURE The compositions of the alloys used in this investigation are shown in Table 1. These alloys have been designed to achieve densities lower than 2.50 g/cm3 and their densitiies are shown in Table 1. Alloys were rapidly quenched from the melt into continuous ribbons by the jet casting process. Article published online by EDP Sciences and available at

3 JOURNAL DE PHYSIQUE TABLE 1 - ALLOY COMPOSITIONS (wt%) AND DENSITIES (a/cm3) A1 loy L i Cu Mg Zr DENSITY These ribbons were mechanically comminuted to -40 mesh powder and then consolidated into bulk compacts by vacuum hot pressing which were then hot extruded to a final rectangular shape of 63.5 mm x 10.2 mm (18:l extrusion ratio, 6:l aspect ratio). Additional alloy 644 was comminuted to -20 mesh powder to study the possible enhancement of mechanical properties by using coarse powder. The extrusions were solutionized for 2 hours at 550 C. Tensile properties were measured using round specimens with 19.1 mm (0.75 in.) gauge length and 3.2 mm (0.125 in.) gauge diameter at a strain rate of 4x10-4/sec. Fatigue crack growth and fracture toughness testing was conducted using compact tension specimens. The dimensions of the specimens were 50 mm x 48 mm x 5 mm, and a load time width of 40 mm. Impact strength was measured using V-notch (notch radius: 0.001") impact specimens. Thin foils for transmission electron microscopy were prepared by jet polishing with an electrolyte of 33% HNO3 and 61% methanol or with one of 10% perchloric acid and 90% ethanol. TEM observations were made at an operating voltage of 120 KV. 1. Microstructure RESULTS Figure 1 is a typical three dimensional composite optical micrograph of as-solutionized alloy. It shows that the microstructure is composed of unrecrystallized grains which are elongated in the extrusion direction. The volume fraction of coarse undissolved constituent particles is minimal. Also visible in Figure 1 is a fine dispersion of particles which appear to be along the prior powder particle boundaries. TEM studies of as-solutionized alloy show that these particles are mostly oxide and/or carbonate (Figure 2). 1. Optical micrograph tionized 550 C/2 hrs. alloy

4 Fig. 2. TEM micrograph of alloy 644 showing oxide stringers. Solutionized 550 C/2 hrs. TEM micrograph of an as-solutionized alloy is shown in Figure 3. It shows that there is a large volume fraction of fine dispersoids within the grain and coarse dispersoids located at the grain boundaries. These dispersoids are all metastable Llq A13Zr which are coherent with the aluminum matrix (6). Figure 4 shows a typical microstructure of an artificially aged alloy. As observed in many alloy systems, an increase in the degree of aging results in a coarsening of 6' precipitates and the development of precipitate free zones. One of the most important microstructural features observed in this alloy system is that most of the 6' is associated with Al3Zr dispersoids, forming composite precipitates. Such homogeneous distribution of a large volume fraction of composite precipitates is quite desirable for the improvement of ductility and toughness. Another important microstructure feature of these alloys is the absence or sluggish precipitation of T1 or S' phases, as was observed in the earlier studies (4,7). Fig. 3. TEM micrograph of alloy 648 Fig. 4. TEM micrograph of alloy 648 showing a distribution of A13Zr. showing composite ~recipitates. Solutionized 550 C/2 hrs. Aged 180 C/16 hrs.

5 C3-312 JOURNAL DE PHYSIQUE 2. Tensile Properties and Fracture Touahness The room temperature tensile properties of aluminum-lithium alloys for various aging conditions are listed in Table 2, along with the selected fracture toughness data. Properties of alloy 644 made from -20 mesh powder are also included. As can be seen in Table 2, good combinations of strength and ductility were obtained in under-aged alloys. It can also be seen that the difference in the particle size (-20 vs. -40 mesh) does not result in any noticeable change in the tensile properties of alloy 644. A comparison between alloy 643 and alloy 644 shows that higher level of Li increases the tensile strength but does not reduce the tensile ductility. Since the underaged alloys show good tensile properties, fracture toughness data were generated for underaged alloys only. Because of the size limitations of the extruded bars, the measured toughness is KQ rather than KI~ (plane strain fracture toughness). It can be seen that alloys 643 and 644 have good fracture toughness in the underaged condition. Although the increase in Li content from 3 wt% in alloy 644 to 3.4 wt% in alloy 643 did not degrade the elongation, it resulted in a decrease in fracture toughness in the L-T orientation. Some of the decrease in fracture toughness might also be due to the higher strength of alloy 643. The change in the Li content had no effect on the fracture toughness in the T-L orientation. Alloys 643 and 644 have similar T-L fracture toughness values which are less than those in the L-T orientation. Fractographic study of fracture toughness specimens showed that failure occurred mainly along the prior powder part.icle boundaries. This indicates that the oxide particles located along the prior powder particle boundaries (Figure 2) are the dominant factor in controlling the T-L fracture toughness of rapidly solidified aluminum-lithium alloys. The change in the powder particle size, from -40 mesh powder to -20 mesh powder, increased the fracture toughness of alloy 644, presumably because of the reduced prior particle boundary area. Table 2 - MECHANICAL PROPERTIES OF RAPIDLY SOLIDIFIED ALUMINUM-LITHIUM ALLOYS Y.S. U.T.S. El KQ Impact Strength Aging MPa MPa % MP~G ~ / m m ~ L-T T-L L-T T-L C/16h (-40 mesh) 160 C/16h C/16h (-20 mesh) 130 C/16h

6 3. Fatigue Crack Growth Fig. 5. Fatigue crack growth rates of Fig. 6. SEM micrographs showing fatigue rapidly solidified A1-Li alloys compared crack profile: (a) alloy 648 and (b) with 7075-T73. alloy 643. Fatigue crack propagation experiments were conducted for selected alloys in under-aged condition using compact tension specimens. The tests were conducted in a laboratory air environment using an R-ratio of 0.5. Crack closure was not measured and thus was not corrected for. Figure 5 shows fatigue crack growth rate, da/dn, as a function of the stress intensity range, AK. The fatigue crack growth rate of 7075-T73 is also included for comparison. Of the alloys studied (643, 644 and 648). alloy 648 has the fastest crack propagation rate and alloy 643 has the slowest crack propagation rate. In all the alloys, the crack path is deflected and tortuous (Figure 6). The tortuousity of the crack path is most severe in alloy 648 (Figure 6a), whose deformation mode is the most planar of the alloys studied. It can also be seen in Figure 6A that alloy 648 shows extrusions and slip line markings along the crack path. Comparison of fatigue crack growth rates of these aluminum-lithium alloys with that of 7075-T73, without considering closure effects, shows that the fatigue performance of alloys 643 and 644 is quite competitive or better than the I/M 7075 alloy. DISCUSSION As has been shown above, good mechanical properties as well as low densities have been obtained in rapidly solidified aluminum-lithium alloys. The improvements in mechanical properties of rapidly solidified aluminum-lithium alloys are mainly due to the beneficial effect of the Zr addition on the microstructure. As shown in Figure 4, the microstructure of aged rapidly solidified aluminum-lithium alloys is comprised of a large volume fraction of homogeneously distributed composite precipitates. Most of the 6' is associated with composite precipitates. These composite precipitates have been shown to significantly affect the deformation behavior of alloys (7-9). It is well known that in aluminum-lithium alloys which contain a small volume fraction of heterogeneously distributed composite precipitates, 6' is sheared by moving dislocations, resulting in severe planar slip. On the other,hand, alloys which contain a large volume fraction of homogeneously distributed composite precipitates are known to exhibit a wavy and homogeneous mode of deformation (7,8) due to the shear-resistant nature of the composite precipitates. The shear-resistant nature of composite precipitates can be understood

7 C JOURNAL DE PHYSIQUE when one considers that the metastable Al3Zr particle in Al-Zr alloys, although it is of the same crystal structure (LIZ) as 6'. is highly resistant to dislocation shear (10). The effect of composite precipitates on modifying the deformation mode can be readily seen in a weak beam dark field image (Figure 7). Superlattice dislocation pairs, which are characteristic of sheared precipitates, are not observed. Single dislocations are seen to move in a wavy mode and many dislocation loops can be seen. By comparing a 6' superlattice dark field image and a weak beam dark field image of the same area, these dislocation loops were found to surround mostly the 6'/AlgZr interface of composite precipitates (6). This means that dislocations by-pass the AlgZr particles leaving the dislocation loops around them, indicating the high resistance of A13Zr particles to dislocation shear. Although the 6' in composite precipitates could be sheared by dislocations, the degree of stress localization imposed by 6' is minimized by the shear-resistant AlgZr which forms the core of composite precipitates. Fig. 7. Weak beam dark f~eld micrograph of deformed alloy 648. The fatigue characteristics of these alloys show some interesting features. As shown in Figure 5, alloy 648 (AT-3.4Li-0.5Zr) has a much faster fatigure crack growth rate than alloy 643 (Al-3.4Li-1Cu-0.5Mg-0.5Zr) which has similar Li content but with additions of Cu and Mg. In general, the high resistance,of aluminum-lithium alloys to fatigue crack progagation has been attributed to the high elastic modulus and to the extensive crack deflection processes induced by shearing of 6' precipitates (11, 12). Several investigations have been shown that a rough fracture surface gives rise to a component of crack closure which has the added effect of reducing crack growth rates. Although the current results do not consider the closure effect, examination of fatigue crack paths of both alloys indicates that the other factors are also playing important roles in controlling the fatigue crack growth behavior of rapidly solidified aluminum-lithium alloys. Alloy 648 exhibits faster crack growth rates than alloy 643, although Figure 6 shows that the crack path tortuosity or crack deflection is more severe in alloy 648. This indicates that roughness induced closure is not the dominant factor which contols the fatigue crack growth of these alloys. The role of slip reversibility in lowering fatigue crack growth rates cannot account for the crack growth behavior of both alloys, since the slip reversibility would be higher in alloy 648 and the plastic strain accumulation in alloy 648 would be lower for a

8 given number of cycles than in alloy 643. On the other hand, it is well known that the plastic zone size ahead of the crack tip with respect to the relevant microstructural unit plays an important role in controlling the fatigue crack propagation behavior of alloys. Alloys 648 and 643 have the identical microstructure. It has been shown that stress-strain behaviors of the two alloys are quite different, alloy 643 showing a much higher degree of strain hardening than alloy 648 (4). Considering that the cyclic plastic zone size is inversely proportional to the square of flow strength (13), altoy 643 may have much smaller plastic zone size than alloy 648. This implies that, for a given AK, alloy 648 should exhibit faster crack growth rates than alloy 643. Observation of extrusions and slip line markings along the crack path of alloy 648 supports this explanation (Figure 6). SUMMARY The results of the present investigation have demonstrated that good combinations of strength, ductility, toughness and fatigue crack growth resistance can be obtained in wt% Li containing rapidly solidified alloys. A fine distribution of non-shearable Alg(Li,Zr) precipitates resulting from a high Zr c6ntent allows optimum mechanical properties to be achieved by simple thermal treatment without any mechanical working. Long standing problems related to the surface oxide layer along powder particle boundaries have been shown to be partially solved by using a coarse powder particle size. An increase in powder particle size resulted in an improvement in toughness without affecting tensile properties. Work is continuing to further improve the transverse-oriented properties by reducing the powder surface oxide layer. REFERENCES T.H. Sanders, Jr., and E.A. Starke, Jr., eds., Aluminum-Lithium Alloys I, TMS-AIME, Warrendale, PA, T.H. Sanders, Jr., and E.A. Starke, Jr., eds., Aluminum-Lithium Alloys 11, TMS-AIME, Warrendale, PA, C. Baker, P.J. Gregson, S.J. Harris and C.J. Peel, eds., Aluminum-Lithium Alloys 111, The Institute of Metals, London, U.K., N.J. Kim, R.L. Bye, A.M. Brown, S.K. Das and C.M. Adam, p. 975 in Aluminum Alloys and Their Physical and Mechanical Properties, eds., E.A. Starke, J., and T.H. Sanders, Jr., EMAS, West Midlands, U.K., N.J. Kim and S.K. Das, p. 213 in Science and Technology of Rapidly Quenched Alloys, eds., M. Tenhover, L.E. Tanner and W.L. Johnson, Materials Research Society, Pittsburgh, PA, N.J. Kim, R.L. Bye, D.J. Skinner and C.M. Adam, p. 367 in Rapidly Solidified Materials, eds. P. W. Lee and R. S. Carbonara, ASM, N.J. Kim, D.J. Skinner, K. Okazaki and C.M. Adam, p. 78 in ref. 3. N.J. Kim and S.K. Das, Scripta Metall., 20, p (1986). F.W. Gayle, J.B. Vander Sande and O.R. Singleton, p. 767 in Aluminum Alloys and Their Physical and Mechanical Properties, eds., E.A. Starke, Jr., and T.H. Sanders, Jr., EMAS, West Midlands, U.K., N. Ryum, Acta Metall., 17, p. 269 (1969). E.J. Coyne, Jr., T.H. Sanders, Jr. and E.A. Starke, Jr., p. 293 in ref. 1. A.K. Vasudevan and S. Suresh, Metall. Trans., 16A, p. 475 (1985). P.C. Paris, p. 107 in Proc. 10th Sagamore Conference, Syracuse University Press, Syracuse, NY, 1962.