HAL Id: jpa-00226605
https://hal.archives-ouvertes.fr/jpa-00226605
Submitted on 1 Jan 1987
HAL is a multi-disciplinary open access archive for the deposit and dissemination of sci- entific research documents, whether they are pub- lished 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.
MICROSTRUCTURE AND PROPERTIES OF
Al-Li-Cu-Mg-Zr (8090) HEAVY SECTION FORGINGS
R. Lewis, E. Starke, Jr., W. Coons, G. Shiflet, E. Willner, J. Bjeletich, C.
Mills, R. Harrington, D. Petrakis
To cite this version:
R. Lewis, E. Starke, Jr., W. Coons, G. Shiflet, et al.. MICROSTRUCTURE AND PROPERTIES OF Al-Li-Cu-Mg-Zr (8090) HEAVY SECTION FORGINGS. Journal de Physique Colloques, 1987, 48 (C3), pp.C3-643-C3-652. �10.1051/jphyscol:1987374�. �jpa-00226605�
JOURNAL DE PHYSIQUE
Colloque C3, supplement au n09, Tome 48, septembre 1987
MICROSTRUCTURE AND PROPERTIES OF Al-Li-Cu-Mg-Zr (8090) HEAVY SECTION FORGINGS
R.E. LEWIS, E.A. STARKE, Jr.* , W.C. COONS* * ( I ) , G. J. SHIFLET* , E. WILLNER, J.G. BJELETICH, C.H. MILLS, R.M. HARRINGTON and D.N. PETRAKIS
Lockheed Missiles and Space Co., Inc., 3251 Hanover street, Palo Alto, CA 94304 and Sunnyvale, CA 94088, U.S.A.
*university of Virginia, Charlottesville, VA 22901, U.S.A.
f *
San Jose, CA 95113, U.S.A.
ABSTRACT
The microstructure and properties of heavy section forgings of the 8090 A1-Li alloy were investigated including the as-cast, homogenized, forged, and heat-treated conditions. The ingots, 305x965~2000-3600 mm in size, were cast by Alcan. Homogen- ization involved 24 to 48 h soaks at 5450C. Ingots were hand-forged by HDAF Ltd. to sizes up to 356x356~1524 mm, then solution-treated at 530°C for 6 h and water quench-.
ed. The material was aged at various times at temperatures of 150, 170, and 1900C.
Microstructures were examined by TEM, CBED, SEM, AES, and optical microscopy. Mech- anical properties were characterized by tensile, fracture toughness, and stress cor- rosion tests. A strong correlation was observed between grain boundary precipitation of a icosohedral (I) phase and certain mechanical properties. The I-phase has been tentatively identified as AlgCuLig, historically called "TzU. When the I-phase was predominant, the fracture toughness, SCC, and tensile ductility were invariably low,.
Factors identified as promoting the formation of I-phase were cooling rates from the solution treatment temperature slower than -10°C s-1, increased aging time, and in- creased aging temperatures in the range 150 to 1900C. Aging conditions which mini- mized the formation of the grain boundary I-phase and, consequently, improved the mechanical properties were determined. Three distinct constituent phases were found in cast ingots having only slightly different chemistries. Remnants of these con- stituent phases were present in every subsequent stage of thermal processing. One, AlLiSi, was discovered to promote surface pitting and to substantially lower the stress corrosion cracking resistance. This phase's reactivity with seawater appears to promote dissolution of the adjacent matrix. Material heat-treated to suppress the formation of the I-phase, but high in silicon, revealed low SCC resistance. In summary, many factors including composition, casting practice, metal-working, and heat-treatment, determine the mechanical properties of 8090 A1-Li in heavy section forgings, and which range from unacceptable to acceptable for high performance aero- space structures.
INTRODUCTION
Aluminum-lithium (Al-Li) alloys are of high current interest for many important, new civilian and military applications. A substantial weight savings may be achieved as a re ult of the lower density, higher specific strength and stiffness of these new
alloy^.^-^ Lockheed Missiles and Space Co. (LMSC) has used one such alloy, 8090, in all product forms including weldments to fabricate structures. An acceptable balance of properties has been readily achieved in product forms characterized by relatively thin sections produced from stock that was substantially reduced from its original ingot dimensions. However, in forms where large reductions are not performed, e.g., heavy section forgings and thick plate, serious limitations are encountered due to inherently low ductility and fracture toughness. Because of the important need for components made from these heavy section product forms, additional work is needed to elucidate the reasons for these unacceptable engineering properties. An understand- ing of the influencing factors is expected to provide the basis for determining if acceptable behavior can be achieved with existing compositions through an optirniz- ation of the processing practices currently performed.
Many explanations exist for the low fracture toughness and tensile ductility in
(''consultant
Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:1987374
C3-644 JOURNAL DE PHYSIQUE
A1-Li alloys. It has been proposed5-7 that they are the direct consequence of inten- se planar slip resulting from dislocation shear of the ordered 6' (A13Li) precipi- tates. This, in turn, leads to the development of a high stress concentration at grain boundaries and causes low energy intergranular fracture. 0thers8y9 have pro- posed that the presence of grain boundary phases, either produced during the casting process or via solid-state precipitation reactions, can lead to low fracture tough- ness by acting as stress risers and preferential sites for microvoid nucleation and growth. Also, grain boundary phases have been proposed to promote the formation of precipitate free zones (PFZ1s). Being much softer than the neighboring age-hardened matrix they accommodate strain localization in the grain boundary regions.6,10 Fin- ally, some invest igatorsl1-l5 contend that the toughness of these materials is dele- teriously affected by elemental segregation of, variously, Na, K, H and S to grain boundaries.
In the present study the objective was to determine the major factors associat- ed with the low toughness and SCC resistance in heavy section forgings of Al-Li-Cu- Mg-Zr alloy 8090 and to formulate the necessary changes in processing or composit- ional control to improve these properties. This was conducted by studying the micro- structure and measuring the properties at different fabrication stages, and evaluat- ing the effect of selected processing variables on these properties.
EXPERIMENTAL PROCEDURES
MATERIALS - Alcan produced the cast ingots by their proprietary direct chill (DC) casting procedures. The nominal ingot sizes were 305x965~2000-3600 mm. These were stress relieved, then homogenized at 545% for 24 to 48 h, using a heating rate of <20°C/h. The ingots were then cut into billets of appropriate size for High Duty Alloys Forging (HDAF) Ltd. to forge into a variety of sizes of large block. After forging, the blocks were solution-treated at 5300C for 6 h, water quenched, then com- pression stress-relieved 3 to 6 X . The blocks were delivered in this condition to Lockheed Missiles & Space Company (LMSC), who performed aging treatments then machin- ed the material into different size specimens for mechanical and physical property tests. A variety of size forgings were selected for property and microstructural characterization from certain Alcan "melt lots". The melt identification, comp- osition, and description of each of the forgings used in this study are presented in Table I. For convenience, subsequent reference will be made to melt nb. by the last three digits, and to forging block size by Roman numeral. Note from Table I that the forgings vary in cross-sectional area from -10,000 to 127,000 mm2, and in length from 127 to 457 mm. In a number of cases, up to four forgings of the same size were prepared from a "log" or multiple in length of the individual forging blocks.
HEAT TREATMENT - Aging of the forging blocks was performed in a recirculating air furnace at 150°C for up to 96 h and at 170°C and 190°C for up to 24 h. Tensile property data provided by Alcan for 1900C aging provided initial guidance for LMSC to concentrate on evaluating the short transverse (ST) orientation as this orientation exhibits the lowest tensile ductility, toughness, and stress corrosion resistance.
MECHANICAL PROPERTY TESTS - The tensile, fracture toughness (FT), and alternate immersion 35% NaCl aqueous stress corrosion (SCC) tests followed appropriate ASTM specifications. The SCC test involved a 10 min. immersion and a 50 min. air drying cycle. The tensile specimens and sustained load SCC specimens have a minor diameter of 6.3 mm. The compact-tension fracture toughness specimens have a thickness ("B"
dimension) of 25.4 mm. All tests were performed in "laboratory air" at 23 to 28OC.
MICROSCOPY - Optical microscopy involved grinding of bakelite-mounted specimens with Sic through 600 grit and rough polishing on nylon cloth with a slurry of 1 pm diamond and oil. This was followed by an intermediate polishing step on a cotton twill hard cloth and a finish polishing step with microcloth, both steps using a slurry of sub-micron ( < 0.02 vm) cerium oxide "Finish-Pol" provided by the Klarnell Corp., San Jose, California. These procedures were developed by one of the authors (W. C. Coons) and have been published by Leco Corp. Etching was performed by ultra- sonic agitation for 20 to 25 s in a solution of 5 ml HNO3, 2 ml HF, 3 ml lactic acid, 50 m l glycerine, 250 ml H20, diluted 10:l with distilled water. Optical micrographs were taken with a Bausch and Lomb Research Metallograph. Scanning electron micro- scopy (SEM) was performed using a Cambridge Stereoscan -Model 11. Auger electron Spectroscopy (AES) was performed on a Perkin-Elmer Model 595 Scanning Auger electron Microscope. Transmission electron microscopy (TEM) was performed on samples mechan- ically thinned, punched to 3 mm dia. discs, then twin-jet polished in 3:l methyl alcoho1:nitric acid at -18% and 25 VDC. The micrographs and convergent beam elect- ron diffraction (CBED) patterns were obtained using a Philips 400T microscope.
RESULTS
COMPOSITION - All of the forgings selected have almost identical composition, see ÿ able I. One main difference to note is silicon, which varies, in wt %, from a low of 0.03 for melt 15/2, to an intermediate level of 0.08 for melt 19/1, and to a high of 0.11 for melt 14/2. These differences, along with general differences in grain size, may be the major reason for the noticeable differences in selected properties, namely, SCC resistance, as discussed below.
MECHANICAL PROPERTIES - The evaluation of forgings after aging at 190 C for 16 or 24 h confirmed the data previously supplied by Alcan, see Table 11. Of particular concern was the low fracture-toughness (values of 9 to 13 ~ ~ a d m ) , and low 'ensile ductility (average values of 2.5 to 3.6% total elongation), see Table 11. Upon aging forging block I at 170°C, a somewhat higher toughness was achieved at 8 and 16 h ag- ing time, but not after 24 h. Only by using an aging temperature of 1500C and aging time of 96 h was a reasonable combination of tensile strength, ductility, and fract- ure toughness achieved,.see Table 11. It is important to note that about the same values of tensile and fracture toughness properties were achieved in all of the (dif- ferent size and composition) forging blocks. It is also important to note the fract- ure toughness of the as-solution treated and compression stress-relieved forging block I, see Table 11, last data line. The toughness well below the surface of this rather large forging was only 18 ~ ~ a d m , but was much higher, 32 ~ ~ a d m , close to the surface. In any case, subsequent aging is not expected to increase the toughness, but rather decrease the toughness of the unaged material.
STRESS CORROSION PROPERTIES - The SCC behavior of the different forging blocks was found to fall into either "good" or "bad" categories, with respect to the maximum tensile stress below which failure did not occur in 30 days, see Table 111. Forging blocks I and V exhibited rather good behavior, and forging blocks 111 and IV failed in about 5 to 9 days at stress levels as low as 103 MPa, well below the acceptance standard of 158 MPa. Note from Table I that all four forging blocks used in these tests have almost identical compositions, but do differ in silicon content. Forging blocks I and V have intermediate and low silicon, 0.08 and 0.03 wt X , respectively, whereas blocks 111 and IV (both from melt 14/2) have the highest silicon, 0.11 wt %.
As discussed below, these differences in silicon are believed to be significant.
MICROSTRUCTURE - The microstructure of the as-cast and the homogenized condit- ions is shown in Figure 1 for three of the melts prepared for the forgings. Note the coarse constituent phases in the grain boundaries of the as-cast structure. The dark-etching phase found in all three melts was identified by TEM and SAD as T2, or A16CuLi3, and is an an icosohedral phase exhibititing a five-fold symmetry diffract- ion pattern, see Figure 2. Using Auger Electron Spectroscopy (AES) in conjunction with laser ion mass spectroscopy (LIMS) to detect lithium, the light-etching grain boundary phase was identified as AlLiSi. The LIMS work was performed by Alcan per- sonnel and is unpublished. The melting point, and presumably the solvus, of this phase is reported as 5770c.16 This temperature is higher than any practicable homo- genization or solution-treatment temperature for the A1-Li alloys. Consequently, upon homogenizing the 8090 alloy ingot castings at 545OC, there is no appreciable change in the amount or shape of the AlLiSi phase, see Figure 1.
In comparison to the AlLiSi phase, the solvus for the T2 is about 5300C, slightly below the 5450C homogenization treatment used for these forging blocks.
Note in Figures l(b), (d), and (f), however, that some of the T2 phase is still present in the grain boundaries following homogenization, apparently because there was inadequate time for its complete re-solution to occur.
Occasionally, other rather coarse constituent phases were found: one is shown in Figure l(d). AES indicated it to be high in titanium, perhaps undissolved rem- nants of Ti-B master alloy particulate added as a grain refiner.
The microstructure in the solution-treated and compression stress relieved blocks reveals significant differences between the near surface and deeper regions.
As shown in Figure 3, complete recrystallization has occurred at and near the forging surface, while only partial recrystallization occurs at midthickness. Similar part- ial recrystallization was observed at quarter thickness. as well. The dislocation density is substantially higher near the surface than it is at midthickness. See Figure 4. Furthermore, the grain boundaries in the proximity of the free surface are relatively free of any large precipitates, see Figure 5(a). Near the free surface of the forgings, 13' exists right up to the grain boundaries, see Figure 5(b). In the same position in the forgings, Sr also precipitates right up to the grain boundaries (not shown). At typical quarter thickness and the midthickness locations such as
C3-646 JOURNAL DE PHYSIQUE
TABLE I. Composition and Description of 8090 Alloy Heavy Section Forgings4 ALCAN Ident. Forging Composition (wt %) Cast Grain
MeltNo. BlockSize(mm) A1 Li Cu Mg Zr Fe Si Ti Size(mm) 9A014/2 I1 26/x267~127 bal. 2.37 1.26 0.62 0.12 0.04 0.11 0.032 -350
111 127x203~762 IV 102x102~406
9A015/2 I 356x356~457 bal. 2.37 1.24 0.57 0.15 0.05 0.08 0.032 70 9A019/1 IV 102x102~406 Bal. 2.35 1.17 0.64 0.13 0.04 0.03 0.025 210 9A022/2 V 254x254~457 bal. 2.44 1.19 0.61 0.13 0.06 0.04 0.025 90
+ data provided by Alcan
TABLE 11. Mechanical Properties of 8090 Alloy Heavy Section Forgings
Forging Ag!ngi UTS TY S Elong.
Block Temp T me ST L T ST L T ST KQ-
ST ("C) (h) (MPa) (MPa) (MPa) (MPa) (MPa) (MPa) (%) (%) (%) ( ~ ~ a d m )
I 150 24 441 283 9
16 24
I+ 190 16 462 452 438 397 375 352 5.0 3.8 2.5
I1 24
III+ 16 462 476 429 390 400 366 2.5.2.7 3.6 24
IV+ 16 458 443 425 395 388 360 6.2 2.5 2.5 V+ 16 507 465 455 450 385 345 6.2 5.0 3.0 I solution treated and compression stress relieved, 1/2t/surface + Alcan tensile data
that displayed in Figure 5(c), the high angle grain boundaries are populated with relatively large, discontinuous T2 precipitates.
In test specimens of more highly aged material the fracture path in both FT and SCC tests preferentially follows the high angle grain boundaries. An example of the intergranular nature of fracture in the SCC specimens is shown in Figure 6. Aging produces additional precipitation of the T2 phase along most high-angle, and some low angle, boundaries. In some instances, the embrittling T2 phase grows along the grain boundary forming long, snake-like precipitates, see Figure 7. In those same regions of the forging blocks, a 6' precipitate free zone (PFZ) exists up to 2-3 pm in width.
Even though Sf precipitates are found in this PFZ, the amount is less than outside of the PFZ. Consequently, this PFZ constitutes a low strength region on both sides of the high angle grain boundaries. The AlLiSi particles formed upon original solidifi- cation of the ingot are retained almost intact in the forged and fully heat-treated microstructure. These particles serve as highly active sites for matrix dissolution and spawn pits on all exposed free surfaces in the SCC test specimens. For example, see the surface region in Figure 8 of a specimen prepared from a forging melt con- taining 0.11 wt % Si. In those forgings prepared from melts containing 0.08 wt % Si or less, the SCC surface pitting either does not occur or is quite minor. See, for example, the surface appearance of the SCC tensile specimen shown on the ieft side in Figure 8(a).
TABLE 111. Alternate Immersion, Short Transverse SCC Test Results+
Aging Temp. Time
("C) (h) 150 60
Spec. Type C-Ring Tensile
Tensile Stress (MPa)
69 103 138 172 207 241 276 310 69 103 138 172 207 241 276 310 69
Forging Block
I I11 IVX v
p/f(t) p/f(t) p/f(t) p/f(t) 2/0
x 215 o/z(i7;31 j
LMSC Acceptance Standard 158
+ p/f(t) = no. that passed >30 days/no. that failed (days to failure)
* block from melt 14/2, see Table I DISCUSSION
It is apparent from the above, that DC castings used to fabricate forgings of heavy section have some intrinsic disadvantages. During solidification of 8090 castings containing >0.08 wt X Si,large amounts of T2 (A16CuLi3) and AlLiSi are both formed as a result of interdendritic solute enrichment (Figure 1). Because of the low toughness associated with the presence of these constituent phases, the ingot castings have a high propensity to crack due to the tensile stresses that develop during cooling followip$ solidification. The T2 phase is also known to form in the 2090 (Al-Li-Cu) alloy. Additionally, it has been identified as being of the same type found in A1-Mn; and can form as a solid-state reaction product either durin slow cooling from solution-treatment temperatures or during artificial aging.17, lf In additidn to the present study, the convergent beam electron diffraction (CBED) method has been used to obtain evidence that this phase is icosohedral (I-phase) and has the same point group, m35, as that determined for the I-phase in the rapidly solidified A1-Mn and Al-Mn-Si alloys .I9
The major factor which tends to limit the toughness, SCC resistance, and tens- ile ductility in 8090 heavy section forgings appears to be the propensity for precip- itation of the T2 phase along high angle boundaries upon (1) slow cooling from solut- ion treatment temperatures, agd (2) aging at temperatures of 150°C and higher. This precipitation is difficult to inhibit. Cooling rate studies performed as a support- ing experiment in the present study involved cold water quenching a Jominy end quench bar prepared from one of the subject heavy section 8090 forgings. In this experi- ment, thermocouples were embedded in a companion test specimen, and the cooling rates measured autographically. After the specimen was quenched, a longitudinal section was made and the microstructures studied metallographically. No precipitation of the T2 phase was detected until a distance of 19 mm from the quenched end. The cooling rate at this location was -lO°C s-1. These results were then used to successfully predict that a 25 mm-thick plate, edge quenched in cold water so both surfaces were simultaneously quenched, would not form T2, but a slightly thicker section would.
C3-648 J O U R N A L D E PHYSIQUE
( a ) AS-CAST HOMOGENI
a@
ZED
LC'
elt 1 5 / 2
, Melt 19/1, 0.03 wt % Si
Fig. 1 - Optical micrographs of A1-2.3Li-1.2Cu-0.7Mg-0.13Zr (8090) in the as-cast and homogenized conditions. Arrows in (a), (b), and (c) indicate AlLiSi constituent phase. Arrow in (d) indicates Ti-rich phase, probably from Ti-B master alloy added for grain refinement
Fig. 2 - TEMs of 8090 alloy ingot, 305x965~2000 mm, melt phase (a) at grain boundary and ( b ) triple point. Note fj
.Px
# - - A
15/2 as-cast condition. I- ve-fold symmetry (CBED)
Fig. 3 - Optical micrographs of 8090 alloy melt 15/2, forging I, 356x356~427 mm, solution treated and compression stress relieved: (a) surface and (b) mid-thickness
Fig. 4 - TEMs of 8090 alloy forging No. I, solution treated and compression stress relieved. Dislocation structure (a) near surface and (b) at mid-thickness
C3-650 JOURNAL DE PHYSIQUE
Fi'g. 5 - TEMs of 8090 alloy, as Fig. 4. Near forging surface (a) grain boundary is relatively free of precipitates and (b) S f existing up to the grain boundaries. ( c ) I-phase commonly present at mid-thickness (CBED- inset)
Fig. 6 - Optical micrograph of stress corrosion specimen from fully aged forging.
Note intergranular nature of cracking. Tensile axis is vertical
This suggests that quenching of forgings having thicker sections will invariably form some amount or a lot of the T2 phase along high angle grain boundaries in all regions of the forging, except a relatively thin layer close to the free surfaces which are quenched. Furthermore, if one can suppress precipitation of T2 on quenching a suit- ably thin section, provided rough machining of the component accomplishes this aim, then aging at temperatures above 150°C still results in the rather profuse precip- itation of T2 along all of the high angle grain boundaries. This precipitation increases with both aging temperature and time, with a concomitant loss in toughness, tensile ductility and SCC,resistance. '
The issue of silicon content and its role in forming AlLiSi is not so clear.
In the present study, an unacceptably low SCC resistance was exhibited by all of the forging blocks produced from the melt .containing 0.11 wt % silicon, but acceptable sCC resistance was exhibited by all the other forging blocks which happen to contain 0.08 or less wt X silicon. The deleterious role of AlLiSi in promoting surface pitting in a salt water environment is rather clear. This pitting can accelerate the incubation period for crack initiation, even though pitting is generally independent of the applied tensile stress. Once initiated, surface cracks experiencing tensile stress proceed to grow in from the free surfaces, resulting eventually in stress cor- rosion failure. This SCC growth rate is expected to be influenced by the grain size of the material and may be influenced by many other factors as well. Unfortunately, we did not have the opportunity to properly resolve the question of grain size effect in the present study. An appropriate experiment would be to compare the SCC resist- ance of specimens from forging blocks both having relatively coarse as-cast grain size, but differing in silicon content. For example, the SCC resistance of specimens from forging block IV prepared from melt 19/1 (low, 0.04 wt X silicon) can be meaningfully compared with the SCC test results of specimens from forging block IV from melt 1412 ("high", 0.11 wt X silicon), shown in Table 111.
CONCLUSIONS
Heavy section forgings of the 8090 alloy composition are currently prepared from large ingot castings. In the as-solidified condition, these castings have a thick, 1 to 3 pm, layer-gf the T2 (A16CuLi3) phase, icosohedral in nature and having the same point group, m35, as previously observed in rapidly solidified A1-Mn alloys. This constituent phase can be taken into solution by heat treatments exceeding -530°C, but re-precipitates preferentially along high-angle grain boundaries upon slow cooling or upon subsequent aging at temperatures above 1500C. Such grain-boundary precipitation provides a low-energy path for crack growth. Thermal treatments which reduce the amount of this grain-boundary precipitate substantially increase the fracture toughness of the forging material.
For compositions containing >0.08 wt. X Si, the equilibrium AlLiSi phase forms upon solidification and is retained through all subsequent thermal mechanical treat- ments. This phase serves as an active nucleation-site for surface pitting upon expo- sure to salt water, and greatly diminishes the material's resistance to SCC. By specifying a maximum allowable silicon content of 0.08 wt. X , this problem should be eliminated or reduced t,o an acceptable level.
Heavy section forgings have an inhomogeneous distribution of recrystallized grain structure, dislocation density, and amount of embrittling I-phase. These microstructural features are the result of varying amounts of plastic deformation and cooling-rate differences. Microstructural variability is intrinsic to heavy section forgings, and nominally has a deleterious effect on the mechanical properties. An improvement in toughness and SCC resistance can be expected by: (a) reducing the size of the forging, (b) increasing the uniformity of deformation during forging and com- pression stress relief, (c) increasing cooling-rate by minimizing the section thick- ness being quenched, and (d) aging at 150°C or less.
REFERENCES
1. Lewis, R. E., et al. (1986), Development of Advanced Aluminun~ Alloys from Rapidly Solidified Powders for Aerospace Structural Applications, Final Report No. AFWAL- TR-86-4108, Contract F33615-78-C-5203, Air Force Wright Aeronautical Laboratories, Wright-Patterson Air Force Base, Ohio.
2. Ekvall, J. C., Rhodes, J. E., and Wald, G. G. (1982), Design of Fatigue and Fracture Resistant Structures, STP 761, Amer. Soc. for Testing and Materials, Philadelphia, Pennsylvania, 328.
C3-652 JOURNAL DE PHYSIQUE
3. Lewis, R. E., Palmer, I. G., Ekvall, J. C., Sakata, I. F., and Quist, W. E.
(1983), Rapid Solidification Processing: Principles and Technologies 111, National Bureau of Standards, Gaithersburg, Maryland, 615.
4. Quist, W. E. and Lewis, R. E. (1986), Rapidly Solidified Powder Aluminum Alloys, STP 890, Amer. Soc. for Testing and Mater., Philadelphia, Pennsylvania, 7.
5. Starke, Jr., E. A., Sanders, Jr., T. H. and Palmer, I. G. (1981), J. Metals, 33
(81, 24-
6. Sanders, Jr., T. H. and Starke, Jr., E. A. (1982), Acta Metall., 30, 927.
7. Noble, B., Harris, S. J. and Dinsdale, K. (1982), Met. Sci., 16, R5,
8. Starke, Jr, E. A . (1983), Strength of Metals and Alloys, ICSMAVI, 3, Pergamon
Press, 1025.
9. Vasudevan, A. K., Ludwiczak, E. A., Baumann, S. F., Howell, P. R., Doherty, R. D.
and Kersker, M. M. (1986), Met. Sci., 2, 1205.
10. Sanders, T. H., Ludwiczak, E. A., and Sawtell, R. A. (1980), Mater. Sci. Eng., 42
247.
11. Sanders, T. El. (June 1979), Factors Influencing Fracture Toughness and Other Properties of Aluminum-Lithium Alloys; Contract No. N62269-764-0271, Final Report, Naval Air Development Center, Warminster, Pennsylvania.
12. Misra, R. D. K., and Balasubramanian, T. V. (1985), Scripta Metall., 19, 1177.
13. Miller, W. S., Thomas, M. P., Lloyd, D. J. and Creber, D. (1986),3ater. Sci.
Tech., 2, 1210.
14. Vasudevan, A. K., Miller, A. C. and Kersker, M. M. (1984), Aluminum-Lithium Alloys 11, TMS-AIME, Warrendale, Pennsylvania, 181.
15. Webster, D. (1986), Aluminum Lithium Alloys 111, The Inst. of Metals, London, 602.
16. Mondolfo, L. (1976), Aluminum Alloys, Butterworths, London, 558.
17. Cassada, W. A., Shiflet, G. J. and Starke, Jr., E. A. (1986), Scripta Met., g ,
751.
18. Cassada, W. A., Shiflet, G. J. and Poon, S.J. (1986), Phys. Rev. Lett., 2 , 2276.
19. Bendersky, L. and Kaufman, M. J. (1986), Philos. Mag., F , L75.
Pig. 7 (left) - TEM of fully heat treated forging. Note elongated, "snake-like"
growth of T2 phase along the high angle grain boundary
Fig. 8 (right) - Stress corrosion test specimens: ( a ) tensile bars after exposure. On left is specimen from melt 15/2, "intermediate" 0.08 wt % silicon level. On right is specimen from forging 111, melt 14/2, llhighu 0.11 wt X silicon level. (b) Pit forming immediately adjacent to AlLiSi particle at free surface of SCC tensile bar
