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The effect of germanium on the precipitation and deformation behavior of Al2Li alloys
Arm metal/. Vol. 34, No. 3, pp. 367-378, 1986 Printed in Great Britain. All rights reserved
Copyright
OOOI-6160/86 $3.00+0.00 Q 1986 Pergamon Press Ltd
THE EFFECT OF GERMANIUM ON THE PRECIPITATION AND DEFORMATION BEHAVIOR OF Al-2Li ALLOYS W. A. CASSADA, G. J. SHIFLET and E. A. STARKE JR Department of Materials Science, University of Virginia, Charlottesville, VA 22901, U.S.A. (Received 26 April 1985; in revised form 26 June 1985)
Abstract-The effect of a small addition of germanium on the precipitation and deformation behavior of an Al-Li alloy was investigated. Alloys of nominal weight composition of AI-2Li and Al-2Li-0.2Ge were solution heat-treated, quenched, and aged for various times at 473 K. The microstructures and deformation behavior were characterized and the results for the two alloys were compared. The solubility of lithium was increased when germanium was in solid solution; however, lithium decreased the solubility of germanium at 473 K resulting in the formation of small germanium precipitates homogeneously distributed throughout the matrix. These precipitates had a very positive effect on the deformation behavior and ductility of the alloy. R&m&Nous avons ttudit l’influence de petites additions de germanium sur la precipitation et la d&formation d’un alliage Al-Li. Nous avons trait6 thermiquement, trempe et recuit pour diverses durkes B 753 K des alliages de compositions nominales Al-2Li et Al-2Li_0,2Ge en poids. Nous avons caract&& les microstructures et la diformation et nous avons cornpark les rtsultats pour les deux alliages. La solubilitC du lithium Ctait augment&e lorsque le germanium ttait en solution solide; cependant, le lithium diminuait la solubilit6 du germanium B 473 K, ce qui conduisait B la formation de petits pr&cipitts de germanium rkpartis de man&e homogene dans la matrice. Ces pr&cipitCs avaient un effet t&s positif sur la d&formation et la ductilitt de l’alliage. ZusammenfassungAer Einflul3 kleiner Zugaben von Germanium auf das Ausscheidungsverhalten in und das Verfestigungsverhalten von Al-Li-Legierungen wurde untersucht. Legierungen mit der nominalen Gewichtszusammensetzung AI-2Li und Al-2Li-0,2Ge homogenisiert, abgeschreckt und fiir verschiedene Zeiten bei 473 K ausgelagert. Mikrostruktur und Verformungsverhalten wurden bestimmt; die Ergebnisse von beiden Legierungen wurden verglichen. Lag Germanium in Lijsung vor, dann war die Liislichkeit des Lithium erhijht. Andererseits senkte Lithium die Liislichkeit des Germaniums bei 473 K, wodurch sich kleine Germaniumausscheidungen in homogener Verteilung in der Matrix bildeten. Diese Ausscheidungen beeinflussten das Verformungsverhalten und die Duktilitiit der Legierung sehr giinstig.
INTRODUCTION Lithium is the lightest metallic element and has a significant effect on density when alloyed with aluminium, producing a 3% decrease for every weight percent added up to the limit of solid solubility [l]. In addition, a concomitant increase in elastic modulus has been reported for both binary and complex alloy systems [2]. Unfortunately, these beneficial effects are usually accomplished by a significant decrease in ductility and toughness making the simple alloys unsuitable for many structural applications. The low ductility has been attributed to strain localization associated with the ordered Ll,, Al, Li(6’) that forms during age hardening [3]. The similarity in structure and lattice parameters of the 6’ and the fee matrix results in a small lattice misfit (- 0.08 f 0.02%) [4] and spherical precipitates which are easily sheared by moving dislocations. Axon and Hume-Rothery found that the lattice of aluminium is contracted by the solution of lithium and expanded by the solution of germanium [5]. Since 367
these elements have opposite effects on the lattice parameter in solid solution, Kujore and Starke [6] suggested that when added in combination they may cluster at the solutionizing temperature in order to minimize lattice distortion and germanium may subsequently effect the shearability of 6’ by modifying S’/u misfit or the nature of the s’/a interface. Similar studies conducted by Gayle [17] and Baumann and Williams [8] investigated the effect of alloying additions of Ag, SC, Cu, Zn, Ga, Mg, Mn, Cr, Zr and Si on the S’ju interface. Of these, only the alloys containing Cu, Ag and Zn had any appreciable effect on either the misfit or segregation of solute between the 6’ and the matrix. Moreover, the observed effect of the ternary additives was not large enough to change the deformation behavior of the alloys as the 6’ remains fully coherent with the matrix and shearable by dislocations. Germanium additions to Al-Li alloys, however, were shown by Kujore and Starke to have a beneficial effect on ductility by decreasing strain localization during tensile tests [6]. In the present investigation carefully controlled iso-
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CASSADA ef al.: DEFORMATION BEHAVIOR OF Al-2Li Table I. Alloy composition Alloy Bl Tl T2
Nominal composition
Actual %Li at. % (wt%)
Actual %Ge at. % (wt%)
AI-2Li AI-2.0Li-0.2 Ge Al-2.4Li-0.2 Ge
8.1 I (2.22) 6.94 (1.88) 9.10 (2.51)
O.lO(O.2) 0.10 (0.2)
thermal heat treatments and detailed examination by transmission electron microscopy (TEM) were used to establish the reasons for the beneficial effect that germanium has on the ductility of Al-Li alloys. EXPERIMENTAL The chemical compositions of the three alloys used in this study are listed in Table 1. Alloys Bl and Tl were prepared by the Reynolds Metals Metallurgical Laboratory where 10 kg heats were induction melted in a graphite crucible and cast into a 6.4cm ID copper mold. All melting and casting were done under a partial pressure of high purity argon. The as-cast ingots were machined in order to remove several millimeters of surface, homogenized at 803 K for 24 h, and plastically deformed at 723 K to provide a 10: 1 area reduction ratio. Alloy T2 was prepared at the Naval Surface Weapons Center. A 5 kg heat was resistance melted in a graphite crucible and cast into a tapered copper mold. Melting and casting were performed under high purity argon at atmospheric pressure. The as-cast ingot was machined and homogenized as before, and plastically deformed at 723 K to a 10: 1 area reduction ratio. All samples of alloys used in these experiments were thermally processed under identical conditions. Samples were solutionized in a high temperature salt-bath at 823 K for 30 min, quenched in water and transferred to the ageing bath. The presence of germanium in the ternary alloys did not have an effect on the grain size. All three alloys exhibited an equiaxed grain structure after solution heat treating with an average grain size ranging from 340 to 360 pm. Two types of ageing treatments were used; either a single ageing treatment for various times at 473 K or a two step ageing treatment which consisted of 12 h at 338 K followed by ageing for various times at 473 K. The single ageing treatment was used for all but one of the experiments. Knoop hardness measurements were made on samples of the aged alloys using a micro-hardness tester with a 1 kg load. Samples used in hardness testing were carefully prepared to remove at least 5 pm of surface by mechanically polishing the aged material through 0.1 pm MgO. TEM foils were prepared by electropolishing 3 mm diameter discs punched from selected microstructures. The discs were electropolished in a twin jet polishing apparatus at 13 V d.c. using a 2: 1 methanol:nitric acid electrolyte cooled to 249 K. Germanium particles were extracted from the ternary alloy samples aged 24 and 110 h at 473 K, using
carbon extraction techniques [9]. Polished specimens were etched for 5 min in a SN NaOH solution at room temperature and approximately 120 nm of carbon was deposited to the etched surface by vapor deposition. The carbon film was subsequently removed from the aluminium surface by a second etch in the NaOH solution and washed several times in distilled water before transferring to a 3 mm copper grid for TEM examination. Tensile testing of the three alloys was performed on a closed loop servohydraulic MTS machine in stroke control at a strain rate of 10-3m s-’ using a 10 mm extensometer mounted on the specimen to measure strain. Selected microstructures were stretched 2% in order to investigate particle-dislocation interactions with TEM. All TEM examinations were conducted using a Philips EM 400T. The fracture surfaces of tensile samples were examined with a JEOL JSM-35 scanning electron microscope (SEM). EXPERIMENTAL RESULTS Isothermal ageing behavior
The age-hardening response for alloys Bl, Tl, and T2 is shown in Fig. 1, where the Knoop Hardness Number (KHN) is plotted as a function of ageing time at 473 K. Comparison of the peak hardness for each alloy demonstrates a direct relationship between the lithium concentration and hardness. Alloy Tl exhibits little hardening during the first 3-5 h, while alloys Bl and T2 show a steady logarithmic increase to peak hardness. After about 5 h alloy Tl begins to
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AGEING
Fig. 1. Variation
100.0
10.0
1.0
TIME
WRS)
in Knoop hardness time at 473 K.
number
with ageing
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Fig. 2. Dark-field transmission electron micrographs of AL, Li precipitate particles in the three alloys after ageing 3, 9 and 12 h (by row) at 473 K. (A)-(C) B1; (D)-(F) Tl; (G)-(I) T2.
harden relatively quickly and reaches peak age after approximately 12 h. Alloy T2 also reaches peak age after about 12 h, while alloy Bl reaches its peak hardness at a somewhat earlier time (about 9 h). The binary alloy Bl overages faster than either of the ternary alloys. Precipitation and coarsening of 6’ Figure 2 is a series of 6’ centered dark-field (CDF) TEM images of the three alloys illustrating the relative 6’ size for each after 3, 9, and 12 h at 473 K. Alloys Bl and T2 show a steady increase in particle size and a steady decrease in number with increasing ageing time. Distinct superlattice spots were observed in the electron diffraction patterns of both after quenching from the solutionizing temperature, and 6’ particles of approximately 2 nm diameter were observed after 30 min ageing at 473 K. The density of particles after short ageing times (OS-l.0 h) was so great that reliable estimates of particle densities were difficult to determine.
The precipitation behavior exhibited by alloy Tl [Fig. 2(D,E,F)] during the first 3-5 h of ageing is very different from the other two alloys. Close examination of 6’ particles during this interval revealed that all of these particles are associated with matrix dislocations [Fig. 2(D)]. The density of 6’ particles increases with ageing time after 3 h. This increase is reflected in the hardness curve as well as in the TEM image [Fig. 2(E)]. Copious 6’ precipitation occurs at some time after the first 3-5 h, and normal coarsening is observed thereafter [Fig. 2(F)]. The micrographs in Fig. 2 also indicate a relationship between the 6’ particle size after a given ageing time and the lithium content of each alloy. The largest particles are seen in alloy Tl which has the lowest percent lithium and, therefore, the lowest nucleation rate. The smallest particles are seen in alloy T2 which has the greatest density of particle and the largest concentration of lithium. The S’ volume fraction determined from the 6’ dark-field images and corrected using the treatment
370
CASSADA
Fig. 3. Variation
et al.:
DEFORMATION
in volume
fraction
of Hilliard [lo], is shown as a function of ageing time in Fig. 3. It is obvious from this plot that nucleation of 6’ in alloys Bl and T2 is over very quickly since the volume fraction of 6 ’ is essentially constant after ageing for 3 h. The differences in precipitation behavior observed in the TEM micrographs (Fig. 2) of alloy Tl are also reflected in the volume fraction ageing curve for this alloy. There is little or no change in volume fraction of 6’ through the first 3 h of 12 h (peakageing. From 3 h to approximately hardness) the volume fraction of 6’ shows a sigmoidal increase similar to that observed for the change in hardness as a function of ageing time (Fig. 1). After 12 h of ageing the volume fraction reaches a constant value. Precipitation and coarsening of germanium
After about 3 h of ageing at 473 K very small (about 2 nm) particles of germanium are observed in the matrix of alloys Tl and T2. An example of these particles is seen in Fig. 4 which is a TEM bright-field (BF) image of alloy Tl aged 3 h. As ageing time is increased the germanium particles coarsen as rods approximately 15-20 nm in length, with an aspect ratio of about 15 [Fig. 5(A) and 5(B)]. These develop into faceted particles [Fig. 5(C) and 5(D)] as ageing progresses. The density of germanium particles in the thin foils was insufficient for obtaining select area electron diffraction patterns. Identification was accomplished by an extraction process in which the germanium particles were removed from the matrix on a carbon replica which was then analyzed by convergent beam electron diffraction in the TEM. Figure 6 shows such a germanium particle extracted from alloy Tl aged 110 h along with its electron diffraction pattern identifying the diamond cubic lattice of germanium (011 zone axis). The extra spots occurring in this pattern are due to twinned areas in the particle which are visible in the micrograph. The normally faceted cor-
BEHAVIOR
of Al,Li
with ageing
OF Al-2Li
time at 473 K.
ners of the germanium particles are rounded by the solutions used to extract it from the matrix. The orientation of the germanium particles in the aluminium matrix can be deduced by systematically producing a series of images from aluminium matrix and germanium precipitate diffracted spots of approximately the same g vector. The occurrence or absence of DF images of the germanium particles along with the orientation relationship of the three forms of germanium particles determined by Koster [l 11, were used to determine the orientation relationship; (11 I)/[0221 Ge parallel to (001)/[020] Al. To determine the effect of incubation time on the precipitation of germanium particles, samples of alloys Tl and T2 were solution heat-treated at 823 K and cold water quenched as before. Half of the
Fig. 4. Transmission electron micrograph nium particles in alloy Tl after ageing
showing germa 3 h at 473 K.
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Fig. 5. Transmission electron micrographs showing germanium particles in alloy Tl. (A) Bright-field image of rod-shaped germanium particles in sample aged 9 h at 473 K; (B) dark-field image of (A); (C) bright-field image of faceted germanium particles in sample aged 110 h at 473 K; (D) dark-field of (C).
samples were immediately transferred to an ageing bath at 473 K, while the other half was pre-aged at 338 K for 12 h before ageing at 473 K. When samples of each were compared for germanium particle density in the TEM no difference could be seen in the number of germanium particles. Preageing, therefore, had no observable effect on the density of germanium particles in any of the alloys investigated.
Tensile properties The tensile properties of the three alloys are given 7(A) and (B) where yield strength and elongation to failure are plotted as a function of ageing time at 473 K. The yield strength of Bl and T2 increase very rapidly, following the rapid increase observed in 6’ volume fractions of the alloys (Fig. 3). Alloy Tl exhibits a slower rate of increase in yield strength with ageing time, also corresponding to
in Fig.
changes in 6’ volume fraction. Alloys Bl and T2 reach similar values of yield strength after 24 h at the ageing temperature (420 and 415 MPa respectively). Alloy Tl, however, reaches a peak yield strength which is considerably less than the other two alloys due to its significantly lower lithium content (about 325 MPa after 24 h at 473 K). Comparing elongation to failure for alloys Bl (8.1 at.% lithium) and Tl (6.9 at.% lithium) demonstrate a decrease in ductility with increasing lithium content [Fig. 7(B)]. After ageing 24 h at 473 K the elongation of Tl is about 9.5% while that of Bl is only about 3.5%. Comparison of elongation to failure between alloys Bl and T2, however, show the opposite trend. When both of these alloys are aged to similar yield strength (24 h at 473 K) the elongation of T2 is about 7% or approximately twice as great as Bl. The lithium content of T2 (9.1 at.%) is about 1 at.% greater than that of Bl.
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Fig. 6. Transmission electron micrograph of a single germanium particle extracted onto a carbon replica film from a specimen of alloy Tl aged 110 h at 473 K. The corresponding convergent-beam electron diffraction pattern [Oil] is shown in the upper right.
A comparison of the fracture surfaces of Bl and T2 aged to a similar yield strength (24 h at 473 K) reveals failure by different mechanisms [Fig. 8(A) and (B)]. Fracture of the binary alloy was predominately intergranular. Slip offsets (arrows) on the grain boundary walls suggest the presence of coarse slip bands within the matrix. The ternary alloy, however, exhibits a
Fig. 7. Variation in the tensile properties of the three alloys with ageing time at 473 K. (A) Yield stress plotted as a function of the ageing time; (B) elongation corresponding with (A) plotted as a function of the ageing time.
Fig. 8. Scanning electron micrographs of fracture surfaces of tensile samples aged 24 h at 473 K; (A) Bl; (B) T2. The yield strength of both alloys was approximately 410 MPa. Arrows in (A) indicate slip band offsets at the grain boundary walls of alloy Bl.
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Fig. 9. Deformation structure near the fracture surface in alloy Bl aged 24 h at 473 K. (A) Strain localization in vicinity of grain boundary (top); (B) dislocation structure in slip bands; (C) sheared A& Li precipitates within slip band. high degree of ductile transgranular failure. This type of failure is consistent with the greater elongation observed for T2. The deformation structures of the tensile specimens are shown in the TEM’s of Figs 9 and 10. At low magnification both Bl and T2 show intense slip bands within the matrix. At higher magnification, however, the ternary alloy reveals a more uniform distribution of dislocations within the matrix between the slip bands. Dark-field images of 6’ in each alloy demonstrate that the binary alloy is subjected to a greater degree of strain localization as sheared 6’ is observed along very long slip lines [Fig. 9(C)]. This type of straight slip line is not observed in the ternary alloy due to a more dispersed dislocation distribution [Fig. 10(C)]. Effect of 2% strain The
ternary
alloy
Tl
aged
48 h at 473 K and
stretched 2% allows a more detailed TEM examination of the deformation behavior of the ternary alloys due to its smaller number fraction of 6’ particles. Figures 1 l(A) and (B) are a bright-field and weak-beam dark-field pair which show the same uniform dislocation distribution seen in the T2 tensile specimen [Fig. 10(B)]. Since the strain is much less, slip bands are not observed in the matrix of this alloy as they are in the T2 tensile specimen [Fig. 10(A) and (B)]. Dislocations are either cross-slipping around 6’ or around germanium particles which lie on common slip planes, or the dislocations are pinned at one of the particle-matrix interphase boundaries. The 6’ dark-field images observed in the T2 tensile specimen [Fig. 10(C)] showed dark wavey lines through the 6’ at or near the locations of dislocations. Furthermore, the S’ particles are obviously distorted from their normally spherical geometry suggesting that these particles are being sheared by the moving dis-
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Fig. 10. Deformation structure near the fracture surface of alloy T2 aged 24 h at 473 K. (A) Deformation in vicinity of grain boundary; (B) dark-field image of dislocation structure; (C) dark-field image of Al, Li precipitates in same region as (B).
locations. This observation tends to support the germanium phase as the agent dispersing the slip in the ternary alloys. The germanium particles are known to be partially incoherent [l l] which is the most important factor in determining the shearability of a precipitate phase. Weak-beam dark-field images [21-231 of germanium particles and dislocations are shown in Fig. 11(D) for the bright-field image in Fig. 1l(C). A region is selected in which a large number of germanium particles are visible with relatively few 6’ particles in order to eliminate any possible influence from 6’ particle-dislocation interaction. The arrows in Fig. 11(D) indicate points where the strain fields of the dislocations are extinguished by pinning at the germanium particles. It is clear from these micrographs that pinning of dislocations at the germanium particles is responsible for homogenizing slip in the ternary alloys.
DISCUSSION OF RESULTS Precipitation
The sequence of precipitation in the ternary alloys and the delay of extensive precipitation of 6’ in alloy Tl can be understood by first considering the aluminium rich side of the Al-Li binary system (Fig. 12) [12]. The relative positions of the alloys Bl, Tl, and T2 on this diagram are indicated at 473 K. From inspection of Fig. 12 it is clear that Bl and T2 lie well under the 6’ solvus boundary where copious precipitation of 6’ is expected to occur (Fig. 3). Alloy Tl on the other hand, lies very close the 6’ solvus. If the ternary addition of germanium had the effect of lowering this solvus a few degrees (S-10 K), the resultant supercooling would be even less. Several ternary additions have been observed to raise the 6’ solvus [8], but to our knowledge, no additions have been observed to lower it. Among the
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Fig. 1I, Deformation structure in the matrix of alloy Tl aged 48 h at 473 K and stretched 2%. (A) Bright-field image; (B) dark-field image of (A) showing homogeneous distribution of dislocations in matrix along with small germanium particles; (C) bright-field image showing spherical Al,Li precipitates and germanium particles; (D) corresponding dark-field image of (C) showing dislocations pinned at the germanium particles. elements found to decrease the solid solubility of lithium in aluminium are: Ag, Zn, Mg, and Mn. Of these Zn and Mn decrease the lattice parameter of aluminium, Mg increases it while Ag has no effect on the lattice parameter of the matrix. Most of these alloying additions will form transitional phases of their own and all of them may also form phases with lithium. These factors probably have the greatest influence on the solubility of lithium in the matrix. Ternary additions of germanium are very different from those additions mentioned above. Covalent solutes in metallic solvents, such as germanium in aluminium, generally do not form transitional phases [13]. An example of this is found in the Al-Si system where no transitional phases exist. Germanium and silicon possess identical valence shells and similar chemistry, so the similarity in behavior of binary alloys of each is not surprising.
The contracting effect of lithium on the aluminium lattice [5, 141 may in fact be the primary reason for the decreased solubility of germanium in aluminium. Conversely, the expanding effect of the germanium on the matrix lattice parameter could increase the solid solubility of the lithium. Within the compositional ranges of the alloys used in this study, germanium expands the aluminium lattice by about 1.08 x 10m4nm/at.% germanium while lithium contracts the aluminium lattice approximately 4.65 x lO~‘nm/at.% lithium [5]. Due to the opposing effects of the solutes on aluminum’s lattice parameter, germanium and lithium atoms may be attracted to one another in order to relieve lattice strain [6]. However, the covalent nature of germanium and the metallic lithium would result in little chemical affinity between the two. The normally high vacancy-binding energy of the germa-
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be limited by the removal of germanium from the solid solution as precipitate particles, thus explaining the relatively slow rate of increase in 6’ volume fraction in Tl observed in Fig. 3. Mechanical behavior
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!
0
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Fig. 12. Aluminum-rich end of the binary Al-Li system after Ceresara et al. [12].
nium in aluminium [ 111is most likely increased as the tra $ ping of a vacancy by a germanium atom in this matrix would release an additional amount of strain energy. The amount of this increase in vacancy binding energy would be related to the decrease in lattice parameter. As the density of Ge-vacancy couples increases there is a concomitant increase in germanium nucleation sites [15]. Because of the decrease in germanium solubility created by the lithium in solid solution, sufficient supersaturation exists for the nucleation and growth of germanium particles. This appears to be happening in alloy T2, since both 6’ and germanium particles have sufficient undercooling for nucleation at early ageing times. The germanium particles are precipitating after 6 ‘, and at about the same time as germanium particles precipitate in alloy Tl. In alloy Tl, however, the composition is at or above the 6’ solvus at 473 K leading to little or no undercooling for precipitation. Figure 3 illustrates that while the volume fraction of 6’ in alloy Tl is nearly zero after three hours, precipitation and growth of 6’ in alloys Bl and T2 is almost complete. As germanium is removed from solution in the form of germanium precipitate particles, the solubility of lithium in solution would be expected to decrease because the lattice parameter of aluminium is contracting to what it would be for a binary of similar lithium concentration. When this happens the 6’ solvus rises sufficiently for 6’ nucleation to occur. Relatively low heterogeneous nucleation rates of 6’ can be expected initially [Fig. 2(D)], but as the solvus becomes more completely restored (that is, most of the germanium has been removed from the solid solution) extensive nucleation would be expected [Fig. 2(E) and (F)]. The kinetics of this process would
The age hardening behavior and volume fraction curves of alloys Bl and T2 are similar to the results obtained by Noble et al. [16] for binary Al-Li alloys. This hardening is attributed to 6’ since they also observed an increase in hardening with increasing lithium content and showed that alloys with insufficient lithium to form 6’ did not harden during ageing. The much wider maxima around the peak hardness values of alloys Tl and T2 relative to the binary (Bl) (Fig. 1) is due to the presence of germanium precipitates. As 6’ coarsens with ageing, rapid softening occurs in the binary alloy, however, the partially coherent germanium particles delay extensive softening by several hours in the ternary alloys, Germanium additions substantially improve the ductility of Al-Li alloys [Fig. 7(B)]. This improvement can be attributed to a change in deformation from intense slip bands in the binary [Fig. 9(B) and (C)] to dispersed slip in alloys Tl and T2 [Fig. 10(B) and (C)l. During plastic strain the 6’ particles are sheared in both the binary and ternary alloys [Figs 10(C) and 1l(C)]; however, in the ternary alloys partially incoherent germanium particles pin dislocations and disperse slip [Fig. 1l(D)]. The shearing of 6’ in both alloys clearly demonstrates that the improved ductility of the ternary is not due to germanium altering the lattice misfit, but due to partially incoherent germanium particles dispersing the slip. Interactions between precipitates and dislocations depend not only on the properties of the precipitate but also on the crystallographic relationship between particle and matrix. The 6 ‘ phase is coherent with the aluminium matrix and the lattice orientation is a cube/cube type [3], which results in all crystallographic planes and directions being identical at the interphase boundary, including the slip systems. Dislocations easily shear 6’ since they are not required to change glide planes at the interphase boundary. Furthermore, the low value of misfit reported [4] suggest a low interfacial energy and, therefore, easy penetration by dislocations. Considerable strain localization is observed in the binary alloy in the regions of slip bands. The ordered 6’ phase is the major source of strength in the binary Al-Li system. As the phase is sheared on a slip plane by a single dislocation an antiphase boundary is created between the sheared regions of the particle [17]. The formation of such a boundary raises the energy of the system by an amount proportional to the created surface area of the antiphase boundary [18]. However, once the boundary has formed, very little additional stress is needed to move a second dislocation on the slip plane since order is restored and the antiphase boundary removed. Once the 6’
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or presence of a germanium PFZ’s; however the high percentage of transgranular fracture [Fig. 8(b)] and the improved ductility observed in the ternary alloys suggest a relatively homogeneous distribution of the germanium precipitates. CONCLUSIONS 1. Ternary addition of 0.1 at.% germanium results in increased solid solubility of lithium in aluminium. Precipitation of 6’ from the saturated solid solution of alloy Tl was observed to be controlled by removal of germanium from solid solution in the form of germanium particles. 2. Germanium has very little if any solid solubility in Al-2Li at 473 K. The solubility of germanium in pure aluminium is 0.5 wt% (or 0.2 at.%) at 450 K
WI. Fig. 13. Precipitate-free region at grain boundary in alloy Bl aged 24 h at 473 K. The large strain fields at the grain boundary
are produced
by the equilibrium
phase
AlLi.
particles intersecting a given slip plane have been sheared by dislocations, the total area of the plane which is ordered 6’ has been reduced. This can be seen in the displaced fragments of 6’ along the slip trace in Fig. 9(C). The fact that less of the hardening phase occupies a slip plane indicates a low energy pathway for dislocation glide that serves to concentrate subsequent deformation on that plane [19]. Intense slip bands in the matrix result in stress concentrations at grain boundaries and intergranular fracture. Strain localization in binary AL-Li alloys is only part of the explanation of the brittle behavior observed in Bl. All of the alloys used in this study contain coarse grain-boundary 6 ’ precipitate-free zones (PFZ), which are considerably weaker and more ductile than the matrix (Fig. 13). Highly localized strain concentrations accumulating in these weak zones from coarse planer slip in the matrix result in local yielding [19,20]. Cracks may then nucleate within these regions at grain-boundary phases or triple junctions and once formed, propagate rapidly through the PFZ. The material fails intergranularly with little macroscopic deformation of the matrix [Fig. 8(A)]. Although the ternary Al-Li-Ge alloys also contain 6’ PFZ’s they exhibit very different deformation behavior because of the presence of small uniformly distributed germanium particles. These particles are strong and incoherent with the aluminium matrix. Consequently, strain is distributed more homogeneously throughout the matrix and the alloy shows greater work hardening and elongation during tensile deformation [Fig. 7(B)]. Previous studies with Al-Ge alloys [24] have demonstrated that germanium PFZ’s may exist in the binary alloy system depending upon the quenching rate from the solid solution. Results of the present study were inconclusive as to the absence
3. The germanium addition has no effect on the coherency of 6’. This observation along with the low solubility of germanium in the matrix indicates that germanium does not alter the nature of a//matrix interface, nor is there any evidence that the addition of germanium has any effect on the misfit between 6’ and the matrix. 4. The germanium addition results in the precipitation of rod-shaped germanium particles approximately 15-20 nm in length at peak-hardness (12 hr/473 K). These particles are observed to precipitate after about 3 h at 473 K and coarsen as partially coherent rods of germanium. At very long ageing times (110 h/473 K) the rods are observed to develop into highly faceted particles. 5. Tensile testing of comparable binary and ternary alloys of equivalent strength (aged 24 h at 473 K) showed a significant increase in the ductility of the germanium containing alloy. 6. TEM studies of deformation behavior in the ternary alloy reveal that nonshearable germanium particles are responsible for the change in slip behavior observed in Al-2Li-0.2Ge alloys. The germanium particles effectively disperse dislocation slip, thereby providing a homogeneous distribution of stresses during deformation. The degree of coarse planer slip characteristic of Al-2Li is greatly reduced and tensile elongation is increased. Acknowledgements-This work was sponsored by the U.S. Army Research Office under governmentcontract DAAG29-83-K-0038. Dr Phillin A. Parrish director. whose support is gratefully acknowledged
REFERENCES S. Balmuth and R. Schmidt, Aluminium-Lithium Alloys (edited by T. H. Sanders Jr and E. A. Starke Jr.) Metall. Sot. A.I.M.E.. Warrendale. Pa (1981). B. Noble, S. J. Harris and K. Dinsdale,‘J. Muter. Sci. 17, 461 (1982). J. M. Silcock, J. Inst. MeraLs 88, 357 (1959/60). D. B. Williams and J. W. Edington, Me/al Sci. 9, 529 (1975).
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Proc. R. Sot., Lond. 5. H. J. Axon and W. Hume-Rothery, 193A, l-24 (1948). and E. A. Starke Jr, The effect of 6. A. 0. Kujore germanium on the mechanical properties of AL-Li-X alloys, unpublished research, Georgia Institute of Technology, 1982 and A. 0. Kujore Ph.D. thesis, Georgia Institute of Technology. Alanta (1982). 7. F. W. Gayle in Aluminium-Lithium Alloys (edited by T. H. Sanders Jr and E. A. Starke Jr), Metall. SOC. A.I.M.E., Warrendale, Pa (1981). 8. S. F. Baumann and D. B. Williams, Aluminium-Lithium Alloys II (edited by T. H. Sanders Jr and E. A. Starke Jr), pp. 17-28. Metall. Sot. A.I.M.E., Warrendale, Pa (1984). Publication 9. R. M. Fisher, ASTM Special Technical No. 155 (1953). Trans. T.M.S.-A.I.M.E. 224, 906 10. J. E. Hilliard, (1962). 11. U. Koster, Mater, Sci. Engng 5, 174 (1969/70). and L. Schiffini, 12. S. Ceresara, G. Cocco, G. Fagherazzi Phil. Mae. 35. 373 (1977). and D. Turnbull, Acta metall. 7, 664 13. H. S. Rosenbaum (1959).
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OF Al-2Li
14. E. D. Levine and E. .I. Rapperport, Trans. T.M.S.-A.I.M.E. 227, 1204 (1963). 15. S. Ceresara and P. Fiorini, Mater. Sci. Engng 7, 168 (1971). 16. B. Noble, S. J. Harris and K. Dinsdale, Metal Sci. 16, 425 (1982). 17. P. M. Kelly, Int. metaN Rev. 18, 31 (1973). 18. J. B. Cohen, .I. Muter. Sci. 4, 1012 (1969). 19. T. H. Sanders Jr and E. A. Starke Jr, Acta metall. 30, 927 (1982). 20. M. Meuris and E. Hornboeen. Pratt. Metall. 13. 160 (1976). 21. D. J. H. Cockayne, I. L. F. Ray and M. J. Whelan, Phil. Msg. 20, 1265 (1969). Proc. R. Sot. A. 22. I. L. F. Ray and D. J. H. Cockayne, 325, 543 (1971). 23. D. J. H. Cockayne, M. L. Jenkins and I. L. F. Ray, Phil. Mug. 24, 1383 (1971). 24. G. W. Lorimer in Precipitation Processes in Soliak (edited by K. C. Russell and H. I. Aaronson), p. 87. Metall. Sot. A.I.M.E. (1976). 25. H. Stohr and W. Klemm, Z. Anorg. Chem. 241, 305 (1939). ”
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Copyright
OOOI-6160/86 $3.00+0.00 Q 1986 Pergamon Press Ltd
THE EFFECT OF GERMANIUM ON THE PRECIPITATION AND DEFORMATION BEHAVIOR OF Al-2Li ALLOYS W. A. CASSADA, G. J. SHIFLET and E. A. STARKE JR Department of Materials Science, University of Virginia, Charlottesville, VA 22901, U.S.A. (Received 26 April 1985; in revised form 26 June 1985)
Abstract-The effect of a small addition of germanium on the precipitation and deformation behavior of an Al-Li alloy was investigated. Alloys of nominal weight composition of AI-2Li and Al-2Li-0.2Ge were solution heat-treated, quenched, and aged for various times at 473 K. The microstructures and deformation behavior were characterized and the results for the two alloys were compared. The solubility of lithium was increased when germanium was in solid solution; however, lithium decreased the solubility of germanium at 473 K resulting in the formation of small germanium precipitates homogeneously distributed throughout the matrix. These precipitates had a very positive effect on the deformation behavior and ductility of the alloy. R&m&Nous avons ttudit l’influence de petites additions de germanium sur la precipitation et la d&formation d’un alliage Al-Li. Nous avons trait6 thermiquement, trempe et recuit pour diverses durkes B 753 K des alliages de compositions nominales Al-2Li et Al-2Li_0,2Ge en poids. Nous avons caract&& les microstructures et la diformation et nous avons cornpark les rtsultats pour les deux alliages. La solubilitC du lithium Ctait augment&e lorsque le germanium ttait en solution solide; cependant, le lithium diminuait la solubilit6 du germanium B 473 K, ce qui conduisait B la formation de petits pr&cipitts de germanium rkpartis de man&e homogene dans la matrice. Ces pr&cipitCs avaient un effet t&s positif sur la d&formation et la ductilitt de l’alliage. ZusammenfassungAer Einflul3 kleiner Zugaben von Germanium auf das Ausscheidungsverhalten in und das Verfestigungsverhalten von Al-Li-Legierungen wurde untersucht. Legierungen mit der nominalen Gewichtszusammensetzung AI-2Li und Al-2Li-0,2Ge homogenisiert, abgeschreckt und fiir verschiedene Zeiten bei 473 K ausgelagert. Mikrostruktur und Verformungsverhalten wurden bestimmt; die Ergebnisse von beiden Legierungen wurden verglichen. Lag Germanium in Lijsung vor, dann war die Liislichkeit des Lithium erhijht. Andererseits senkte Lithium die Liislichkeit des Germaniums bei 473 K, wodurch sich kleine Germaniumausscheidungen in homogener Verteilung in der Matrix bildeten. Diese Ausscheidungen beeinflussten das Verformungsverhalten und die Duktilitiit der Legierung sehr giinstig.
INTRODUCTION Lithium is the lightest metallic element and has a significant effect on density when alloyed with aluminium, producing a 3% decrease for every weight percent added up to the limit of solid solubility [l]. In addition, a concomitant increase in elastic modulus has been reported for both binary and complex alloy systems [2]. Unfortunately, these beneficial effects are usually accomplished by a significant decrease in ductility and toughness making the simple alloys unsuitable for many structural applications. The low ductility has been attributed to strain localization associated with the ordered Ll,, Al, Li(6’) that forms during age hardening [3]. The similarity in structure and lattice parameters of the 6’ and the fee matrix results in a small lattice misfit (- 0.08 f 0.02%) [4] and spherical precipitates which are easily sheared by moving dislocations. Axon and Hume-Rothery found that the lattice of aluminium is contracted by the solution of lithium and expanded by the solution of germanium [5]. Since 367
these elements have opposite effects on the lattice parameter in solid solution, Kujore and Starke [6] suggested that when added in combination they may cluster at the solutionizing temperature in order to minimize lattice distortion and germanium may subsequently effect the shearability of 6’ by modifying S’/u misfit or the nature of the s’/a interface. Similar studies conducted by Gayle [17] and Baumann and Williams [8] investigated the effect of alloying additions of Ag, SC, Cu, Zn, Ga, Mg, Mn, Cr, Zr and Si on the S’ju interface. Of these, only the alloys containing Cu, Ag and Zn had any appreciable effect on either the misfit or segregation of solute between the 6’ and the matrix. Moreover, the observed effect of the ternary additives was not large enough to change the deformation behavior of the alloys as the 6’ remains fully coherent with the matrix and shearable by dislocations. Germanium additions to Al-Li alloys, however, were shown by Kujore and Starke to have a beneficial effect on ductility by decreasing strain localization during tensile tests [6]. In the present investigation carefully controlled iso-
368
CASSADA ef al.: DEFORMATION BEHAVIOR OF Al-2Li Table I. Alloy composition Alloy Bl Tl T2
Nominal composition
Actual %Li at. % (wt%)
Actual %Ge at. % (wt%)
AI-2Li AI-2.0Li-0.2 Ge Al-2.4Li-0.2 Ge
8.1 I (2.22) 6.94 (1.88) 9.10 (2.51)
O.lO(O.2) 0.10 (0.2)
thermal heat treatments and detailed examination by transmission electron microscopy (TEM) were used to establish the reasons for the beneficial effect that germanium has on the ductility of Al-Li alloys. EXPERIMENTAL The chemical compositions of the three alloys used in this study are listed in Table 1. Alloys Bl and Tl were prepared by the Reynolds Metals Metallurgical Laboratory where 10 kg heats were induction melted in a graphite crucible and cast into a 6.4cm ID copper mold. All melting and casting were done under a partial pressure of high purity argon. The as-cast ingots were machined in order to remove several millimeters of surface, homogenized at 803 K for 24 h, and plastically deformed at 723 K to provide a 10: 1 area reduction ratio. Alloy T2 was prepared at the Naval Surface Weapons Center. A 5 kg heat was resistance melted in a graphite crucible and cast into a tapered copper mold. Melting and casting were performed under high purity argon at atmospheric pressure. The as-cast ingot was machined and homogenized as before, and plastically deformed at 723 K to a 10: 1 area reduction ratio. All samples of alloys used in these experiments were thermally processed under identical conditions. Samples were solutionized in a high temperature salt-bath at 823 K for 30 min, quenched in water and transferred to the ageing bath. The presence of germanium in the ternary alloys did not have an effect on the grain size. All three alloys exhibited an equiaxed grain structure after solution heat treating with an average grain size ranging from 340 to 360 pm. Two types of ageing treatments were used; either a single ageing treatment for various times at 473 K or a two step ageing treatment which consisted of 12 h at 338 K followed by ageing for various times at 473 K. The single ageing treatment was used for all but one of the experiments. Knoop hardness measurements were made on samples of the aged alloys using a micro-hardness tester with a 1 kg load. Samples used in hardness testing were carefully prepared to remove at least 5 pm of surface by mechanically polishing the aged material through 0.1 pm MgO. TEM foils were prepared by electropolishing 3 mm diameter discs punched from selected microstructures. The discs were electropolished in a twin jet polishing apparatus at 13 V d.c. using a 2: 1 methanol:nitric acid electrolyte cooled to 249 K. Germanium particles were extracted from the ternary alloy samples aged 24 and 110 h at 473 K, using
carbon extraction techniques [9]. Polished specimens were etched for 5 min in a SN NaOH solution at room temperature and approximately 120 nm of carbon was deposited to the etched surface by vapor deposition. The carbon film was subsequently removed from the aluminium surface by a second etch in the NaOH solution and washed several times in distilled water before transferring to a 3 mm copper grid for TEM examination. Tensile testing of the three alloys was performed on a closed loop servohydraulic MTS machine in stroke control at a strain rate of 10-3m s-’ using a 10 mm extensometer mounted on the specimen to measure strain. Selected microstructures were stretched 2% in order to investigate particle-dislocation interactions with TEM. All TEM examinations were conducted using a Philips EM 400T. The fracture surfaces of tensile samples were examined with a JEOL JSM-35 scanning electron microscope (SEM). EXPERIMENTAL RESULTS Isothermal ageing behavior
The age-hardening response for alloys Bl, Tl, and T2 is shown in Fig. 1, where the Knoop Hardness Number (KHN) is plotted as a function of ageing time at 473 K. Comparison of the peak hardness for each alloy demonstrates a direct relationship between the lithium concentration and hardness. Alloy Tl exhibits little hardening during the first 3-5 h, while alloys Bl and T2 show a steady logarithmic increase to peak hardness. After about 5 h alloy Tl begins to
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Fig. 1. Variation
100.0
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1.0
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WRS)
in Knoop hardness time at 473 K.
number
with ageing
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Fig. 2. Dark-field transmission electron micrographs of AL, Li precipitate particles in the three alloys after ageing 3, 9 and 12 h (by row) at 473 K. (A)-(C) B1; (D)-(F) Tl; (G)-(I) T2.
harden relatively quickly and reaches peak age after approximately 12 h. Alloy T2 also reaches peak age after about 12 h, while alloy Bl reaches its peak hardness at a somewhat earlier time (about 9 h). The binary alloy Bl overages faster than either of the ternary alloys. Precipitation and coarsening of 6’ Figure 2 is a series of 6’ centered dark-field (CDF) TEM images of the three alloys illustrating the relative 6’ size for each after 3, 9, and 12 h at 473 K. Alloys Bl and T2 show a steady increase in particle size and a steady decrease in number with increasing ageing time. Distinct superlattice spots were observed in the electron diffraction patterns of both after quenching from the solutionizing temperature, and 6’ particles of approximately 2 nm diameter were observed after 30 min ageing at 473 K. The density of particles after short ageing times (OS-l.0 h) was so great that reliable estimates of particle densities were difficult to determine.
The precipitation behavior exhibited by alloy Tl [Fig. 2(D,E,F)] during the first 3-5 h of ageing is very different from the other two alloys. Close examination of 6’ particles during this interval revealed that all of these particles are associated with matrix dislocations [Fig. 2(D)]. The density of 6’ particles increases with ageing time after 3 h. This increase is reflected in the hardness curve as well as in the TEM image [Fig. 2(E)]. Copious 6’ precipitation occurs at some time after the first 3-5 h, and normal coarsening is observed thereafter [Fig. 2(F)]. The micrographs in Fig. 2 also indicate a relationship between the 6’ particle size after a given ageing time and the lithium content of each alloy. The largest particles are seen in alloy Tl which has the lowest percent lithium and, therefore, the lowest nucleation rate. The smallest particles are seen in alloy T2 which has the greatest density of particle and the largest concentration of lithium. The S’ volume fraction determined from the 6’ dark-field images and corrected using the treatment
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of Hilliard [lo], is shown as a function of ageing time in Fig. 3. It is obvious from this plot that nucleation of 6’ in alloys Bl and T2 is over very quickly since the volume fraction of 6 ’ is essentially constant after ageing for 3 h. The differences in precipitation behavior observed in the TEM micrographs (Fig. 2) of alloy Tl are also reflected in the volume fraction ageing curve for this alloy. There is little or no change in volume fraction of 6’ through the first 3 h of 12 h (peakageing. From 3 h to approximately hardness) the volume fraction of 6’ shows a sigmoidal increase similar to that observed for the change in hardness as a function of ageing time (Fig. 1). After 12 h of ageing the volume fraction reaches a constant value. Precipitation and coarsening of germanium
After about 3 h of ageing at 473 K very small (about 2 nm) particles of germanium are observed in the matrix of alloys Tl and T2. An example of these particles is seen in Fig. 4 which is a TEM bright-field (BF) image of alloy Tl aged 3 h. As ageing time is increased the germanium particles coarsen as rods approximately 15-20 nm in length, with an aspect ratio of about 15 [Fig. 5(A) and 5(B)]. These develop into faceted particles [Fig. 5(C) and 5(D)] as ageing progresses. The density of germanium particles in the thin foils was insufficient for obtaining select area electron diffraction patterns. Identification was accomplished by an extraction process in which the germanium particles were removed from the matrix on a carbon replica which was then analyzed by convergent beam electron diffraction in the TEM. Figure 6 shows such a germanium particle extracted from alloy Tl aged 110 h along with its electron diffraction pattern identifying the diamond cubic lattice of germanium (011 zone axis). The extra spots occurring in this pattern are due to twinned areas in the particle which are visible in the micrograph. The normally faceted cor-
BEHAVIOR
of Al,Li
with ageing
OF Al-2Li
time at 473 K.
ners of the germanium particles are rounded by the solutions used to extract it from the matrix. The orientation of the germanium particles in the aluminium matrix can be deduced by systematically producing a series of images from aluminium matrix and germanium precipitate diffracted spots of approximately the same g vector. The occurrence or absence of DF images of the germanium particles along with the orientation relationship of the three forms of germanium particles determined by Koster [l 11, were used to determine the orientation relationship; (11 I)/[0221 Ge parallel to (001)/[020] Al. To determine the effect of incubation time on the precipitation of germanium particles, samples of alloys Tl and T2 were solution heat-treated at 823 K and cold water quenched as before. Half of the
Fig. 4. Transmission electron micrograph nium particles in alloy Tl after ageing
showing germa 3 h at 473 K.
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Fig. 5. Transmission electron micrographs showing germanium particles in alloy Tl. (A) Bright-field image of rod-shaped germanium particles in sample aged 9 h at 473 K; (B) dark-field image of (A); (C) bright-field image of faceted germanium particles in sample aged 110 h at 473 K; (D) dark-field of (C).
samples were immediately transferred to an ageing bath at 473 K, while the other half was pre-aged at 338 K for 12 h before ageing at 473 K. When samples of each were compared for germanium particle density in the TEM no difference could be seen in the number of germanium particles. Preageing, therefore, had no observable effect on the density of germanium particles in any of the alloys investigated.
Tensile properties The tensile properties of the three alloys are given 7(A) and (B) where yield strength and elongation to failure are plotted as a function of ageing time at 473 K. The yield strength of Bl and T2 increase very rapidly, following the rapid increase observed in 6’ volume fractions of the alloys (Fig. 3). Alloy Tl exhibits a slower rate of increase in yield strength with ageing time, also corresponding to
in Fig.
changes in 6’ volume fraction. Alloys Bl and T2 reach similar values of yield strength after 24 h at the ageing temperature (420 and 415 MPa respectively). Alloy Tl, however, reaches a peak yield strength which is considerably less than the other two alloys due to its significantly lower lithium content (about 325 MPa after 24 h at 473 K). Comparing elongation to failure for alloys Bl (8.1 at.% lithium) and Tl (6.9 at.% lithium) demonstrate a decrease in ductility with increasing lithium content [Fig. 7(B)]. After ageing 24 h at 473 K the elongation of Tl is about 9.5% while that of Bl is only about 3.5%. Comparison of elongation to failure between alloys Bl and T2, however, show the opposite trend. When both of these alloys are aged to similar yield strength (24 h at 473 K) the elongation of T2 is about 7% or approximately twice as great as Bl. The lithium content of T2 (9.1 at.%) is about 1 at.% greater than that of Bl.
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Fig. 6. Transmission electron micrograph of a single germanium particle extracted onto a carbon replica film from a specimen of alloy Tl aged 110 h at 473 K. The corresponding convergent-beam electron diffraction pattern [Oil] is shown in the upper right.
A comparison of the fracture surfaces of Bl and T2 aged to a similar yield strength (24 h at 473 K) reveals failure by different mechanisms [Fig. 8(A) and (B)]. Fracture of the binary alloy was predominately intergranular. Slip offsets (arrows) on the grain boundary walls suggest the presence of coarse slip bands within the matrix. The ternary alloy, however, exhibits a
Fig. 7. Variation in the tensile properties of the three alloys with ageing time at 473 K. (A) Yield stress plotted as a function of the ageing time; (B) elongation corresponding with (A) plotted as a function of the ageing time.
Fig. 8. Scanning electron micrographs of fracture surfaces of tensile samples aged 24 h at 473 K; (A) Bl; (B) T2. The yield strength of both alloys was approximately 410 MPa. Arrows in (A) indicate slip band offsets at the grain boundary walls of alloy Bl.
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Fig. 9. Deformation structure near the fracture surface in alloy Bl aged 24 h at 473 K. (A) Strain localization in vicinity of grain boundary (top); (B) dislocation structure in slip bands; (C) sheared A& Li precipitates within slip band. high degree of ductile transgranular failure. This type of failure is consistent with the greater elongation observed for T2. The deformation structures of the tensile specimens are shown in the TEM’s of Figs 9 and 10. At low magnification both Bl and T2 show intense slip bands within the matrix. At higher magnification, however, the ternary alloy reveals a more uniform distribution of dislocations within the matrix between the slip bands. Dark-field images of 6’ in each alloy demonstrate that the binary alloy is subjected to a greater degree of strain localization as sheared 6’ is observed along very long slip lines [Fig. 9(C)]. This type of straight slip line is not observed in the ternary alloy due to a more dispersed dislocation distribution [Fig. 10(C)]. Effect of 2% strain The
ternary
alloy
Tl
aged
48 h at 473 K and
stretched 2% allows a more detailed TEM examination of the deformation behavior of the ternary alloys due to its smaller number fraction of 6’ particles. Figures 1 l(A) and (B) are a bright-field and weak-beam dark-field pair which show the same uniform dislocation distribution seen in the T2 tensile specimen [Fig. 10(B)]. Since the strain is much less, slip bands are not observed in the matrix of this alloy as they are in the T2 tensile specimen [Fig. 10(A) and (B)]. Dislocations are either cross-slipping around 6’ or around germanium particles which lie on common slip planes, or the dislocations are pinned at one of the particle-matrix interphase boundaries. The 6’ dark-field images observed in the T2 tensile specimen [Fig. 10(C)] showed dark wavey lines through the 6’ at or near the locations of dislocations. Furthermore, the S’ particles are obviously distorted from their normally spherical geometry suggesting that these particles are being sheared by the moving dis-
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Fig. 10. Deformation structure near the fracture surface of alloy T2 aged 24 h at 473 K. (A) Deformation in vicinity of grain boundary; (B) dark-field image of dislocation structure; (C) dark-field image of Al, Li precipitates in same region as (B).
locations. This observation tends to support the germanium phase as the agent dispersing the slip in the ternary alloys. The germanium particles are known to be partially incoherent [l l] which is the most important factor in determining the shearability of a precipitate phase. Weak-beam dark-field images [21-231 of germanium particles and dislocations are shown in Fig. 11(D) for the bright-field image in Fig. 1l(C). A region is selected in which a large number of germanium particles are visible with relatively few 6’ particles in order to eliminate any possible influence from 6’ particle-dislocation interaction. The arrows in Fig. 11(D) indicate points where the strain fields of the dislocations are extinguished by pinning at the germanium particles. It is clear from these micrographs that pinning of dislocations at the germanium particles is responsible for homogenizing slip in the ternary alloys.
DISCUSSION OF RESULTS Precipitation
The sequence of precipitation in the ternary alloys and the delay of extensive precipitation of 6’ in alloy Tl can be understood by first considering the aluminium rich side of the Al-Li binary system (Fig. 12) [12]. The relative positions of the alloys Bl, Tl, and T2 on this diagram are indicated at 473 K. From inspection of Fig. 12 it is clear that Bl and T2 lie well under the 6’ solvus boundary where copious precipitation of 6’ is expected to occur (Fig. 3). Alloy Tl on the other hand, lies very close the 6’ solvus. If the ternary addition of germanium had the effect of lowering this solvus a few degrees (S-10 K), the resultant supercooling would be even less. Several ternary additions have been observed to raise the 6’ solvus [8], but to our knowledge, no additions have been observed to lower it. Among the
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Fig. 1I, Deformation structure in the matrix of alloy Tl aged 48 h at 473 K and stretched 2%. (A) Bright-field image; (B) dark-field image of (A) showing homogeneous distribution of dislocations in matrix along with small germanium particles; (C) bright-field image showing spherical Al,Li precipitates and germanium particles; (D) corresponding dark-field image of (C) showing dislocations pinned at the germanium particles. elements found to decrease the solid solubility of lithium in aluminium are: Ag, Zn, Mg, and Mn. Of these Zn and Mn decrease the lattice parameter of aluminium, Mg increases it while Ag has no effect on the lattice parameter of the matrix. Most of these alloying additions will form transitional phases of their own and all of them may also form phases with lithium. These factors probably have the greatest influence on the solubility of lithium in the matrix. Ternary additions of germanium are very different from those additions mentioned above. Covalent solutes in metallic solvents, such as germanium in aluminium, generally do not form transitional phases [13]. An example of this is found in the Al-Si system where no transitional phases exist. Germanium and silicon possess identical valence shells and similar chemistry, so the similarity in behavior of binary alloys of each is not surprising.
The contracting effect of lithium on the aluminium lattice [5, 141 may in fact be the primary reason for the decreased solubility of germanium in aluminium. Conversely, the expanding effect of the germanium on the matrix lattice parameter could increase the solid solubility of the lithium. Within the compositional ranges of the alloys used in this study, germanium expands the aluminium lattice by about 1.08 x 10m4nm/at.% germanium while lithium contracts the aluminium lattice approximately 4.65 x lO~‘nm/at.% lithium [5]. Due to the opposing effects of the solutes on aluminum’s lattice parameter, germanium and lithium atoms may be attracted to one another in order to relieve lattice strain [6]. However, the covalent nature of germanium and the metallic lithium would result in little chemical affinity between the two. The normally high vacancy-binding energy of the germa-
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be limited by the removal of germanium from the solid solution as precipitate particles, thus explaining the relatively slow rate of increase in 6’ volume fraction in Tl observed in Fig. 3. Mechanical behavior
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Fig. 12. Aluminum-rich end of the binary Al-Li system after Ceresara et al. [12].
nium in aluminium [ 111is most likely increased as the tra $ ping of a vacancy by a germanium atom in this matrix would release an additional amount of strain energy. The amount of this increase in vacancy binding energy would be related to the decrease in lattice parameter. As the density of Ge-vacancy couples increases there is a concomitant increase in germanium nucleation sites [15]. Because of the decrease in germanium solubility created by the lithium in solid solution, sufficient supersaturation exists for the nucleation and growth of germanium particles. This appears to be happening in alloy T2, since both 6’ and germanium particles have sufficient undercooling for nucleation at early ageing times. The germanium particles are precipitating after 6 ‘, and at about the same time as germanium particles precipitate in alloy Tl. In alloy Tl, however, the composition is at or above the 6’ solvus at 473 K leading to little or no undercooling for precipitation. Figure 3 illustrates that while the volume fraction of 6’ in alloy Tl is nearly zero after three hours, precipitation and growth of 6’ in alloys Bl and T2 is almost complete. As germanium is removed from solution in the form of germanium precipitate particles, the solubility of lithium in solution would be expected to decrease because the lattice parameter of aluminium is contracting to what it would be for a binary of similar lithium concentration. When this happens the 6’ solvus rises sufficiently for 6’ nucleation to occur. Relatively low heterogeneous nucleation rates of 6’ can be expected initially [Fig. 2(D)], but as the solvus becomes more completely restored (that is, most of the germanium has been removed from the solid solution) extensive nucleation would be expected [Fig. 2(E) and (F)]. The kinetics of this process would
The age hardening behavior and volume fraction curves of alloys Bl and T2 are similar to the results obtained by Noble et al. [16] for binary Al-Li alloys. This hardening is attributed to 6’ since they also observed an increase in hardening with increasing lithium content and showed that alloys with insufficient lithium to form 6’ did not harden during ageing. The much wider maxima around the peak hardness values of alloys Tl and T2 relative to the binary (Bl) (Fig. 1) is due to the presence of germanium precipitates. As 6’ coarsens with ageing, rapid softening occurs in the binary alloy, however, the partially coherent germanium particles delay extensive softening by several hours in the ternary alloys, Germanium additions substantially improve the ductility of Al-Li alloys [Fig. 7(B)]. This improvement can be attributed to a change in deformation from intense slip bands in the binary [Fig. 9(B) and (C)] to dispersed slip in alloys Tl and T2 [Fig. 10(B) and (C)l. During plastic strain the 6’ particles are sheared in both the binary and ternary alloys [Figs 10(C) and 1l(C)]; however, in the ternary alloys partially incoherent germanium particles pin dislocations and disperse slip [Fig. 1l(D)]. The shearing of 6’ in both alloys clearly demonstrates that the improved ductility of the ternary is not due to germanium altering the lattice misfit, but due to partially incoherent germanium particles dispersing the slip. Interactions between precipitates and dislocations depend not only on the properties of the precipitate but also on the crystallographic relationship between particle and matrix. The 6 ‘ phase is coherent with the aluminium matrix and the lattice orientation is a cube/cube type [3], which results in all crystallographic planes and directions being identical at the interphase boundary, including the slip systems. Dislocations easily shear 6’ since they are not required to change glide planes at the interphase boundary. Furthermore, the low value of misfit reported [4] suggest a low interfacial energy and, therefore, easy penetration by dislocations. Considerable strain localization is observed in the binary alloy in the regions of slip bands. The ordered 6’ phase is the major source of strength in the binary Al-Li system. As the phase is sheared on a slip plane by a single dislocation an antiphase boundary is created between the sheared regions of the particle [17]. The formation of such a boundary raises the energy of the system by an amount proportional to the created surface area of the antiphase boundary [18]. However, once the boundary has formed, very little additional stress is needed to move a second dislocation on the slip plane since order is restored and the antiphase boundary removed. Once the 6’
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or presence of a germanium PFZ’s; however the high percentage of transgranular fracture [Fig. 8(b)] and the improved ductility observed in the ternary alloys suggest a relatively homogeneous distribution of the germanium precipitates. CONCLUSIONS 1. Ternary addition of 0.1 at.% germanium results in increased solid solubility of lithium in aluminium. Precipitation of 6’ from the saturated solid solution of alloy Tl was observed to be controlled by removal of germanium from solid solution in the form of germanium particles. 2. Germanium has very little if any solid solubility in Al-2Li at 473 K. The solubility of germanium in pure aluminium is 0.5 wt% (or 0.2 at.%) at 450 K
WI. Fig. 13. Precipitate-free region at grain boundary in alloy Bl aged 24 h at 473 K. The large strain fields at the grain boundary
are produced
by the equilibrium
phase
AlLi.
particles intersecting a given slip plane have been sheared by dislocations, the total area of the plane which is ordered 6’ has been reduced. This can be seen in the displaced fragments of 6’ along the slip trace in Fig. 9(C). The fact that less of the hardening phase occupies a slip plane indicates a low energy pathway for dislocation glide that serves to concentrate subsequent deformation on that plane [19]. Intense slip bands in the matrix result in stress concentrations at grain boundaries and intergranular fracture. Strain localization in binary AL-Li alloys is only part of the explanation of the brittle behavior observed in Bl. All of the alloys used in this study contain coarse grain-boundary 6 ’ precipitate-free zones (PFZ), which are considerably weaker and more ductile than the matrix (Fig. 13). Highly localized strain concentrations accumulating in these weak zones from coarse planer slip in the matrix result in local yielding [19,20]. Cracks may then nucleate within these regions at grain-boundary phases or triple junctions and once formed, propagate rapidly through the PFZ. The material fails intergranularly with little macroscopic deformation of the matrix [Fig. 8(A)]. Although the ternary Al-Li-Ge alloys also contain 6’ PFZ’s they exhibit very different deformation behavior because of the presence of small uniformly distributed germanium particles. These particles are strong and incoherent with the aluminium matrix. Consequently, strain is distributed more homogeneously throughout the matrix and the alloy shows greater work hardening and elongation during tensile deformation [Fig. 7(B)]. Previous studies with Al-Ge alloys [24] have demonstrated that germanium PFZ’s may exist in the binary alloy system depending upon the quenching rate from the solid solution. Results of the present study were inconclusive as to the absence
3. The germanium addition has no effect on the coherency of 6’. This observation along with the low solubility of germanium in the matrix indicates that germanium does not alter the nature of a//matrix interface, nor is there any evidence that the addition of germanium has any effect on the misfit between 6’ and the matrix. 4. The germanium addition results in the precipitation of rod-shaped germanium particles approximately 15-20 nm in length at peak-hardness (12 hr/473 K). These particles are observed to precipitate after about 3 h at 473 K and coarsen as partially coherent rods of germanium. At very long ageing times (110 h/473 K) the rods are observed to develop into highly faceted particles. 5. Tensile testing of comparable binary and ternary alloys of equivalent strength (aged 24 h at 473 K) showed a significant increase in the ductility of the germanium containing alloy. 6. TEM studies of deformation behavior in the ternary alloy reveal that nonshearable germanium particles are responsible for the change in slip behavior observed in Al-2Li-0.2Ge alloys. The germanium particles effectively disperse dislocation slip, thereby providing a homogeneous distribution of stresses during deformation. The degree of coarse planer slip characteristic of Al-2Li is greatly reduced and tensile elongation is increased. Acknowledgements-This work was sponsored by the U.S. Army Research Office under governmentcontract DAAG29-83-K-0038. Dr Phillin A. Parrish director. whose support is gratefully acknowledged
REFERENCES S. Balmuth and R. Schmidt, Aluminium-Lithium Alloys (edited by T. H. Sanders Jr and E. A. Starke Jr.) Metall. Sot. A.I.M.E.. Warrendale. Pa (1981). B. Noble, S. J. Harris and K. Dinsdale,‘J. Muter. Sci. 17, 461 (1982). J. M. Silcock, J. Inst. MeraLs 88, 357 (1959/60). D. B. Williams and J. W. Edington, Me/al Sci. 9, 529 (1975).
378
CASSADA
et al.:
DEFORMATION
Proc. R. Sot., Lond. 5. H. J. Axon and W. Hume-Rothery, 193A, l-24 (1948). and E. A. Starke Jr, The effect of 6. A. 0. Kujore germanium on the mechanical properties of AL-Li-X alloys, unpublished research, Georgia Institute of Technology, 1982 and A. 0. Kujore Ph.D. thesis, Georgia Institute of Technology. Alanta (1982). 7. F. W. Gayle in Aluminium-Lithium Alloys (edited by T. H. Sanders Jr and E. A. Starke Jr), Metall. SOC. A.I.M.E., Warrendale, Pa (1981). 8. S. F. Baumann and D. B. Williams, Aluminium-Lithium Alloys II (edited by T. H. Sanders Jr and E. A. Starke Jr), pp. 17-28. Metall. Sot. A.I.M.E., Warrendale, Pa (1984). Publication 9. R. M. Fisher, ASTM Special Technical No. 155 (1953). Trans. T.M.S.-A.I.M.E. 224, 906 10. J. E. Hilliard, (1962). 11. U. Koster, Mater, Sci. Engng 5, 174 (1969/70). and L. Schiffini, 12. S. Ceresara, G. Cocco, G. Fagherazzi Phil. Mae. 35. 373 (1977). and D. Turnbull, Acta metall. 7, 664 13. H. S. Rosenbaum (1959).
BEHAVIOR
OF Al-2Li
14. E. D. Levine and E. .I. Rapperport, Trans. T.M.S.-A.I.M.E. 227, 1204 (1963). 15. S. Ceresara and P. Fiorini, Mater. Sci. Engng 7, 168 (1971). 16. B. Noble, S. J. Harris and K. Dinsdale, Metal Sci. 16, 425 (1982). 17. P. M. Kelly, Int. metaN Rev. 18, 31 (1973). 18. J. B. Cohen, .I. Muter. Sci. 4, 1012 (1969). 19. T. H. Sanders Jr and E. A. Starke Jr, Acta metall. 30, 927 (1982). 20. M. Meuris and E. Hornboeen. Pratt. Metall. 13. 160 (1976). 21. D. J. H. Cockayne, I. L. F. Ray and M. J. Whelan, Phil. Msg. 20, 1265 (1969). Proc. R. Sot. A. 22. I. L. F. Ray and D. J. H. Cockayne, 325, 543 (1971). 23. D. J. H. Cockayne, M. L. Jenkins and I. L. F. Ray, Phil. Mug. 24, 1383 (1971). 24. G. W. Lorimer in Precipitation Processes in Soliak (edited by K. C. Russell and H. I. Aaronson), p. 87. Metall. Sot. A.I.M.E. (1976). 25. H. Stohr and W. Klemm, Z. Anorg. Chem. 241, 305 (1939). ”
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