STRUCTURE AND PROPERTIES OF RAPIDLY SOLIDIFIED ALUMINUM-LITHIUM ALLOYS

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STRUCTURE AND PROPERTIES OF RAPIDLY SOLIDIFIED ALUMINUM-LITHIUM ALLOYS

N. Kim, R. Bye, S. Das

To cite this version:

N. Kim, R. Bye, S. Das. STRUCTURE AND PROPERTIES OF RAPIDLY SOLIDIFIED ALUMINUM-LITHIUM ALLOYS. Journal de Physique Colloques, 1987, 48 (C3), pp.C3-309-C3-315.

�10.1051/jphyscol:1987335�. �jpa-00226566�

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JOURNAL DE PHYSIQUE

Colloque C 3 , supplgment au n09, Tome 48, septembre 1987

STRUCTURE AND PROPERTIES OF RAPIDLY SOLIDIFIED ALUMINUM-LITHIUM ALLOYS

N.J. K I M , R . L . BYE and S . K . DAS

Metals and Ceramics Laboratory, Allied-Signal Incorporation, P.O. Box 1021 R, Morristown, NJ 07960, U.S.A.

Considerable improvements have been made during the last decade in the properties of A1-Li alloys for aerospace applications. However, the development of A1-Li alloys for applications that require higher stiffness and lower density than those achieved with ingot casting technology necessitates the use of rapid solidification processing to minimize microsegregation related property deficiencies. The present paper discusses the current status of rapidly solidified A1-Li alloy development at Allied-Signal. The main emphasis of the program has been placed on the achievement of optimum properties by simple heat treatment process, which appears to be essential for the development of high-performance near net-shape A1-Li alloy forgings.

INTRODUCTION

Aluminum-lithium alloys, because of their low density and high elastic modulus, have received great attention for aerospace applications (see

e.g., 1-3). The addition of one wt% lithium (- 3.5 at%) to aluminum decreases the density by

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3% and increases the.elastic modulus by

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^6%,hence giving a substantial increase in specific modulus. Their tendency to brittleness, however, has been a major factor which hinders the development of high

performance aluminum-lithium alloys. Although there has been much progress in the development of ingot metallurgy (I/M) aluminum-lithium alloys in recent years, the problem of segregation encountered in ingot casting limits the maximum lithium content to

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2.7 wt%, giving weight saving potential to approximately 8%. Increasing the lithium contents to further improve the density and elastic modulus usually results in a reduction in ductility and fracture toughness. Moreover, some of the approaches utilized in ingot cast aluminum-lithium alloys to improve toughness, such as stretching, cannot be utilized for forging applications. In aluminum-lithium alloys prepared by rapid solidification, however, possible weight savings of 10-15% can be achieved and elastic modulus can be increased further by using higher lithium contents ( > 3 wt%). In addition, rapidly solidified aluminum-lithium alloys do not require a stretching operation, thereby making them particularly well suited for forging applications.

The present paper discusses the structure and properties of the recently developed rapidly solidified aluminum-lithium alloys with high lithium ( > 3 wt%) and zirconium (0.5 wt%) contents. The main principles of alloy design and rapid solidification processing utlilized in our development program have been discussed in previous papers (4,5)

MATERIALS AND EXPERIMENTAL PROCEDURE

The compositions of the alloys used in this investigation are shown in Table 1. These alloys have been designed to achieve densities lower than 2.50 g/cm3 and their densitiies are shown in Table 1. Alloys were rapidly quenched from the melt into continuous ribbons by the jet casting process.

Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:1987335

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TABLE 1

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ALLOY COMPOSITIONS (wt%) AND DENSITIES (a/cm3)

A1 l o y L i Cu Mg Zr DENSITY

These ribbons were mechanically comminuted t o -40 mesh powder and then con- s o l i d a t e d i n t o b u l k compacts by vacuum hot pressing which were then h o t extruded t o a f i n a l rectangular shape o f 63.5 mm x 10.2 mm (18:l e x t r u s i o n r a t i o , 6 : l aspect r a t i o ) . A d d i t i o n a l a l l o y 644 was comminuted t o -20 mesh powder t o study t h e p o s s i b l e enhancement o f mechanical p r o p e r t i e s by using coarse powder. The extrusions were s o l u t i o n i z e d f o r 2 hours a t 550°C.

Tensile p r o p e r t i e s were measured using round specimens w i t h 19.1 mm (0.75 i n . ) gauge length and 3.2 mm (0.125 i n . ) gauge diameter a t a s t r a i n r a t e o f 4x10-4/sec. Fatigue crack growth and f r a c t u r e toughness t e s t i n g was conducted using compact tension specimens. The dimensions o f t h e specimens were 50 mm x 48 mm x 5 mm, and a load time w i d t h o f 40 mm. Impact s t r e n g t h was measured using V-notch (notch radius: 0.001") impact specimens. Thin f o i l s f o r transmission e l e c t r o n microscopy were prepared by j e t p o l i s h i n g w i t h an e l e c t r o l y t e o f 33% HNO3 and 61% methanol o r w i t h one o f 10% p e r c h l o r i c a c i d and 90% ethanol. TEM observations were made a t an operating v o l t a g e o f 120 KV.

RESULTS 1. M i c r o s t r u c t u r e

Figure 1 i s a t y p i c a l t h r e e dimensional composite o p t i c a l micrograph o f as-solutionized a l l o y . I t shows t h a t t h e m i c r o s t r u c t u r e i s composed o f u n r e c r y s t a l l i z e d g r a i n s which are elongated i n t h e e x t r u s i o n d i r e c t i o n . The volume f r a c t i o n o f coarse undissolved c o n s t i t u e n t p a r t i c l e s i s minimal. Also v i s i b l e i n Figure 1 i s a f i n e d i s p e r s i o n o f p a r t i c l e s which appear t o be along t h e p r i o r powder p a r t i c l e boundaries. TEM s t u d i e s o f as-solutionized a l l o y show t h a t these p a r t i c l e s are mostly oxide and/or carbonate (Figure 2).

1. O p t i c a l m i c r o g r a p h t i o n i z e d 550°C/2 hrs.

alloy

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Fig. 2. TEM micrograph of alloy 6 4 4 showing oxide stringers. Solutionized 550°C/2 hrs.

TEM micrograph o f an as-solutionized a l l o y i s shown i n Figure 3. I t shows t h a t there i s a l a r g e volume f r a c t i o n o f f i n e dispersoids w i t h i n the g r a i n and coarse dispersoids located a t t h e g r a i n boundaries. These dispersoids are a l l metastable L l q A13Zr which are coherent w i t h the aluminum m a t r i x

( 6 ) . Figure 4 shows a t y p i c a l microstructure o f an a r t i f i c i a l l y aged a l l o y . As observed i n many a l l o y systems, an increase i n the degree o f aging r e s u l t s

i n a coarsening o f 6' p r e c i p i t a t e s and t h e development o f p r e c i p i t a t e f r e e zones. One o f the most important m i c r o s t r u c t u r a l features observed i n t h i s a l l o y system i s t h a t most o f t h e 6 ' i s associated w i t h Al3Zr dispersoids, forming composite p r e c i p i t a t e s . Such homogeneous d i s t r i b u t i o n o f a l a r g e volume f r a c t i o n o f composite p r e c i p i t a t e s i s q u i t e d e s i r a b l e f o r t h e improvement o f d u c t i l i t y and toughness. Another important microstructure f e a t u r e o f these a l l o y s i s the absence or sluggish p r e c i p i t a t i o n o f T1 o r S' phases, as was observed i n the e a r l i e r studies ( 4 , 7 ) .

Fig. 3. TEM micrograph of alloy 6 4 8 Fig. 4 . TEM micrograph of alloy 6 4 8 showing a distribution of A13Zr. showing composite ~recipitates.

Solutionized 550°C/2 hrs. Aged 180°C/16 hrs.

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2. Tensile Properties and Fracture Touahness

The room temperature t e n s i l e p r o p e r t i e s o f aluminum-lithium a l l o y s f o r various aging c o n d i t i o n s are l i s t e d i n Table 2, along w i t h the selected f r a c t u r e toughness data. Properties o f a l l o y 644 made from -20 mesh powder are a l s o included. As can be seen i n Table 2, good combinations o f s t r e n g t h and d u c t i l i t y were obtained i n under-aged a l l o y s . I t can a l s o be seen t h a t t h e d i f f e r e n c e i n t h e p a r t i c l e s i z e (-20 vs. -40 mesh) does not r e s u l t i n any noticeable change i n t h e t e n s i l e p r o p e r t i e s o f a l l o y 644. A comparison between a l l o y 643 and a l l o y 644 shows t h a t higher l e v e l o f L i increases the t e n s i l e s t r e n g t h but does not reduce the t e n s i l e d u c t i l i t y .

Since the underaged a l l o y s show good t e n s i l e properties, f r a c t u r e toughness data were generated f o r underaged a l l o y s only. Because o f t h e s i z e l i m i t a t i o n s o f t h e extruded bars, t h e measured toughness i s KQ r a t h e r than K I ~ (plane s t r a i n f r a c t u r e toughness). I t can be seen t h a t a l l o y s 643 and 644 have good f r a c t u r e toughness i n the underaged condition. Although the increase i n L i content from 3 w t % i n a l l o y 644 t o 3.4 w t % i n a l l o y 643 d i d not degrade the elongation, i t r e s u l t e d i n a decrease i n f r a c t u r e toughness i n t h e L-T o r i e n t a t i o n . Some o f t h e decrease i n f r a c t u r e toughness might a l s o be due t o t h e higher s t r e n g t h o f a l l o y 643. The change i n t h e L i content had no e f f e c t on the f r a c t u r e toughness i n the T-L o r i e n t a t i o n . A l l o y s 643 and 644 have s i m i l a r T-L f r a c t u r e toughness values which are l e s s than those i n the L-T o r i e n t a t i o n . Fractographic study o f f r a c t u r e toughness

specimens showed t h a t f a i l u r e occurred mainly along the p r i o r powder part.icle boundaries. This i n d i c a t e s t h a t t h e oxide p a r t i c l e s located along the p r i o r powder p a r t i c l e boundaries (Figure 2) are the dominant f a c t o r i n c o n t r o l l i n g the T-L f r a c t u r e toughness of r a p i d l y s o l i d i f i e d aluminum-lithium a l l o y s . The change i n the powder p a r t i c l e size, from -40 mesh powder t o -20 mesh powder, increased t h e f r a c t u r e toughness o f a l l o y 644, presumably because of the reduced p r i o r p a r t i c l e boundary area.

Table 2

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MECHANICAL PROPERTIES OF RAPIDLY SOLIDIFIED ALUMINUM-LITHIUM ALLOYS

Y.S. U.T.S. E l KQ Impact Strength

Aging MPa MPa %

M P ~ G

~ / m m ~

L-T T-L L-T T-L

644 130°C/16h 430 534 6.7 32 23 0.0322 0.0074

(-40 mesh) 160°C/16h 487 581 6.0

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190°C/16h 456 537 5.6

- - - -

644

(-20 mesh) 130°C/16h 461 574 5.9 38

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^0.0557 ^0.0189

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3. Fatigue Crack Growth

Fig. 5. Fatigue crack growth rates of Fig. 6. SEM micrographs showing fatigue rapidly solidified A1-Li alloys compared crack profile: (a) alloy 648 and (b)

with 7075-T73. alloy 643.

Fatigue crack propagation experiments were conducted f o r selected a l l o y s i n under-aged c o n d i t i o n using compact tension specimens. The t e s t s were conducted i n a laboratory a i r environment using an R - r a t i o o f 0.5. Crack closure was not measured and thus was not corrected f o r . Figure 5 shows f a t i g u e crack growth rate, da/dN, as a f u n c t i o n o f t h e s t r e s s i n t e n s i t y range, AK. The f a t i g u e crack growth r a t e o f 7075-T73 i s a l s o included f o r comparison. O f the a l l o y s studied (643, 644 and 648). a l l o y 648 has t h e f a s t e s t crack propa- g a t i o n r a t e and a l l o y 643 has the slowest crack propagation r a t e . I n a l l the a l l o y s , t h e crack path i s deflected and tortuous (Figure 6). The t o r t u o u s i t y o f the crack path i s most severe i n a l l o y 648 (Figure 6a), whose deformation mode i s t h e most planar o f t h e a l l o y s studied. I t can also be seen i n Figure 6A t h a t a l l o y 648 shows extrusions and s l i p l i n e markings along t h e crack path. Comparison o f f a t i g u e crack growth r a t e s o f these aluminum-lithium a l l o y s w i t h t h a t o f 7075-T73, without considering c l o s u r e e f f e c t s , shows t h a t t h e f a t i g u e performance o f a l l o y s 643 and 644 i s q u i t e competitive o r b e t t e r than the I / M 7075 a l l o y .

DISCUSSION

As has been shown above, good mechanical p r o p e r t i e s as w e l l as low den- s i t i e s have been obtained i n r a p i d l y s o l i d i f i e d aluminum-lithium a l l o y s . The improvements i n mechanical p r o p e r t i e s o f r a p i d l y s o l i d i f i e d aluminum-lithium a l l o y s are mainly due t o the b e n e f i c i a l e f f e c t o f t h e Zr a d d i t i o n on the microstructure. As shown i n Figure 4, t h e microstructure o f aged r a p i d l y s o l i d i f i e d aluminum-lithium a l l o y s i s comprised o f a l a r g e volume f r a c t i o n o f homogeneously d i s t r i b u t e d composite p r e c i p i t a t e s . Most o f t h e 6' i s asso- c i a t e d w i t h composite p r e c i p i t a t e s . These composite p r e c i p i t a t e s have been shown t o s i g n i f i c a n t l y a f f e c t the deformation behavior o f a l l o y s (7-9). It i s w e l l known t h a t i n aluminum-lithium a l l o y s which contain a small volume f r a c t i o n o f heterogeneously d i s t r i b u t e d composite p r e c i p i t a t e s , 6 ' i s sheared by moving d i s l o c a t i o n s , r e s u l t i n g i n severe planar s l i p . On the other,hand, a l l o y s which contain a l a r g e volume f r a c t i o n o f homogeneously d i s t r i b u t e d composite p r e c i p i t a t e s are known t o e x h i b i t a wavy and homogeneous mode o f deformation (7,8) due t o t h e shear-resistant nature o f the composite p r e c i p i - tates. The shear-resistant nature of composite p r e c i p i t a t e s can be understood

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when one considers t h a t the metastable Al3Zr p a r t i c l e i n Al-Zr a l l o y s , although i t i s o f the same c r y s t a l s t r u c t u r e ( L I Z ) as 6'. i s h i g h l y r e s i s t a n t t o d i s l o c a t i o n shear (10). The e f f e c t o f composite p r e c i p i t a t e s on modifying t h e deformation mode can be r e a d i l y seen i n a weak beam dark f i e l d image

(Figure 7 ) . S u p e r l a t t i c e d i s l o c a t i o n p a i r s , which are c h a r a c t e r i s t i c o f sheared p r e c i p i t a t e s , are not observed. Single d i s l o c a t i o n s are seen t o move i n a wavy mode and many d i s l o c a t i o n loops can be seen. By comparing a 6' s u p e r l a t t i c e dark f i e l d image and a weak beam dark f i e l d image o f t h e same area, these d i s l o c a t i o n loops were found t o surround mostly t h e 6'/AlgZr i n t e r f a c e o f composite p r e c i p i t a t e s ( 6 ) . This means t h a t d i s l o c a t i o n s by-pass t h e AlgZr p a r t i c l e s l e a v i n g the d i s l o c a t i o n loops around them, i n d i c a t i n g t h e high resistance o f A13Zr p a r t i c l e s t o d i s l o c a t i o n shear. Although the 6 ' i n composite p r e c i p i t a t e s could be sheared by d i s l o c a t i o n s , the degree o f s t r e s s l o c a l i z a t i o n imposed by 6' i s minimized by the shear-resistant AlgZr which forms t h e core o f composite p r e c i p i t a t e s .

F i g . 7 . Weak beam d a r k f ~ e l d m i c r o g r a p h o f deformed a l l o y 648.

The f a t i g u e c h a r a c t e r i s t i c s o f these a l l o y s show some i n t e r e s t i n g features. As shown i n Figure 5, a l l o y 648 (AT-3.4Li-0.5Zr) has a much f a s t e r f a t i g u r e crack growth r a t e than a l l o y 643 (Al-3.4Li-1Cu-0.5Mg-0.5Zr) which has s i m i l a r L i content but w i t h a d d i t i o n s o f Cu and Mg. I n general, the h i g h r e s i s t a n c e , o f aluminum-lithium a l l o y s t o f a t i g u e crack progagation has been a t t r i b u t e d t o t h e h i g h e l a s t i c modulus and t o the extensive crack d e f l e c - t i o n processes induced by shearing o f 6 ' p r e c i p i t a t e s (11, 12). Several i n v e s t i g a t i o n s have been shown t h a t a rough f r a c t u r e surface gives r i s e t o a component o f crack closure which has the added e f f e c t o f reducing crack growth rates. Although the c u r r e n t r e s u l t s do not consider t h e closure e f f e c t , exa- mination o f f a t i g u e crack paths o f both a l l o y s i n d i c a t e s t h a t the other fac- t o r s are a l s o p l a y i n g important r o l e s i n c o n t r o l l i n g the f a t i g u e crack growth behavior o f r a p i d l y s o l i d i f i e d aluminum-lithium a l l o y s . A l l o y 648 e x h i b i t s f a s t e r crack growth r a t e s than a l l o y 643, although Figure 6 shows t h a t t h e crack path t o r t u o s i t y o r crack d e f l e c t i o n i s more severe i n a l l o y 648. This i n d i c a t e s t h a t roughness induced closure i s not t h e dominant f a c t o r which con- t o l s t h e f a t i g u e crack growth o f these a l l o y s . The r o l e o f s l i p r e v e r s i b i l i t y i n lowering f a t i g u e crack growth r a t e s cannot account f o r t h e crack growth behavior o f both a l l o y s , since the s l i p r e v e r s i b i l i t y would be higher i n a l l o y 648 and t h e p l a s t i c s t r a i n accumulation i n a l l o y 648 would be lower f o r a

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g i v e n number o f c y c l e s t h a n i n a l l o y 643. On t h e o t h e r hand, i t i s w e l l known t h a t t h e p l a s t i c zone s i z e ahead o f t h e crack t i p w i t h r e s p e c t t o t h e r e l e v a n t m i c r o s t r u c t u r a l u n i t p l a y s an i m p o r t a n t r o l e i n c o n t r o l l i n g t h e f a t i g u e c r a c k propagation behavior o f a l l o y s . A l l o y s 648 and 643 have t h e i d e n t i c a l m i c r o - s t r u c t u r e . I t has been shown t h a t s t r e s s - s t r a i n b e h a v i o r s o f t h e two a l l o y s a r e q u i t e d i f f e r e n t , a l l o y 643 showing a much h i g h e r degree o f s t r a i n harden-

i n g than a l l o y 648 ( 4 ) . Considering t h a t t h e c y c l i c p l a s t i c zone s i z e i s i n v e r s e l y p r o p o r t i o n a l t o t h e square o f f l o w s t r e n g t h ( 1 3 ) , a l t o y 643 may have much s m a l l e r p l a s t i c zone s i z e t h a n a l l o y 648. T h i s i m p l i e s t h a t , f o r a g i v e n AK, a l l o y 648 should e x h i b i t f a s t e r c r a c k growth r a t e s t h a n a l l o y 643.

Observation o f e x t r u s i o n s and s l i p l i n e markings a l o n g t h e c r a c k p a t h o f a l l o y 648 supports t h i s e x p l a n a t i o n ( F i g u r e 6 ) .

SUMMARY

The r e s u l t s o f t h e present i n v e s t i g a t i o n have demonstrated t h a t good combinations o f s t r e n g t h , d u c t i l i t y , toughness and f a t i g u e c r a c k growth r e s i s t - ance can be o b t a i n e d i n 3-3.4 w t % L i c o n t a i n i n g r a p i d l y s o l i d i f i e d a l l o y s . A f i n e d i s t r i b u t i o n o f non-shearable A l g ( L i , Z r ) p r e c i p i t a t e s r e s u l t i n g from a h i g h Zr c 6 n t e n t a l l o w s optimum mechanical p r o p e r t i e s t o be achieved by simple thermal treatment w i t h o u t any mechanical working. Long s t a n d i n g problems r e l a t e d t o t h e s u r f a c e o x i d e l a y e r along powder p a r t i c l e boundaries have been shown t o be p a r t i a l l y s o l v e d by u s i n g a coarse powder p a r t i c l e s i z e . An increase i n powder p a r t i c l e s i z e r e s u l t e d i n an improvement i n toughness w i t h o u t a f f e c t i n g t e n s i l e p r o p e r t i e s . Work i s c o n t i n u i n g t o f u r t h e r improve t h e t r a n s v e r s e - o r i e n t e d p r o p e r t i e s by reducing t h e powder s u r f a c e o x i d e l a y e r .

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T.H. Sanders, J r . , and E.A. Starke, J r . , eds., Aluminum-Lithium A l l o y s I, TMS-AIME, Warrendale, PA, 1981.

T.H. Sanders, Jr., and E.A. Starke, J r . , eds., Aluminum-Lithium A l l o y s 11, TMS-AIME, Warrendale, PA, 1984.

C. Baker, P.J. Gregson, S.J. H a r r i s and C.J. Peel, eds., Aluminum-Lithium A l l o y s 111, The I n s t i t u t e o f Metals, London, U.K., 1986.

N.J. Kim, R.L. Bye, A.M. Brown, S.K. Das and C.M. Adam, p. 975 i n Aluminum A l l o y s and T h e i r P h y s i c a l and Mechanical P r o p e r t i e s , eds., E.A. Starke, J., and T.H. Sanders, Jr., EMAS, West Midlands, U.K., 1986.

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P.C. P a r i s , p. 107 i n Proc. 1 0 t h Sagamore Conference, Syracuse U n i v e r s i t y Press, Syracuse, NY, 1962.