EXTRUSION OF AN Al-Li-Zr ALLOY PREPARED FROM ATOMIZED POWDER

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EXTRUSION OF AN Al-Li-Zr ALLOY PREPARED FROM ATOMIZED POWDER

M. Mahmoud, H. Mcshane, T. Sheppard

To cite this version:

M. Mahmoud, H. Mcshane, T. Sheppard. EXTRUSION OF AN Al-Li-Zr ALLOY PREPARED FROM ATOMIZED POWDER. Journal de Physique Colloques, 1987, 48 (C3), pp.C3-327-C3-334.

�10.1051/jphyscol:1987337�. �jpa-00226568�

EXTRUSION OF AN Al-Li-Zr ALLOY PREPARED FROM ATOMIZED POWDER

M.S. MAHMOUD, H.B. McSHANE and T. SHEPPARD

D e p a r t m e n t o f M a t e r i a l s , I m p e r i a l C o l l e g e , G B - L o n d o n , SW7 2 B P , G r e a t - B r i t a i n

A b s t r a c t

The consolidation and subsequent heat treatment of an inert gas atomized Al-Li-Zr alloy are discussed. The consolidation process used was hot extrusion. The process parameters investigated include extrusion temperature, solution heat treatment and ageing temperature. Attempts are made to correlate the resulting mechanical properties with the extruded structure observed by transmission electron microscopy.

Introduction

Interest in the Aluminium-Lithium family of alloys continues to grow because o f the exciting possibilities these alloys offer in areas where specific strength and specific modulus properties are important. The potential strengthening mechanism which is available in these alloys, namely the controlled homogeneous precipitation of the metastable, ordered, coherent A13Li (6') phase is well documented. The inherent poor fracture toughness properties due to the propensity for localised slip in this system are also familiar as are the'attempts to improve fracture toughness by the use of various alloying additions. Typically magnesium and copper have been added in order to produce a distribution of A1 CuMg (S) pWase which causes dispersion of slip ana improves resistance to fracture.

In the present investigation the problem of localised slip is attacked by the addition of significant quantities of Zirconium (0.66wt%) a dispersoid forming element which has be n used in aluminium alloys as a graln ref inerl, to promote strengthening9, ' to faci 1 i tate superplastic behaviour3 and as a recrystal 1 isation inhibitor4. In binary A1 -Zr a1 loys the Al Zr dispersoid forms but in the ternary Al-Zr system a composite al[Al f ~ i ~ r ) ] ordered phase is reported5 to precipitate at temperatures

> ^45d°C. The production of the alloy used in this investigation by

atomization, a rapid solidification technique, permits the retention of the Zirconium in supersaturated solid solution. This, in turn, provides the opportunity to study the influence of processing conditions on the subsequent precipitation of the Zirconium bearing phase. Post-extrusion heat treatment to facilitate 6' precipitation both homogeneously and on existing Zirconium dispersoids is investigated and the resulting tensile and fracture toughness properties are determined.

Experimental

The powder was manufactured in an up-draught inert gas atomizer6 using argon and a melt temperature of 950°C. The powder was stored under argon to minimize the particle surface contamination. It was then cold compacted unidirectionally in a hardened steel, stearic acid lubricated ram and die compaction set giving 75mm diameter compacts of 88% Theoretical Density. The compacts were stored under argon in the presence of silica gel

Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:1987337

C3-328 JOURNAL DE PHYSIQUE

and phosphorous pentoxide. Extrusion was carried out on a 5MN hydraulic press. A constant extrusion ratio of 20:l and a constant ram speed of 9mm/sec were used. The billets were heated by induction to various extrusion temperatures ranging from 250°C to 500°C. Typical heating times of 29

-

TABLE I a) As extruded and aged at : 150°C

170°C 190°C

b) Solution soaked at 54O0C/+hr, C.W.Q. and aged at : 150°C 170°C 190°C c) Heated at 420°C/1 hr, C.W.Q. and aged at : 150°C

170°C 190°C

The response was monitored by hardness testing. Tensile test and short rod fracture toughness test specimens were machined from the under-aged, peak-aged and over-aged extrudes. The atomized powder and the consolidated extrudes were examined by scanning and electron microscopy respectively.

Results and Discussion The Atomized Powder

The chemical analysis of the powder used in this investigation is given in Table 11.

TABLE I1

Wt.% L i Zr Fe S i A1

2.54 0.66 <0.05 0.03 Balance The morphology of the powder w s essentially spherical as would be expected from Argon atomised powder6 with very small satel 1 i te particles occasionally adhering to the larger particles. This is a fairly common occurrence in atomization caused by collisions between particles in flight before solidification is complete. The powder is fairly fine and the size distribution of the particles is given in Table 111.

TABLE I11

Response to Heat Treatment

The structure of the as-extruded material is shown in Figures 1 and 2 for extrusion temperatures of 250°C and 400°C respectively. In both cases a fairly well established subgrain structure, due to dynamic recovery, is apparent. This is typical of the structure f rmed in many extruded, rapidly solidified aluminium powder a1 loys7r8. In the higher temperature extrude there is a considerable amount of fine (0.02pm) precipitate ~ r e s e n t This precipitate is probably ZrAl which has

p a r t i c l e s i n p e r i o d i c p a r a l l e l " s t r i n g s " c l o s e l y resembles t h e interphase p r e c i p i t a t i o n o f MZ3C6 a t t h e y / a i n t e r f a c e which ~ o n e ~ c o m b e ~ has r e p o r t e d d u r i n g t h e t r a n s f o r m a t i o n o f chromium s t e e l s . He r e p o r t e d t h a t these " s t r i n g s " o f p r e c i p i t a t e are o f t e n associated w i t h p l a n a r low energy

i n t e r f a c e s . ~ e s l O has a l s o i n v e s t i g a t e d t h e behaviour o f A13Zr p r e c i p i t a t e s d u r i n g h i g h temperature deformation and has found t h a t they can be d i s s o l v e d by moving boundaries i n what he terms s t r a i n induced continuous r e c r y s t a l l i s a t i o n . C l e a r l y t h e n a t u r e o f t h e f o r m a t i o n o f t h e p r e c i p i t a t e shown i n Figure 2 may be complex and w i l l r e q u i r e d e t a i l e d i n v e s t i g a t i o n . I n t h e present i n v e s t i g a t i o n considerable amounts o f t h i s p r e c i p i t a t e were found i n t h e extrudes produced a t temperatures o f 400°C and above. However, t h e r e was no evidence o f t h i s p r e c i p i t a t e i n t h e lower temperature extrudes. O f p a r t i c u l a r i n t e r e s t i s t h e apparent absence o f g r a i n boundary p r e c i p i t a t e s o r indeed t h e 6 p r e c i p i t a t e which has been r e p o r t e d t o l e a d t o reduced f r a c t u r e r o p e r t i e s i n these a l l o y s by t h e process o f g r a i n boundary d u c t i l e f r a c t u r e 1 ?

The response t o t h e v a r i o u s heat treatments i n v e s t i g a t e d i s t y p i f i e d f o r t h e 450°C extrude aged a t 190°C i n t h e hardness curves shown i n F i g u r e 3. As expected t h e value o f peak hardness increases s l i g h t l y (from 157 t o 161) as t h e ageing temperature decreases from 190°C t o 150°C.

However, t h e ageing t i m e r e q u i r e d f o r peak hardness increases from 23hours a t 190°C t o 100 hours a t 150°C. It was t h e r e f o r e decided t o use 190°C as t h e standard ageing temperature i n t h e i n v e s t i g a t i o n because o f t h e r a p i d response a t t h i s temperature and because the peak hardness obtained a t 190°C was o n l y 2-3% l e s s than t h a t obtained a t 150°C. Although t h e peak hardness i s obtained q u i c k l y a t 190°C, no s i g n i f i c a n t overaging occurs before 10 hours. I t i s a l s o obvious t h a t t h e g r e a t e s t response i n t h e 450°C extrude was obtained by ageing t h e as-extruded m a t e r i a l , w i t h o u t a preceding s o l u t i o n soak treatment. The ageing curves f o r t h e as-extruded m a t e r i a l f o r a l l s i x e x t r u s i o n temperatures a r e shown i n Figure 4. C l e a r l y t h e t h r e e h i g h e r temperature extrudes a r e s i g n i f i c a n t l y harder. However, t h e value o f (peak hardness - as extruded hardness) f o r a l l s i x e x t r u s i o n temperatures i s v e r y s i m i l a r ( o f t h e order o f 43 VPN) and thus approximately t h e same amount o f l i t h i u m must have been a v a i l a b l e i n a l l cases f o r p r e c i p i t a t i o n as 6 ' ( A l L i ) . The h i g h e r e x t r u s i o n temperature specimens do have s i g n i f i c a n t ? y higher as-extruded hardness. Any d i f f e r e n c e s i n hardness due t o subgrain s i z e and p e r f e c t i o n v a r i a t i o n s caused by d i f f e r e n t e x t r u s i o n temperatures would tend towards i n c r e a s i n g hardness w i t h decreasing e x t r u s i o n temperature, the reverse o f t h e t r e n d shown i n F i g u r e 4. Therefore t h e h i g h e r as-extruded hardnesses o f t h e h i g h e r temperature extrudes i s most probably due t o t h e presence o f t h e A13Zr p r e c i p i t a t e mentioned e a r l i e r .

The e f f e c t o f i n t r o d u c i n g a 54O0C/*-hour s o l u t i o n soak before ageing a t 190°C i s shown i n Figure 5. A l l o f t h e curves have converged although t h e r e i s a s l i g h t i n d i c a t i o n t h a t t h e lower temperature extrudes e x h i b i t s l i g h t l y g r e a t e r hardness. The value o f (peak-aged hardness - s o l u t i o n soaked hardness) i s s t i l l o f t h e order o f 43 VPN f o r a l l extrudes showing t h a t t h e s o l u t i o n soak has n o t p r o v i d e d any more L i t h i u m f o r p o t e n t i a l p r e c i p i t a t i o n as 6 ' . T h i s i s n o t s u r p r i s i n g as t h e r a p i d s o l i d i f i c a t i o n process ensured r e t e n t i o n o f L i t h i u m and Zirconium i n super-saturated s c l i d s o l u t i o n and t h e r e was no evidence o f p r e c i p i t a t i o n o f any L i t h i u m bearing p r e c i p i t a t e s d u r i n g even t h e low temperature e x t r u s i o n s (when t h i s a l l o y composition i s w i t h i n t h e 2 phase a + 6 f i e l d ) . The increase i n hardness i n t h e low temperature extrudes caused by the 540°C soak p r i o r t o ageing ( i . e . t h e s o l u t i o n soaked hardness i s 10-16 VPN g r e a t e r than t h e as-extruded hardness) can be e x p l a i n e d by t h e p r e c i p i t a t i o n o f t h e A13Zr phase.

F i g u r e 6 i s an e l e c t r o n micrograph o f t h e 350°C extrude which has been subjected t o t h e 540°C s o l u t i o n soak and then water quenched. There i s considerable evidence o f discontinuous p r e c i p i t a t i o n o f AlgZr as

C3-330 JOURNAL DE PHYSIQUE

described by Gayle and vandersande5 who also investigated microstructural evolution in a rapidly solidified Al-Li-Zr alloy. The higher temperature extrudes exhibit a slight decrease in hardness ( = 7-8 VPN) on s o l u t i i ~ soaking and this can be ascribed to static recovery processes. In thest higher temperature extrudes the AlgZr whase is still present in the solution soaked and aged condition in the form of the periodic parallel

"strings" previously mentioned (Figure 7a - 450°C extrude).

In general the 6 ' precipitates nucleated both homogeneous1 y and also on existing A13Zr particles during ageing o f this alloy. Figure 7b is the corresponding Dark Field image o f the micrograph representing the 450°C extrude, solution soaked and aged. Some o f the 6 ' can clearly be seen to have formed "composite" particles by nucleating on A13Zr particles as reportef previously by Gayle and vandersande12 and Gregson and Flower . Further evidence of the presence of "composite" particles present in the solution soaked and aged 250°C extrude are shown in the Bright and Dark Field images in Figures 8 a and 8b. Here the "composite"

phase is present in both the discontinous precipitate form and also as discrete spherical particles.

Part of this investigation involves optimization of the heat treatment for this Al-Li-Zr alloy and s o heat treatments specifically aimed at precipitating the Zr are also being examined. Preliminary results are shown in Figure 9 which depicts the 190°C ageing behaviour of extrudes subjected to a 420°C/1 hr heat treatment. Clearly the hardness of the aged specimens has been reduced when compared t o Figure 5. This is probably due to precipitation of equilibrium 6 (Al-Li) phase during the 420°C treatment although evidence has not yet been obtained. No additional precipitation of Zr was detected after this 1 hour soak and longer times are being investigated.

In conclusion, the higher temperature extrudes in the T 5 condition exhibit hardnesses remarkably similar to those o f the lower temperature extrudes in the T 6 condition, Figure 10.

Mechanical ropert ties

In order to investigate the mechanical properties under-aged, over-age!

(both to 80% of the peak) and peak-aged specimens were prepared in the "T5 condition for the 500, 450 and 400 OC extrudes and in the "T6" condition for the 350, 300 and 250 OC extrudes.

Ultimate tensile strength and yield strength data were obtained and Figure 11 depicts the yield strength as a function of ageing condition.

Both UTS and YS vary typically from 520-560 MPa for the peak-aged condition. The higher temperature extrudes in the T 5 condition are typically 20 MPa weaker than the lower temperature extrudes in the T 6 condition. Gayle and vandersande12 achieved peak-aged strengths o f 500 MPa (UTS) and 450 MPa (YS) for an A1-2.34%Li - 1.07%Zr atomized alloy. The achievement of peak-aged strength in 2 4 hours ageing at 190°C in the alloy presently being investigated can be compared with the time to peak of 1 hour for the Gayle and VanderSande alloy. This acceleration of the ageing behaviour may be due to the greater amount of Zirconium which may accelerate nucleation. The greater strength achieved in the present alloy may be attributed to a greater amount of homogeneous precipitation of 6' 2ue to the increased Lithium content and the reduced Zirconium content.

The % elongation and short rod fracture toughness values are given in Figures 1 2 and 13 respectively. Unfortunately, the material exhibits relatively poor ductility in the heat treated conditions with. peak-aged elongation values o f = 5% compared to = 15% for the as-extruded material.

d u c t i l i t y = 6% elongation. As would be expected from t h e UTS and YS data t h e higher temperature extrudes a r e s l i g h t l y more d u c t i l e (+ 3%

e l o n g a t i o n ) . The f r a c t u r e toughness dependence on e x t r u s i o n temperature i s s l i g h t l y g r e a t e r w i t h peak-aged Kv values o f = 8 MPa Jm f o r t h e higher e x t r u s i o n temperature T5 specimens compared t o values o f = 5 MPa Jm f o r t h e lower e x t r u s i o n temperature T6 specimens. The under ? n @ over-aged specimens show h a r d l y any improvement i n t r a c t u r e toughness i n general.

I t i s however worth n o t i n g t h a t toughness values o f about 10 MPa Jm have been achieved f o r t h e 500°C extrude i n t h e underaged and peak-aged c o n d i t i o n i n d i c a t i n g t h a t o p t i m i s a t i o n o f processing c o n d i t i o n s and heat treatment may y i e l d acceptable f r a c t u r e p r o p e r t i e s .

Concl usions

An A1-2.5XLi-0.6%Zr i n e r t gas atomized a l l o y has been s u c c e s s f u l l y f a b r i c a t e d by h o t e x t r u s i o n . I t was found t h a t extrudes produced a t temperatures o f > 400°C e x h i b i t e d t h e i r maximum mechanical s t r e n g t h i n t h e T5 c o n d i t i o n whereas extrudes produced a t lower temperatures were s t r o n g e s t i n t h e T6 c o n d i t i o n . T h i s dependence on e x t r u s i o n temperature i s thought t o be due t o p r e c i p i t a t i o n o f a Zirconium bearing phase d u r i n g e x t r u s i o n a t t h e higher temperatures. For lower e x t r u s i o n temperatures a subsequent h i g h temperature soak i s r e q u i r e d t o p r e c i p i t a t e t h i s phase b u t we should r e c a l l t h a t any p r e c i p i t a t e formed concomitant w i t h s t r a i n i s l i k e l y t o e x h i b i t a f i n e r morphology.

On ageing this a l l o y a t 190°C peak hardness i s obtained i n o n l y 2$

hours. The 6 ' phase p r e c i p i t a t e s b o t h homogeneously and a l s o on e x i s t i n g A13Zr d i s p e r s o i d s t o g i v e t h e composite [A13(LiZr)] a ' phase.

The heat t r e a t e d a l l o y has a maximum YS o f 560 MPa which i s promising i n v i e w o f t h e r e l a t i v e l y l o w a l l o y i n g a d d i t i o n s o f Z r and L i used. The d u c t i l i t y and f r a c t u r e toughness a r e somewhat d i s a p p o i n t i n g (6% e l o n g a t i o n and maximum Kv = 8 MPa Jm). The e f f e c t o f e x t r u s i o n temperature i s once again e v i d e n t i n t h a t t h e higher e x t r u s i o n temperatures g i v e somewhat b e t t e r d u c t i l i t y and f r a c t u r e toughness.

Acknowledgement :

The p r o v i s i o n o f a research g r a n t by the Aluminum Conpany of America (ALCOA) i s g r a t e f u l l y acknowledged.

Re~eretlces

1. L. Mondolfo, A1 uminium Alloyus, publ . Butterworth, 1976, p.413.

2. T. Ohashi and R. Ichikawa, Met.Trans.A., 12A (1981), p.546.

3. R.H. B r i c k n e l l and J.W. Edington, Met.Trans.A., 10A (1979), p.1257.

4. Aluminium-Properties and Physical M e t a l l u r g y , J.E. Hatch, ed.

(ASM, Metals Park, Ohio) 1984.

5. F.W. Gayle and J.B. VanderSande, S c r i p t a Met., 1984, vo1.18, 473.

6. A. Unal , proc. A1 uminium Technology '86, ed. T.Sheppard, p.673, I n s t . o f Metals, London, 1986.

7. T. Sheppard and H.B. McShane, Powder Met., 1976, 19, 126.

8. T. Sheppard e t a l . , Powder Met., 1983, 26, 10.

9. R.W.K. Honeycombe, " S t e e l s - M i c r o s t r u c t u r e s and P r o p e r t i e s " , p.72, publ. Edward Arnold, 1981.

10. E. Nes, J.Mat.Sci. ( L e t t e r s ) , vo1.13, 2052, 1978.

11. A.K. Vasudevan and R.D. Doherty, Acta.Met., 1987, 35 ( 6 ) , p.1193.

12. F.W. Gayle and J.B. VanderSande, proc. "Rapidly S o l i d i f i e d Powder Aluminium A l l o y s " , ASTM, 1984.

13. P.J. Gregson and H.M. Flower, proc. Aluminium Technology '86, ed. T. Sheppard, I n s t . o f Metals, London, 1986, p.423.

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FIG. 1 250°C extruded structure FIG. 2 400°C extruded structure

- 3 450C extrude aged 190C Fig

100 j . . . ^, . . . . . . . , . . . . I . . . ^, ....

/

1

b.

0 10 10

Ageing Time (hr) Fig 5 S.T 540C/0.5hr, a g e d

160 ^{I ' '} .'....' ^' . . ' . . . . l .

As extruded aged 190C

I . ^{' '} " " I ' ' ' I " " 1

1 oo

0 10

Ageing Time (hr)

FIG. 6. DiscontinuousA1,Zr precipitate in solution soaked extrude

(350°C extrusion temp. ) .

I00

1

1 oO

0 10 ^I LO

Ageing Time (hr)

showing parallel "strings" of showing composite a' A1,Zr ~ r e c i ~ i t a t e . ~ r e c i ~ i t a t e and homoqeneous

FIG. Sa. Solution soaked and peak aged FIG. 8b. As 8a. Dark Field Image.

250°C extrude showing both discontinous and spheroidal composite a' precipitate.

EMRUSION TEMP

250C

Fig 9 S.T 420C/l h r , a g e d Fig 10 Ageing summary

I , . . I .... ¹ . . . I . . L L -

A

X

A .

n ₀ H0 ..

T

EMRUSON TEMP

)

1 0 0 . 0 - . . ^{1 .} . . . , . . . . I ' . . , . . . . ^100, . . . , . . . . I . . . , . . . . I . . " " " I

0 1 0 t Z

0 10 10 10 0 10 10 10

Ageing Time ( h r ) Ageing Time (hr)

160

1 4 5 -

I

1 5 0 . 0 -

1 3 7 . 6 -

' ' " " " I ' ' " " ' " ' ' " " "

8 $ 8 2 '

8 2 m $ T 8 T

$ 0 + d

0 O X %

C) 6, n

JOURNAL DE PHYSIQUE

Fig 12 Elongation vsAgeing

1 8 . 0 -

Ageing condition

EYIRVSON TEMP

35w

EXTRUSION TEMP

sow 250C

Ageing condition

Ageing condition