GRAIN REFINEMENT AND SUPERPLASTICITY IN A LITHIUM-CONTAINING Al-Mg ALLOY BY THERMOMECHANICAL PROCESSING

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GRAIN REFINEMENT AND SUPERPLASTICITY IN A LITHIUM-CONTAINING Al-Mg ALLOY BY

THERMOMECHANICAL PROCESSING

S. Hales, S. Oster, B. Sanchez, T. Mcnelley

To cite this version:

S. Hales, S. Oster, B. Sanchez, T. Mcnelley. GRAIN REFINEMENT AND SUPERPLASTICITY IN A LITHIUM-CONTAINING Al-Mg ALLOY BY THERMOMECHANICAL PROCESSING. Journal de Physique Colloques, 1987, 48 (C3), pp.C3-285-C3-291. �10.1051/jphyscol:1987332�. �jpa-00226563�

J O U R N A L D E P H Y S I Q U E

Colloque C3, supplkment au n09, Tome 48, septembre 1987

GRAIN REFINEMENT AND SUPERPLASTICITY IN A LITHIUM-CONTAINING A1-Mg ALLOY BY THERMOMECHANICAL PROCESSING

S.J. H A L E S , S.B. O S T E R , B.W. S A N C H E Z and T.R. M c N E L L E Y M a t e r i a l s E n g i n e e r i n g G r o u p , D e p a r t m e n t of M e c h a n i c a l E n g i n e e r i n g , Code 6 9 , Naval P o s t g r a d u a t e S c h o o l , M o n t e r e y , CA 9 3 9 4 3 - 5 0 0 0 , U . S . A .

ABSTRACT The refined microstructures and superplastic properties resulting from controlled thermomechanical processing of an A1-8Mg-0.5Li-0.23Zr alloy were evaluated.

Rolling between the solvus temperatures for Ng and Li in the alloy allowed for grain refinement by precipitation of 6 (AlgMg5) during deformation and subsequent precipitation of G1(Al3Li) on cooling. Increasing the rolling strain enhance the su erplastic ductility of the alloy at 573K in the strain-rate regime of 10-q-10-2 S

- '

. Elongations in excess of 500 pct. , without cavitation, and a corresponding strain-rate sensitivity coefficient of approximately 0.5, were obtained. TEM

investigations of the microstructural characteristics responsible for the mechanical behavior revealed that a more uniformly refined grain structure (3-5um) evolved by continuous recrystallization (CRX) in material experiencing the larger rolling strain.

INTRODUCTION Al alloys containing Wg and Li exhibit both solid solution and precipitation hardening but are also prone to the formation of brittle phases on grain boundaries such as A12MgLi (1). Grain refinement may be expected to reduce the susceptibility of these alloys to such embrittlement. Aging studies conducted on Al-Mg-Li alloys (1-3) have revealed that Mg precipitates independently of Li and follows the sequence that occurs in an Al-Mg system. Li co-precipitates as 6' (an intermetallic based on Al3Li) which is a metastable, ordered phase with an LIZ-type structure (1 5). This precipitation reaction is unaffected by the presence of Mg in solid solution (3). The 6'- phase forms as coherent, spheroidal precipitates unifonly distributed within the matrix usually

-

^7Onm in size after peak-aging treatments (1,3). In AL-Li-X alloys the 6' is primarily responsible for strengthening (2), thus it may be anticipated that Li additions to AL-Mg alloys create the potential for precipitation hardening treatments in otherwise non-heat treatable material.

Research on Al-Mg-X (X = Zr,Mn,Cu) alloys containing 8-10 wt.pct.Mg (5-9) has revealed that the materials can behave in a superplastic manner in an essentially non-recrystallized condition, i.e., without recrystallization heat treatments prior to deformation. The themechanical processing ('IMP) used has, as its central feature, mechanical mrking by rolling at 573K and since this temperature is below the @,-solvus , the intermetallic 6 -phase (AQMgg) precipitates concurrently with the generation of dislocations (5). Superplastic elongations in excess of 500 pct. have been obtained at 573K and strain rates of 2-5 X 10-3

s - ~

(7,8). The corresponding homologous temperature, 0.7 Tm, is considerably lower than generally reported for Al alloys and the strain-rate higher by an order of magnitude (10,ll).

It is also possible to correlate apparent strain hardening during superplastic flow with grain growth and a decrease in the value of the strain-rate sensitivity coefficient, m (= dlnoldlni) (8).

The microstructural transformation from the as-rolled, heavily deformed structure to a refined (sublgrain structure, capable of sustaining superplastic deformation, has been interpreted in terms of CRX occuring both preceding and concurrent with defamation (9,12). 'Ihis mechanism is characterized by the

development of moderate-angle boundaries by extensive dislocation rearrangement and the absence of high-angle boundary migration (1 3,14)

.

Especially noteworthy is the observation that these warm-rolled alloys do not cavitate during superplastic flow at 573K. It is believed that this is a result of the highly refined grain structure

(2- 5 retained during deformation, and the relatively low temperature gnployed (7)

.

Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:1987332

C3-286 JOURNAL DE PHYSIQUE

Current research i s focussing on the effect of the t o t a l s t r a i n introduced into the material during warm rolling and the effects of Li additions t o these alloys.

The processing parameters that can be controlled during rolling are temperature, reduction per pass, reheat time between passes and t o t a l rolling strain. Since CRX r e l i e s on extensive rearrangement of dislocations, it might be expected that an increase i n dislocation density, created by increasing the rolling s t r a i n , may enhance the development of a refined microstructure by t h i s mechanism. Finally, the

6

'

solvus temperature for a 0.5 wt.pct .Li addition w i l l be

-

400K (4,16) and the @

s o l w s temperatire for a 8 wt.pct.Mg addition

-

590K (5). Thus, grain refinement during deformation m y be achieved via CRX, as observed in A1-Mg-X alloys, and precipitation hardening may then be exploited by subsequent lower temperature a p , i n ~ treataents.

[?X'ERIIENTAL PROCEDURE 'Ihe nominal camposition 9f the alloy in t h i s investigation was A1-8Mg-0.5Li-0.23Zr (wt.pct.) and details of the 'IMP used have been described previously (6). The essential features are solution treatment and hot forging a t 713K, quenching t o ambient temperature, followed by reheating and rolling a t 573K.

The multiple pass rolling was conducted using constant reductions per pass of lm, corresponding t o an approximate reduction of 4% on the f i r s t pass and the reheat time between passes kept constant a t 4 minutes. One batch of material was subject t o a f i n a l reduction of 85% (€true = 1.9) corresponding to a reduction on the l a s t pass of 24%

.

^A second batch of material was processed to a f i n a l reduction of 92%

(ztme

T

2.6) corresponding t o a reduction ^on the l a s t pass of 37%. Hereafter the

processing to cr = 1.9 w i l l be referred t o as 'IMP I and t o c, = 2.6 as 'IMP 11.

Elevated temperature mechanical properties were evaluated using samples machined from the as-rolled sheet with tensile axes parallel t o the rolling direction. Simulated superplastic forming a t 573K was accomplished utilizing constant crosshead speeds and the data obtained corrected t o compensate for the decrease in true strain-rate with increasing s t r a i n (8). The data were reduced t o t n l e s t r e s s and ductility versus s t r a i k r a t e plots for the various i n i t i a l s t r a i n rates. Transmission electron microscopy (TIN) investigations were conducted on specimens extracted from material such that f o i l normals were parallel t o the sheet normal direction.

RESULTS Ihe steady increase in dislocation density with increased rolling s t r a i n produces an extremely distorted structure and, thus, microstructural comparison of the material a f t e r different strains i s impractical. The rolling temperature (573K) i s such that gradual precipitation of the @-phase concurrently with dislocation generation occurs while the Li in the alloy remains i n solid solution. The as-rolled microstructure does not represent the condition of the alloy a t the comnencement of elevated temperature testing because the material experiences the equivalent of a 45-60 minute s t a t i c anneal as the furnace equilibrates a t the t e s t temperature. Fig. 1 shows the condition of the microstructure a t t h i s point. The 'IMP I material, Fig. l ( a ) , exhibits a somewhat banded structure with the 0-phase distributed in adjacent areas of high and low volume fractions of precipitates.

F I G 1. The microstructure after a 45-60 minute s t a t i c anneal showing well-defined

boundaries i n ( a ) TMP I and ( b ) TMP 11w i t h a banded structure discernible i n TMP I.

direction are also apparent. In comparison, the 'IMP I1 material, Fig. l(b),

exhibits a much more uniform distribution of 6 precipitates and the grain structure appears more homogeneous. The size and morphology of the precipitates is similar for the two conditions, the B being heavily faulted, approximately equiazed inshape and 0.5-2.0m is size.

A closer examination of the struc,ture of the matrix reveals the effect of the differing 6 distribution. The micrographs in Fig. 2 show representative areas from ^' the 'IMP I material. Fig. 2(a) shows an example of a 0 -rich region which exhibits 6 precipitates interlaced with well-defined boundaries of apparent moderateangle as well as a distinct absence of individual dislocations. 'Ihis is to be compared with Fig. 2(b) hich shows a 6 -lean region where it can be seen that the microstructure approximates to a recovered substructure with individual dislocations present as regular arrays in low-angle boundaries. In both micrographs it is evident that the

6'-phase has precipitated as a uniform distribution of - 2 5 m particles. It may be- inferred that air cooling from the testing temperature occurred at a sufficiently slow rate in this case to allow precipitation to occur.

The microstructure typical of the 'IMP I1 material following the static anneal can be seen in Fig. 3. ?he uniform distribution of 6 precipitates has produced a structure consisting of well-defined, moderateangle boundaries throughout the mic'rostructure. The apparent grain size is 1-3 l.r~~ with 0 residing at grain

boundaries and is very similar to that observed in Fig. 2(a). The 6'-phase is not observed in this sample which may reflect on the cooling rate after deformation.

Surmnarizing, the 'IMP I1 material has a fine grain size which is of the order of the spacing between the 6 precipitates. In the 'IMP I material this is also the case in the 6 -rich regions, but a recovered structure consisting of predominantly lowangle

boundaries exists in the B-lean regions.

of m determined to be

-

^0.5. In contrast, the 'IMP I material exhibits only moderate

8 o versus In E m e is not sigmoidal and a t y p i c a l m i c r o s t r u c t u r e i n TMP II. 'IMP 11 material is considerably lower (25-

C3-288 J O U R N A L DE PHYSIQUE

1 6 ~ 1f3 16* 40%) suggesting that the larger rolling

600-

,

^{I I 5 I} "'"I l 1 + strain has produced a softer material which, in turn, is capable of greater

-

ductility.

Microstructures typifying the condition of

-

the two materials at the completion of deformation at a strain rate of 6.7 X 10-3

-

^S-1 are presented in Fig. 5. Figs. 5(a) and 5(b) show an example of a 6 -rich and

- ^e

-lean region respectively in the 'lMP I material. The apparent grain size in Fig.

AI-B ~q-0.5 ~i-0.23 zr

-

⁵ (a) is 1-3 p~ whereas in Fig. 5(b) it is

warn r o l w 01 3-5 p suggesting that larger grains

573 K ( 3 0 0 . ~ 1 develop in 8 -lean regions during

O

-

_O-TMPI I R , = I . ~ I deformation. In comparison with Fig. 2(a) ,

0 - T M P E I R , - Z . ~ I Fig. 5(a) shows that little change has

occurred during straining. Tne 6-phase has not changed in size or morphology and is

-

present at grain boundary triple points.

In contrast to these micrographs a

-

comparison of Fig. 5(b) with Fig. 2(b)

-

reveals that the lowangle boundaries present at the onset of deformation have

TEST TEMPERATURE transformed into high-angle boundaries

D a 573 K (300°c) pinned by 6 particles. The relatively

k lo1

- -

wider spacing of the dictates .that the

I I * ^{a 8} *.*rn#l

-

grain size is correspondingly larger,

10- 1 c i 3 16' lo' 3-5 p. It may be inferred from these

STRAIN RATE C , S-' observations that wen though the 6 phase

FIG 4 . M e c h a n i c a l t e s t d a e a f o r TMP I is present as relatively coarse and TMP I I . D u c t i l i t y i s much h i g h e r precipitates it still aids in the

w h i l e s t r e n g t h i s l o w e r f o r TMP I I . retardation of microstructural coarsening during deformation. The microstructure in Fig. 6 shows the ?MP I1 material after deformation to fracture at the same strain rate and it is apparent that the same microstructural features seen in Fig. 5(a) are observed. Tne grains are essentially dislocation free, although some grains contain subgrain boundaries in which individual dislocations are discernible. The overall low dislocation density in grain interiors suggests that the material. exhibiting the peak ductility (556 pct. elongation) has deformed by grain boundary sliding (GBS).

A similar conclusion can thus be drawn from the 6 -rich regions in 'IMP I material observed in Fig. 5(a).

The curves in Fig. 7 show that the stress for a given strain-rate increases with strain during simulated superplastic forming for both material conditions.

-rich r e g i o n and (b) a B - l e a n r e g i o n . N o t e the d i f f e r e n c e i n a p p a r e n t g r a i n - s i z e .

micrographs in Fig. 8 show that microstruc- tural coarsening has occurred to a greater extent at lower strain rates than at higher strain-rates. The lmer strain-rate has allowed grain growth such that the grain- size is n m

-

^{5p1, in ?MP} I material, Fig.

8 (a) , and 'LMP I1 material, Fig. 8 (b)

.

^The

plots in Fig. 7 show that that extent of strain hardening is greater at 10-4 s-1 strain-rate, and thus the effect may be correlated with microstructural coarsening

" during deformation. Again, the uniform distribution of 6' is observed within grains and the presence of 6 at grain boundaries is noted. The 6' precipitates can be seen more clearly in the dark field micrograph in Fig. 9. The area is the same as that presented in Fig. 8(b) and reveals

11 s h o w i n g a v e r y s i m i l a r s t r u c t u r e . a 6' precipitate size of 10-25nm.

DISCUSSION Utilizing the ?MP of this study it has been shown that the equilibrium volume fraction of 8 is attained after an approximate true strain of 1.0 is achieved during rolling at 573K (5). Thus, the rolling strains achieved in this study ensure that the maximum amount of 0 is chlt of solution at the initiation of the elevated temperature testing. The presence of the 6 controls microstructural coarsening such that a stable structure capable of superplastic elongations is obtained. The transformation of the microstructure from the as-rolled condition to that at the onset of deformation has previously been interpreted in terms of CRX (9). There is no evidence in this study for the formation of new grains by nucleation and growth mechanisms pertaining to discontinuous recrystallization. This supports the claim that microstructural evolution occurs via CRX such that the material exhibits a fine-grain superplastic response at 573K. The precise role of in affecting grain refinement is uncertain but it may be assumed that even though the precipitates are relatively coarse they are capable of pinning moderate-angle boundaries. The resultant grain size is thus of the order of the interparticle spacing. It has been observed previously that AlgZr mrticles pin grain boundaries (1 0,ll) tut in this alloy this would not appear to be the case since they are not present in sufficient numbers to produce the grain sizes seen. The development of a finer microstructure by the extensive rearrangement of dislocations in the 'IMP I1 material is consistent if one considers that the larger rolling strain, with a higher rolling strain-rate on final pass, provides a greater dislocation density and further homogenization of the 6 distribution. It has been shown (17) that the presence of coarse, second phase precipitates can increase the stored energy attained from 2IP through additional dislocation generation during the f ina

1 I 1 I r I I I ~ I I

(a) AI-,kg-0.5 ~ i - 0 2 3 Zr

I T W I (e, = 1.9)

TEST TEMPERATURE 573 K (300.C)

10' - -

I 1 I 1 1 1 1 1 1 1 1 I

16' 6' 6'

STRAIN RATE i . T'

F I G 7 . M e c h a n i c a l t e s t d a t a f o r ( a ) TMP I and ( b ) TMP I I . T h e m v a l u e i s l a r g e s t a t s m a l l s t r a i n s and d e c r e a s e s w i t h d e f o r m a t i o n d u e t o m i c r o s t r u c t u r a l c o a r s e n i n g .

C3-290 JOURNAL DE PHYSIQUE

TMP I I . The s t r u c t u r e s h a v e c o a r s e n e d d u r i n g d e f o r m a t i o n t o a s i m i l a r g r a i n - s i z e . manifests itself in the development of a structure containing a larger area fraction of moderate-angle boundaries produced by the CRX mechanism. In addition, a more uniform distribution of 6 leads to a more hogeneous microstructure with respect to grain size distribution.

In the 'IMP I material the existence of @ -lean regions leads to the development of only a recovered structure in these areas. However, in the @-rich regions, where the @ precipitates have aided the CRX process, a microstructure consistent with a 2- 5~ grain structure has developed. During elevated temperature deformation it is to be expected that the fi -lean regions deform by a solute drag-type mechanism where- as the @ -rich regions by GBS (18). The mechanical property data for 'IMP I material reveals an m value of

-

0.3 which is consistent with deformation by solute drag. The microstructural data for material which has attained peak ductility reveals a consi-

derable dislocation density within constituent grains to support this hypothesis.

However, in the 'IMP I1 material the extent of fi -lean regions is sufficiently small such that the microstructure is one of continuously recrystallized grains.

The mechanical property data for this material reveals an m value of -0.5 with a considerably lower flow stress and a shift of peak ductility to higher strain-rates.

The microstructure shows a marked absence of dislocations in grain interiors suggesting that GBS is predminantly responsible for deformation.

In comparing the microstructures before and after elevated temperature testing it is apparent that structural coarsening has occured in both material conditions during deformation. Analysis of the mechanical property data leads to the correlation of strain hardening with microstructural coarsening (8). If grain growth occured at similar rates in IMP I and 1MP I1 material then it may be inferred

ced deformation with a coarser

(18,19) suggest that material starting with a finer grai~size will be weaker and able to sustain grain coarsening longer before dislocation processes intervene and limit attainable ductility. The curves

indicating strain hardening show that the TMP I1 material is more susceptible to this effect which reflects the fineness of the structure at the onset of deformation. A rise in flow stress results particularly at lower strain rates and this is related to microstructural coarsening. The results may be interpreted interms of the ?MP I material deforming by a dislocation gener- ation and recovery process consistent with a solute drag mechanism and an m value of

-

^{0.3. The IMP} I1 material exhibits a

FIG 9.Dark f i e l d TEM m i c r o g r a p h showing fine-grain superplastic response with an t h e d i s t r i b u t i o n o f t h e 6'-phase. effective grait?-siZe of 3-!im and an m

value of

-

0.5. The larger rolling strain has allwed the structure to evolve by CRX such that it deforms by diffusional mechanisms such as

as.

The difference in microstructural evolution between ?MP I and 'IMP I1 material can be attributed to the distribution of the 6 phase and its ability to impede derate-angle boundary migration.

In comparison with Al-Mg-X alloys, processed in the same manner (5-9), the elevated temperature mechanical properties are apparently unaffected by the presence of 0.5 wt .pct.Li. h i s is accaunted for by considering that the rolling and testing temperature is above the 6' solvus (4,16), thus, the Li remains in solid solution.

On subsequent slow cooling from the test temperature the 6' phase precipitates, independently of the Mg in the alloy, to a fine dispersion. Thus, there is great potential for decreasing the density of Al-Mg-X alloys and also increasing the ambient temperature mechanical properties using peak aging treatments dter deformation.

CONCLUSIONS

1 Increasing the rolling from 1.9 to 2.6 during 'IMP enhances the su

P

erplastic dActility of the alloy at 573K between strain-rates of 10-3 and 1 G-

s - ~ .

^{h e}

strain-rate sensitivity coefficient, m is increased from 0.3 to 0.5 and the flow stress decreased by 2540% allowing peak ductilities in excess of 500% to be achieved in this temperature/strain-rate regime.

2)TEM reveals that the increase in rolling strain enhances grain refinement both preceding and concurrent with superplastic deformation. Amore uniformly refined grain structure evolves by CRX dependent on the distribution of the @phase. The grain-size attained (2-5p) facilitates superplastic behavior &en ccnnpared to material experiencing a lesser rolling strain.

3)The grain refinement effect observed is consistent with the mechanical property data obtained. ?he strength and ductility data form elevated temperature are similar to those reported for othe Al-Mg-X alloys processed using the same 'IMP. The Li addition forms 6' precipitates independently of the Mg in the alloy.

A C K N This research activity has been supported ~ ~ by the Naval Air Systems Command with Dr. Laris Sloter as program mitor.

EzlzFmmcES

1 .Noble, B., Harris, S. J. and Harlow, K., ALuminmLithium Alloys 11, Sanders, T.

H. and Starke, E. A., eds., QG-AIME, Warrendale, PA (1984), pp. 65-77.

2.Harris, S. J., Noble, B. and Dinsdale, K., ibid., pp. 219-233.

3.Sanders, T. H., Aluminum-Lithium Alloys, Sanders, T. H. and Starke, E. A., eds., TW-APE. Warrendale. PA (1 981

.

^{vv. 63-67.}

4.williams; D. B., ibid., pp. 87-100:

5,McNelley, T. R. and Garg A., Scripta Metall., Vol. 18 (1984), pp. 917-920.

6.PlcNelley, T. R., et al., &tall. Trans. A, Vol. 17A (1986), pp. 1035-1041.

7.Lee, E.-W., et a?. , &tall. Trans. A, Vol. 17A (1986), pp. 1043-1050.

8.Lee, E.-W. and McNelley, T. R., Mater. Sci. Eng. J., in press.

9.Hales, S. J. and McNelley, T. R., Acta Metall., in press.

lO.Watts, B. M., Stowell, M. J., et al., k t . Sci. J., Vol. 10 (1976), pp. 189-197.

1 1 .Wert, J. A., et al., Metall. Trans. A, Vol. 12A (1981), pp. 1267-1276.

12.Nes, E., Su e lasticit Bauelet, B. and Suery, M., eds., Edition du C.N.R.S., Paris (I 9i8- 4.

13.Haessner, F., Recrystallization of &tallic Materials, Haessner, F., ed., Dr.

Riederer Verlae: QnbH. Stuttgart (1978). _. . vv. 1-10.

14.Nes, E.

,

~ecq%allization &d Gain Growth inWlti-phase and Particle Containing Materials, Hansen, N., Jones, A. R. and Leffers, T., eds., Riso National Laboratory, Roskilde, Denmark (1 980)

,

pp. 85-95.

15.Thompson, G. E. and Noble, B., J. Inst. Metals, Vol. 101 (1973), pp. 111-115.

16.Sigli, C. and Sanchez, J. M., Acta Metall., Vol. 14 (1986), pp. 1021-1028.

17.Sheppard, T., Zaidi, M. A., et al., Microstructural Control of Aluminum Alloys, Chia, E. H. and McQueen, H. J., eds., D E - A N , Warrendale, PA (1986), pp. 19-43.

18.Sherby, 0. D. and Wadsworth, J., Deformation process in^ and Structure, Krauss, G., ed., ASM, Paetals Park, OH (1984), pp. 355-390.

19.Gifkins, R. C., Superplastic Forming of Structural Alloys, Paton, N. E. and Hamilton, C. H., eds., IMS-AIME, Warrendale, PA (1982), pp. 3-26.