The cyclic hardening of Al3Mg alloy

Scripta

METALLURGICA

Vol. 18, pp. 981-984, 1984 Printed in the U.S.A.

Pergamon Press Ltd. All rights reserved

THE CYCLIC H A R D E N I N G OF A I - 3 M g ALLOY S° I. Kwun* and M. E. Fine Dept. of M a t e r i a l s Science & E n g i n e e r i n g and M a t e r i a l s Research Center N o r t h w e s t e r n University, Evanston, IL 60201 (Received May (Revised July

31, 1984) 5, 1 9 8 4 )

Introduction It is well k n o w n that pure A1 c y c l i c a l l y hardens not only at room t e m p e r a t u r e but even more at low t e m p e r a t u r e (1,2). The p r e s e n t r e s e a r c h i n v e s t i g a t e d the strain c o n t r o l l e d fatigue b e h a v i o r of Ai-3wt% Mg alloy b o t h at r o o m t e m p e r a t u r e and at liquid n i t r o g e n temperature including t r a n s m i s s i o n e l e c t r o n m i c r o s c o p y of thin foils taken from the specimens. Enormous cyclic h a r d e n i n g was o b s e r v e d in this alloy c o m p a r e d to pure AI. A p o s s i b l e e x p l a n a t i o n of cyclic h a r d e n i n g is p r o p o s e d on the basis of the m i c r o s t r u c t u r a l changes from strain cycling. W h i l e the m i c r o s t r u c t u r e of A I - M g solid s o l u t i o n before and after fatiguing (3,4) and the strain response under c o n t r o l l e d stress of c o m m e r c i a l A I - M g alloy (5) at room temperature had been studied, the effect of Mg a d d i t i o n s in solid solution on strain c o n t r o l l e d fatigue b e h a v i o r had not been studied at r o o m or at low temperature. Experimental The alloy e x a m i n e d in this study was high purity Ai-3wt% Mg (3.01 Mg, 0.006 Si, 0.002 Fe, 0.0003 Mn, 0.003 Cu, 0.003 Zr, 0.003 Ti and b a l a n c e AI) d o n a t e d by A L C O A Technical Center. The specimens of this alloy were s o l u t i o n i z e d at 430°C for 5 hrs., Water q u e n c h e d followed by r e h e a t i n g at 250°C for i0 hrs. and then water quenched. The resulting grain diameters were a p p r o x i m a t e l y 200 ~m. The cyclic s t r e s s - s t r a i n curves were o b t a i n e d by the incremental step test technique (6). In a d d i t i o n constant total strain a m p l i t u d e tests were also done. The low temperature e x p e r i m e n t at 77°K were p e r f o r m e d by immersing the specimens in liquid nitrogen. The cyclic flow stress responses were m e a s u r e d by cycling at various constant total strain amplitudes until failure. Results and D i s c u s s i o n s The m o n o t o n i c and cyclic s t r e s s - s t r a i n curves of AI-3Mg alloy at 298°K and 77°K are shown in Fig. I. The cyclic s t r e s s - s t r a i n curves o b t a i n e d from the saturation stresses of the constant strain a m p l i t u d e tests c o i n c i d e d with those o b t a i n e d by the incremental step test. The cyclic 0.2% o f f s e t y i e l d stresses were 172 M N / m ~ and 236 M N / m 2 at 298°K and 77°K r e s p e c t i v e l y compared to m o n o t o n i c yield stresses of 51 MN/m 2 and 63 M N / m ~. The cyclic h a r d e n i n g response, (Oc-Om), to s a t u r a t i o n in this alloy at given strain a m p l i t u d e s are c o m p a r e d with those for pure A1 in Fig. 2. Here O c is the m a x i m u m cyclic flow stress and Om is the m a x i m u m stress for a given strain amplitude at the end of the first quarter cycle w h i c h was in compression. While pure A1 c y c l i c a l l y hardens slowly, the alloy hardens much more rapidly, e s p e c i a l l y at the low temperature. However, the initial h a r d e n i n g rate in the alloy at 77°K was very close to that at 298°K. This result implies that the initial d e f o r m a t i o n m e c h a n i s m s at 77°K and 298°K are not so different. The m i c r o s t r u c t u r e s after fatigue at 298°K and 77°K with 0.6% of strain amplitude, Fig. 3, 4, show large numbers of d i s l o c a t i o n loops and d i s l o c a t i o n dipoles. The as heated (i.e. before fatigue) specimen, Fig. 5, c o n t a i n e d a very low density of d i s l o c a t i o n s and d i s l o c a t i o n loops. The latter r e s u l t e d from collapse of q u e n c h e d - i n v a c a n c y clusters. On the other hand, after fatigue there were many d i s l o c a t i o n loops, the smallest loops were about 30A ° in diameter. A few of the loops were elongated, some to a length of 0.15 ~m. Similar m i c r o s t r u c t u r a l changes were o b s e r v e d by W a l d r o n (3) who fatigued a similar alloy under c o n s t a n t stress at room temperature. *Permanent address is Dept. of M e t a l l u r g i c a l Engineering, Sungbuk-Ku, Seoul, Korea.

0036-9748/84 Copyright (c) 1 9 8 4

981 $3.00 + Pergamon

Korea University,

.00 Press

Ltd.

Anam-Dong,

982

CYCLIC

HARDENING

OF

AI-3Mg

Vol,

18,

No.

9

In the present work at a strain amplitude of 0.6%, the ordinary dislocation density also increased but not nearly as much as the dislocation loop density. Occasionally areas with dislocation cells were observed in the 298°K fatigued specimens but not at 77°K. The density of dislocation loops and dipoles was higher for the specimens fatigued at 77°K than those fatigued at 298°K. The specimens fatigued at 0.3% strain amplitude also showed a similar increase in dislocation density but not as much as at 0.6% strain amplitude (compare Fig. 3 and Fig. 4). However, dislocation cells were not detected a t a l l at this lower strain amplitude. In contrast to the AI-3Mg alloy after fatigue cycling at 298°K, pure A1 showed a well defined dislocation cell structure (7). The alloy contained very few cells and then only at higher strain amplitude at room temperature. The dislocation density, mostly dislocation tangles, was much higher in pure A1 at 77°K (7) than in the alloy. This difference originates from the well known fact that dissolved Mg in A1 makes cross slip more difficult (8). The above observations suggest that the cyclic hardening in this alloy has resulted from the increased dislocation loop density and, to a lesser extent, the increase in general dislocation density. During initial fatigue cycling, the dislocation loops might form by the interaction between the shuttling mobile dislocations and the dislocation loops already present as well as the forest dislocations. The Mg atoms and vacancies are suggested to form complexes (9) and pin the mobile dislocations (i0) facilitating loop formation and increasing the friction stress. As the deformation proceeds, the mobility of the dislocations decreases due to interaction with the increased dislocation debris density resulting in cyclic hardening as well as increased vacancy production and Mg-vacancy complexes. The effect of temperature is suggested to be through its effect on thermally activated c r o s s slip of screw dislocations. In pure AI, tangles thereby result at 77°K rather than cells as at 298°K. In the AI-3% Mg alloy, the loop production increases on lowering the temperature because cross slip becomes more difficult with less chance of mobile dislocations avoiding existing loops and Mg-atom clusters by the cross slip mechanism. Because the solubility of Mg in A1 is about 1% at room temperature, it was necessary to check whether part of the cyclic hardening might result from either precipitates or preprecipitates (ii). One specimen was prepared with the same heat treatment as the others but furnace cooled from 250°C instead of being quenched. The monotonic and cyclic response of this specimen to straining was essentially the same as for the quenched specimens proving precipitation during strain cycling had not played a role in the cyclic hardening. Acknowledgment One of the authors (S.I.K.) would like to thank the Korean Government for their financial support during his leave of absence at Northwestern University. The use of facilities of Northwestern University's Materials Research Center sponsored under NSF-MRL Grant DMR-8216972 is acknowledged. References i. 2. 3. 4. 5. 6. 7. 8. 9. i0. ii.

P. K. Liaw, Ph.D. thesis, Northwestern University, 1980. C. E. Feltner, Phil. Mag. 28, 1229 (1965). G. W. J. Waldron, Acta Metl. 13, 897 (1965). B. Ramaswami and T. W. F. Lau, Mat. Sci. & Engr. 46, 221 (1980). C. Laird and A. R. Krause, pp. 691-715 in Inelastic Behavior of Solids, eds., M. F. Kanninen, W. F. Adler, A. R. Rosenfield and R. I. Jaffee, McGraw Hill, 1970. D. T. Raske and JoDean Morrow, ASTM STP 465, i, 1969. P. K. Liaw, M. E. Fine, M. Kiritani and S. Ono, Scripta Metl. ii, 1151 (1977). J. T. McGrath and G. W. J. Waldron, Phil. Mag. 9, 249 (1964). C. Panseri, T. Federighi and S. Ceresara, Trans AIME 227, 1122 (1963). T. Tabata, H. Fujita and N. Nakajima, Acta Met. 28, 735 (1980). T. Sato, Y. Kojima and T. Takahashi, Met. Trans. A. 13A, 1373 (1982).

Vol.

18, No.

9

CYCLIC HARDENING

OF AI-3Mg

983

2S(~ -

'//

200

/

•

CYCLIC

I

®

/ 15G

/"iiiill

N

/

Z

Ig u') IOC u~ bJ n~ l'Oq

o. 5

;/

....

0.3 STRAIN

Fig. i. Monotonic and cyclic stress-strain curve obtained by incremental step test and constant strain amplitude test at 298°K and 77°K. Each curve represents two or more tests.

..~

0.45 (%)

(16

2OO . . . . 77"K 298"K

~'P"

•

~ ~/2-0.G%

AI*3Hg ALLOY

Z

p/

/

~

p/S

I

ff 5O

•

~ . . . A , / 2 • 0.3 %

/s

s SS s'ls Sr /

I0 DO DO0 NO. OF REVERSALS (ZN)

Fig. 2.

DOO0

Cyclic hardening behavior before saturation stress at constant strain amplitudes. °c: cyclic flow stress m: monotonic flow stress

984

CYCLIC

HARDENING

OF AI-3Mg

(a) Fig.

Vol.

(b)

3. Microstructures after fatigue at 298°K. (a) 0.6% strain amplitude (b) 0.3% strain amplitude

(a)

(b)

Fig. 4. Microstructures after fatigue at 77°K. (a) 0.6% strain amplitude (b) 0.3% strain amplitude

Fig.

(weak beam)

5. Microstructure of AI-3Mg alloy before fatigue.

18, No.

9