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Correlation of microstructure with tensile deformation behaviour in dispersion strengthened aluminum Surana, Narendra Singh 1967

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CORRELATION OF MICROSTRUCTURE WITH TENSILE DEFORMATION BEHAVIOUR IN DISPERSION STRENGTHENED ALUMINUM BY NARENDRA SINGH SURANA B. Sc. (Met-Engg.), Banaras, H. Un i v e r s i t y , 1965 A THESIS.SUBMITTED IN PARTIAL FULFILMENT OF THE REQUIREMENTS FOR THE•DEGREE OF MASTER OF APPLIED SCIENCE i n the Department of METALLURGY We accept t h i s thesis as conforming to the standard required from candidates for the degree of MASTER OF APPLIED SCIENCE. THE UNIVERSITY OF BRITISH COLUMBIA May, 1967 In p r e s e n t i n g t h i s t h e s i s i n p a r t i a l f u l f i l m e n t o f t h e r e q u i r e m e n t s f o r an advanced d e g r e e a t the U n i v e r s i t y o f B r i t i s h C o l u m b i a , I a g r e e t h a t t h e L i b r a r y s h a l l make i t f r e e l y a v a i l a b l e f o r r e f e r e n c e and s t u d y . I f u r t h e r a g r e e t h a t p e r m i s s i o n f o r e x  t e n s i v e c o p y i n g o f t h i s t h e s i s f o r s c h o l a r l y p u r p o s e s may be g r a n by the Head o f my D e p a r t m e n t o r by h i s r e p r e s e n t a t i v e s . I t i s u n d e r s t o o d t h a t c o p y i n g o r p u b l i c a t i o n o f t h i s t h e s i s f o r f i n a n  c i a l g a i n s h a l l no t be a l l o w e d w i t h o u t my w r i t t e n p e r m i s s i o n . D e p a r t m e n t o f M e t a l l u r g y  The U n i v e r s i t y o f B r i t i s h C o l u m b i a V a n c o u v e r 8, Canada Date M a y 5, 1967 ABSTRACT The t e n s i l e behaviour of an A l - l O S b a l l o y , conta ining a d i s p e r s i o n of the i n t e r m e t a l l i c compound A l S b , has been compared w i t h that of an overaged Al-4Cu a l l o y i n the temperature range of -100 to 200°C. As an a i d to the i n t e r p r e t a t i o n of the a l l o y r e s u l t s , pure a l  uminum has a l s o been examined, over the range of gra in s ize represented by the two a l l o y s . A coarse d i s p e r s i o n of AlSb was not found to increase the y i e l d strength of aluminum f o r a given i n i t i a l g r a i n for subgrain) s i z e . •However, the d i s p e r s i o n markedly increased .the r e s i s t a n c e of aluminum to recovery and r e c r y s t a l l i s a t i o n , at the same time i n c r e a s i n g the true ul t imate s t rength and the uniform e l o n g a t i o n . The observed t e n s i l e p r o p e r t i e s of pure aluminum have been explained by a c e l l formation and migrat ion -argument-. The t e n s i l e deformation behaviour of A l - l O S b and A l - U C u has been explained i n the l i g h t of current theor ies and as an extension of the i n t e r p r e t a t i o n of the behaviour of pure aluminum. ACKNOWLEDGEMENTS The author wishes to express h i s s incere g r a t i t u d e , ' to D r . J . A . Lund, f o r h i s advice and a s s i s t a n c e , through the preparat ion of t h i s t h e s i s . He a l s o wishes to express h i s thanks to D r . R. Warda f o r h i s h e l p f u l suggestions during t h i s work. The author wishes to thank the members of the f a c u l t y and f e l l o w graduate students of the Department of Meta l lurgy f o r many h e l p f u l d i s c u s s i o n s . S p e c i a l thanks are extended to Messers . A . L a c i s and P. M u s i l f o r t h e i r ass is tance i n metallography, and to the t e c h n i c a l s t a f f of the department f o r t h e i r va luable a s s i s t a n c e . F i n a n c i a l ass is tance provided by the Defence Research Board and N a t i o n a l Research C o u n c i l i s g r a t e f u l l y acknowledged. TABLE OF CONTENTS Page I. INTRODUCTION 1 Dispersion Strengthening Approaches 1 • Scope of the Present Investigation h II. EXPERIMENTAL 7 Materials 7 Melting and Solidification . 7 Fabrication 11 Tensile Specimens 11 Tensile Tests . 12 Metallography 13 Measurement of Grain-Size Ik- III. OBSERVATIONS AND RESULTS 17 Metallographic Observations 17 Tensile Data-General ' 32 Yield Stress of Polycrystalline aluminum kO Ultimate Tensile Strength h2 Percent Elongation h5 Work Hardening Rates hi Discontinuous Yielding h9 IV. DISCUSSION . 53 A. Pure Aluminum 53 (a) Deformation Behaviour 53 (b) Structure of 200°C Extruded and Annealed Aluminum 5^  (c) Tensile Behaviour 55 (i) Maximum true stress 55 ( i i ) Work hardening behaviour. 62 ( i i i ) Yield Stress 63 (iv) Percent uniform elongation 63 B. Dispersion Strengthened Aluminum Alloys 6h (a) Deformation and Recovery Processes 6h (b) Structure of Al-lOSb extrusions 6j • (c) Tensile Behaviour of Al-lOSb 6j (d) Tensile Behaviour of overaged Al-i+Cu 69 (e) Tensile Behaviour of Solution Treated Al-l+Cu at room temperature 71 Table.of Contents (Cont.) Page C. E f f e c t of S t r a i n Rate on Strength Properties at Room Temperature 71 V . CONCLUSIONS . ... 73 VI. BIBLIOGRAPHY ' 76 LIST OF FIGURES Page 1. • Al-Sb phase .diagram 5 2. Sketch of the shotting equipment used 10 3. Sketch of.a rod shaped grain 15 h. Sketch of possible chords i n a semi-circle ... 19 5. Transmission electron micrograph of 200°C extruded and annealed aluminum from shot • . . . 19 6. Transmission electron.micrograph of 200°C extruded and annealed aluminum from shot 19 7. Transmission electron micrograph of 200°C extruded and annealed aluminum from. shot 20 8. Transmission electron micrograph of 200°C extruded and annealed aluminum from shot 20 9. Transmission electron micrograph of 200°C extruded and annealed aluminum- from cast b i l l e t 22 10. Transmission electron micrograph.of 200°C extruded and annealed aluminum from cast b i l l e t ' 22 11. Transmission electron micrograph of 200°C extruded and annealed aluminum; from . cast b i l l e t 23 12. Transmission electron ;micrograph of 200°C extruded and annealed aluminum from cast b i l l e t 23 13. O p t i c a l micrograph of 300°C extruded and annealed aluminum . . . 24 ik. O p t i c a l micrograph of 300°C extruded and annealed. aluminum 24 15. O p t i c a l micrograph of 400°C extruded and annealed aluminum 26 16. Op t i c a l micrograph of 400°C extruded and annealed aluminum 26 17. O p t i c a l micrograph of 400°C extruded and annealed aluminum • 27 L i s t of Figures (Con't) P a g e 1 8 . Transmission electron micrograph of Al-lOSb extruded and annealed at 300°C ; . . 27 19* Transmission electron micrograph of Al-lOSb extruded and annealed at 300°C .. .. 28 20. Transmission electron micrograph of Al-lOSb extruded and annealed at 300°C 2 8 21. Transmission, electron micrograph of Al-lOSb extruded and annealed at 300°0 ...... 29 22. Transmission electron micrograph of Al-lOSb extruded and annealed at 300°C. Shoving AlSb p a r t i c l e s 29 23. Transmission electron .micrograph ofAl-10Sb extruded and annealed at.300°C; Showing AlSb p a r t i c l e s 30 24. Replica micrograph.showing the si z e of AlSb p a r t i c l e s , compared to Figure 25 30 2 5 . Replica micrograph showing the si z e of AlSb p a r t i c l e s , compared to Figure .24 31 2 6 . O p t i c a l micrograph of overaged Al-^Cu showing grain s i z e 31 27. . O p t i c a l micrograph of overaged Al-l+Cu showing f i n e nature of CuAl2 i n r e l a t i o n to grain s i z e 33 28. O p t i c a l micrograph of s o l u t i o n treated Al-4Cu showing grain s i z e •••• ........ .. 33 2 9 . V a r i a t i o n df y i e l d strength with temperature for d i f f e r e n t grain sized aluminum. 41 - 1 / 2 30. V a r i a t i o n of y i e l d strength with (grain s i z e ) - for aluminum at d i f f e r e n t t e s t temperatures h3 31. , V a r i a t i o n of ultimate t e n s i l e strength with temperature for aluminum,and a l l o y s ' 44 3 2 . V a r i a t i o n of percent uniform elongation with temperature for aluminum and a l l o y s . . . . k6 33. A discontinuous y i e l d i n g curve 51 L i s t of Figures (Cont.) Page 3k. True stress-.strain. curve f o r aluminum and a l l o y s at -100°C (0.19Tm) .. 56 35. True s t r e s s - s t r a i n curve for aluminum and al l o y s at 20°C (0.32Tm) .... 57 36. T r u e - s t r e s s - s t r a i n curve for aluminum and a l l o y s at 100 9 C (O.ltOTm) 58 , 37. True s t r e s s - s t r a i n curve f o r aluminum and al l o y s at 200°C (0.51Tm) 59 LIST OF TABLES Page 1. Analysis of aluminum used 8 2. Analysis of antimony used •••• 8 3. Grain s i z e of materials investigated 18 k. Tensile data for 2.7 p aluminum 3^  5. T e n s i l e data for 3.5 unr aluminum •• 35 6. Tensile data for 10 ym aluminum 36 7. Tensile data for 98 urn aluminum 37 8. Tensile data for Al-lOSb . 38 • 9. T e n s i l e data for overaged A1ThCu 39 10. Tensile data for s o l u t i o n treated Al-l+Cu 39 11.. Work hardening exponent (n) , hQ 12. P l a s t i c s t r a i n s f or drops and Dc- i n 300°C annealed aluminum • 52 I '' INTRODUCTION D i s p e r s i o n Strengthening Approaches P r e c i p i t a t i o n hardened a l l o y s owe much of t h e i r s t rength t o a f i n e d i s p e r s i o n of a s u b - p r e c i p i t a t e , or "zones" . Many of these a l l o y s continue to e x h i b i t r e l a t i v e l y high strength even a f t e r coherency s t r a i n s have been e l iminated by overaging. D i s p e r s i o n strengthened a l l o y s of t h i s group are " s o l u t i o n - t r e a t a b l e " , i . e . , the dispersed second phase must be soluble i n the s o l i d matrix at some elevated temperature. Even at temp eratures below that at which they d i s s o l v e completely , the second phase p a r t i c l e s w i l l have a tendency to reduce t h e i r surface area by coarsening, or coalescence. Since the strengthening e f f e c t of a d i s p e r s i o n i s s t rongly dependent upon the p a r t i c l e spacing, such coarsening i n v a r i a b l y causes a decrease i n a l l o y s t rength . Therefore , the s t ructure and mechanical p r o p  e r t i e s of d i s p e r s i o n strengthened a l l o y s produced by solute r e j e c t i o n from s o l i d s o l u t i o n are i n h e r e n t l y unstable at elevated temperatures. Various methods f o r the product ion of d i s p e r s i o n strengthened 1 a l l o y s conta ining a more s table second phase have been discussed by Grant Of these , two methods have rece ived considerable a t t e n t i o n : <fi) Incorporat ing an oxide i n a metal matrix by powder m e t a l l u r g i c a l t echniques .or i n t e r n a l o x i d a t i o n of a s u i t a b l e a l l o y . Since the discovery 2 of " S i n t e r e d Aluminum Powder" f SAP ) by Irmann , there has been considerable a c t i v i t y i n t h i s f i e l d . if i i ) Incorporat ing f i n a metal matrix) a d i s p e r s i o n of an " i n s o l u b l e " i n t e r m e t a l l i c compound formed dur ing the s o l i d i f i c a t i o n of a s u i t a b l e a l l o y . High rates of s o l i d i f i c a t i o n are d e s i r a b l e to obtain a f i n e d i s p e r s i o n . Thus, powders obtained by a melt fragmentation technique, e . g . , a tomising , are a - 2 - suitable starting material. Towner undertook a general comparative study of the two methods, applied to aluminum. He reported that alloys having a dispersion of an inter- metallic compound during rapid solidification of the alloy have somewhat lower strength properties but significantly higher d u c t i l i t y (especially at elevated temperatures) as compared to alloys containing a dispersion of oxide. Alloys containing fine dispersions of "insoluble" intermetallic compounds offer the promise of higher strength levels at elevated temperatures than age-hardening alloys by virtue of their greater structural s t a b i l i t y . -Probably the most effective way to ensure that a uniform, fine intermetallic compound dispersion is provided by solidification i s to employ an alloy of eutectic or near eutectic composition. It is a basic requirement that the eutectic be formed between an intermetallic compound of unique composition and a pure metal. Solubility of the two eutectic phases in each other should be negligibly low, i f possible, at a l l temperatures below the eutectic isotherm. This is essential to structural stability. Moreover, the volume fraction of the intermetallic dispersion in the alloy should not be more than about 15$, or the du c t i l i t y of the alloy is unlikely to be sat isfactory. Among the other requirements of a suitable alloy system are the following: C I ) The melting point of the intermetallic should be high so that i t does not lose strength appreciably at elevated temperatures. (2) The melting point of the eutectic should not be much lower than that of the pure metal on which the alloy is based; otherwise the working temperature range of the alloy would be reduced substantially relative to - -3 - the pure s o l v e n t . (3) The morphology of the e u t e c t i c should be of the divorced t y p e 1 0 , wi th e s s e n t i a l l y equiaxed or s p h e r o i d a l p a r t i c l e s of the i n t e r m e t a l l i c phase, u n i f o r m l y dispersed i n the metal matr ix . 11 Spengler has proposed that divorced e u t e c t i c s are formed when two parameters and cj? are both l e s s than 0 .1 where 6 = T 2 - T E T x = M e l t i n g temperature of the lower melt ing phase T s . = M e l t i n g temperature of the higher melt ing phase TJJ, = E u t e c t i c temperature and (j> = <j) p f i 4> 2 = Volume f r a c t i o n of the higher melt ing phase = Volume f r a c t i o n of the lower melt ing phase Although the r e l i a b i l i t y of these e m p i r i c a l c r i t e r i a i s 12-13 open t o quest ion , they can be employed as rough guides t o the s e l e c t i o n of s u i t a b l e a l l o y systems. • In fact the tendency to the formation of a divorced type of e u t e c t i c i s increased at the r a p i d rates of s o l i d i f i c a t i o n which are of i n t e r e s t i n the present work. Many systems ( e . g . Aluminum-Nickel , Aluminum-Antimony, Aluminum-Iron, Aluminum-Cobalt , Magnesium-Silicon) s a t i s f y most o f the aforementioned requirements. - k - Of these systems, the A l - S b system has been chosen f o r i n c l u s i o n i n the present study. Part of the A l - S b phase diagram i s reproduced i n Figure Ik 15 1 from Hansen , i n c o r p o r a t i n g a c o r r e c t i o n given by E l l i o t t . The s o l i d s o l u b i l i t y of antimony i n aluminum i n the range 200 - 6k^,°C i s l e s s than ^ '16 0.10$ antimony. The compound AlSb. ' i s reported t o have a cubic s t ructure of the ZhS type w i t h a l a t t i c e parameter of 6.1J55 - 0.0001 A ° . Scope of the Present I n v e s t i g a t i o n Powder m e t a l l u r g i c a l procedures were i n v o l v e d i n f a b r i c a t i n g an a l l o y conta ining a f i n e d i s p e r s i o n of an i n t e r m e t a l l i c compound. I t was found.that the matrix g r a i n s ize was i n v a r i a b l y very f i n e i n such m a t e r i a l , p a r t i c u l a r l y when compared to the matrix g r a i n s i z e i n a s o l u t i o n - t r e a t e d and overaged a l l o y . As a r e s u l t , i t was necessary to t r y to i s o l a t e the e f f e c t of g r a i n boundaries o n . t e n s i l e behaviour i n the temperature range of i n t e r e s t . This was done by studying the behaviour of pure aluminum i n a f a i r l y wide range of g r a i n s i z e . The scope of the present work thus embodied the f o l l o w i n g : (1) E f f e c t of matrix g r a i n s i z e on the t e n s i l e p r o p e r t i e s and work- hardening behaviour i n the temperature range 0.2 to 0.5 Tm (Tin, the melt ing point of matrix metal i n degrees K e l v i n ) f o r a d i s p e r s i o n strengthened mat e r i a l i n v o l v i n g a d i s p e r s i o n of an i n t e r m e t a l l i c compound formed during rapid s o l i d i f i c a t i o n . (2) V a r i a t i o n of the t e n s i l e p r o p e r t i e s and work-hardening behaviour of pure p o l y c r y s t a l l i n e aluminum with g r a i n s i z e i n the temperature range of 0.2 to 0.5 Tm. Figure 1. - A l - S b Phase Diagram - 6 - (3) Comparison of a f i n e grained a l l o y conta ining a coarse d i s p e r s i o n of a thermal ly s table i n t e r m e t a l l i c compound with a coarse grained a l l o y c o n t a i n i n g a f i n e d i s p e r s i o n of a thermally unstable compound with regard to t h e i r t e n s i l e p r o p e r t i e s and work-hardening behaviour i n the temper ature range of -100°C to +200°C. - 7 " II EXPERIMENTAL Materials Superpurity aluminum was supplied by the Aluminum Company of Canada in the form of 25 lb ingots, the analysis of which is shewn in Table 1. Antimony was supplied by J.T. Baker Chemical Co., Phillipsburg, N.J., in the form of lumps, the analysis of which i s shown in Table 2. Reagent grade copper shot from Fisher Scientific (meeting A;C.S. specifications) containing C u m i n 99'99$> w a s used for alloying with superpurity aluminum. Melting and Solidification Aluminum and alloys were melted in sil i c o n carbide crucibles. In preparing an alloy, pure aluminum was melted f i r s t , the alloying element was added, and the melt was stirred. The temperature of the melt was maintained at 50 to 70°C above the liquidus temperature of the alloy. The melt was given occasional stirring to enhance dissolution. In the case of pure aluminum, the temperature of the furnace was maintained at 710 to 730°C following a gradual heating of the charge. Rapid solidification was brought about by two different methods: Table I Analysis of Aluminum Used Cu Fe S i Mg A l (by dif f e r e n c e ) .001 .002 .001 .004 99-992 Table 2 Analysis,of. Antimony Used Fe. As Pb Cu Sb (by difference) i 00 i ,002 .004 .004 .002 99.988 - 9 - (1) C h i l l C a s t i n g : In a 1" diameter by 3" long s p l i t copper mould. (2) S h o t t i n g : • To get s o l i d i f i c a t i o n rates higher than those p o s s i b l e i n c h i l l c a s t i n g , a s h o t t i n g technique was employed. Melts of the d e s i r e d composition were made to pass through a f i n e o r i f i c e under pressure and the i s s u i n g stream was quenched i n water. The equipment used f o r t h i s purpose i s sketched i n Figure 2. Graphite was used f o r the container mainly because of i t s good m a c h i n a b i l i t y . A replaceable graphite bottom was machined to s c r e w - f i t i n t o the c r u c i b l e , and had a 0.020" diameter hole d r i l l e d i n i t . • The melt ing furnace was set c lose to the shot t ing equipment. As soon as the melt was ready, the l i d of the container was l i f t e d and t h e - u n e l t poured i n t o i t . Argon gas pressure on the melt was g r a d u a l l y increased to a l e v e l where good s p h e r i c a l shot was found to be obtained by experience . This occured at a gauge pressure of about 25 pounds per square i n c h . Shot was recovered from the water, washed repeatedly i n a l c o h o l , and stored i n an evacuated d e s s i c a t o r t o minimise surface ox idat ion p r i o r to u s e . - 1 0 - Figure 2. Shotting Equipment - l l - F a b r i c a t i o n Shot was converted to wrought m a t e r i a l i n two stages, compacting and ex t ruding , whereas cast mater ia ls were extruded d i r e c t l y . Compacting: The shot was s t a t i c a l l y co ld compacted to obtain b i l l e t s of 3 A " diameter . Pressures of about 100,000 p . s . i . were employed. A t h i n l a y e r of l u b r i c a n t ( conta ining Molybdenum >d'isulphide) was a p p l i e d to components of the compacting d i e s , but due care was taken to minimise entrapment of l u b r i c a n t i n the compacts. Compacts were degreased and throughly cleaned i n chlorothene ( t r i - c h l o r o - e b h y l e n e ) u s i n g a supersonic v i b r a t o r y c leaner . E x t r u s i o n : Cast or compacted b i l l e t s of 3 A i n c h diameter were hot extruded to wires of 0.15 incfr diameter . The ex t rus ion : . ra t io was 25 to 1. The conta iner , together wi th d i e , plugs and ram was preheated " i n s i t u " by an e l e c t r i c res is tance-heated s leeve . To minimise the o x i d a t i o n of aluminum b i l l e t s , p a r t i c u l a r l y the i n t e r n a l oxidat ion of shot compacts, b i l l e t s were put i n the container only when the whole assembly had reached the d e s i r e d ext rusion temperature. The f i r s t few inches of each ex t rus ion was discarded and the l a s t l/2 to 3 A inches of each b i l l e t was not extruded. T e n s i l e Specimens A small p r e c i s i o n lathe was used f o r machining t e n s i l e specimens from extruded w i r e . -The specimens were machined to a -1'2 - reduced diameter of 0.100" i n a gauge length of 1 i n c h . Next the specimens were placed i n n. i i u u r * on the top surface of a f i r e b r i c k , and annealed i n a muffle furnace . The specimen-brick assembly was allowed to c o o l i n a i r a f t e r removal from the furnace . In the case of aluminum - copper, s o l u t i o n treatment was done + + i n a tube furnace at 535 - 5° C . Aging was then c a r r i e d out at 500 - 5°C f o r 60 minutes i n a tube furnace . T e n s i l e Tests A l l t e n s i l e t e s t s were performed on a Instron t e n s i l e t e s t i n g machine. Threaded "tap wrench" gr ips were used i n most cases . In the case of the coarses t -gra ined aluminum, where sl ippage i n the mechanical g r i p s was observed, specimens were cemented i n t o hollow-ended g r i p s u s i n g an epoxy r e s i n . Four t e s t temperatures were used: - 1 0 0 ° C , Room Temperature, 100°C and 2 0 0 ' ° C , cover ing the range of 0.19 to 0.51 Tm. (of aluminum). Betroleum ether , cooled w i t h l i q u i d n i t r o g e n was used f o r the lower temp- e r t a u r e . For higher temperature t e s t s specimens and g r i p s were immersed i n a heated s i l i c o n o i l ba th . A c o n t r o l l e r incorporated i n the heating c i r c u i t maintained temperatures of 100' - 1°C and 200 - 2 ° C . Tension t e s t s at a l l temperatures were performed at a s t r a i n rate of 0.33 x 10 sec . : - In a d d i t i o n some room temperature t e s t s were -h -1 -2 -1 performed at s t r a i n rates of 0.33 x 10 sec . - . and 0.33 x 10 sec . . . . Metallography 3 O p t i c a l Microscopy: Specimens to be examined were mounted i n "Koldmount" s e l f - c u r i n g r e s i n and p o l i s h e d . o n a s e r i e s of emery papers . The specimens were then e l e c t r o p o l i s h e d i n a s o l u t i o n of composi t ion- N i t r i c a c i d - 1 p a r t , Methyl a l c o h o l - 1 p a r t , Hydrochlor ic a c i d - 1 ml per 50 ml of the mixture . The e l e c t r o p o l i s h i n g s o l u t i o n was prevented from heating by means of water- c o o l i n g c o i l s around the beaker, and was s t i r r e d by a magnetic s t i r r e r . A s t a i n l e s s s t e e l beaker served as the cathode. A good p o l i s h was obtained at 3 v o l t s . D i f f e r e n t etchants were necessary to r e v e a l the microstructure of d i f f e r e n t m a t e r i a l s . Coarse grained aluminum was e f f e c t i v e l y etched by immersing the specimen f o r 1 5 - 2 0 seconds i n a 10$ sodium hydroxide s o l u t i o n heated to 7O0 C . Aluminum - *4-$ copper (Overaged, as w e l l as s o l u t i o n treated) was etched e l e c t r o E y t i c a l l y i n the p o l i s h i n g s o l u t i o n simply by t u r n i n g down the vol tage to 1 . 5 V . The s t ructure of medium g r a i n s i z e d aluminum was best revealed by anodis ing the e l e c t r o p o l i s h e d surface i n a 1 . 8 $ s o l u t i o n of f l u o b o r i c a c i d (HBF ) at 5 0 V and u s i n g an aluminum cathode. The a n i s o - 4 t r o p i c but e p i t a x i a l oxide l a y e r was then examined under p o l a r i s e d l i g h t . 17 The d e t a i l s . o f t h i s technique are discussed by Barker . E l e c t r o n Microscopy: - E l e c t r o n miscroscopy was used to study the microstructures of very f i n e grained aluminum and the aluminum - 10$ antimony a l l o y . S t r i p s of about 0 . 0 0 7 " to 0 . 0 1 0 " thickness were spark machined from extruded wires p a r a l l e l to the ex t rus ion d i r e c t i o n . S t r i p s were e l e c t r o  p o l i s h e d and t h i n n e d . u s i n g the same p o l i s h i n g s o l u t i o n as described above. T h i n f i l m s were examined u s i n g an H i t a c h i HU-11A E l e c t r o n Microscope. - 14 - A r e p l i c a technique was used w i t h some specimens. A neg a t i v e c e l l u l o s e acetate r e p l i c a was d r y - s t r i p p e d from a p o l i s h e d and e t  ched specimen s u r f a c e . Chromium was then deposited at an angle of 45 degrees t o the surface of the c e l l u l o s e acetate , followed by the d e p o s i t i o n of a t h i n f i l m of carbon normal to the s u r f a c e . The c e l l u l o s e acetate was then d i s s o l v e d away us ing acetone, l e a v i n g the carbon r e p l i c a wi th the chromium shadowing i n t a c t . • Measurement of G r a i n - S i z e F u l l m a n 1 R has shown that the diameter & of s p h e r i c a l gra ins i s he la ted to the mean l i n e a l t raverse length 1 at any s e c t i o n as f o l l o w s : d = 1 . 5 1 For the purpose of determining the s i z e of equiaxed grains the above r e l a t i o n s h i p should be a p p l i c a b l e reasonably w e l l and was the one used g e n e r a l l y i n the present work. •Some of the " g r a i n s " were found t o be elongated i n low- 1 8 temperature e x t r u s i o n s . The e s s e n t i a l arguements of Fullman can be used to develop the f o l l o w i n g r e l a t i o n s f o r t h i s case: •In three dimensions the " g r a i n s " can be approximated to r o d s . The longer dimension i s i n the ex t rus ion d i r e c t i o n as w e l l as the t e n s i l e d i r e c t i o n , and i s perpendicular to the d i r e c t i o n of examination. Figure 3 i s a sketch of a rod-shaped g r a i n . PQRS and ABCD are sec t ions p a r a l l e l t o the longer dimension of the g r a i n . I f the e l e c  t ron miscroscope f o i l s are prepared p a r a l l e l to the longer dimension of the rods then sect ions v i s i b l e i n the microstructure would be those p a r a l l e l to PQJRS and ABCD. t - 15 - Extrusion and Tensile D i r e c t i o n Figure 3 . Sketch of a rod shaped-.grain - 16 - In Figure k the average width of a l l the p o s s i b l e chords e . g . AB, CD, E F , PQ, e t c . i n a s e m i c i r c l e can be shown mathematically t o be : Average width = 7T<^ 4~ where d i s the rod diameter . Since the p r o b a b i l i t y of seeing any p a r t i c u l a r chord as the shorter dimension of a rec tangular s e c t i o n i n the microstructure i s the same as that of any other chord, by averaging the shorter dimensions of the " g r a i n s " from the microstructure the a c t u a l diameter of the " g r a i n s " i s three dimensions could be estimated reasonably w e l l . The d i r e c t i o n along which s l i p occurs w i t h i n any g r a i n i s assumed to be at an angle of to the load a x i s . Thus the average width L of the elongated grains i s r e l a t e d to the e f f e c t i v e diameter of the grains a s : ^ e f f e c t i v e = 1 - IT - III OBSERVATIONS AND RESULTS Metallographic Observations: Results of the measurement of "grain" si z e on d i f f e r e n t materials are c o l l e c t e d in.Table 3 . Some features of the microstructure are dealt with i n the following sections, whereas more d e t a i l e d i n t e r p r e t a t i o n s follow i n the Discussion Section. ( l ) 2.7 um Pure Aluminum (extruded from shot) T y p i c a l transmission electron micrographs of the 2.7 Pun pure aluminum are shown as Figures 5 through 8. In these micrographs only a few "grains" exhibit any s i g n i f i c a n t d i f f e r e n c e i n t h e i r transparency to electrons compared to other "grains" in.the area represented. This indicates that most.of the "grains" i n each area were misoriented at ..relatively small angles to one another. This observation was confirmed by a l i m i t e d number of selected area d i f f r a c t i o n patterns on adjacent."grains". The misorientation was l e s s than 5° i n each case. The term "grains" w i l l be used henceforth for any regions i n the microstructure that are separated by distinguishable boundaries, regardless of the magnitude of the misorientation across the boundaries. Grains i n Figure 5 have r e l a t i v e l y sharp boundaries and few c e l l s of dense d i s l o c a t i o n s , whereas the boundaries i n Figures 6, 7 and 8 are l e s s d i s t i n c t and there are more d i s l o c a t i o n s evident, mainly near the boundaries. - 18 - Table 3 Grain S i z e of M a t e r i a l s Invest igated M a t e r i a l . S o l i d i f i c a t i o n , Working Grain S ize and Thermal H i s t o r y i n ym Pure A l Shot;, Cold Compacted, 2.7 Extruded at 200 °C Annealed at 200 °C For 1 h r . Pure A l C h i l l C a s t , Extruded 3.5 at 200 °C and annealed at 200 °C f o r 1 h r . Pure A l C h i l l Cas t , Extruded 10 at 300°C and annealed at 300°C f o r 1/2 h r . Pure A l C h i l l Cas t , Extruded 98 at 400 C and.annealed at 400°C f o r 1/2 ,hr. A l - l O S b Shot , Cold Compacted, 2.9 Extruded at 300°C and annealed at 300°C f o r 1/2 h r . A l - U C u C h i l l Cast , Extruded at 220 300 °C, S o l n t r e a t e d at 535°C f o r 4 h r . and overaged at 300 °C f o r l . h r . A l - U C u C h i l l Cas t , Extruded at 270 300°C and s o l u t i o n t rea ted at 535 °C X 4Tirs. 25 mts. - 19 - Figure 6. Transmission e l e c t r o n micrograph of annealed aluminum from shot . 200°C extruded and - 20 - Figure 8. Transmission e l e c t r o n micrograph of 200°C extruded and annealed aluminum from shot . - 21 - Some of the grains appear to he elongated i n Figure 5 , ( d i r e c t i o n marked with the arrow) whereas they appear to be comparatively equiaxed i n Figures 6 through 8 . This has placed a l i m i t a t i o n of the accuracy with which grain diameter could be estimated. Any entrapped aluminum oxide ( o r i g i n a t i n g on the surface of. shot) was apparently dispersed s u f f i c i e n t l y coarsely in,the microstructure that i t could not be detected as such by transmission microscopy. I t i s therefore u n l i k e l y that oxide i n t h i s material contributed s i g n i f i c a n t l y to ..deformation behaviour. (2) 3• 5 Wim Pure Aluminum (extruded from b i l l e t ) T y p i c a l t h i n f i l m electron micrographs of the 3 . 5 yra pure aluminum are shown i n Figures 9 through 12. Contrast differences between adjacent, grains were not d i s t i n c t i n a l l areas but on the average were greater than those observed i n the.2.7 yim aluminum. S i m i l a r l y the grain boundaries were sharper, and fewer d i s l o c a t i o n aggregates were seen i n the microstructure. In the areas examined, the-grains were found to be r e l a t i v e l y equiaxed. (3) 10 Mim Pure Aluminum Typical o p t i c a l micrographs of the 10 micron aluminum. (a f t e r anodising and using p o l a r i s e d l i g h t ) are shown as Figures 13 and Ik). The grains were a l l equiaxed. The contrast between adjacent grains was sharp. This suggests that r e c r y s t a l l i s e d grains had formed 19 which i s expected as a r e s u l t of a 300°C anneal - 22 - Figure 10. Transmission e l e c t r o n micrograph of 200°C extruded and annealed aluminum from cast b i l l e t . Figure 11. Transmission e l e c t r o n micrograph of 200°C extruded and annealed aluminum from cast b i l l e t . - 2h - Figure 14. O p t i c a l micrograph of 300°C extruded and annealed aluminum using p o l a r i s e d l i g h t on anodised s u r f a c e . - 25 - (U) 98 ym Pure Aluminum Ty p i c a l o p t i c a l micrographs of 98 urn pure aluminum are.shown i n Figures 15 through 17.. M o s t o f ; t h e grains are equiaxed, although some exceptions are apparent. The grain boundaries were generally sharply revealed. (5) Aluminum - 10% Antimony Figures 18 through 21 are transmission electron .micrographs of areas which reveal the matrix grain boundaries i n the Al-10$ Sb. The lack of contrast between adjacent grains suggest that many of the boundaries are probably of low angle. The boundaries are not sharp, and some grains contain d i s l o c a t i o n sub-structures. The p a r t i c l e s of. AlSb appear black where they were retained i n the f o i l at the time of examination. White areas are believed t o - correspond to.where AlSb p a r t i c l e s at the surface have been extracted i n the e l e c t r o p o l i s h i n g operation. A l l the AlSb p a r t i c l e s were found to l i e either on, or very close,to the matrix grain boundaries. Figures 22 and 23 are.from areas i n which the i n t e r m e t a l l i c particles, of AlSb were more c l e a r l y detectable. Most of the pa r t i c l e s , did.not have any one dimension p a r t i c u l a r l y l arger than any other, dimension. The p a r t i c l e s tended.to be,angular i n shape, although the corners of some were rounded, possibly during e l e c t r o p o l i s h i n g . Figures - 2 k and 25 are micrographs of r e p l i c a s taken from d i f f e r e n t areas to emphasise the observation that the AlSb dispersion was not very uniform - p a r t i c u l a r l y with respect to the size of the AlSb p a r t i c l e s . - 26 - ilR^ iiiililiillllilllillla^ ^ Figure 15. Optical micrograph of 1+00°C extruded and annealed aluminum. 19c i Figure 16. Optical micrograph of U00°C extruded and annealed aluminum. - 27 - O p t i c a l micrograph of 400°C extruded and annealed aluminum. and annealed at 3 0 0 ° C . - 28 F i g u r e 2 0 . 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 g r a p h o f A l - l O S b e x t r u d e d and a n n e a l e d a t 3 0 0 ° C . F i g u r e 22. 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 g r a p h o f A l - l O S b e x t r u d e d a n d a n n e a l e d a t 3 0 0 ° C , s h o w i n g A l S b p a r t i c l e s . - 30 - Figure 23. Transmission electron micrograph of Al-lOSb extruded and annealed at 300°C, showing AlSb p a r t i c l e s . cr r ft.. • e>*^  Oaf" r % f» - 10 k Figure 2k. Replica micrograph showing the siz e of AlSb p a r t i c l e s , compared to Figure 25. - 31 - - 2> " •* Figure 25. R e p l i c a micrograph showing the s i z e of AlSb p a r t i c l e s , compared to Figure 2k. Figure 26. O p t i c a l micrograph of overaged A l - U C u showing g r a i n s i z e . - 32 - Aluminum - k% Copper Figure 26 i s an o p t i c a l micrograph of the overaged Al-l+Cu a l l o y . The matrix grains are equiaxed and large (220urn). Figure 27 i s an enlarged o p t i c a l micrograph of the a l l o y which reveals the extremely fin e nature of the dispersion of CuAlg p a r t i c l e s i n r e l a t i o n to the grain s i z e . Figure 28 i s an o p t i c a l micrograph of the s o l u t i o n treated Al-l+Cu.alloy. The grains are equiaxed and coarse (270 ym). Tensile Data - General Values of y i e l d stress (0.2% o f f s e t ) , ultimate t e n s i l e strength, t o t a l percent elongation and percent uniform elongation are c o l l e c t e d i n Tables k through 10 for d i f f e r e n t s t r a i n 'rates and temperatures, and f o r a l l the materials investigated. A y i e l d stress at le s s than 0.2$ o f f s e t was not obtainable.from the machine curves with adequate r e p r o d u c i b i l i t y . In most of.the work two tests were performed under each set of conditions to check the r e p r o d u c i b i l i t y " of each r e s u l t . The v a r i a t i o n i n the strength parameter was generally l e s s than 2% and was les s than 5% i n a l l cases. There was more scatter i n the r e s u l t s f o r the. elongation parameters, which was not unexpected because of the inherent l i m i t a t i o n s of such data. - 33 Figure 2 7 . O p t i c a l micrograph of overaged Al-i+Cu showing f i n e nature of CuAlg d i s p e r s i o n i n r e l a t i o n to g r a i n s i z e . Figure 2 8 . O p t i c a l micrograph of s o l u t i o n t rea ted A l - U C u showing g r a i n s i z e . Table k - 3h - Tensile Data for 2 . 7 urn Aluminum Temperature S t r a i n Y i e l d Strength of Test Rate (0.2% o f f s e t ) °G sec 1 p . s . i . Ultimate T o t a l per-. Uniform Tensile cent percent Strength Elongation Elongation p . s . i . 100 0 . 3 3 x l 0 - 3 14,900 1 5 , 5 0 0 17,000 17,000 2k 21 1 9 . 2 16.7 20 -k 0 .33x10 1 3 , 5 0 0 13,700 l U . 1 0 0 1 3 , 8 0 0 5.k 5.6 1.2 1 . 3 - 0 . 3 3 x l 0 - 3 lk,100 lit,700 l U , 8 0 0 l i t , 8 0 0 5.k 5 . 8 1.1 1.3 0 . 3 3 x l 0 " 2 l i t i 100 - l U . 1 0 0 l U j l O O l i+ ,100 5.9 6.8 1 . 1 2.6 100 0 . 3 3 x l 0 - 3 13,000 12,800 1 3 , 1 0 0 1 3 , 2 0 0 6.1 6.3 1.1 1.2 200 ' 0 . 3 3 x l 0 ~ 3 10,100 10,300 1 0 , 5 0 0 1 0 , 5 0 0 7 . 9 7 . 1 . • l . U 1 . 3 Table 5 - 35 - Tensile Data for 3.5 u m Aluminum Temperature S t r a i n Y i e l d Strength Ultimate. T o t a l per- Uniform of Test Rate {0,2% o f f s e t ) T e n s i l e cent percent °C sec--*- p . s . i . Strength Elongation Elongation p . s . i . •rlOO 0.33xl0~ 3 15,100 15,100 1 6 , 9 0 0 17,800 23.2 2 6 . 6 15-7 1?.4 20 0,33x10"^ 13,300 12,500 13*300 12 , 9 0 0 9-7 11.1 2.0 3,0 0,33xl0~ 3 13,200 12 , 9 0 0 13 , 6 0 0 13,300 11.8 14.8 3.8 6 . 6 0.33xl0~ 2 12 , 6 0 0 12,500 13,500 14,000 15.2 17.7 5-8 6.9 100 0.33xl0~ 3 12,200 12,400 12,400 12,800 9.7 10.2 1.0 1.0 200 0.33xl0" 3 10,500 • 9 , 9 0 0 10,500 9,900 10.0 10.0 1.1 1.0 - 36 - Table 6 Tensile Data.for 10 u m Aluminum Temperature S t r a i n Y i e l d Strength Ultimate T o t a l per- Uniform of Test Rate {0.2% o f f s e t ) Tensile cent percent ° C sec"^- p s i Strength Elongation Elongation p s i -100 0.33xl0 _ : s 9,500 9,^00 20 0.33xl0 _ 1 + 8,600 10,000 0.33xlO - 3 8,600 9,000 0.33xl0" 2 8,000 8,900 100 0.33xl0 - 3 7,800 7,900 200 0.33xl0~ 3 6,000 6,000 1 7 , 0 0 0 3 5 . 9 25.2 1 5 , 7 0 0 3 9 - 5 3 1 . 5 10,200 27.k. iQ.k 1 0 , 9 0 0 22.8 1 3 . 6 11,1+00' 3 1 . 1 21.1 1 1 , 1 0 0 3 0 . 2 2 1 . 1 1 1 , U 0 0 3 1 . 8 23.8 1 1 ^ 3 0 0 3 1 . 1 22.8 8,800 26.6 1 3 . 8 9 , 0 0 0 2 6 . 3 1 2 . 8 6,000 1 3 . 6 1.8 6 , 1 0 0 1 3 . 6 1.0 - 37 - Table 7 Tensile Data f o r 98 Mm Aluminum Temperature S t r a i n Y i e l d Strength Ultimate T o t a l per- Uniform of Test Rate (0.2% o f f s e t ) Tensile : cent.. percent sec - 1 p s i Strength Elongation Elongation p s i -100 0.33x10' -3 3,800 20 0 .33x10 -3 2,800 10,000 1+1.4 32.6 100 0 .33x10 -3 2,600 7,800 1+2.8 28.6 200 0 .33x10 -3 1,600 4 ,300 1+1.9 20.2 - 38 - Table.8 Tensile Data for Al-lOSb Temperature S t r a i n Y i e l d Strength Ultimate T o t a l per- Uniform of Test Rate {0,2% o f f s e t ) Tensile cent Percent °C sec" p s i Strength Elongation Elongation p s i -100 0.33xl0 - 3 16,400 26,200 22.6 17.7 16,100 26,300 25.0 21.6 20 0.33xlO _ i + 13,400 18,500 23.8 14.0 14,000 18,800 24.6 1U.5 0.33xl0 - 3 15,100 20,00.0 23.4 17.1 15,000 19,700 25<0 19.3 0.33xl0 - 2 13,900 20,500 21.8 17.4 14,200 20,300 20.9 15-7 100 0.33xl0 - 3 12,600 15,200 20.6 12.0 12,900 15,000 21.0. 13.1 200 0.33xl0 - 3 10,600 10,700 19-7 1.2 10,800 10,900 16.3 1.1 Table 9 T e n s i l e Data f or Overaged Al-l+Cu - 3 9 - Temperature S t r a i n Y i e l d Strength Ultimate Total; per- Uniform of Test Rate (0.2% o f f s e t ) Tensile cent percent C S e c - 1 p s i Strength Elongation Elongation p s i -100 0.33xlO~ 3 15,600 3*+,000 11.7 10.1+ 15, H 00- 36,000 12.2 11.2 20 0.33xl0 _ 1 + l!+,600 35,100 . 11.3 10.5 0.33xl0 - 3 16,600 36,800 11.5 .9.6 16,200 36,100 11.8 10.1+ 0.33xl0 - 2 16,800 36,300 9-8 8.1 100 0.33xl0 - 3 15,^00 31,200 12.5 7-9 200 0.33xl0" 3 13,1+00 • 19,^00 11.3 6.2 Table 10 Tensile Data f or Solution Treated Al-l+Cu Temperature, S t r a i n Y i e l d Strength Ultimate T o t a l per- Uniform, of Test Rate (0.2% o f f s e t ) Tensile cent.. percent C sec -^" p s i Strength Elongation Elongation p s i . 20 0 . 3 3 x l 0 - 3 21,800 39,1+00 21.7 18.5 - ko - Because.of the r e l a t i v e l y good r e p r o d u c i b i l i t y observed i n the strength r e s u l t s f o r the other materials only one tes t was performed under each set of conditions for 98 \xm pure aluminum and i n some.instances for A l - 4 C u . Tests at d i f f e r e n t s t r a i n rates were not performed on the 98 km aluminum as the r e s u l t s obtained e a r l i e r with pure aluminum d i d not suggest enough dependence on s t r a i n rate at room temperature to j u s t i f y a continuation of the p r a c t i c e . It was not.possible to obtain the ultimate t e n s i l e strength and elongation values i n the case of 98 ym A l at -100°C because of the p r a c t i c a l d i f f i c u l t y of avoiding grip-slippage with t h i s material. Y i e l d Stress of P o l y c r y s t a l l i n e Aluminum Figure 29 shows.the v a r i a t i o n of 0.2% o f f s e t y i e l d stress f o r aluminum with t e s t temperature. 20 Carreker et a l have investigated the t e n s i l e behaviour of two sets, of p o l y c r y s t a l l i n e aluminum which had the following analysis: (AC) S i 0.006%, Cu 0.015%, Fe 0.006%, A l 99-975% (BS) S i 0.006%, Cu 0.002%, FeO.003%, A l 99-987% Lot BS was reported to show no grain s i z e dependence of y i e l d strength, although there was considerable scatter i n the data. Reported r e s u l t s f o r l o t AC are reproduced i n Figure 29 for comparison with the 20 present r e s u l t s . The range of grain si z e studied by Carreker e t a l - hi - _i 1 _ _ ( ! 200 kOO 600 800 _ Temperature Figure 29. V a r i a t i o n of Y i e l d Strength w i t h Temperature f o r D i f f e r e n t G r a i n Sized Aluminum - 1+2 - was l i m i t e d , t h e i r f i n e s t material having 21 ym grains. For comparable grain sizes Garreker et al' have observed a higher y i e l d stress value than i n the present work. This was not unexpected because the flow stresses i n the former work were measured at 0 . 5 % o f f s e t as opposed to the 0 . 2 % o f f s e t used i n the present work. The most important f i n d i n g of the present work i n t h i s regard i s the marked increase i n t h e - y i e l d strength of pure p o l y c r y s t a l l i n e A l , as a r e s u l t of :refinement i n the g r a i n . s i z e , at a l l the temperatures investigated within the range 0.19 Tm to 0.51 Tm. Figure 30 shows a plo t of y i e l d stress (0.2% o f f s e t ) against -1/2 1 where ± i s the i n i t i a l g ram siz e or sub g r a i n . s i z e , for pure aluminum at the four t e s t temperatures. Ultimate Tensile Strength Figure 31 i s a plot of ultimate t e n s i l e strength as a function of temperature for a l l the materials investigated. The f i n e r the grain s i z e , the higher was the U.T.S. for pure aluminum at a l l temperatures. A gradual decrease can be seen i n the U.T.S. values for pure aluminum with increasing temperature for a l l grain s i z e s . The U.T.S. of Al-lOSb was higher than that of pure aluminum of comparable grain s i z e over the e n t i r e temperature range investigated. The difference was more marked at•lower temperatures. oTi oA c U ~ o?8 r.o 1 - l / 2 ( vm)-1'2 1»- Figure 30. Y i e l d Strength V/S (gra in s ize ) For Aluminum at D i f f e r e n t Test Temperatures. Figure 31. V a r i a t i o n of Ult imate T e n s i l e Strength w i t h Temperature For Aluminum and A l l o y s - 1*5 - A r e l a t i v e l y rapid.drop i n the U.T.S. of overaged Al-l+Cu at 200°C was observed. Percent Elongation Because of the marked dependence of t o t a l elongation (to fracture) on such factors as specimen.geometry and gauge length, i t i s not meaningful to compare present values with others previously reported for aluminum. The amount of uniform elongation ( i . e . p r i o r to l o c a l necking) i s generally a more useful parameter although i t i s found to be d i f f i c u l t to assess t h i s quantity i f the machine curve does not show a well-defined peak. Percent uniform elongation i s . p l o t t e d as a function of temperature, for a l l the materials investigated, i n Figure 32. An increase i n the value of t o t a l elongation with decreasing temperature o 20 below 300 K was reported by Carreker et a l . Similar behaviour i s observed with respect to t o t a l elongation as well as uniform elongation i n the present work. A l l the finer.grained materials appeared to reach a l i m i t i n g lower value of percent uniform.elongation with r i s i n g temperature. This was reached for 2.7 ym. aluminum at room.temperature, for 3 5 P aluminum at 100°C and for 10 ym aluminum at 200°C. Al-10% Sb, which maintained i t s d u c t i l i t y with increasing temperature compared with pure aluminum of comparable grain s i z e , did reach the lower l i m i t i n g value at 200°C. Figure 32. V a r i a t i o n of Percent Uniform Elongat ion w i t h Temperature f o r Aluminum and A l l o y s - 47 - Overaged Al -4Cu loses d u c t i l i t y with increasing temperature at an appreciably lower rate than any other of the materials investigated i n the temperature range.of 0.2 to 0.5 Tm. Pure aluminum of 98 Mm grain s i z e i s more d u c t i l e than any of the other materials over the temperature range investigated,'but i t loses i t s d u c t i l i t y at higher temperatures f a s t e r than the overaged Al-UCu a l l o y . Work Hardening Rates During homogeneous p l a s t i c deformation i n tension, the true stress a and true s t r a i n e are often found to be empi r i c a l l y r e l a t e d by anexpression of the form n a = K e where n i s defined as the s t r a i n hardening exponent*. and K i s the strength c o e f f i c i e n t It should be noted that there i s nothing basic about the r e l a t i o n s h i p , and i t s v a l i d i t y i s open to question. However, the equation has been found by many other i n v e s t i g a t o r s ' to be,useful i n comparing the work hardening rates of d i f f e r e n t materials. In many cases n has been found to change with s t r a i n , and comparisons are possible only within those l i m i t e d ranges of s t r a i n where n may be reasonably assumed to be constant. For the present work log a has been p l o t t e d against loge In most cases the value of n was found to change gradually over.the complete s t r a i n region. However, an attempt was made to compare average values of n over the small p l a s t i c range of 0.2% to 2% (Table l l ) . - 1+8 - Table 11 Table of Work Hardening Exponent (n) Material Value of n ( e p l > 0 Q 2 _ Q > 0 2 ) at -100°C 20°C 100 °C 200° C (0.19Tm) (0.32Tm) (O.UOTm) (0.51Tm) 2.7 ym A l 0.028 * * * 3.5 ym A l 0.021 0.028 * * 98 ym A l 0.29 0.35 0.30 0.33 2.9 ym Al-lOSb O.lU • 0.079 0.064 * Overaged Al-l+Cu 0.28 0.30 0.30 0.18 Solution Treated Al-HCu - 0.13 * Ma t e r i a l necks down at very low st r a i n s No t e s t was performed - 49 - It was not possible to estimate n values for the 10 ym aluminum specimens because of the discontinuous y i e l d i n g behaviour t h i s material exhibited i n the s t r a i n range of i n t e r e s t . Fine grained materials necked a f t e r very small s t r a i n s at higher temperatures and no value of n was reported for these cases. From Table 11 the following observations can be made about, the work hardening rates of various materials i n the s t r a i n range 0.2 to 2%. (1) The two f i n e s t grained aluminum extrusions showed a very low rate of work hardening at temperatures where work hardening could , i be observed at a l l . (2) The f i n e grained Al-lOSb a l l o y work hardened much more ra p i d l y at 0.19 Tm than at 0.32 Tm or 0.40 Tm. (3) The coarsest grained aluminum and overaged Al-l+Cu exhibited a comparatively high rate of work hardening at a l l temperatures, although the value of n for Al-4Gu.dropped appreciably at 0.51 Tm. (4) Solution treated and quenched Al-4Cu had a s u b s t a n t i a l l y lower work hardening rate than overaged Al-UCu at room temperature. Discontinuous Y i e l d i n g An i n c i d e n t a l observation during t h i s work was that a number of drops i n stress occurred during the e a r l i e r stages of deformation of aluminum which had been annealed at 300°C and which had a grain s i z e of 10 microns. - 50 - The phenomenon was never observed i n aluminum specimens that were annealed at 200°C or lj-00°C . For the 300°C annealed aluminum, -k the phenomenon was observed at a l l the threestralni rates (0.33x10 , -3 -2 _ i 0.33x10 , 0.33x10 sec ) i n v e s t i g a t e d at room temperature and at a l l the t e s t temperatures where there was any work-hardening at a l l , namely -100°C, 20°C and 100°C. The re levant p o r t i o n of a l o a d - e l o n g a t i o n curve showing -3 -1 the load drops i n 300°C extrusions tes ted at 20°C and 0.33x10 ^ sec i s reproduced i n Figure 33. The drops i n a l l cases were d i s t i n c t . F ive such drops could be seen i n Figure 33- •The number of drops v a r i e d from 3 to 5 i n the range of s t r a i n rate and temperature i n v e s t i g a t e d . A marked r e g u l a r i t y was noted regarding drops and D,_ which a l s o were detected i n a l l the t e s t s . The s t r a i n s at which these l a s t two drops occurred are c o l l e c t e d i n Table 12. A trend toward i n c r e a s i n g valnaes of the p l a s t i c s t r a i n at which drops D. and D occur w i t h i n c r e a s i n g temperature or decreasing s t r a i n rate can be seen i n Table 12. 20 Carreker et a l , a l s o reported that they observed a y i e l d phenomenon at t e s t temperatures below 195°K i n pure p o l y c r y s t a l l i n e a l  uminum which was annealed at' 300°C and had a g r a i n s i z e of 21 microns. They have reported i t s absence i n aluminum annealed at k00°C and 500°C. How ever , no d e t a i l s of the load drops were documented. Present work e s t a b l i s h e s the absence of t h i s phenomenon i n aluminum annealed at 200°C. - 51 - - ko Load f i b s . - 20 Percent E l o n g a t i o n Area of the Specimen = 0.0073 i n . Gauge length = 1 " Figure 33. Discontinuous Y i e l d i n g - 52 - Table 12 Tabulation of P l a s t i c Strains at Which Drops D^ and D, Occur of.Test Rate °C s e c " 1 D4 D5 Percent Percent -10.0 0.33X10""3 1.2 1.3 1.1 1.3 20 0.33xl0 _ 1 + 1.5 2.6 2.8 4.0 0.33xl0 - 3 2.0 2.5 1.6 2.1 0.33xl0" 2 0.6 1.7 1.3 1.7 100 0.33xl0 - 3 2.1 2.5 3.5 h.5 200 0.33xl0 - 3 IV DISCUSSION - 53 - A Pure Aluminum (a) Deformation Behaviour 21 H e i d e r r e i c h , u s i n g a t ransmission e l e c t r o n irmicroscopy technique, showed that the grains of high p u r i t y p o l y c r y s t a l l i n e aluminum are broken up during deformation i n t o u n i t s which he descr ibed as an arrangement of s l i g h t l y misoriented " c r y s t a l domains" (now c a l l e d " c e l l s " ) . 22 -25 A s e r i e s of X- ray i n v e s t i g a t i n n s by Gay et a l showed that a c e l l s t ructure i s present i n other deformed face-centered cubic metals (copper, n i c k e l , lead) as w e l l as i n aluminum and some body centered cubic metals . They concluded that the c e l l s i z e i s independent of the i n i t i a l g r a i n s i z e and decreases to a . - l i m i t i n g v a l u e a f t e r a c e r t a i n s t r a i n . These r e s u l t s have since been confirmed by t ransmission e l e c t r o n . 26 microscopy of t h i n metal f o i l s . The c e l l s were found to be r e l a t i v e l y f ree of d i s l o c a t i o n s but were separated by wal ls of h i g h d i s l o c a t i o n d e n s i t y . 19 Weissmann et a l , have s tudied the d e t a i l s of the d i s l o c a t i o n s t ructure of co ld worked h i g h - p u r i t y p o l y c r y s t a l l i n e aluminum and the processes which occur i n r e c r y s t a l l i s a t i o n . They have suggested that the d i s l o c a t i o n tangles (which represent the i n i t i a l stage of formation of a d i s l o c a t i o n c e l l s t ructure) are formed as a r e s u l t of the mutual a t t r a c t i o n of d i s l o c a t i o n s i n primary and secondary s l i p systems. They found that i n c r e a s i n g deformation givetjrise to a c e l l s t ructure which e x h i b i t s w e l l - d e l i n e a t e d d i s l o c a t i o n boundaries a f t e r about 30$> r e d u c t i o n . A d d i t i o n a l co ld working increased the d i s l o c a t i o n d e n s i t y as w e l l as the twist and asymmetry; of the low-angle boundaries separat ing the c e l l s . No appreciable increase i n d i s l o c a t i o n den s i t y w i t h i n the c e l l s was found, nor was any decrease i n c e l l s i z e observed - 54 - w i t h f u r t h e r co ld working. The increase i n d i s l o c a t i o n d e n s i t y w i t h i n the c e l l w a l l s w i t h i n c r e a s i n g deformation was explained as r e s u l t i n g from the a t t r a c t i o n which screw d i s l o c a t i o n s i n the c e l l w a l l s exert on d i s l o c a t i o n s ed generated w i t h i n ; the c e l l . I t was suggest that the formation of dense tangles i n the c e l l w a l l s was f a c i l i t a t e d by the ease w i t h which d i s l o c a t i o n s can circumvent obstacles w i t h i n the c e l l by means of c r o s s - s l i p . 19 Weissmann et a l f u r t h e r suggest that there i s an interconnect ion between the mechanical and thermal i n s t a b i l i t y of asymmetrical low angle boundaries . The strengthening e f f e c t of the low angle boundaries was a t t r i b  uted to the e f f e c t i v e s t r e s s f i e l d of screw d i s l o c a t i o n s , which are mechan i c a l l y s t a b i l i s e d by the edge d i s l o c a t i o n s . The -stress-, f i e l d of the screw d i s l o c a t i o n s was a l s o considered t o be responsible f o r the thermal i n s t a b i l i t y of the low-angle boundaries , which on annealing are g r a d u a l l y converted i n t o pure t i l t boundaries . (b) St ructure of 200°C'Extruded and Annealed Aluminum S t r u c t u r a l changes t a k i n g place i n aluminum as a . r e s u l t of heavy . deformation (extrusion r a t i o of 25 t o 1) at 2 0 0 , C are l i k e l y to be the r e s u l t of cooperative processes . C e l l s t ructure formation (expected due to deformation) probably competed w i t h simultaneous recovery processes . However, because the deformation i n v o l v e d was at a very h i g h r a t e , a c e l l s t ructure more t y p i c a l of c o l d deformed aluminum probably formed. Subsequent annealing at 200°C has r e s u l t e d i n the microstructures 19 seen i n Figures 5 through 12. Weissmann et a l , pointed out that screw d i s  l o c a t i o n s are responsible f o r the thermal i n s t a b i l i t y of asymmetrical low angle boundaries , which on annealing are g r a d u a l l y converted i n t o pure t i l t boundaries . He observed that upon annealing i t i s the screw d i s l o c a t i o n s which are f i r s t expel led from asymmetrical boundaries (and d i s t o r t e d l a t t i c e domains. ) .The appearance of a c h a r a c t e r i s t i c p r o f u s i o n of i n d i v i d u a l jogged 19 d i s l o c a t i o n s (as at B i n Figure 6) are termed by Weissmann et a l as "disentanglement of d i s l o c a t i o n s . " Sharp boundaries (as at A i n Figures 5 through 12) are supposed to be p r i n c i p a l l y t i l t boundaries , and the 19 r e s u l t i n g s t ructure i s termed subgrains . (c) T e n s i l e Behaviour of P o l y c r y s t a l l i n e Aluminum ( i ) Maximum True Stress True s t r e s s - t r u e s t r a i n ( p l a s t i c only) curves f o r a l l the m a t e r i a l s tes ted at four d i f f e r e n t temperatures are p l o t t e d i n Figures 34 through 37. In Figure J>k, i t can be seen that at -100°C a l l the pure p o l y c r y s t a l l i n e aluminum extrusions work hardened at d i f f e r e n t r a t e s , yet reached e s s e n t i a l l y the same peak value of t rue s t r e s s (where necking com menced). The same observat ion could be made f o r a l l these extrusions i n the case of 20°C deformation (Figure 35) except that the f i n e s t grained (2.7 um) extrusions necked a f t e r very small s t r a i n , and supported a s l i g h t l y 3'S higher maximum true s t ress than, the o thers . At 100°C (Figure 36) -3-r5" Vim ex t r u s i o n s a l s o s t a r t e d to neck a f t e r l i t t l e s t r a i n . The two coarsest grained extrusions supported n e a r l y the same l e v e l of t rue s t ress before necking commenced, al though the 10 um aluminum d i d support s l i g h t l y more s tress than the 98 um aluminum. At 200°C (Figure 37) there was a s u b s t a n t i a l dependence of the maximum a t t a i n e d true s t r e s s on g r a i n s i z e , al though i t should be noted that only the coarsest grained ex t rus ion deformed e x t e n s i v e l y before s t a r t i n g to neck. 1+0,000 "I <~ 1 16 2k 32 True S t r a i n Figure 5%. True Stress-Slzain Curves f o r Aluminum and A l l o y s at -100°C (0.19 Tm) . : g \£ " ^ 32 ^True Strain (%) Figure 35. True stress-strain !.j Curves for Aluminum and Alloys at 20°C (0.32 Tm) 40,000 Figure 36. True s t ress - s t ra in Curves f o r Aluminum and A l l o y s 100°C (0.40 Tm) - 59 - 30,ooo -, 98 y ro A l _, , , 8 16 2 4 True S t r a i n ($) F igure 37. iTrue: s t r e s s - s t r a i n . Curves f o r Aluminum and A l l o y s at 200°C (0.51 Tm) -6o- The f a c t that the onset of necking was s t ress -dependent and apparantly g r a i n s ize - : ' independent at low temperatures may be explained p a r t i a l l y on the b a s i s of a c e l l formation agruement. The suggestion i s that no matter what the s t a r t i n g g r a i n s i z e , i f s u f f i c i e n t uniform deformation occurred;, a l i m i t i n g c e l l s i z e would be a t t a i n e d . Since flow s tress i s de pendent on c e l l s i z e 2 ^ , the flow s t ress would r i s e to a value determined by the l i m i t i n g c e l l s i z e , and not determined.by the o r i g i n a l g r a i n s i z e . 28 •Warrington has found that the l i m i t i n g c e l l s i z e of copper decreases w i t h decreasing deformation temperature. T y p i c a l values reported f o r copper were 0.6 urn at 0.06 Tm and "J \m at 0.68 Tm. I t has a l s o been found that the l i m i t i n g c e l l s i z e f o r aluminum i s decreased by deforming at 27 temperatures below•room temperature . Therefore , an increase i n the l i m i t i n g c e l l s ize of aluminum can be' expected by deforming at temperatures above room .temperature as i s the case-with copper. •At -100°C a l l the pure aluminum extrusions^ a t t a i n the minimum c e l l s i z e f o r t h i s temperature a f t e r s t r a i n i n g d i f f e r e n t amounts. At room temperature, i t may be argued that the 2.7 aluminum has an i n i t i a l c e l l s i z e approaching the lower l i m i t i n g s i z e f o r 20°C. A f t e r very small s t r a i n s the c e l l s i z e reaches the minimum and p l a s t i c i n s t a b i l i t y sets i n because the m a t e r i a l no longer work hardens w i t h i n c r e a s i n g s t r a i n . In f a c t , the p r o v i s i o n of a d d i t i o n a l s t r a i n energy probably promotes the migrat ion of c e l l boundaries (growth of subgrain boundaries from c e l l s ) to produce a form of " w o r k - s o f t e n i n g . " This argument" i s equivalent to one which says that recovery processes are capable of occurring l o c a l l y i n aluminum at 20°C i n regions of very h i g h s t r a i n energy. - 61 - S i m i l a r l y the 3-5 Mm extrusion at 20°C work hardens as f a s t as smaller c e l l s are produced. Because of the small i n i t i a l subgrain s i z e , the minimum c e l l s i z e f o r 20°C i s reached a f t e r only 2$> s t r a i n (Figure 35)- By c o n t r a s t , the coarser grained extrusions undergo considerably more s t r a i n before the minimum c e l l s i z e i s a t t a i n e d f o r 20°C. A l s o , the maximum true s t ress reached decreases s l i g h t l y with i n c r e a s i n g g r a i n s i z e . T h i s p r o b  a b l y r e s u l t s from the f a c t that thermal ly a c t i v a t e d processes contr ibute to . recovery , and that more time i s a v a i l a b l e f o r such processes to operate while the i n i t i a l l y coarse grained extrusions are deformed. Thus competi t ion between, s t r a i n - a c t i v a t e d c e l l - w a l l migrat ion may r e s u l t i n the l i m i t i n g c e l l s i z e (and maximum true s t ress ) being s l i g h t l y g r a i n - s i z e dependent at 20°C. Reference to Tables k through 6 i n d i c a t e s an i n c r e a s i n g value of percent uniform elongat ion with i n c r e a s i n g s t r a i n r a t e . This may be a t t r i b u t e d to an i n c r e a s i n g amount of dynamic recovery or c e l l w a l l mig r a t i o n w i t h decreasing s t r a i n rate ( i . e . i n c r e a s i n g time a v a i l a b l e f o r recovery) . At 100°C, the arguments; of the previous paragraphs are p r o b  a b l y s u f f i c i e n t t o account f o r the observed t e n s i l e behaviour of pure a l  uminum i n Figure 36. The r e l a t i v e lack of dynamic recovery i n the f i n e s t - grained e x t r u s i o n s , and the occurence of i t i n the coarser grained extrusions can be used to account f o r the s u b s t a n t i a l l y higher value of maximum true s t ress a t t a i n e d i n the former m a t e r i a l s . At 200°C, the la rge dependence of maximum true s t ress on i n i t i a l g r a i n or subgrain s i z e may again be a t t r i b u t e d to the e f f e c t s of dynamic recovery processes on c e l l formation and growth. - 62 - The c o n t r i b u t i o n of g r a i n boundary s l i d i n g i s known to be greater at higher temperatures. However, i t s c o n t r i b u t i o n t o t o t a l deform a t i o n i n the present work i s not l i k e l y t o be more than 10$ at 200°C 29 f o r any of the mater ia ls under i n v e s t i g a t i o n . ( i i ) Work Hardening Behaviour An explanat ion of the dependence of work hardening rate on g r a i n s i z e and temperature fo l lows l o g i c a l l y from the arguments ; used above t o i n t e r p r e t maximum, true s t ress v a l u e s . The e f f e c t of i n c r e a s i n g s t r a i n on c e l l s t ructure has been studied by K e l l y u s i n g an X-ray technique . I t was found that the c e l l s i z e i s diminished sharply at f i r s t and the rate of decrease becomes smaller as the l i m i t i n g c e l l s i z e i s approached. The r e s u l t has been con- 27 f irmed by Swann u s i n g t ransmission e l e c t r o n microscopy. A c c o r d i n g l y , one sees from Figures jk- through 37 that aluminum extrusions work harden more r a p i d l y i n the e a r l i e r stages of deformation. •Reference to Table 11 and Figures 3^  through 37 reveals that i n the present work the rate of work hardening increased w i t h i n c r e a s i n g i n i t i a l g r a i n s i z e . With f i n e i n i t i a l grains or subgrains l i t t l e a d d i t i o n a l refinement of the substructure r e s u l t s from deformation, and so work h a r d  ening from t h i s source i s s m a l l . The reverse i s true f o r coarse i n i t i a l g r a i n s . ( i i i ) Y i e l d Stress (0.2$ O f f s e t ) The o f f s e t y i e l d s t ress data f o r pure aluminum extrusions were / -1/2 presented g r a p h i c a l l y i n Figure 30 as p l o t s of 60.2$ v . ' s . 1 where 1 - 63 - i s the i n i t i a l g r a i n or subgrain s i z e . With the exception of r e s u l t s f o r the extrusions of f i n e s t i n i t i a l g r a i n s i z e , the data appear t o obey a H a l l - P e t c h type of r e l a t i o n s h i p at each of the t e s t temperatures. However, i t i s f e l t that the apparant l i n e a r i t y of the p l o t s f o r 3.5 ym t o 98 Vini ex t r u s i o n s i s f o r t u i t o u s , rather than meaningful , e s p e c i a l l y at temperatures above - 1 0 0 ° C . A f t e r 0.2$> s t r a i n , appreciable m i c r o s t r u c t u r a l changes have occurred i n these m a t e r i a l s , depending on tes t temperature and i n i t i a l g r a i n s i z e . Thus coarse grained extrusions have work hardened considerably wi thin- the f i r s t 0.2$ s t r a i n , whereas f i n e - g r a i n e d extrusions have n o t . At the higher temperatures, even.dynamic recovery processes may be i n f l u e n c i n g the 30 ^1 flow s t ress at 0..2$ s t r a i n . Thus, i n terms of the o r i g i n a l arguments • ' J .. 1 upon which the H a l l - P e t c h r e l a t i o n s h i p was based, i t i s u n l i k e l y that i t can be meaningful ly a p p l i e d to the present o f f s e t y i e l d s t ress r e s u l t s f o r pure aluminum. ( iv) Percent Uniform -Elongation During work hardening the c e l l s i z e of a deforming specimen gets f i n e r a n d . f i n e r , and necking sets i n when the l i m i t i n g c e l l s i z e i s reached f o r the p a r t i c u l a r t e s t temperature. T h i s attainment of l i m i t i n g c e l l s ize would happen a f t e r l e s s deformation i f the d i f f e r e n c e between the i n i t i a l gra in s ize and f i n a l c e l l s ize i s l e s s . Therefore , a coarse grained ex t rus ion i s expected to have a higher value of percent uniform elongat ion than the f i n e r grained one at any temperature (Figure-32). •As the work hardening rate of aluminum i s not found t o be s i g n i f i c a n t l y dependant on temperature (Table 11) the value of percent uniform e longat ion i s expected to decrease w i t h i n c r e a s i n g temperature because of an increase i n the l i m i t i n g c e l l s i z e . - 6h - B. Dispersion Strengthened Aluminum Al l o y s (a) Deformation and Recovery Processes 32 Brimhall et.-.al i n t h e i r transmission electron microscopic studies of the deformation of i n t e r n a l l y oxidised a l l o y s found that a high d i s l o c a t i o n density was developed a f t e r r e l a t i v e l y small amounts of s t r a i n . The d i s l o c a t i o n s were found to be arranged randomly rather than, i n c e l l walls i f the p a r t i c l e d i s t r i b u t i o n was s u f f i c i e n t l y fine., A very f i n e c e l l structure was produced i f the d i s t r i b u t i o n of p a r t i c l e s was somewhat coarser . These observations were interpreted as meaning that p a r t i c l e s can act as sources as well as b a r r i e r s to d i s l o c a t i o n s . 32 I t was suggested by Brimhall et a l that p a r t i c l e s may serve as b a r r i e r s to d i s l o c a t i o n s i n several ways. Screw d i s l o c a t i o n s may' be forced to c r o s s - s l i p when meeting a p a r t i c l e , r e s u l t i n g i n the formation of jogs a n d d i p o l e s . Edge.dislocations, which cannot c r o s s - s l i p , tend to bow around the p a r t i c l e s . The screw segments which r e s u l t from bowing can c r o s s - s l i p , however, leaving loops and a Jog i n the d i s l o c a t i o n 33 l i n e as described by Ashby et a l The idea that p a r t i c l e s may act as primary sources of d i s l o c a t i o n s was a r r i v e d at only on the basis of experimental 32 observations, and no s a t i s f a c t o r y t h e o r e t i c a l model was proposed However, the p a r t i c l e s may s e r v e r s secondary sources a f t e r d i s l o c a t i o n motion i s i n i t i a t e d , by inducing c r o s s - s l i p and forming dipoles as a r e s u l t . - 6 5 - Dispersions are also.known to have a d e f i n i t e retarding influence on annealing processes and t h i s i s manifested i n the greatly improved t h e r m a l . s t a b i l i t y of the mechanical properties of many metals which have been dispersion strengthened. The i n h i b i t i o n of annealing has been generally interpreted 3k i n terms of p a r t i c l e r e s t r a i n t to the motion of boundaries . One 3 5 example-of t h i s type of model i s that proposed by Gatt'i'. and Fullman to explain.the lack of macroscopic r e c r y s t a l l i s a t i o n i n i n t e r n a l l y oxidised aluminum a l l o y s . They assumed that r e c r y s t a l l i s a t i o n n u c l e i formed i n a structureless matrix ( i . e . portion of the matrix not containing any di s l o c a t i o n s ) as a r e s u l t of thermal f l u c t u a t i o n s . These n u c l e i , bounded by mobile i n t e r f a c e s , were depicted', as growing at the expense of the surrounding strained matrix. The growth process stops i f the r e s t r a i n i n g force due to surface tension of.the i n t e r f a c e or that due to p a r t i c l e s , i s equal to or,greater than .the d r i v i n g force for boundary migration due to the stored energy difference between the r e c r y s t a l l i s e d volume and the strained matrix. When considering the retention of strength, and of an apparently cold worked microstructure to unusually high temperatures i n 3 6 i n t e r n a l l y oxidised copper extrusions, Preston and Grant proposed that annealing was i n h i b i t e d s o l e l y as a r e s u l t of the interference by dispersed p a r t i c l e s with the migration of sub-boundaries formed during deformation. - 6 6 - In general, annealing could be considered as occuring i n two stages:-(a) the formation, from.the as-deformed substructure, of a sub-grain or nucleus which i s surrounded by a boundary capable,of migration under thermal a c t i v a t i o n , and (b) the growth of the sub-grain: by boundary migration through the remaining defect structure. 37 Brimhall et a l , were f i r s t to u t i l i s e transmission electron microscopy to study the annealing behaviour of dispersion strengthened a l l o y s . They concluded that annealing i n h i b i t i o n by dispersions i s involved i n both of the above stages. They argued that i n al l o y s containing f i n e dispersions, any c e l l s formed would have a much smaller average misorientation and lower d i s l o c a t i o n density, i n the c e l l walls as compared to,pure metals. Sub—grain-formation/ i t s e l f would thus be much less l i k e l y . In addition to the influence of p a r t i c l e s on the deformation mode, and therefore the type and density of di s l o c a t i o n s i n the c e l l w alls, there can also be,direct p a r t i c l e involvement i n the d i s l o c a t i o n rearrangement process within the walls. In cases where sub grain boundaries were observed to form i n such a l l o y s , t h e i r growth was i n h i b i t e d because of the lack of. a c e l l u l a r d i s t r i b u t i o n (or, concentration) of di s l o c a t i o n s i n the surrounding matrix. In other words the .driving f orce-for sub-grain growth i s lacking when the disl o c a t i o n s , a r e d i s t r i b u t e d uniformly, rather than i n . c e l l s . - 67 - (b) Structure of Al-lOSb Extrusions The. Al-lOSb a l l o y was - extruded and "annealed" at 300°C with the r e s u l t i n g microstructure of Figures 18 to 21. Pure aluminum, of s i m i l a r mechanical and thermal h i s t o r y , had the much coarser micro- structure revealed i n Figures 13 and lh. There i s s u f f i c i e n t evidence that annealing processes have proceeded, to an.advanced s'tage i n the 300 °C extruded and annealed pure metal, which can be described as. r e c r y s t a l l i s e d . By contrast, i n the Al-lOSb extrusion there i s evidence that some of the energy, of the farming process has been "stored'.' or retained even a f t e r .prolonged heating at 300 V. I t i s s i g n i f i c a n t also that the AlSb p a r t i c l e s (Figures 18 to 21) are concentrated on c e l l or sub-grain boundaries i n the microstructure. 32 In the l i g h t of the work of Bri m h a l l . e t ' a l . i t may be argued that-during extrusion AlSb p a r t i c l e s could have acted as sources of d i s l o c a t i o n s to produce a r.ather f i n e c e l l s i z e at 300°C. But , during subsequent annealing AlSb p a r t i c l e s presumably acted as b a r r i e r s to c e l l w all migration r e t a i n i n g the c e l l structure rather w e l l . (c) Tensile .Behaviour of Al-lOSb The 0.2% o f f s e t y i e l d stress values for the Al-lOSb (at . a l l t e s t temperatures)' are comparable, to those of pure aluminum of . s i m i l a r grain.size.. No d i r e c t contribution of the dispersion t o y i e l d strength by an Orowan mechanism would a c t u a l l y be expected i n t h i s material because the AlSb p a r t i c l e s are too widely spaced. However, i t i s clear that the dispersion, coarse as i t may be, has res u l t e d i n a marked -- 68 - increase i n the resistance of deformed aluminum to recovery processes at 300°C. Thus the y i e l d strength of Al-lOSb i s appreciably greater than that of pure aluminum of i d e n t i c a l mechanical and thermal h i s t o r y . The Al-lOSb a l l o y was also observed to work-harden to a greater extent than pure aluminum,of comparable-initial g r a i n . s i z e at a l l t e s t temperatures (Figures 3h to 37) except 200°C,where neither of the two materials work hardened. This was c l e a r l y r e f l e c t e d i n the much higher,maximum true stress attained i n Al-lOSb compared to.the pure aluminum. At-100°C, where dynamic recovery e f f e c t s are l i k e l y to be small, the r o l e of the AlSb dispersion i s interpreted as follows. The p a r t i c l e s of AlSb act as sources of d i s l o c a t i o n s . This, plus the . i n h i b i t i n g e f f e c t of the p a r t i c l e s on the migration of d i s l o c a t i o n s to the early-formed c e l l w alls, encourages the formation of smaller c e l l s than are possible i n the pure metal. Expressed otherwise,- the, l i m i t i n g c e l l s i z e a t t a i n a b l e due to deformation i s smaller i n the-. presence of a dispersion. Thus the Al-lOSb a l l o y i s able to work- harden to a higher true stress level-before the d r i v i n g force for sub grain: formation and migration becomes great enough to end the hardening process. The rate.of work hardening of Al-lOSb decreases with increase i n the t e s t temperature (Figures 3*+ to 36). As the temperature? i s r a i s e d there i s a decrease i n the effectiveness of AlSb p a r t i c l e s i n preventing migration of d i s l o c a t i o n s to c e l l walls formed early i n the deformation process. Thus the l i m i t i n g minimum c e l l s i z e increases with temperature, with an e f f e c t s i m i l a r to that•described previously f o r pure aluminum. Although the i n i t i a l substructure of the Al-lOSb was.produced . as a r e s u l t of 300°C deformation, i t i s apparently not appreciably coarser (or ."weaker") than that produced by 200°C t e n s i l e deformation. This i s i n f e r r e d from the t e n s i l e ;test behaviour of the a l l o y at 200°C, wherein necking commenced.after very l i t t l e p l a s t i c s t r a i n . ( F i g u r e 37). The uniform elongation r e s u l t s . f o r Al-lOSb (Figure 32) are explainable i n the same terms as were used previously to int e r p r e t the d u c t i l i t y of pure aluminum, and i t s dependence on grain size ,and temperature. Whereas pure aluminum.of 2.7 or 3.5 micron grain s i z e was able to s t r a i n v e r y . l i t t l e . a t 20 °C before a t t a i n i n g the l i m i t i n g c e l l s i z e , and therefore.beginning to neck, the Al-lOSb a l l o y of si m i l a r grain s i z e was not so l i m i t e d . Thus, the incorporation of a coarse AlSb dispersion has resu l t e d i n simultaneous improvements i n the strength and d u c t i l i t y of pure aluminum. (d) Tensile Behaviour of. Overaged Al-ltCu Extrusions of Al-l+Cu, a f t e r solution-treatment and overaging, possessed a coarse grained r e c r y s t a l l i s e d matrix, i n which there was a very f i n e CuAlg dispersion (Figure.27). Any contribution which the presence of. CuAlg p a r t i c l e s might have made.to the refinement of the-matrix sub-structure during extrusion was subsequently destroyed by the high temperature s o l u t i o n treatment.. - 70 - The increment i n y i e l d strength (at a low o f f s e t ) for t h i s . material over pure aluminum.of comparable grain si z e must be a t t r i b u t e d p a r t i a l l y to s o l i d s o l u t i o n strengthening and p a r t i a l l y to strengthening by an Orowan mechanism. 32 During deformation i n - tension, the work of Brimhall et a l suggests that l i t t l e or no c e l l formation would occur.at low or ordinary temperatures i n the overaged Al-4Cu a l l o y . The very high- i n i t i a l rate of work-hardening (Figures 3b to 37) i n t h i s material can be a t t r i b u t e d to a rapid rate of d i s l o c a t i o n .generation,- the p a r t i c l e s themselves contributing by acting as sources. The concept of l i m i t i n g c e l l s i z e i n t h i s material probably has no meaning, so that i t was possible to. obtain a high.true stress l e v e l without p l a s t i c i n s t a b i l i t y ensuing; due to c e l l - w a l l migration and the r e l a t e d processes discussed above. In fact r e l a t i v e l y l i t t l e necking preceded.fracture of the overaged Al-UCu extrusions,, suggesting that the fracture stress of the a l l o y was approached.throughout the specimen. The effectiveness of the CuAlg pa r t i c l e s , i n preventing any dynamic recovery processes at t e s t temperatures as high as 100°C i s also evident.from the data. Only at the highest t e s t temperature (200°C) was there an appreciable decline i n strength and rate of work hardening. In f a c t , t h i s might be a t t r i b u t a b l e to coarsening of the CuAlg dispersion due to the combined a c t i v a t i n g e f f e c t s of s t r a i n and temperature. - 71 - (e) Tensile Behaviour.of Solution-treated AI-UCu at Room Temperature An Al-4Cu a l l o y quenched from a so l u t i o n treatment temperature 27 of 550°C has been studied by Swann i n the electron microscope a f t e r 10% elongation. In contrast to pure aluminum, t h i s a l l o y d i d not form any c e l l structure. In the absence of a c e l l structure and i n the presence of a supersaturation of point defects produced by quenching, the a l l o y can be expected to work harden r a p i d l y , although not quite as r a p i d l y as i n the presence of a f i n e dispersion of a,second phase. The s o l u t i o n treated a l l o y has a higher y i e l d strength than the overaged a l l o y presumably because of the .much larger amount.of copper i n s o l i d s o l u t i o n . The maximum true stress which was attained i n so l u t i o n treated Al-i+Cu was higher than that observed i n overaged Al-UCu, This was consistent with the l e s s e r necking that preceded fracture i n the former material. (C) E f f e c t of S t r a i n Rate on Strength Properties at Room Temperature 38 Moon and.Campbell have reported that y i e l d strength and ultimate t e n s i l e strength were not.influenced or only s l i g h t l y a f f e c t e d , 39 by s t r a i n rate at room temperature i n aluminum a l l o y s . Trozera et a l have seen some increase i n y i e l d and t e n s i l e strengths when the s t r a i n ^ ' * ' 0 " 6 - 4 _! ^ _ l - l rate was increased from <S&s&s3fr~ sec. to 1 x 10 sec. , for pure aluminum. The s t r a i n rate dependence was very small at 20°C or lower temperatures. - 72 - In the .present i n v e s t i g a t i o n no d i s t i n c t dependence of y i e l d and ultimate t e n s i l e strength could he .detected over,the s t r a i n -k -1 • -2• -1 • rate range of 0.3 x 10 sec.. to 0.3 x 10 sec. , at room temperature : for any of-/the aluminum and aluminum alloys , tested. This r e s u l t could be expected because of the small range of s t r a i n rate involved. - 73 - V. CONCLUSIONS. 1. A c e l l formation and growth .theory due to Gay , K e l l e y , Weissmann.and others can be used to explain s a t i s f a c t o r i l y the t e n s i l e deformation behaviour of pure aluminum as a function of i n i t i a l microstructure and'test temperature. 2. At low temperatures, where thermally activated recovery processes are slow, the true ultimate strength of pure aluminum i s e s s e n t i a l l y g r a i n - s i z e independent. This i s consistent with the concept of a l i m i t i n g minimum size of the c e l l s produced by deformation, t h i s s i z e determining the limit'.of work hardening, and being independent of i n i t i a l microstructure i n the pure metal.. 3. At temperatures above about 0.25 Tm, and i n the s t r a i n rate range of the present work, dynamic recovery l i m i t s the amount,of work- hardening which can occur with deformation. In t h i s case, the true ultimate strength becomes lower as the i n i t i a l g r a i n , s i z e increases. In e f f e c t , the minimum-; c e l l . s i z e attained during deformation i s l a r g e r , the larger the s t a r t i n g grain s i z e , because recovery (sub grain formation) occurs simultaneously with c e l l formation. h. The:0.2% o f f s e t y i e l d strength of pure aluminum, appeared at a l l temperatures to obey a Hall-Petch r e l a t i o n s h i p with i n i t i a l grain si z e i n the range 3.5 to 98 microns. However, i n view of.the extensive occurpecl c e l l formation which must have ooourod i n the f i r s t 0.2% of s t r a i n , and because dynamic recovery probably affected the rate of work hardening at the upper end of the temperature range, i t i s believed that the apparent agreement with the Hall-Petch equation was l a r g e l y f o r t u i t o u s . - 7 4 - 5. The observed dependence on grain s i z e and temperature of the rate of work-hardening and uniform elongation i n pure aluminum are consistent with the c e l l formation and migration arguments: used to i n t e r p r e t true ultimate strengths. 6 . A coarse dispersion of an i n t e r m e t a l l i c compound i n aluminum has been found to retard markedly the recovery processes which r e s u l t i n sub-grain formation and growth. It i s probable that.the AlSb p a r t i c l e s i n the Al-lOSb extrusions of the present work, acted .as d i s l o c a t i o n sources, and thereby contributed to a f i n e r c e l l s i z e than was obtainable i n pure alumium. In addition, the AlSb p a r t i c l e s r e s t r i c t e d the c e l l w all migration which i s necessary to sub-grain formation and the nucleation of r e c r y s t a l l i s a t i o n . 7 . Al-lOSb containing a coarse dispersion of AlSb, has an appreciably higher o f f s e t y i e l d strength than pure aluminum of s i m i l a r mechanical and thermal h i s t o r y , f o r the reasons outlined i n Conclusion ( 6 ) . However, t h e : y i e l d strength was comparable to pure aluminum of s i m i l a r grain s i z e , i n d i c a t i n g that n o . s i g n i f i c a n t Orowan-type contribution to strength was provided.by the dispersion. 8. The Al-lOSb a l l o y work hardened more r a p i d l y , and showed greater uniform elongation than aluminum of comparable i n i t i a l grain s i z e . This i s interpreted as r e f l e c t i n g the lower l i m i t i n g c e l l s i z e attainable i n the presence of the coarse dispersed phase. 9 . The high y i e l d strength and work hardening behaviour of overaged Al-HCu i s probably best explained i n terms of.an Orowan-^Ashby mechanism, invol v i n g the,fine dispersion of ;CuAl_ as primary b a r r i e r s to d i s l o c a t i o n s . - 75 r C e l l formation i s not expected i n such a material, and has not been observed by others. Because the s t r a i n energy of deformation i s uniformly d i s t r i b u t e d i n . t h i s case, instead of being concentrated i n c e l l walls, normal annealing processes are evaded. Thus, the Al-UCu a l l o y i s able-to work-harden almost to i t s fracture, s t r e s s , and l i t t l e necking precedes f r a c t u r e . 10. The solution-treated Al-ltCu,alloy i s also believed to deform without any c e l l formation. Thus, l i k e the overaged a l l o y , i t work-hardens to a high true ultimate strength, and l i t t l e necking precedes f r a c t u r e . 11. Pronounced discontinuous y i e l d i n g can occur under c e r t a i n conditions i n p o l y c r y s t a l l i n e pure aluminum. From the present work, and that of others reported previously, i t appears that the occurence of the : phenomenon i s dependent on annealing temperature. Only 300°C annealed aluminum has shown discontinuous y i e l d i n g . It i s not known what-is the o r i g i n of t h i s behaviour. - 76 - VI BIBLIOGRAPHY I . Grant , N. J ; The Strengthening of Metals (Reinhold P u b l i s h i n g C o . ) 1964, 174. .2. Irmann, R; M e t a l l u r g i a 4 6 , 1952, 125. 3. L y l e , J . P . J r ; Metal Progress , 62, 1952, 109. 4. Nock, J . A . J r ; S . A . E . T r a n s . , 6l, 1953, 209. 5. L y l e , J . P. J r ; M a t e r i a l s and Methods, 4_5, 1956, 106. 6. D i x , E . H . J r ; Aero' Engg. Rev; J a n . 1956, 106. 7. B l o c k , E . A . , and Hug. H ; Sympo. D i s p e r s i o n Strengthening, I n t e r n a t i o n  a l Powder M e t a l l u r g y , June i960, New York . 8. Block E . A . , - Met. Rev. 6, 196l, No. 22. 9. Towner, R. . J ; Metals Engg. Q u a r t e r l y , 1, 1961, 24. 10. C o l l i n s , W. T . J r ; and Mondolfo, L . F ; T r . A . - I . M . E ; 253, 1965, I67I. I I . Spengler , H ; M e t a l l , 1 1 , 1957, 384. 12. Chadwick, G . A; Progr . M a t e r . - S c i . , 12, 1963, 97. 15. Davies V de L ; J l . I n s t . Metals 93, 1964-65, 10. 14. Hansen. M; ' " C o n s t i t u t i o n of Bina-ry A l l o y s " (McGraw,Hill ) 1958, 130. 15. E l l i o t t , R . - P ; " C o n s t i t u t i o n of Binary A l l o y s , F i r s t Supplement, : .(McGraw H i l l ) 1965, 54. 16. Giesecke . G; and P f i s t e r , H ; Acta C r y s t , 11,1958, 369. 17. Barker , L . . J . T r . A . S . M. 42, 1950, 347. 18. Fullman, 'R. L ; J l . Metals 5_, 1953, 447. 19. Weissmann, S; Imura, T ; and Hosokawa, N ; Recovery and R e c r y s t a l l i s a t i o n of Metals ( Interscience) 1963, 2 4 l . 20. Carreker R. P. J r ; and Hibbard W. .R. J r ; Trans . A . I . M. E . 209, 1957, 1157. -11- 21. H e i d e n r e i c h , R. D ; J . A p p l . Phys. 20, 19^ 9, 993. 22. Gay, P; and K e l l y , A ; Acta C r y s t . 6, 1953, 185. 23. Gay, P ; and K e l l y A; Acta C r y s t ; 6, 1953, 172. 2k. Gay, P; H i r s c h , P. B ; and K e l l y , A;. Acta C r y s t J_, 195^ , kl. 25. K e l l y , A ; Acta C r y s t , 19^k, 55^ . 26. . H i r s c h , P. B; H o m e , R. W;-and Whelan, M. J ; P h i l M a g . , 1, 1956, 677- 27. Swann, P. R; " E l e c t r o n Microscopy and Strength of C r y s t a l s " ( I n t e r - science) I963,'131. .28. Warr ington, D. .H; Proceedings of the European Regional Conference on E l e c t r o n Microscopy, D e l f t , I96I, 35'1" 29. Stevens, R. N ; Met. Rev. 11, 1966, No. 108. 30. Armstrong, R; Codd, I , Douthwaite, R . M ; and Petch, N . J ; P h i l Mag 7x 1962, U5. 31. Warrington, D. H ; J.I-.'B.-L. ,201, I963, 610. 32. B r i m h a l l , J . L ; and Huggins R. A ; T r . A . I .M.E . ' . 253, 1965, IO76. 33. Ashby, M. F ; and Smith, G. C ; P h i l Mag.. % i960, 298. 3k. • Smith C . S . Trans A . I . M . E . 175, 191+8, 15. 35. G a t t i , A ; , a n d Fullman, R. L ; T r . A . I . M . E . 215, 1959, 762. 36. Pres ton, 0; and Grant , N . J ; T r . A . I . M . E . .221, I96I, l6k. 37. B r i m h a l l , J . L ; K l e i n , M. J ; and Huggins, R. A ; Acta Met lh, ±966, k^9. 38. Moon, D. P; and Campbell , J . E ; L i g h t Metal Age 20, 1962, 18. 39. T r o z e r a , T . A ; Sherby, 0. D; Dora, J . E ; T r . A. S . M. k_9_, 1957, 173-

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