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UBC Theses and Dissertations

Phase transformations in the ag-20.1 to 27 at o/o a1 alloys. Cambal, Ludvik 1972

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PHASE TRANSFOKMATIONS IN THE Ag - 20.1 TO 27 AT % Al ALLOYS by LUDVIK CAMBAL • A.Sc., Technical University, Brno, Czechoslovakia VOLUME 1 ( T E X T ) A THESIS SUBMITTED IN PARTIAL FULFILMENT OF THE REQUIREMENTS FOR THE DEGREE OF MASTER OF APPLIED SCIENCE in the Department of METALLURGY We accept this thesis as conforming to the required standard THE UNIVERSITY OF BRITISH COLUMBIA April, 1972 In present ing th i s thes is in pa r t i a l f u l f i lmen t o f the requirements for an advanced degree at the Un ivers i t y of B r i t i s h Columbia, I agree that the L ibrary sha l l make it f r ee l y ava i l ab le for reference and study. I fu r ther agree that permission for extensive copying o f t h i s thes i s fo r scho la r l y purposes may be granted by the Head of my Department or by his! representat ives . It is understood that copying or pub l i c a t i on of th i s thes i s f o r f i nanc i a l gain sha l l not be allowed without my tten permiss ion. w r i Department of Metallurgy  The Un ivers i ty o f B r i t i s h Columbia Vancouver 8, Canada Date May 31, 1972 - i i -ABSTRACT The anisothermal and isothermal phase transformations i n the Ag-20.1 to 27 at% A l a l l oys were examined. The coo l ing rates employed during continuous coo l ing had a d i s t i n c t e f f e c t on the phase t ransformat ions. In the 24, 25, 26 and 27 at% A l a l l oys the low temperature transformations required coo l ing rates of 0 .3-0.6 °C per minute i n order to obta in the phases pred ic ted from the phase diagram. The 20.1 and 23 at% A l a l l oys exh ib i ted much.less y-phase at room temperature a f te r coo l ing at 0 .3 °C per minute than an t i c ipa ted from the phase diagram. The y-phase developed wi th in the (5 + p) f i e l d i n the 24 and 25 at% a l l oys resembles Widmannsfatten p l a t e s . The WidmannstStten s t ruc ture was a lso present i n the 25 at% a l l o y at room temperature. During the transformat ion £ (£ -»- y) i n the two phase f i e l d the parent £ matrix develops a polygonized substructure . This substructure i s in te rpre ted i n terms of d i s l o c a t i o n rearrangements i n the accommodation region adjacent to the >i-phase p l a t e s . I t .was found that the quenched B-phase exh ib i t s p r e c i p i t a t i o n hardening during low temperature ageing. At 200°C the y-phase grows i n the 3-matrix as sphe r i c a l p a r t i c l e s . The radius of these p a r t i c l e s i s a l i n e a r funct ion of t ime. Thus the 3 -> y transformation was c l a s s i f i e d as an in te r f ace con t ro l l ed process . The k i n e t i c data obtained during the isothermal t ransformat ion of the quenched £-phase in to the y-phase at 160 and 200°C ind i ca tes tha t the o v e r a l l growth rate of th i s reac t ion i s a lso in te r face con t ro l l ed . However, the growth cha rac t e r i s t i c s and the surface d i s t o r t i o n accompanying the E, -> y transformat ion ind ica te that more than one atomic process operates during the 5 -»• y r eac t ion . •  The X-ray and e l ec t ron microscopy analyses suggest that there i s not a random d i s t r i b u t i o n of - i i i -the Ag and A l atoms i n the y - p h a s e . I t .appears t h a t the y-phase c o n t a i n s l a y e r s r i c h i n Ag o r A l atoms. The r e g u l a r s t a c k i n g of: t h e s e l a y e r s may g i v e r i s e to a l o n g range s u p e r l a t t i c e as suggested by the X-ray d i f f r a c t i o n p a t t e r n s . The d e s i g n a t i o n o f the y-phase as an i s o m o r p h i c 8-Mo s t r u c t u r e seems i n c o r r e c t . E l e c t r o n t r a n s m i s s i o n m i c r o s c o p y of t h e . i s o t h e r m a l l y produced y +/ £ s t r u c t u r e s i n the 24 a t % a l l o y e s t a b l i s h e d t h a t t h e r e i s a. c r y s t a l l o g r a p h i c r e l a t i o n s h i p between the p a r e n t and the p r o d u c t . The s t r u c t u r e r e l a t i o n s h i p between the Widmannstatten y-phase and the ^ - m a t r i x i n the 25 a t % a l l o y was a l s o o b t a i n e d u s i n g X-ray ^La/ue t e c h n i q u e s . S-phase s i n g l e c r y s t a l s t r a n s f o r m e d i n t o y - p h a s e p o l y c r y s t a l s . The o r i e n t a t i o n of the y - c r y s t a l l i t e s was n o t c o m p l e t e l y random. A l t h o u g h , a h i g h degree of randomness was i n d i c a t e d when u s i n g X-ray D.S. t e c h n i q u e s on f i n e r o d s . "Fragmentation" of the £-phase s i n g l e c r y s t a l i n t o a p o l y c r y s t a l t a k e s p l a c e d u r i n g room temperature d e f o r m a t i o n of the quenched £-phase. R e l a t i v e l y s m a l l amounts o f d e f o r m a t i o n i n t r o d u c e s numerous twin l a m e l l a e i n t o the £-phase. A l t h o u g h the d u c t i l i t y o f the deformed S-phase d e c r e a s e d r a p i d l y w i t h i n c r e a s i n g d e f o r m a t i o n , no p r e s e n c e of a s t r a i n i n d u c e d t r a n s f o r m a t i o n p r o d u c t was d e t e c t e d u s i n g X-ray methods. The 5 + y i s o t h e r m a l t r a n s f o r m a t i o n b e a r s a s t r o n g resemblance t o the 3 ^ 5 t r a n s f o r m a t i o n i n the Ag-Zn a l l o y . - iv «. TABLE OF CONTENTS Page INTRODUCTION 1 REVIEW OF PUBLISHED WORK ± PURPOSE OF THE PRESENT INVESTIGATION 4 1. EXPERIMENTAL PROCEDURES 4 1.1 Alloys Preparation 5 1.2 Homogenizing 6 1.3 Hot Rolling 6 1.4 Cold Rolling 6 1.5 Cold Drawing of Wire 7 1.6 Heat Treatment , 7 1.7 Preparation for Optical Metallography 8 1.8 Preparation for Quantitative Metallography 8 1.9 Preparation for Electron Microscopy 9 2. GENERAL PROPERTIES AND MORPHOLOGY OF PHASES IN COMPOSITION RANGE Ag-20.1 TO 27 AT % Al 10 2.1 Preliminary Results 10 2.2 Hardness of the c, and u-Phases 12 2.3 X-Ray Phase Fields Determination 12 X-Ray D.S. Patterns Evaluation 14 2.4 Optical Metallography of PHases in the Composition Range Ag-20.1 to 27 at % Al 20 «- V -Page 3. TWO-PHASE EQUILIBRIUM STRUCTURES 24 3.1 (?+y) Mixture in 24 at % Alloy 24 3.2 (s+y) Mixture in 25 at % Alloy 28 4. COMMENT ON SURFACE METALLOGRAPHY 29 5. DECOMPOSITION OF THE QUENCHED g-PHASE 30 5.1 Metallography of the g^-Phase Decomposition in 24 at % Alloy 30 5.2 Pre-Precipitation Phenomena During g^-Phase Decomposition 31 5.2.1 X-Ray D.S. Sequence of the Pre-Precipitation Stage 32 5.2.2 X-Ray Laue Transmission Sequence of the g^ Decomposition 34 6. DILATOMETRY OF ANISOTHERMAL PHASE TRANSITIONS IN 24 AT % ALLOY 36 7. ANALYSIS OF THE <;-PHASE - 24 AT % Al 37 7.1 ^-Single Crystal Analysis 39 7.2 Some Aspects of Heat Treatment 40 - v i «• Page 8. TRANSFORMATION OF QUENCHED Z, -PHASE IN 24 AT % A l ALLOY . 41 8.1 Surface D i s t o r t i o n During — a » u Transformation 45 8.2 Dilatometry Under Isothermal Conditions 47 8.3 Deviations from Isothermal Conditions 48 8.4 Isothermal Transformation Z —=» y^ i n 24 at % Q 1 A l l o y 50 8.4.1 X-Ray D.S. Sequence 51 8.4.2 X-Ray Transmission Laue 53 8.4.3 O p t i c a l Metallography of £Q *-Transformation 54 8.4.4 Q u a n t i t a t i v e Metallography 55 9. KINETICS OF PHASE TRANSFORMATIONS 57 9.1 General Theory of Phase Transformations 57 9.2 Comments ^o"3^ ^ \i± K i n e t i c s 62. 9.3 K i n e t i c s of £Q — y ^ Transformation i n 24 at % A l l o y 6 3 9.4 Growth C h a r a c t e r i s t i c s of y^-Phase 6 ^ 10. ISOTHERMAL TRANSFORMATION IN 25 AT % A l ALLOY 6 8 10.1 O p t i c a l Metallography and X-Ray D.S. Pat t e r n s ... 6 8 10.2 C r y s t a l l o g r a p h i c R e l a t i o n s h i p c//yjr i n 25 at % A l A l l o y 6 9 10.3 Transformation Sequence f o r Si n g l e C r y s t a l s ..... 71 10.4 ?-Phase N u c l e a t i o n on Free Surfaces 72 - v i i -r> Page 10.5 Effect of Deformation on r^-Phase 73 11. u-PHASE CRYSTAL STRUCTURE ANALYSIS 80 11.1 Experimental Details of X-Ray Analysis (24 at % Al) 80 11.2 Conclusions from X-Ray Analysis of u-Phase 82 12. ELECTRON MICROSCOPY OF ? Q AND u-PHASE IN 24 AT % Al ALLOY , 88 12.1 Cq-Phase 88 12.2 Structural Changes of ?^-Phase i n El-Microscope Stage 90 12.3 Transformation — > y^ 90 12.4 Double Diffraction Effects 93 12.5 y^-Phase After Long Isothermal Tempering 95 12.6 y -Phase Structure 96 13. REPLICAS FROM FRACTURE SURFACES 97 14. DISCUSSIONS OF RESULTS 98 Comments £0"Transformation Mechanism 115 CONCLUSIONS 119 - v i i i -LIST OF FIGURES (T?.134 - 223 FIGURE ALLOY DESCRIPTION AT % AL 1 Phase diagram of Ag-Al binary system (1). 2 Thermodynamic properties of Ag-Al alloys at 550°C. 3 Lattice parameters versus electron concentration. 4 - . . . . .. A t % A 1 ( 7 ) > 5 Melting furnace. 6 Heat treatment furnace. 7 Operating conditions of the furnace. 8 Electropolishing apparatus. 9 24 B i l l e t slowly cooled from 700°C to room temperature. 10 24 B i l l e t Quenched from 550°C and drawn to wire. 11 24 Quenched from 550°C and tempered at 200°C for 1 minute. 12 Hardness versus Al-content for different heat treatments. X-RAY D.S. PATTERNS OF PHASE FIELDS AFTER QUENCHING FROM TEMPERATURES (°C): 13 20.1 500, 400, 300, 200 14 23 500, 400, 300 15 24 500, 435, 400, 300, 200 16 25 500, 300 17 25 375 - Cooled at 0.6°C/min. from 500°C 18 23 Room temperature - Cooled at 0.6°C/min. from 500°C. - i x -FIGURE ALLOY DESCRIPTION AT % AL X-RAY OF y?L - STRUCTURES: Kl 21 24 Laue back r e f l . Cu. . rad., cooled at 0.4°C/inln. (w) 22 24 " " " Mo. . " (w) 23a 25 " " " Cu. , " " 7°C/min. (w) 23b 25 " " " Cu. , " " 0.4°C.mln. (w) OPTICAL METALLOGRAPHY OF STUCTURES AT TEMPERATURES: 25 20.1 500°C (Quenched). 26 20.1 R.T. - cooled at 0.3°C/min. 27 23 500°C (Quenched) 28 23 R.T. - cooled at 0.3°C/min. 29 24 500°C (Quenched) 30 24 435°C " 31 24 435°C 32 24 400°C 33 24 300°C 34 24 R.T. - cooled at 0.6°C/min. fracture surface (polarized l i g h t ) . 35 24 R.T. - cooled at 0.1°C/min. ele c t r o p o l i s h e d . 36 24 37 24 R.T. - cooled at 0.1°C/min. (polarized l i g h t ) . FIGURE ALLOY AT % AL DESCRIPTION 38 25 500°C 39 25 R.T. - cooled at 0.4°C/min. 40 25 41 26 42 27 11 " TWO-PHASE EQUILIBRIUM STRUCTURE (24 AT %) : 43 24 435°C (annealed 185 hours and quenched). 44 24 " " " " " 45 24 11 " " " " and tempered at 200°C for 30 minutes. 46 24 X-Ray D.S. Pattern of 435°C struct. CuK rad. 47 24 X-Ray Laue " " " C u ( w ) 11 48 24 11 " " " 11 Mo ( w ) 11 X-RAY LAUE PHOTOGRAPHS AFTER COOLING AT 5°C PER MINUTE FROM 500°C TO TEMPERATURES: 49 24 425°C (Quenched) 50 24 415°C (held 15 minutes and Quenched) 51 24 400°C 52 24 300°C (held 7 hours and Quenched) - xi -FIGURE ALLOY DFSCRIPTION AT % AL TWO-PHASE STRUCTURE IN THIN FOIL: 53 24 Cooled at 7°/min. from 700°C to 435°C, held 30 min. and quenched. 54 24 Electron diffraction pattern of Fig. 53 TWO-PHASE EQUILIBRIUM STRUCTURE (25 AT %) 55 25 400°C (annealed 200 hours and quenched) Chemically etched. 56 25 Same treatment as Fig. 55, electropolished (-etched). 57 25 Same treatment as Fig. 55 D.S. pattern CuK^ rad. 58 25 " " " electron Microprobe scan. 59 (deleted) SURFACE METALLOGRAPHY: 60 25 Quenched from 500°C, not electropolished after quenching, polarized light. 61 25 Same as Fig. 60, lightly electropolished, polarized light. - x i i -FIGURE DESCRIPTION S-PHASE: 62 24 Quenched from 700° C into brine at 0°C. Sample thickness 0,030" 63 24 Quenched from 700°C into brine - 4°C. sample thickness 0,005" 64 24 - Phase tempered 43 hours at 100°C. 65 24 it I I I I I I I I 66 24 " " " 20 seconds at 200°C. 67 24 I I I I I I gQ " " 68 24 I I I I ll 210 " ll 69 24 Precipitate radius versus time (CQ—> y? at 200°C) 70 (deleted) 71 24 Precipitate hardening of the B^-Phase 72 24 X-Ray D.S. sequence of the B^-Phase decomposition 73 24 X-Ray Laue Sequence of the Bp-Phase decomposition ANISOTHERMAL DILATOMETRY: 74 24 R.T. to 680°C 75 • 24 R.T. to 600°C - x i i i -AT TOY FIGURE AT % AL DESCRIPTION £ -PHASE: 76 24 X-Ray D.S. patterns of Cq-Phase (66=Powder, 67=Solid Rod 77 24 X-Ray Laue Pattern of C^-Phase (500°C) 78 24 " 11 " 11 (460°C) 79 25 Undissolved R.T. Structure after 5 min. at 500°C. 80 25 Structure as i n Fig. 79 after 75 min. at 500°C. 81 25 X-Ray Pinhole Photograph (CrK^) of deformed t Q , annealed 15 min. at 500°C. 82 25 Enlarged detail of Fig. 81. TRANSFORMATION C Q—»y: 83 24 Dilatation curve. Heating at 0.33°C/min. 84 24 11 11 " 2°C/min. 85 23 D.S. pattern of CQ tempered 60 min. at 200°C 86 24 " " " " 8 " 11 " 87a 24 Laue Pattern, C^-Phase 87b 24 " " " " tempered 5 min. at 187°C. 88a 24 11 " " " 88b 24 " " " " tempered 20 sec. at 240°C. - xiv -FIGURE A T % A L DESCRIPTION 88c 24 " " " " tempered 40 sec. at 240°C. 88d 24 11 " 11 " 100 " 89b 24 Structure after Tempering 20 sec. at 240°C. 89c 24 " " " " 40 " 90A 24 Surface Relief, Cg-Phase after 20 min. at 160°C. 90B 24 " " " " " 35 " " 90C 24 " " " " 50 " 90D 24 As Fig. 90C, electropolished 91 24 Surface Relief, tg-Phase after 40 min. at 160°C. 92 24 Isothermal dilatometry of CQ y transformation 93 24 Sample temperature variation at 201°C 94 24 " 11 " 11 231°C 95 24 " " " " 304°C 96 24 X-Ray D.S. pattern (Cu), Cg-Phase 97 24 " " " " after 2 min. at 200°C gg 24 " 11 " " " 5 " " 99 24 " " " " " 9 " " 100 24 11 " " " deformed 15% 101 24 " " " " " " and tempered 2 min. at 200°C 102 (deleted) 103 104 - X V -AT T DY FIGURE AT % AL DESCRIPTION 105 24 Laue transmission, CN after tempering 22 min. at 200 °C as thin f o i l 4 106 24 Laue transmission, C, after tempering 22 min. at 200 °C as bulk material. 107 24 Microstructure - CQ after 2 min. at 200 °C 108 24 11 " 6 11 109 24 " " 9 " " 110 24 y-Phase internal structure(treatment as for Fig. 108). 111 24 112 24 " Morphology in three dimensions 113 24 Quant. Metallography, CQ after 4 min. at 200 °C 114 24 11 " " " 6 " 11 H5 24 " " " " 7 " " 116 24 C —»»• y Kinetics 117 24 118 24 119 24 119-1 24 CQ after 30 min. at 160 °C (slowly heated up) 119-2 24 The same as Fig. 119-1 plus tempering 10 min. at 160 °C. 119-3 24 CQ after 1 + 1 min. at 200 °C 119-4 24 11 1 1 1 + 1 + 2 min. at 200 °C 119-5 25 CQ -Phase mechanically ground and chemically etched 119-6 24 CQ-?hase after 3 min. at 210 °C. Mechanically prepared surface and chemically etched. - xvi -FIGURE A T ^ A L DESCRIPTION 119-7 120 121 122 123 124a 124b 125 126 127 128 129 130 131 132 133 134 135 136a 136b 136c 136d 24 25 25 25 25 25 25 25 25 25 25 25 25 25 26 27 26 25 25 25 25 25 Same treatment as for Fig. 119-6 (different sample) CQ after 5 hours at 200 °C. 50 100 11 7 11 80 " Relationship between C-matrix and u-plates D.S. Pattern Laue Pattern, Cg-Phase i t i t II t i II tr after 2 hours at 200 C " 4 X-Ray D.S. Pattern C Q " 32 11 " CQ electropolished " " after 30 hours at 200°C t l t l It CjQ It t l i i i i i i ]_oo " " Structure as in Fig. 132 after re-polishing C N mechanically polished, after 15 min. at 200°C (surface) Laue Pattern of the Cg-Phase Laue Pattern of the Cg-Phase deformed 3% II i t t i t t i t 177 Structure of Cg-Phase deformed 3% - x v i i -AT T OY FIGURE rZ B, A T DESCRIPTION AT % AL 136e 25 Structure of Cg-Phase deformed 17% 136f 25 Laue Pattern " " 10% 136g 25 Structure of " 11 11 137-1 Neumann Bands in Bronze and Brass (22) 137-2 Lattice reorientation during twinning (22) 137-3 25 Local deformation of the t^-Phase. LOCALLY DEFORMED CQ-PHASE AFTER 75 MIN. AT 200°C: 138a 25 Close to Indentation 138b 25 1 mm. from Indentation 138c 25 3 mm. from Indentation X-RAY D.S. PATTERNS OF THE y-PHASE: 139 24 yD„-Phase, pattern 68{Cu rad.) 25 " 11 11 72 " " 140 24 y^-Phase (Isothermally produced) Cu rad. 141 24 " " ^ n <^ j^-RT " Phas* F e " 142 24 " " " " " Cr " ELECTRON TRANSMISSION MICROSCOPY: 143 24 Cn-Phase 5 days after quenching - x v i i i -FIGURE ALLOY AT % AL DISCRIPTION 144 24 Cg-Phase less than 4 hours after quenching. 145 24 n ti it ti 2 " " " 146 24 S.A.D. of Fig. 145 147 24 Cg after 2 hours in the beam 148 24 I I I I ^ it ti ti ti 149 24 ti I I 2 "" " 11 ti 150 24 " 150 sec. at 200°C. 151 24 " " " " (enlarged) 152 24 " 11 6 min. " (as bulk) 153 24 (deleted) 154 24 Cg after 210 sec. at 200°C 155 24 y-Phase of Fig. 154 at large magnification 156 24 I I it it I I ti ti 157 24 S.A.D. of Fig. 155 158 24 CQ after 30 min. at 200°C. Thin fragment. 159 24 S.A.D. of Fig. 158 160 24 CQ after 210°C at 200°C (y + Q 161 24 S.A.D. of Cg i n Fig. 160 162 24 S.A.D. of (CQ + y) boundary in Fig. 160 163 24 S.A.D. of (CQ + y) layers i n Fig. 160 164 24 y-fragments i n carbon 165 24 CQ after 2 hours at 200°G 166 24 it it -j I I I I - xix « AT T DY FIGURE AT % AL DISCRIPTION 167 24 y as cooled from 650°C at 0.4°C per min. to R R l REPLICAS OF THE y-PHASE: 168 24 Cg after 30 min. at 200°C. Fracture surface 169 24 " " 15 11 " 300°C. j_70 24 " " " " " " " " 171 24 y as cooled from 650°C at 0.4°C per minute. K l 172 Proposed Model of Clustering - XX -ACKNOWLEDGEMENT The author wishes to thank his research director, Dr. E.B. Hawbolt for the advice given throughout the period of study and during the preparation of this thesis. Thanks, are also extended to other members of the faculty, fellow graduate students, and technicians for their helpful cooperation. The financial support from the National Research Council of Canada (Grant No. 8-6155) i s gratefully acknowledged. - 1 -INTRODUCTION The Ag-Al binary phase diagram is shown in Fig. l.V It contains two terminal solid solutions a. , a... and three intermediate Ag' Al phases g, t,, and y. Ag-Al alloys have limited practical importance. The system, however, exhibits several overlapping phase fields making i t suitable for examining composition invariant phase transformations. REVIEW OF PUBLISHED WORK The thermodyanmic properties of Ag-Al alloys at 550°C are shown 2 M in Fig. 2. The positive value of the heat of mixing, H , at the Al-rich side implies clustering (CP. zones). The negative value of M H at the Ag-rich side would imply short-range order. Al-Base Alloys: The phase transformations that occur on the Al-rich side of the binary system are relatively well understood. Upon continuous cooling, the h.c.p. £-phase precipitates from 3 the f.c.c. a-phase i n the form of Widmannstatten plates. The phases have the following orientation relationship: (lll)a//(0001) ? • [llp]d//[1120] ? In the a supersaturated solid solution, spherical G.P. zones form 4 5 during quenching and low temperature aging. * The clustering of Ag in the solid solution was interpreted in terms of a miscibility gap - 2 -existing in the two phases region, (a,Z,).^ Nicholson and Nutting^ suggested the following sequence for the transformation: 1. Spherical G.P. zones, rich in Ag. 2. Partial ordering in the G.P. zones. 3. Development of Ag-rich stacking faults by subsequent shear on close packed planes. 4. ? ordered precipitate. 5. Z, formed by discontinuous precipitation. Ag-Base Alloys: Massalski and Cockayne^ measured the variation of the lattice parameters with Al content for the f.c.c. a and the h.c.p. z, as quenched from 550°C. Fig. 4 shows a continuous decrease of the "a" parameter in the a-phase and an increase of this parameter in the Z, -phase with addition of A l . The c-parameter and the ratio^i — in h.c.p. S-phase decrease linearly with the Al content (Fig. 3). These observations have been interpreted as Brillouin-zone effects^ 8 though Newmann suggested an explanation of the effect by the presence of alternating Ag-rich and Al-rich layers. Massalski and 9 King described the £ phase in the Ag-Al system as a typical electron compound with ratio — = 1.48, similar to that present in Ag-Cd and Ag-Zn alloys, where e = number of valency electrons and a = number of atoms in the-compounds. Hawbolt and Brownfound that the g-phase transformed either to massive Z, or martensite, depending on the quenching conditions. Arias and K i t t l " ^ obtained the following phases on quenching from 750°C: - 3 -at % A l : Ice Brine: Ice Water: Air 23.2 Mart. + C* (traces) a + £ + Mart. £+ct m m m 24 a + r + Mart. E+a m m^ 24.6 ?. + Mart. + 6 % m 25.1 5 + B ' 5+y m 25.9 K, £+y 27.9 C m C+y 12 Hawbolt and Massalski suggested the Ag-Al system as being ideal for the study of massive reactions. The following reactions were ,12 investigated ?m (c irregular grains often striated) B — * y m (y m spherical) z, — » - y m (y very irregular patches) To retain(i-phase at room temperature, cooling rates of the order r -1 /\5 j 12 of 10 deg. per. sec. were necessary. 13 Spencer and Knight found that the peritectoid reaction a + t, —»• y in the 23% a l was extremely sluggish. The reaction was not completed after 17 hours at 428°C and the alloy consisted of two phases a + y, contrary to the equilibrium phase diagram requiring a single y-phase only. The observation implies that the equilibrium phase diagram cannot be used without caution when an attempt i s made to predict certain phases at certain temperatures after annealing periods of the order, of - 4 -minutes or a few hours respectively. PURPOSE OF PRESENT INVESTIGATION The present investigation was initiated to examine thejnorphology and kinetics of the phase transformations that occur in the composi-tion range Ag-20 at % Al to Ag-27 at % Al. Both, continuous cooling and isothermal transformations were to be examined. 1. EXPERIMENTAL PROCEDURES The following alloys were used in the investigation: Ag-20.1 at % Al Ag-23 at % Al Ag-24 at % Al Ag-25 at % Al Ag-26 at % Al Ag-27 at % Al Alloys 20.1 and 24 at % Al provided a convenient means of assessing the composition control throughout the experimental work. Any loss of Al in the 20.1 % alloy should result in the disappearance of the C phase. Whereas any loss of Al in the 24% alloy should give rise to the appearance of some a-phase for temperatures between 610-448°C. - 5 -Most of the work was done on the alloys containing 24 and 25 at % Al. 1.1 Alloys Preparation ! Alloys were prepared by induction melting the appropriate weights of silver (99.999%) and aluminum (99.999%) in a graphite crucible. The i n i t i a l melting produced a nonhomogenous ingot 0.6 inch in diameter. This material was induction remelted and c h i l l cast into, a water cooled 0.7" x 2" x 0.2" copper mold. A l l melting and casting was performed under a helium atmosphere. The design of the melting furnace and water cooled mold i s shown in Fig. 5. The weight of the charge before melting and the weight of the ingot after were used as a measure of the composition control. In a l l cases the weight loss was no more than 0.005%. The composition was also checked by a gravimetric measurement of the Al content. This analysis involved a dissolution of the alloy in an HNO^  solution and re-precipitation of Al^O^ in NH^ OH. The resulting analyses were in good agreement with the nominal compositions. Flat ingots were desirable for several reasons: (a) C h i l l casting was more effective and macrosegregation was minimized. (b) Sheets rolled from f l a t ingots had a constant degree of total deformation. - 6 -1.2 Homogenizing Before homogenizing, the ingot was hot rolled at 700°C (25% reduction). The rolled product was then covered with fine graphite chips and annealed^in an argon atmosphere for 30 hours. After annealing, the samples were cooled to 450°C at approximately 300°C/hour and from 450°C to room temperature at approximately 100°C/hour. 1.3 Hot Rolling The homogenized ingots were hot rolled at 650°C, employing reductions of 30% per pass. The f i n a l sheet thickness was limited for hot rolling as the rapid cooling of the thin sample resulted in cracking. The optimum f i n a l thickness for hot rol l i n g was 0.100". 1.4 Cold Rolling Cold rolling was possible when alloys were quenched from either the t, or the 8-phase fields. The quenched £-phase could not be deformed at room temperature by more than 20% without cracking. To continue cold r o l l i n g i t was necessary to anneal at 500°C and requench the sample. Most samples were fabricated by cold rolling of the homogenized slabs. - 7 -1.5 Cold Drawing of Wire Wires of diameter 0.060" were produced by repeated drawing after quenching from the £-phase field.- To produce a wire of a given diameter from a b i l l e t 0.140" x 0.140", required 8-10 draws with intermediate heating at 550°C followed by quenching. 1.6 Heat Treatment A vertical tube furnace was designed to permit: (a) rapid change of the temperature . (b) temperature control + 2°C (c) protection against oxidation and contamination (d) rapid quenching from elevated temperature. The tube furnace i s shown in Fig. 6. Temperature control was 1 maintained by welding a control thermocouple on the external surface of the tube (NO. 4, Fig. 6). A second thermocouple was inserted through the open valve 7 and positioned adjacent to the sample laying on the supporting graphite disc (No. 6). The temperature obtained from the inserted thermocouple was considered as a true temperature of the sample during heat treatment. Heat treatment and quenching of the specimen were possible without transfer of the specimen from the furnace. The heat treatment to obtain the required structure included heating, short evacuation period and quenching (Fig. 7). Quenching was accomplished by a jet of water directed onto the specimen in the centre of the tube. The water jet did not contact the tube wall. Upon impact of the water jet the flexible supporting disc (No. 6, Fig. 6) was instantly released from the specimen holder (No. 5). The disc and the specimen f e l l into the bottom container which was then f i l l e d up with a liquid from the top container. 1.7 The Preparation for Optical Microscopy The quenched C-phase could not be polished mechanically as this structure was strongly affected by deformation. The y-phase could be mechanically polished without d i f f i c u l t y . A l l samples for metallographic examination were electropolished using the apparatus shown i n Fig. 8. Optimum conditions: AC current Electrolyte: 10% KCN Voltage: 60-95 V Distance B: 10-15 mm No additional etching was necessary as the resulting structures were already lightly etched. The size of the sample was limited to 0.125" in diameter as over-etching occurred at larger distances from the centre of the jet. 1.8 Preparation of Surfaces for Quantitative Metallography A distinct contrast between the z, and y-phases was an essential requirement for kinetic measurements using the IMANCO QUANTIMET 720. - 9 -The procedure described in Section 1.7 did not produce^necessary optical contrast between the two phases. The following procedure gave good results: The surfaces were ground on emery papers (up to the finest, grade 4). The distortion of the £-matrix phase was heavy but did not affect the contrast at the £-y interfaces. The sample was then etched in the solution: 20 gm chromium trioxide 1.5 gm sodium sulfate 3 ' 100 cm water The etching time varied from 1 to 5 minutes. 1.9 Preparation of Foils for Electron Transmission Microscopy £-phase: Bulk f o i l s were thinned using the electropolishing apparatus previously described (Fig. 8). Prior to perforation, the f o i l was repeatedly turned to expose both sides to the jet polish. Mixtures of £+y and pure y: Good results were obtained when using a window el-polishing method and.the following polishing conditions: Electrolyte: 3.0 gm KCN 1.5 gms sodium hydroxide 0.5-2 gms silver cyanide -lo-co. 3 gm potassium carbonate 3 100 cm water temperature 0°C DC current - 5-10 volts stainless steel crucible as a cathode. 2. GENERAL PROPERTIES AND MORPHOLOGY OF PHASES IN COMPOSITION  RANGE Ag-20.1 TO 27 AT % Al 2.1 Preliminary Results In order to specify standard procedures for specimen treatment, i t was necessary to know how different alloys responded to heat treatment and deformation. Symmary of Preliminary Observations: 1. The £-phase was retained to room temperature by quenching in water. It was very ductile i n the quenched state though exhibited rapid work hardening and became b r i t t l e when cold deformation exceeded approximately 20%. Cold deformation was accompanied by sound effects similar to "crying" during twinning in t i n . 2. The y-phase as produced from the £-phase by slow cooling to room temperature, was extremely b r i t t l e in the alloys 23 and 24 at % Al. The alloy 25 at % Al cooled at a rate of 100-200°C/hour to room temperature exhibited an intercrystalline fracture, but was relatively ductile compared to 23 and 24 at % alloys. 3. The 25 at % alloy became as b r i t t l e as alloys 23 and 24 at % when i t was cooled from the t, f i e l d to room temperature at 40 deg. per - 11 -hour (and slower). 4. The quenched r, phase of the 23 and 24 at % Al alloys became very b r i t t l e under the following conditions: Annealed Cold rolled sheet, cast or cold drawn wire hot rolled b r i t t l e after material at b r i t t l e after 200° 1 min. 8 min. 400°C 10 sec. 1 min. In a l l cases the heat treatment prior to tempering was the same: Heating at 700°C for 15 minutes, cooled i n air , reheating at 500°C for 20 minutes and quenched in water. Differences recited to the fabrication procedure were not completely eliminated by the short annealing at 700°C. This may explain the different behavior of the structures during tempering. The drastic change of properties during the transition t, -*• u in the 24 at % alloy are demonstrated in the following sequence of photographs: Fig. 9 - Material cooled from 700°C at 40°C per hour to room temperature - broken, as dropped on the floor. Fig..10 - Material quenched from 550°C, drawn to wire. Fig. 11 - Wire produced in the quenched x, state and subsequently tempered at 200°C for 1 minute; fragmented by dropping on the floor. Fragments could be mechanically crushed into fine powder, - 12 -without obvious plastic deformation of the particles. 2.2 Hardness of the g and y-phases It was found that the hardness increased by X4 to X5 when the 24 at % alloy was transformed from the £ to the y-phase. Generally in a l l alloys except 27 at % the presence of the y-phase was accompanied by an increase in hardness. It was impossible to establish i f some precipitation hardening of the ^-matrix took place as micro-hardness indentations on the ^-grains were distorted. For this reason, fine grained specimens and a macro-load of 5 kg. were used to obtain an average value of the hardness of the two-phase mixtures. The variation of hardness with Al content for structures produced by different heat treatments i s shown on Fig. 12. A l l quenched samples had a grain size of approximately 200 microns. Cast, homogenized and slowly cooled (at 0.4°C/min.) samples had a non-uniform grain size. The results shown in Fig. 12 reveal a peak hardness at 24 at % Al. An examination of the phase diagram (Fig. 1) would indicate that i f hardness is related to the presence of the y-phase, the alloy 23 at % Al should have at least a hardness equivalent to that exhibited by the 24 at % Al alloy. 2.3 X-Ray Phase Fields Determination This part of the investigation was conducted to establish the phases that were present when the annealing times (and cooling rates) - 13 -were different from those used to construct the equilibrium phase diagram. Experimental Details: Polycrystalline rod specimens produced by electropolishing cold drawn wire were employed. The diameter of the rod was 0.020" max. Debye-Scherrer patterns (D.S.) were obtained using the 57.4 mm diameter cylindrical camera and CuKa radiation. Heat treatments for the different alloys are given in Table 1. A l l structures were produced by continuous cooling either from 700°C or 500°C. Continuous cooling from 700°C i n alloys 24, 25, 26, and 27 at % Al resulted in a large grained structure unsuitable for D.S. patterns. Samples from these alloys were annealed at 700°C for 40 minutes, quenched in water, re-heated to 500°C and continuously cooled to lower temperatures as indicated in Table 1. Prior to quenching, the samples were held at temperatures for varying periods of time. TheJow temperature structures result therefore from phase transformations taking place during continuous cooling from high temperatures and isothermal heating at the given lower temperature. Cooling rates were: 700-500°C, 10°C/min; 500-200°C, 7°C/min. - 14 -TABLE 1 At % Al 700°C min Quenched from temperature 500°C 435°C 400°C 300°C 200°C min min min min min min 20,1 301 ^ 30 », * AO ' V - — • — 60 120 23 30 i s» 30 " 1- *• 40 1 — — s — 60 24 30 i 30 h- 30 a » 40 H 5 ^ 4 0 i- » ^ 6 0 H " — » 1 2 0 25 r- 30-* »_ _ ^ 40 (- " =-40 H " *~ 60 H • r- : B—120 X-Ray D.S. Pattern Evaluation The phases to be expected over the given composition range at temperatures below 605°C are: a-phase f.c.c. structure £-phase h.c.p. structure y-phase 8-Mn structure - 15 -a-Phase: ' 0 The lattice parameter, a = 4.0528 A at the s o l u t i l i t y limit 20 At % Al, aid temperature 500°C was used for indexing this phase."^ The calculated values of the 26 angles and the corresponding reflecting planes are list e d i n the Appendix. £-Phase: The reported la t t i c e parameters of the 2;-phase at 27 at % Al are"'": a = 2.877 A c = k .662 A ±= 1.62k a A l l quenched ^-phases (500°C) gave comparable patterns within the composition range 20.1-27 at % Al. The £-phase pattern from the alloy 24 at % Al was used as a standard for the determination of this structure in the other alloys. The pattern was analyzed by a graphical method (Bunn chart). Angles 20 and corresponding planar indices are listed in Appendix. (Used notation h, k, 1 which i s equivalent to a notation h, k, 1, 1 where h + k = - i ) . u-Phase: This phase has been described in the literature as complex cubic, isomorphic with B-manganese structure, having the lattice parameter - 16 -15 2 a = 6.92. The calculated values of sin 0 for this type of structure are list e d in Appendix-Tab. 2, column III. The values correspond to CuKct radiation. Alloy 20.1 at % Al The set of D.S. patterns is shown in Fig. 13. 500°C: Distinct a lines and weak £ lines are evident, in agreement with the phase diagram. Weak ? lines can be found at 20 = 41°, 53.5°, 71°...etc. 400,300,200°C: The intensity of the a lines does not change with . decreasing temperature. This suggests that the ct-phase does not participate in the y-phase formation. The y-phase appeared at 400°C as weak lines at 20 = 43.5°, 49°, 56°, 59°...etc. According to the phase diagram, the amount of y-phase should increase as the temperature i s lowered and the y pattern should gain in intensity. Patterns from 400, 300, and 200°C, show that this did not happen under employed conditions. Alloy 23 at % Al The set of D.S. patterns i s shown in Fig. 14. 500°C: The structure i s a mixture of the t, and the a-phase as expected from the phase diagram. - 17 -S lines at 29 = 36°, 38.5°, 41°, 54°...etc. a lines at 26 = 45°, 98.5°, 1110...etc. 40O,30O°C: The ft.-phase appears in the 400°C pattern, u-lines at 29 = 22°, 26°, 29°, 31°, 43.5°...etc. The a-lines coincide approximately with the position of the y-lines and thus no comment can be made as to the decomposition of the a-phase. The C-phase changed readily; at 400°C the z, lines for the (100) and (302) planes disappeared (29 = 36 and 160 degrees, respectively). There i s some evidence of a preferred orientation in the y-phase (arcs instead of continuous diffraction lines are v i s i b l e in the u-patterns). Alloy 24 at % Al The set of D.S. patterns i s shown in Fig. 15. 500°C: Only lines of the•£ were present. 435°C: A sharp and relatively serong y-pattern has developed with l i t t l e change in the ^-pattern. y lines at 29 = 29°, 43.5°, 49°...etc. 400°C: The (100) C-line at 36.5 degrees is not visible on this pattern. This could indicate that the structure has already transformed to the y-phase. No changes were observed on patterns taken from temperatures below 400°C. This implies that this temperature is in the single phase f i e l d . - 18 -Alloys 25, 26, 27 at % Al Employing a cooling rate of 7°C per minute from the £-phase f i e l d , only ^-patterns were present at a l l temperatures below 448°C (Fig. 16). Thus the cooling rates and isothermal treatments did not allow the expected transformations to occur. It was found that the change of cooling rate had a more pronounced effect on the £—s»-y transformation in the 25, 26, 27 at % alloys than prolonged heating periods. Cooling of the 25 at % alloy from 500°C at a rate 6f 0.3°C per minute to 375°C resulted in the structure corresponding to the two-phase mixture (£ + y) , Fig. 17. The diffraction patterns of the room temperature structures were taken from samples cooled from 500°C to room temperature at 0.6°C per minute. Alloy 23 at % - Fig. 18 Alloy 24 at % - Fig. 19 Alloy 25 at % - Fig. 20 In the alloy 25 at % Al (Fig. 20) the presence of £ is indicated i. by the strong widely separated arcs and the presence of the y-phase by the more continuous weaker lines. Alloys 24, 25, 26, and 27 at % Al developed large grains of the £-phase upon slow continuous cooling from' 700°C. The grain size of this £ was 3 to 5 mm. in diameter. Thus i t was possible to examine the phase changes in single grains comparable to single crystals. The single crystals of the £-phase in the 24 at % alloy always transformed during continuous cooling to room temperature into the y-phase structure having a large number of - 19 -crystallites. X-ray photographs from the C-single crystals after anisothermal transformation to the y-phase indicated that the y-crystallites were not randomly oriented. The Laue pattern of the V (equilibrium y-phase at room tempera-ture) produced by cooling at 0.4°C/min. of the alloy 24 at % is shown in Fig. 21. On the background of the Laue photograph (Fig. 21) a diffuse single crystal pattern i s vi s i b l e . As the lattice parameter of the y-phase is large, by using the Mo radiation the prominent Debye rings formed by characteristic radiation should be concentrated at low diffraction angles - i.e., absent on the back reflection photograph. Fig. 22 shows the same sample as Fig. 21 but obtained using Mo white radiation. The faint single crystal zones indicate that the original £-phase persists in some small amount in this structure. The retained £-phase i s , however, broken down into a fine substructure. In order to detect the "background" patterns in y structures RT i t was necessary to operate the Mo-tube at least at 38 KV and 15 mV. The marked effect of the cooling rate on phase transformations in 25 at % alloy was also observed on the single crystals. Fig. 23a shows a Laue pattern of the sample cooled from 700°C at 7°C per minute to room temperature. The pattern belongs to the untransformed t,-phase. Upon cooling at a rate of 0.3°C per minute the transformation took place - as seen in Fig. 23b. The background t;-patterns were more obvious in the room tempera-ture structures of the alloy 25 at %, as shown in Fig. 24 (Mo white radiation). The difference in intensities of the background patterns - 20 -on the Laue photographs taken by Cu and Mo radiation may be attributed to the different operating conditions (Cu-tube operated at 33 KV and 14 mV). The results obtained using D.S. and Laue X-ray techniques showed that the x, y transformation occurs rapidly in the alloy 24 at % Al and extremely slowly in the alloy 25 at % A l . 2.4 Optical Metallography of Phases In the Composition Range  Ag-20.1 to 27 at % Al To reveal the typical phase morphology for the experimental conditions the following structures were produced: (a) Structure at 500°C (b) Structure of two-phase mixture (c+y) (c) Room temperature structures (R.T.) The following heat treatments were used to obtain the required structures: (a) Heating at 700°C for 30 minutes, slow cooling to 500°C,. holding at temperature for 30 minutes and quenching in water. (b) Heating at 700°C for 30 minutes, continuous cooling to 435, 400 and 300°C, holding for one hour at the specified temperature and quenching i n water (only for alloy 24 at % Al). (c) Heating at 700°C for 30 minutes, continuous cooling at 0.3°C per minute to room temperature. For the 24 at % alloy,cooling rates of 0.6 and 0.1°C per minute were used to obtain the y-phase. - 21 -Alloy 20.1 at % Al 500°C structure: Consists of twinned a-grains and the £-phase arranged along the r o l l i n g direction (R.D.) -Fig. 25. R.T. structure: The y-phase forms a continuous network on the a-grain boundaries - Fig. 26. The amount of the y-phase in the R.T. structure is substantially less than i t would be anticipated from the equilibrium phase diagram. The amount of y can be roughly estimated as equal to the ("-phase amount present at 500°C. This observation agrees with X-ray results (intensities of t, and y lines on D.S. patterns). Alloy 23 at % Al 500°C structure: The a-phase in the S-phase matrix formed non-quiaxed grains - Fig. 27. Lenticular grains were also frequently observed. R.T. structure: Some a-phase particles were internally twinned -Fig. 28. Alloy 24 at % Al 500°C structure: (C+y) at 435°C: C-phase only - Fig. 29. The structure i s shown in Fig. 30 and 31. The - 22 -structural features in Fig. 30 are i n a more defined crystallographic relationship than those in Fig. 31. (£+u) at 400°C: The structure s t i l l consists of two phases -Fig. 32, and i s comparable with the partially transformed areas in Fig. 30 and Fig. 31. The observation indicates that this small amount of the S-phase was not detected by the X-ray D.S. technique. 300°C structure: The micrograph shown in Fig. 33 represents this structure. No second phase (?) can be identified in the photograph. R.T. structure: It was found, that the 24 at % always transformed into the b r i t t l e y-phase when cooled from 500°C in air or in a furnace cooled at 7°C per minute to room temperature. Metallography established that e the y structure in 24 at % alloy i s built up of Kl plates or layers arranged into larger blocks. The evidence is presented in the following figures: Fig. 34 - Cooling rate 0.6°C per minute, fractured surface unpolished, polarized light/ Fig. 35 - Cooling rate 0.1 deg. per minute, electro-polished (-etched). Fig. 36 - Cooling rate 0.1 deg. per minute, electro-polished (-etched). Bands of y-plates with different orientation with respect to the sample surface. - 23 -Figure 37 - Same sample as in Fig. 36 (different place), using polarized light. The illumination of polarized light reveals mutual overlapping of several plates. The structure responded to polarized light because i t was deliberately over-etched for this purpose. Alloy 25 at % Al 500°C structure: S-phase only - Fig. 38. R.T. structure: The transformation product after cooling from f i e l d at 0.3-0.6°C per minute resembled a typical Widmanstatten structure - Fig. 39. Fig. 40 shows a partial section in the plane parallel to the Wi dmanstatten plates. Upon cooling at 7°C per minute, the t, phase was retained to room temperature (already shown using X-ray methods). Alloy 26 and 27 at % Al 500°C structures: Identical with those in 24 and 25 at % alloy. R.T. structures: After cooling at 0.4°C per minute to room tempera-ture, the structures contained a small amount of Widmannstatten-like y-phase. This i s shown in Fig. 41 (26 at % alloy) and Fig. 42 (27 at % alloy). 3. TWO-PHASE EQUILIBRIUM STRUCTURES 3.1 ( g + u ) Mixture in 24 at % Alloy The following heat treatment schedule was employed to obtain the equilibrium 2-phase structure. 700°C - 1 hour The following figures show the resulting structure obtained in two different samples: Fig. 43 - Sample having uniformly distributed £-islands within u-matrix. Fig. 44 - Microstructure show^ing long y-plates and larger areas of untransformed C . - 25 -Large patches of the S-phase readily transformed at 200°C into another phase which was not comparable to any phase obtained under continuous cooling. Hardness measurement showed that the product i s the y-phase. This structure i s shown in Fig. 45. Phases are marked e as: SQ =» s-phase quenched to room temperature; y = equilibrium y-phase at 435°C; y* = y-phase obtained by isothermally heating the Sq-phase at 200°C for 4 minutes. Small separated s^-islands shown in Fig. 43 did not transform to y^ even after 30 minutes at 200°C. This implies that these areas are enriched in H. and that s—»- V transformation i s much slower as the Al content increases. X-Ray Analysis: A 0.020° diameter rod was cut from a sample having the structure shown in Fig. 43. The rod was electropolished and D.S. and Laue patterns were obtained using a cylindrical camera. The patterns are shown in Figures 46, 47, and 48. The inspection of Laue patterns (Fig. 47, 48) shows several important details. There i s a similar single crystal pattern to that found in the e y^-structures. This presumably belongs to the s-phase. The y-phase is present as a larger number of smaller crystallites giving rise to the Debye rings. A l l intense smeared arcs seen i n Fig. 47 are common to both the S-zones and the Debye rings of the y-phase. This is evidence that the y-phase may be nucleated from the S by some mechanism of atoms - 26 -re-shuffling within the r, crystal. The Laue pattern in Fig. 47 cannot be explained i f long range diffusion i s considered as the only mechanism operating in this stage of the transformation. As the y-phase appears within the £-phase, the £-phase undergoes homogenous breakdown into a very fine substructure. Mo-radiation is suitable for revealing the substructured ^-pattern as the interfering y spots are effectively removed from the pattern (Fig. 48). The experimental results showed that two processes occur when the X transforms to y by continuous cooling. (a) A breaking of the perfect single crystal t, into a fine substructure. (b) Nucleation of the y-phase within the ?-matrix (presumably with some coherency) and a development of a polycrystalline y-structure as the transformation proceeds. t • To establish the true sequence for these processes, the following experiments were conducted: Samples were continuously cooled through the two-phase f i e l d , cooling being stopped for short periods of time at certain temperatures and the sample was then quenched. By quenching from progressively , lower temperatures the z, — » - y transformation could be examined. The following heat treatment was employed: - 27 -700 °C cooling at 5 deg. per min. 15 min. 415°C 15 min. 400°C 300°C 1 hour r Room Temperature f f B C D 7 hours Laue patterns were obtained from the samples using Cu white radiation. A cooling rate of 5 deg. per minute suppressed the trans-formation - start i n the two-phase region as sample A showed the C-pattern only (Fig. 49). The homogenous breakdown of the £-phase into a substructure obviously preceded y-phase formation in sample B as shown in Fig. 50. The Debye arcs of the y-phase are absent, i.e. , the y-phase starts to develop at the later stage. The pattern from sample C (400°C) indicates a stage of the y-nucleation within the substructured ^-matrix -(Fig. 51). Metallographlc examination of sample B showed the r, .structure only. Samples D and E as quenched from 300°C after 1 and 7 hours had similar X-ray diffraction patterns and morphology. The Laue pattern - 28 -and microstructure of sample E (held 7 hours at 300°C) i s shown in Fig. 52. This structure i s a single phase and is comparable with the structure shown in Fig. 33. The Laue pattern is very similar to y -patterns obtained from 24 at % alloy. It also shows a diffuse single crystal pattern in the background. In order to obtain information on the substructure in the ^-matrix the sample was heated to 700°C, cooled to 435°C at 7 deg. per minute, held at this temperature 30 minutes and quenched. A thin f o i l from this sample was examined using transmission electron microscopy. The structure exhibited parallel fringes but showed no evidences of equiaxed subgrains - Fig. 53. The selected area diffraction pattern (Fig. 54) confirmed that the structure i s similar to that of the substructured ^-matrix as detected by X-ray photographs. Each diffraction spot is s p l i t into several spots. The conclusion from this observation is that the substructure can be qualitatively described as a small displacement of thin layers within the ^-structure. 3.2 (g+y) Mixture in 25 at % Alloy An equilibrium two-phase structure was produced by c ooling from 500°C to 400°C at 0.3°C/min., holding at this temperature for 200 hours and quenching. The resulting microstructure is shown in Fig. 55 (chemically etched) and Fig. 56 (electropolished). A Debye-Scherrer X-ray pattern is included in Fig. 57. The structure is comparable to that obtained by continuous cooling to room temperature (Fig. 39). The spacings between the lamellae is however, larger. - 29 -An electron microprobe analysis (Fig. 58) showed that the lamellae are rich in aluminum - i.e., they are the £-phase. Equilibrium heating for 200 hours at 700°C has allowed the y-plates to grow together to form the continuous matrix phase. The D.S. X-ray pattern in Fig. 57 shows only a few t, spots consistent with the large grain . -phase size of the original £ (arcs at 26 = 36, 54, 78 degrees). The resulting y-phase i s more randomly oriented giving rise to more continuous Debye rings (26 = 29, 43, 56,...degrees). This suggests that i f the morphology of the y-phase results from the presence of a unique crystallographic orientation relationship with the ?-phase, several variants are available. 4. COMMENT ON SURFACE METALLOGRAPHY OF STRUCTURES The transformation phenomena i s many alloys can be conveniently followed on free surfaces, polished prior to the transformation event. However, for some alloys the observations on surfaces may not be representative of the processes occurring in the bulk of the material. A sample from the 25 at % alloy was smoothly electropolished, heated at 500°C for 20 minutes and quenched. The structure as seen on the prepolished surface using polarized light is shown in Fig. 60. The surface was then lightly electropolished and the same area was examined using polarized light. The structure i s shown in Fig. 61. * The JXA-3A Electron Probe was used. - 30 -From a comparison of Fig. . 60 and 61 i t is evident that the boundaries of the structural features were not exactly reproduced on the pre-polished surface. This implies that Ag-Al alloys may not be suitable for an evaluation of transformation processes from surface observations. 5. DECOMPOSITION OF THE QUENCHED B PHASE The structure of the 24 at % alloy as quenched from 700°C into brine at 0°C i s shown in Fig. 62. The B-phase was not retained by quenching. The thickness of this sample was 0.030". A comparison 12 with published structure of the quenched B suggests that the structure in Fig. 62 i s a massive c;-phase. However, i t was found that when the sample thickness did not exceed 0.005", the g-phase could be retained to room temperature by quenching into brine (-4°C). 5.1 Metallography of the g^-Phase Decomposition - Alloy 24 at % The structure of the quenched g-phase, four days after quenching, i s shown in Fig. 63. The second phase on the grain boundaries is the £-phase that was present in the as-quenched structure. When the it structure was heated to 100°C remained unchanged after a heating period of 2-3 hours. After 10-15 hours, distinct dark patches appeared and grew with increasing time. Fig. 64 and 65 show colonies - 31 -of the product-phase growing from the quenched B ( B Q ) after 43 hours at 100°C. When the Bp-phase underwent tempering at 200°C, the decomposition proceeded rapidly. The product phase had a regular spherical shape when growing in the grain interior. This morphology is identical with the 12 published results. The progress of the transformation at 200°C i s shown in the sequence of micrographs - Fig. 66, 67 and 68. Fig. 66 - BQ + 20 seconds at 200°C Fig. 67 - Bp + 90 seconds at 200°C Fig. 68 - Bp + 210 seconds at 200°C. The product of the BQ decomposition at 200°C responds to etching quite differently f rom that produced at 100°C. The regular spherical shape of the growing precipitates allowed the measurement of radial growth with respect to time. The radius (R) versus the annealing time at 200°C i s shown in Fig. 69. The R values represent an average size of 3-4 of the largest precipitates found in the micrographs after the indicated annealing times. The growth i s a linear function of time. It was observed, that thin layers of the u-product formed on prepolished surfaces during the 8—*- y transformation (Figs. 66 to 68 represent structures in the sample interior). 5.2 Pre-Precipitation Phenomena During Bn~Phase Decomposition The experimental work conducted in the present study revealed - 32 -several significant characteristics of the transition from the high temperature quenched g-phase to the equilibrium y-phase. The gg-phase was found to be unstable at room temperature. The * hardness increased from HV115 to HV140 within 4 days after quenching. No second phase was visible in this structure. The age hardening of the g^-matrix was established by tempering at 100°C. The hardness change versus ageing time is shown in Fig. 71 (bottom curve). No precipitates of the second phase were detected before the peak hardness (HV = 240) was reached. When metallographically visible precipitates started to grow in the matrix, the hardness of the matrix decreased. The microhardness of the precipitates was comparable to that previously found in the y-phase. At 200°C the decomposition of the g^-phase was so rapid that hardening of the matrix was precluded by the development of the much harder product phase. An i l l u s t r a t i o n of the two different hardening processes that take place during the decomposition of thehigh temperature g^-phase in the 24 at % alloy i s shown in Fig. 71. The hardness of the y-precipitate is plotted i n the upper part of the diagram. 5.2.1. X-Ray D.S. Sequence of the Pre-Precipitation Stage To further examine the pre-precipitation hardening process, D.S. patterns were taken from the sample as ageing progressed at 100°C. The hardness was measured on the Tukon hardness testing apparatus using a load of 50 gram - 33 -The sample was a 0.003" diameter rod quenched from 700°C into brine at -4°C. D.S. X-ray patterns obtained after consecutive ageing periods are shown in Fig. 72. Pattern a - g^-phase b - gp-phase 60 sec. at 100°C c - gg-phase 1 hour at 100°C d - gg-phase 49 hours at 100°C e - gg-phase 49 hours at 100°C plus 45 minutes at 200°C. For an evaluation of the structural changes occurring in the pre-precipitation stage, only patterns, a, b, and c can be used ( i t was established metallographically, that at longer ageing periods the true precipitation took place). From patterns a, b, and c the following conclusions can be drawn. (1) The g-phase does not exhibit any long-range ordering during quenching. (Pattern "a" in Fig. 72 i s the b.c.c. structure and is analyzed in the Appendix). (2) During pre-precipitation the g-matrix is significantly distorted and undergoes some structure changes. The source of distortion cannot be the presence of a visible second-phase as no such phase was detected microscopically. It could be due to processes occurring within the original matrix on a very fine scalej i.e. clustering or G.P. zones formation. The possibility of the existence of a transition phase in the early stage of the g -decomposition is shown in the pattern c (Fig. 72). - 34 -The nature of this structure (designated as g) may be investigated in terms of i t s similarity to the y and ^-phases respectively or may be analyzed independently on these phases. In pattern c (Fig. 72) the 3 line at 40 degrees is accompanied by two s a t e l l i t e lines (39 and 41 degrees). However, these lines are also present in the u-pattern (e). The strong line observed at 79 degrees in pattern c i s missing in the u-pattern and thus no obvious relationship exists between the g and the y-phase. A comparison with the ^-pattern shows that the s a t e l l i t e lines at 39, 41, and 79 degrees can be s-lines. A geometrical comparison of the patterns does not permit any conclusions to be made about the nature of the g-phase. There are several alloy systems reported injthe literature which. exhibit precipitation hardening and form transition structures due to the sinusoidal variations in composition or lattice dimensions within 16 the matrix phase. The cluserting of one element in an Ag-Al binary alloy would introduce a variation in the scattering power and the lattice parameter (provided the clusters are large). Thus the g could have some similarity to a modulated structure. 5.2.2. X-Ray Laue-Transmission Sequence of the g^ Decomposition Transmission Laue photographs were taken from f l a t single crystals of the g-phase diring decomposition into the y-phase. Photographs were obtained using Fe white radiation. The thickness of samples was 0.003". - 35 -Fig. 73 shows the transmission Laue patterns corresponding to the following treatments: Patter a - Bp-phase " Bq-phase 60 sec. at 100°C c - Bg-phase 1 hr.'at 100°C d - Bp-phase 43 hrs. at 100°C and 45 min. at 200°C. The sequence shows that changes prior to the visible precipitation of the y-phase are associated with a homogenous distortion of the 8Q-matrix. The streaking of a l l spots can be seen in pattern b (after 60 seconds of ageing). Patterns b and c were taken from the same sample in an identical position. Several zones are common to both patterns. This implies a close similarity between the original structure and the transition phase. Numerous streaked spots around the centre in pattern d (Fig. 73) indicate that a great number of y-crystallites nucleated within a single crystal of the B-phase. Originally i t was proposed that i n the composition range 20-27 at % Al, only the four equilibrium phases a, B> S> and y are involved.' The possibility of the formation of a transitional, metastable phase and precipitation hardening prior to decomposition of thejhigh temperature phase into the equilibrium y-phase was not considered. The separation of one element into zones under equilibrium conditions depends on the thermodynamic properties of a given alloy at a particular temperature. The implication for the Ag-Al alloy i s that clusters might form either during the reaction B—>• v or c , y . - 36 -The kinetics aid mechanism may be quite different when different structures are involved. 6. DILATOMETRY OF ANISOTHERMAL PHASE TRANSITIONS IN ALLOY 25 AT % Al To obtain information about the volume changes associated with the equilibrium phase transformations, dilatation of the sample was recorded during continuous heating and cooling. This technique also provided a measure of the temperature associated with the onset of the transformations for various cooling and heating rates. A THETA dilatometer with a horizontal furnace was used. The sample size was 0.225" x 0.975" x 1.869". A constant heating-cooling rate df 30°C/hour was employed. The starting structure was the u D r r-phase obtained by cooling the material from 700°C to room temperature at 0.4°C per minute. On heating to 680°C at 30°C/hour (Fig. 74) the smooth curvature in the i n i t i a l portion of the curve is indicative of the smooth transition from the y to the s-phase. It appears to occur over the, temperature range starting approximately at 295°C (2 mV). Just above 6058C the S —^- 3 reaction takes place rapidly with a marked increase in the specimen length. On cooling from the 3~phase the 3 transforms back to the £ at just below 605°C. The small displacement of the transformation temperature of the £ reaction implies that the reaction can occur without d i f f i c u l t y . From an inspection of the phase diagram i t follows that the £ and 3-fields are not separated by a wide - 37 -two-phase f i e l d at the composition 24 at % A l . Therefore the composition invariant (polymorphic) transformation is possible. This would explain the shape of the dilatometric curve at 605°C. The sharp change in length at 392°C during cooling suggests that the s-phase w a s supercooled and remained untransformed while passing temperatures corresponding to the two-phase f i e l d . Upon reaching the boundary (s+y)/y (392°G) the rapid, composition invariant transition s —y occurs as shown by the sudden AL change over a small temperature range. When the dilatometric experiment was repeated under identical conditions except the heating stopped just below the s — t r a n s i t i o n temperature, the cooling curve did not exhibit the sharp change indicating the s y transition (Fig. 75). This observation was interpreted as evidence that the s-phase when produced by two different thermal treatments may not be exactly the same. The difference could be associated with the presence of inhomogeneities which may not be completely dissolved at temperatures within the S-field. 7. ANALYSIS OF THE s-PHASE - 24 AT % Al The crystal structure of the s~phase w a s examined using both powder and rod shaped specimens. The powder was heated at 700°C for 30 minutes, cooled to 500°C, held 15 minutes at temperature and quenched. F\ D.S. diffraction - 38 -pattern from the powder i s shown in Fig. 76 (pattern 66). The solid polycrystalline rod (0.020" diameter) was quenched from 630°C, reheated to 500°C - held 1 hour at 500°C and quenched. This procedure yielded the reasonably small grain size required for satisfactory Debye rings for lattice parameters measurements. The diffraction pattern from the polycrystalline solid rod is shown in Fig. 76 (pattern 67). Both patterns in Fig. 76 were taken using CuKa radiation. An inspection of pattern 66 (powder) shows lines at 20 - 45 and 99 degrees. These lines should not be present in the (".-phase pattern and were always missing on patterns obtained from the electropolished solid rod samples. The spurious lines were attributed to the loss of aluminum from the large surface area of the powder particles (the lines at 45 and 99 degrees are consistent with a-phase lines). For this reason, powder specimens were not used for D.S. patterns. Pattern 67 is indexed in Appendix, Table 1. The lattice parameters calculated from pattern 67 were: o a = 2.8692 A o c = 4.6706 A - = 1.6278 a The u/fpublished data for an alloy 25 at % Al quenched from 550°C . 7 i s : o a = 2.870 A 0 c = 4.665 A - 39 -7.1 fr-Single Crystal Analysis Rods 0.020" diameter of the alloy 24 at % were cooled slowly from 700°C to the £-field, held at temperature for 30 minutes and quenched. This treatment produced large single grains of the £-phase suitable for Laue X-ray examination. The shape of the samples permitted the use of a cylindrical camera for Laue photographs. On these photographs the continuous reflecting zones of the crystal were recorded. The analysis of the cylindrical Laue patterns was done using an appropriate scale converting linear coordinates into angular coordinates. The single crystal patterns obtained on quenching from 500°C and 460°C are shown in Fig. 77 and Fig. 78 respectively. Both patterns were obtained using Cu white radiation. Three strong dark spots and one spot exhibiting a circular halo are visible in Fig. 77. Similar halos around two prominent spots are also present in Fig. 78. Neither very intense spots, nor diffuse halos were observed in t\-patterns obtained using Mo-white radiation. . When optimum exposure conditions * and V:>.- samples with perfectly smooth, unetched were used, surfaces^, intense spots with well separated scattered halos^ were only found in the Laue diffraction patterns taken with Cu radiation (Fig. 78). Intense spots can be found on Laue pattern when a certain set of planes is accidentally in the Bragg condition for the strong Ka^ i radiation. Such an excess of radiation may produce enlarged,, intense Laue spots but cannot producev'circular halos^ surrounding the spots. The scattering was generally observed around reflections from planes - 40 -of the type: (0006), (1015), (1124) and in the back reflection region only (the intensity of radiation i s less i n the transmission region due to absorption). The effect, however, could be related to the presence of internal inhomogeneities within the ("-phase. Guinier has shown that small spherical clusters (G.P. zones) in the matrix can give rise to a scattering of radiation around each Laue spot. The scattering is very weak and can be detected only around very intensely diffracted beams. This approach, realized as a low angle X-ray scattering method, was successfully used for the it S analysis of G.P. zones in the Ag-Al system on the Al-rich side. ' )i • • 7.2 Some Aspects of Heat Treatment The dilatometry showed (Fig. 74,75) that there are differences between transformations occurring during different heating cycles though the phases involved in the reactions are expected to be the same. It was shown metallographically that the equilibrium low temperature structure in the alloy 25 at % dissolved relatively slowly at 500°C. jFig. 79 shows the structure of the 25 at % alloy after heating i n i t i a l l y at 0.5°C/min to 450°C, then heating rapidly to 500°C, holding for 5 minutes at this temperature and quenching. Within the (".-grains, there are s t i l l v i s i b l e features resembling an original Widmannstatten equilibrium structure. The experiment was repeated using identical heating conditions, except the sample was at 500°C for 75 minutes. The resulting structure is shown in Fig. 80. - h i -lt was experimentally proved that disturbances introduced into the structure during cold deformation were not eliminated by annealing within the £-field. Fig. 81 is a back reflection pinhole X-ray photograph (CrKa) of the 25 at % Al alloy quenched from 500°C, cold rolled 10% then annealed at 500°C for 15 minutes and quenched. The radial displacement of the diffraction spots i s shown i n the magnified portion of the Debye rings (Fig. 82). This observation implies that the perfection of the lattice was not restored. In order to remove the effects of prior heat treatments or deformation i t was necessary to heat the sample into the g-phase f i e l d . The standard structures of the (--phase were obtained either by continuous cooling from the g-field or by quenching from the g-field and subsequent heating to the £-field. 8. TRANSFORMATION OF THE QUENCHED c~-PHASE IN 24 AT % A1-ALL0Y Preliminary investigations established that the ("-phase as quenched (£p) is very ductile but becomes extremely b r i t t l e when tempered for short times at temperatures 200 to 420°C (Fig. 11). This drastic change in properties is associated with the transformation CQ V • The dilatometric changes recorded during continuous heating of the quenched v, from room temperature up to 600°C revealed an important * characteristic of this reaction. * The dilatometer and specimen previously described were used for this test. - 42 -Fig. 83 shows the dilatation curve for heating at a rate of 0.33°C per minute. The distinct shift of the curve at 165°C corresponds to the t,^ —»- y reaction. The transition was raised to higher temperatures by increasing the heating rate. At 0.5°C per minute the reaction occurred at 171°C. Fig. 84 i s a dilatometry curve obtained- during heating at 2°C per minute. This higher heating rate resulted in two effects seen on the heating curve: (a) A shift of the reaction to approximately 220°C. (b) An increase of the sample temperature due to the released latent heat of the transformation. The crystal structure change occurring during the £^ decomposition was established by X-ray techniques. The resulting structure i n a l l alloys investigated showed the presence of y lines after a certain period of tempering at 200°C. However, the speed of the reaction appeared to be high in the alloys 20,1, 23, and 24 at % Al and very slow in 25, 26, 27 at % alloys. Fig. 85 shows D.S. pattern of the alloy 23 at % Al quenched from 500°C and tempered at 200°C for 60 minutes. The spotty lines at 45 and 65 degrees belong to the untransformed a , and y-lines are continuous. Fig. 86 shows a D.S. pattern of the alloy 24 at % Al quenched from 500 and tempered at 200°C for 8 minutes. The pattern corresponds to a partially transformed structure (c^-lines at 36, 54, and 65, etc. degrees). The appearance of the Debye rings of the y-phase in the patterns in Fig. 85 and 86 implies that the product of the reaction t, —»- y - 43 -is randomly distributed and i t s crystallites are very small. It was shown previously that during continuous cooling of the <;-phase, the y-phase developed within a single grain of the t, as a large number of small crystallites, these crystallites not being randomly oriented (Fig. 14 shows intensity maxima on the Debye rings). However, D.S. patterns obtained from samples of small diameter and rotated during exposure may not give explicit information about the degree of randomness of the y-phase arising from the z, ^. y transformation i n bulk crystals. Transformations in fine rods and thin f o i l s which have a large surface to volume ratio may differ from the reactions occurring in bulk samples. This assumption was verified using single crystal of the £-phase (thickness 0.040") and X-ray Laue technique. Fig. 87a shows a single crystal pattern of the t^-phase. After tempering at 187°C for 5 minutes the diffraction pattern changed in a manner demonstrated in Fig. 87b. If the additional faint spots are due to a phase transformation then the regular displacement of the central zone indicates that the product phase developed by some coordinated atom displacement taking place during the transformation. In order to examine the relationship between X-ray diffraction effects and the structural changes observed metallographically during &>e — y transformation, a ^ - s i n g l e crystal of the 24 at % Al alloy was transformed isothermally at 240°C. The size of the sample was 0.250" x 0.050" x 0.300". The X-ray Laue patterns and microstructures were obtained from the same area following three interrupted tempering cycles. The Laue patterns and microstructures corresponding to different stages of the transformation are list e d in the following table. - 44 -Figures Treatment X-Ray Laue Metallography Quenched (--phase (c^) 88 r + 20 sec/240°C 88b 89b ^ 2.0 ? Q + 20 + viec/240 oC 88c 89c £ Q + 20 + 20 + 60 sec/240°C 88d The area irradiated by the X-ray beam is outlined by the large circ l e shown in Figures 89b and 89c. The diameter of the beam was approximately 0.8 mm. The central c i r c l e in Fig.s 89b, c positions the fiducial mark introduced by using a pointed needle. This mark served as a nucleation site for the transformation. The deformation in the v i c i n i t y of the mark did not affect the s-phase pattern (Fig. 88a). Light electropolishing was necessary following each treatment. The block of y-phase in the center ofthe irradiated area in Fig. 89b i s rot a single crystal. If i t was a single crystal grown indie (--matrix randomly, then the corresponding Laue pattern would show the essentially unchanged ^-pattern and the weaker but distinct y-phase single crystal pattern having no relationship to £-Laue spots. The Laue pattern of this area (Fig. 88b) shows a similar effect as seen in Fig. 87b - i.e. , diffraction spots of the second phase appear to be produced by splitting of the original <;-phase sppfots (three pairs are shown in the area marked by the c i r c l e ) . As - 45 -the transformation proceeded new nuclei appeared (Fig. 89c) and the number of new Laue spots increased (Fig. 88c). Fig. 88d represents pattern of the microstructure consisting of the ^ pif-phase only (after 20 + 20 + 60 seconds annealing). The arrows in Figures 88b and 88c show several strong £-phase spots which clearly decomposed into groups of faint spots located on the potential Debye rings of the y-phase drawn in.the pattern. This implies, that some reflecting planes are identical for the parent and the product providing the £ spots are produc ed by Ka radiation. It i s obvious from Fig. 88d that the distribution of reflections on Debye rings i s not as uniform as appeared to be In the D.S. patterns obtained from fine rods. More general implications of the observation made on single crystal — y isothermal transformation means that the crystallo-graphic relationship between these two phases might exist. 8.1 Surface Distortion During — y Transformation The quenched £-phase in the 24 at % Al alloy transformed to y-product either by isothermal annealing at a constant temperature or during continuous heating from room temperature. The mechanism of the 2;^  y transformation should be the same for both isothermal and anisothermal conditions, providing the chosen reaction temperatures are within the y-phase f i e l d . Upon slow heating of the t^-phase from room temperature to approximately 165°C distinct areas transformed to the y-phase. Surface rumplings associated with the transformation could be observed on the - 46 -prepolished surface. Geometrically these areas resembled highly coordinated atomic displacements similar to those observed on martensitic transformations. Fig. 90-A - C"Q heated at 1 ° C per minute from room temperature to 1 6 5 ° C , held for 20 minutes at this temperature and quenched. Fig. 90-B - Same treatment as A plus 15 minutes at 1 6 5 ° C . Fig. 90-C - Same treatment as A+B, plus 15 minutes at 1 6 5 ° C . : The block of the product phase as revealed by ligh t l y electro-etching the surface has approximately the same external shape as was outlined by the surface distortions. Fig. 90-D i s the same area as shown in Fig. 90-C following electroetching. The relationship between the internal details and the crystallographic features of the distortions are apparent. Etching revealed fine cracks at the intersection of two distinct directions. The twin-like distortions appeared to grow from one center giving rise to the formation of star-like irregular blocks of the y-phase. Fig. 91 shows several different configurations of twin-like r e l i e f s on the prepolished surface of a partially transformed sample. The distinct geometry of the surface distortions could only be detected by using oblique illumination and a reaction temperatures no higher than 1 7 0 ° C . (A more rapid growth of the y-nuclei at higher temperatures restricted their maximum size prior to impingement. - 47 -8.2 Dilatometry Under Isothermal Conditions The measurable change of the sample l e n g t h during the c,^ ^ u transformation i n the 24 at % a l l o y offered*convenient means f o r determining the k i n e t i c s at a constant temperature. However, the d i l a t a t i o n r e l a t e d to the g r i p p i n g arrangement d i d not a l l o w an exact continuous reading of the sample extension as the r e a c t i o n proceeded. This d i f f i c u l t y was p a r t i a l l y overcome by i n t e r r u p t e d h e a t i n g and quenching of the sample. The AL corresponding to the sample transformation was recorded a f t e r short annealing times at 200°C. A p l o t of AL versus time f o r 4 heating and quenching c y c l e s i s shown In F i g . 92. The f o l l o w i n g c o n d i t i o n s were employed: Sample diameter 0.082", length 2", s i l i c o n - o i l bath at 200°C. The AL was obtained using the v e r t i c a l c r e e p - t e s t i n g machine and a constant load of 150 grams to ensure the s t a b i l i t y of the setup. (The load d i d not cause any measurable extension due to creep.) The graph i n F i g . 92 shows th a t s u b s t a n t i a l p r o p o r t i o n s of the m a t e r i a l have transformed w i t h i n the f i r s t two c y c l e s ( t o t a l time at 200°C being approximately 8 minutes). The t o t a l d i f f e r e n c e i n the length was: AL^ + AL 2 + A L 3 = 0.0026" The average volume change accompanying the £ i s : / u t r a n s f o r m a t i o n - 48 -2 2 TTD TTD V - V L - L ? x 100 - ^—2 5 x 100 4 C As the value of D and D (the diameters of the wire prior to and x, y after the transformation) remained approximately the same within experimental error, the volume change was: \ " ^ x 100 = ^0026 x 1 ( ) 0 L 0 > 1 3 % L c 2 where: L is the length after transformation to the JLL-phos e and is the length prior to the transformation. The sharp steps on the AL curve (Fig. 92) recorded during the reaction at 200°C (upper part of the curve) are indicative of a discontinuous process of consecutive shear displacements occurring within the material during the course of the transformation. This would be consistent with the twin-like surface markings described in Section 8.1. 8.3 Deviations from Isothermal Conditions From the dilatometric measurements i t was concluded that under certain conditions the t, — ^ y reaction can occur with sufficient heat evolution to overheat.the sample to a temperature higher than the isothermal temperature of the surroundings. The magnitude of this temperature difference was examined. - 49 -i Thermal Analysis Experiments: Chromel-Alumel wires were welded to the ends of a sample approxi-mately 1 inch long. The average temperature change at two junctions could be thus recorded during the r, u transformation. A s i l i c o n o i l bath maintained at temperatures 201, 220 and 304°C was used to obtain the effect of the reaction temperature on the true temperature change of the sample. Experimental Procedure: The specimens were quenched from 500°C into water to reain the S-phase and were then immersed into o i l . The transformation was considered to be completed when ithe temperature of the specimen became constant and coincided with the isothermal temperature of the bath. The mass of the sample was small and thus the released heat; did not markedly change the temperature of the bath. After the transformation cycle, the sample was quenched to R.T. and the same cycle was repeated on the f u l l y transformed material. Thus each run yielded two curves, one i l l u s t r a t i n g the internal thermal changes due to the SQ . u transformation and the other the heating of the same sample without a phase transformation. The results are shown in Fig. 93, Fig. 94, and Fig. 95. The magnitude of the temperature increase was: - 50 -Bath temperature Sample overheating Fig. 201°C 5°C 93 221°C 10°C 94 304°C 20°C 95 The reaction at 201°C exhibits an incubation period of approximately 85 sec. This implies that the high heating rate was sufficient to permit retention of the untransformed £-phase during the heating cycle. The incubation period vanished at the higher temperatures. At 304°C, the transformation had already started during the heating cycle. The results showed that at 200°C the transformation is controllable, i.e., i t starts with 100% £Q and ends with 100% y within a convenient period of time. However, the most important observation is that the reactions £^ . ^ y does not occur in a truly isothermal., manner. When the bath at 200°C was vigorously stirred, the heating curve including the reaction z,q — j » - y was identical to the heating curve from the already transformed sample. Under these circumstances the transformation taking place at 200°C could be considered as an iso-thermal event. A l l isothermal experiments were carried out in a stirred bath at temperatures not exceeding 200°C. r 8.4 Isothermal Transformation Za~~*" u. in Alloy 24 at % Al Both hardness and volume change can be considered as direct, though qualitative evidence that the transformation z, — i — V has - 51 -occurred. Quantitative data concerning the structural changes and the kinetics of the £ —»- y transition were obtained using: (1) D.S. X-ray method (polycrystals) (2) X-ray Laue method (single crystals) (3) Quantitative metallography. 8.4.1 X-Ray D.S. Sequence D.S. Patterns from 0.020" diameter sample were obtained using the CuKa radiation after various annealing times at 200°C. The resulting photographs (Figs. 96 to 99) represent a complete transformation of a single £^ grain into the polycrystalline y-phase structure (y^). A l l exposures were obtained with the sample in the same position in the camera but- rotated around an axis perpendicular to the beam. Fig. 96 - £p structure, one grain. Fig. 97a - £Q (as Fig. 96) annealed 2 minutes at 200°C. Fig. 97b - The same as Fig. 97a, with scale. Fig. 98 - The same as Fig. 97 and 98 plus 3 minutes at 200°C. Fig. 99 - The same as Fig. 96, 97 and 98 plus 4 minutes at 200°C. The evaluation of the D.S. patterns: Fig. 97a and 97b show fine, but sharp and continuous y^ -lines - 52 -at the 29 angles 38°, 68°, 72.5° etc. Some £-spots s p l i t into two diffraction spots as can be seen at 29 = 114°, 110°, and 89° (pattern in Fig. 97a). Several lines of the y-phase developed at positions common to £-reflections (lines at 147°, 137°, 130°, 114°, 41° and 39° in Fig. 98). This suggests that the product and parent phase crystal structures may be closely related, providing the identity of several planar d-spacings i s not accidental.^ £-Spots at 113° and 109° which sp l i t during the transformation^are in the positions corresponding to the Debye rings of the developing u-pattern. The X-ray sequence in Fig. 96 to 98 confirmed the sp l i t t i n g of some c"-phase spots as was observed using the Laue technique on bulk samples. The high degree of randomness of the u-product as demonstrated in the D.S. patterns should be related to the experimental conditions (small diameter, sample rotated during exposure). The surface transformation product was not removed after annealing as i t would change the sample diameter. The diffraction lines of the y-phase at large diffraction angles are broadened and the Kct^, Ko^ doublets cannot be resolved. Two possible reasons could explain these effects: (1) Line broadening due to the small size of the y-crystallites. (2) Line broadening due to the lattice distortions. The kinetics of the £Q—»- y^ transformation was substantially faster when the £^-phase underwent some cold deformation. This i s shown in Figs. 100 and 101. Fig. 100 - D.S. pattern of the £^ deformed 15% at room temperature. The undeformed pattern was similar to that in Fig. 96. - 53 -Fig. 101 - The same sample as i n Fig. 100 after tempering 2 minutes at 200°C. Fig. 100 shows that cold deformation caused not only la t t i c e distortions (which would give rise to streaking of the original spots) but i t broke single crystals into fine f a i r l y misoriented regions (almost continuous Debye rings after deformation). The sharp and almost complete pattern of the y-phase i s apparent after only 2 minutes annealing following 15% cold deformation. Whereas in the undeformed sample the reaction had just initiated after the same period of time (Fig. 97a). The drastic change in the kinetics of the £—=»• y reaction i n the deformed sample suggests that both the deformation of the matrix (which produces a . structure of misoriented .grains) and the change of the h.c.p. £-phase to the cubic y-phase are directly related as two steps necessary to accomplish the £Q—s*. y transformation. 8.4.2. X-Ray Transmission Laue The Laue transmission photographs obtained from pa r t i a l l y transformed Cp-phase thin crystals indicated similar s p l i t t i n g of the original c^-phase spots as observed in back reflection Laue patterns, providing conditions for an exposure were satisfactory. Blurred and weak diffraction spots associated with the transformation were not recorded i f either the sample was thick or the intensity of the beam was low (for example when a non-focused beam of 0.2 mm diameter was used). - 54 -A striking characteristic of a l l transmission photographs obtained from material containing the y-phase were radial streaks concentrated in the central region. They represent Laue reflections from numerous y-phase crystallites affected by a la t t i c e strain. The y-phase Debye rings appeared in the transmission patterns after longer isothermal annealing. However, they were not continuous as observed in D.S. patterns of the fine rods (Figs. 96 to 99). Fig. 105 shows a Laue photograph of s^-phase f o i l transformed after 30 minutes at 200°C. (The fcil thickness was 0.005"). The same treatment was performed on a bulk sample 1/8" x 1/8" x 0.200". The f o i l 0.005" cut from this sample yielded the Laue pattern shown in Fig. 106. 8.4.3 Optical Metallography of S^—» y^ Transformation The size of samples used for optical microscopy was 0.500^ x 0.250" x 0.040". Prior to the metallographic examination a layer approximately 0.015" thick was removed from the surface. The progress of the transformation as observed in the electropolished metallographical samples is shown in the following figures: Fig.. 107 - 2 minutes at 200°C. Fig. 108 - 6 minutes at 200°C. Fig. 109 - 9 minutes at 200°C. The etching accompanying the electropolishing revealed the complex internal structure of the y-phase. Blocks of the y-phase in partially transformed samples resembled internal geometrical characteristics comparable to the surface relief s observed on a prepolished surface. - 55 -The internal structure of the y-phase blocks is shown in Figures 110 and 111. Fig. 110 documents distinct ridges and pointed outlines of the y-blocks similar to that seen in Fig. 90. Fig. I l l shows twinning within an area partially transformed to the y-phase. It was established that the irregular separate blocks of the y-phase growing injthe t^-matrix during isothermal transformation cannot be single Jli-grains. X-ray methods showed that the number of y-crystallites must be much larger than the number of optically visible y-blocks i n partially transformed samples (Figures 108, 109). However, the real size and shape of the y-crystallites could not be determined using an optical microscope. The three-dimensional morphology of the y-phase in the fu l l y transformed samples was examined by a consecutive removal of the surface layers. Irregularities and a marked difference in the y-block size is demonstrated in the following sequence of photographs: r Fig. 112-A - The y^-structure after transformation 30 minutes at 200°C. Fig. 112-A-l - As A, 25 microns removed Fig. 112-A-2 - As A, 50 microns removed. A l l micrographs were obtained from the same area on the sample. 8.4.4. Quantitative Metallography The QUANTIMET 720 was used to measure the fraction transformed during the isothermal t,^—»• y^ reaction. The 20 to 36 readings at 290 times magnification gave an average value of the transformed volume fraction. The viewing area of each reading was 0.573 mm . If A = area observed on the microscope screen Au = area covered by the product V = volume of the sample Vy = volume of the product transformed, then the fraction transformed i s : Vy Ay X = V- = A " ; The QUANTIMET gave values of Ay and A was constant. The accuracy of the A readings was less in samples transformed for short times. The small patches of the product formed in the i n i t i a l stage were not "dense" and had a fuzzy outline. The appearance of the analyzed . structures present after tempering at 200°C is demonstrated in Fig. 113 (4 min.), Fig. 114 (6 min.) and Fig. 115 (7 min.). A comparison of the photographs shown in Fig. 114 and 108 (both representative of 6 min. at 200°C) shows a difference in the amount of the product. This i s due to the removal of the smaller y particles during the electropolishing procedure (Fig. 108). To retain a l l the y-phase present and to produce sharp contrast between the phases, the procedure described in Section 1.8 was employed. Fig. 113 documents one of the general features of the £ — y transformation-preferred surface nucleation. A continuous layer of the y-phase is shown on the two surfaces. Quantitative data was obtained only from the sample interior. 9. KINETICS OF PHASE TRANSFORMATION^ 17 18 9.1 General Theory of Phase Transformations ' A phase transformation i s the rearrangement of atoms from one metastable configuration to another of lower free energy. In the rearrangement two differeng basic atomicprocesses may be involved: (1) Independent movement of single atoms. (2) Cooperative translation of many atoms. The f i r s t process includes both long-range diffusion and short-range diffusion. The second process involving a cooperative movement of the atoms (shear and twinning) is a martensitic type of reaction. The direct and necessary experimental evidence that a structure transforms martensitically i s the appearance of distortion on the surface. This reaction most frequently requires rapid cooling but can take place isothermally as well. The martensitic transformation i s necessarily a composition-invariant reaction. In the metallic systems phase-transformations occur almost exclusively as heterogenous reactions, i.e., they start at identifiable centres in the original phase. If long-range diffusion i s the rate controlling mechanism during phase transformation then the linear growth of the product phase follows the relationship: x « /5T where: x = linear dimension; D = diffusion coefficient; t.= time. - 58 -The experimental kinetic data on interface controlled, non-martensitic transformations can be evaluated in terms of the empirical Johnson-Mehl-Avrami equation: X = 1 - exp(-kt) n (1) where x = fraction transformed k = reaction constant t = time n = constant dependent on the nucleation and growth processes. The equation (1) can be used only when k and n are constants independent of x (and t) at a constant temperature. The time exponent n and the rate constant k are useful empirical parameters providing a concise description of the isothermal reaction kinetics, k has a dimension of time and may take any positive value. Rearranging equation (1) and taking logarithms of both sides gives: In ^ = ( k t ) n (2) log In -r^— = n log t + n log k (3) When values of log In - r ^ — are plotted against log t the plot should give a straight line, providing that equation (1) is valid for the process. The slope is equal to the parameter n. For growth processes occurring without a composition variation between ftie parent and the product, i.e, processes not controlled by - 59 -long range diffusion, n w i l l vary according to the following nuclea-tion conditions: n Conditions 4 Increasing nucleation rate 4 Constant nucleation rate 3-4 Decreasing nucleation rate 3 Zero nucleation rate (sites saturated) 2 Grain edge nucleation after saturation 1 Grain boundary nucleation after saturation The heterogenous reactions in metallic systems frequently occur in two or more consecutive steps when different atomic processes are involved. The activation energy for transformation calculated from experimental kinetic data by no means specifies unambiguously the atomic process of the reaction. More precisely i t is an average value for a l l processes taking place during the transformation event. For calculations of the activation energy, Q, several modifications of the Arrhenius rate equation are used. —0 /RT Arrhenius equation: k = Be (5) k = rate constant B = frequency factor * Thisks true for prosesses taking place consecutively.When different processes occur .simultaneously,then the value of the activation energy corresponds to the rate controlling process. - 60 -1. The Time f o r S p e c i f i c Fraction Transformed Method The rate equation f o r an isothermal process i s : !f - kf(x) C6) where k = rate constant f(x) = empirical function describing reaction with respect to time. Rearrangement of (6): x=X t x = k" 1 / f _ 1 ( x ) d x (6a) x=0 where t = time f or X f r a c t i o n transformed at a s p e c i f i c temperature, x The value of the i n t e g r a l i s a constant, providing that the function f "'"(x) does^not vary with temperature over the chosen temperature range. Hence: t x - k 1 (7) (7) . _ - l Q/RT and X a or t v = A e Q / R T (7a) A. where A incorporates B ^ and the p r o p o r t i o n a l i t y constant. - 61 -Taking logarithms: r n t x . + ! . • 1 (8) Thus a plot of versus — yields a straight line for a series of temperatures when A and Q are independent of temperature. Q can be calculated from the slope of the plot (Slope = % ) . When this procedure i s repeated for different values of X and a set of parallel lines are obtained, i t can be concluded that Q i s also independent of the transformed fraction. 2. The Change of Rate Method The specimen is partly transformed at the temperature T^, then transferred into a bath at temperature T^ (the difference T^ - T^ being no more than 10 degrees). The reaction goes on at different temperatures with different rates measured from gradients on the X-time plot just before and immediately after the change. Since the fraction transformed, X, i s the same at the instant of change the rates are: Hence - 62 -6 A G * A G k h 9.2 method offers several advantages: (1) Only one specimen i s necessary. (2) No completion of the reaction i s required. (3) No differences in structural changes may be expected within a narrow temperature range. Growth Rate for Interface-Controlled Transformations. 18 A general equation for the growth rate is : = interface thickness = free-energy change per g-atom for the g — a reaction = activation energy per mole = Boltzmann's constant = Planck's constant. S Comments on g^—*~ y_^  Kinetics in 24 at % Al Alloy 8 Spherical precipitates of u£ (Fig. 67, 68) grew at 200°C with a tant rate in a l l directions. Isotropic growth rate is to be cted for an interface controlled, composition invariant reaction, growth rate obtained from the slope of the line in Fig. 69 is Y exp (10) Y = 2.5 x 10 -5 -1 cm sec - 63 -However, i t was established that the g^-phase exhibits a preprecipita-tion reaction and probably forms a transition phase g at lower tempera-tures. It is thus more r e a l i s t i c to describe the process of the g -decomposition as: The kinetic data obtained at 200°C is then related to the jjj transition. 9.3 Kinetics of c^  —*** Transformation in 24 at % Al Alloy Quantitative metallography yielded the fraction transformed (x) in a specific time (t). Values for^transformation at 200°C and 160°C are shown in Figures 116 and 117, respectively. A plot of log In versus log t (equation (4)) is shown in Fig. 118. Straight lines imply that equation (1) is applicable for this transformation. The slope n = 3 means that the nucleation sites are rapidly exhausted in the early stagejpf the transformation. The activation energy was calculated using the time to a fixed transformed fraction (equation (8)). From the slope of the lines corresponding to 0.3, 0.5, and 0.7 transformed (Fig. 119) the activation energy was found to be: Q = 25.6 kcal/mole £ ^ _ During heating to 200°CvTnay occur~ja reaction): g — s * . g - 64 -The use of only two temperatures (Fig. 119) is unsufficient to prove that the frequency•factor A and the activation energy Q are | independent of temperature. However, r e s i s t i v i t y measurements of the 19 transformation for a series of temperatures showed that log t vs. Y relationship i s linear and a comparable value of Q was obtained. Neither activation energy for volume diffusion nor grain boundary ' diffusion in the alloy 24 at % Al was available for direct comparison with experimentally found value corresponding to v £ — y transforma-tion process. An approximate comparison can be made using data related to the £-phase of the Ag-35 at % Al alloy. The activation energy of volume diffusion in this alloy as obtained from the diffusion 90 coefficients at 497°C and 535°C was: Q = 42.5 kcal/mole 2 1 A ratio 2 5 • 6 - n fi 423 " ° -6 suggests that the transformation process £ y requires an fie activation energy comparable to the value of^activation energy for grain boundary diffusion (or interface controlled process). £ 9.4 Growth Characteristics of y . Phase : *x  The activation energy 25.6 kcal/mole would imply an interface controlled mechanism for the £ — y transformation. This is in - 65 -agreement with the fact that the Johnson-Mehl-Avrami equation is obeyed. Twin-like reliefs observed on a prepolished surface during the r, —»• y transforma-tion indicate that a coordinated atom movement may also be involved in the r, — ^ y reaction. This means that the activation energy cannot specify the transformation mechanism unambiguously. Furthermore, the classification of the reaction as an interface controlled process whenever the JMA equation is obeyed may not be j u s t i f i e d in some cases. The JMA equation was derived for the following conditions: (1) The product grows by an advancement of the product interfaces uniformly in a l l directions. The growing product is thus approximately spherical and does not alter i t s shape during growing. (2) Growth rate of a l l (planar or curved) interfaces should be constant. In order to classify the transformation reaction as an interface controlled process, the observed growth geometry must exhibit character-i s t i c s given by conditions 1 and 2. Experiments were conducted to establish the external shape of the growing y patches and the differences in mobility of individual interfaces (qualitatively). Samples of thickness 0.050" were slowly cooled from. 700°C to 500°C and quenched. The quenched £ was transformed for certain periods of time at 160° and 200°C and the external shape of the y-blocks was examined metallographically. The results are shown - 66 -in the following figures: Fig. 119-1 - £Q immersed in the bath at 140°C, heated at 1 degree per minute to 160°C, held 30 minutes at temperature and quenched. Fig. 119-2 - The same sample as in Fig. 119-1 tempered an additional 10 min. at 160°C. The surface was lightly electro.etched to reveal the displacement of the original interfaces. This figure^ shows the composite photograph. Fig. 119-3 - CQ> tempered at 200°C for two consecutive 1-minute treatments. The shape of the original nuclei developed during the f i r s t minute is outlined i n the centre. The growth in the second minute^ corresponds to the external boundary revealed by electroetching. Fig. 119-4 - The same sample as in Fig. 119-3, tempered 2 minutes at 200°C (electroetching was employed to reveal the i n i t i a l growth of the y-phase nuclei and two consecutive displacements of the y-phase boundary). The change of t h e - b l o c k shapes as observed on (or near) the surface cannot be attributed to an accidental protrusion of they -crystallites from the interior of the sample. It was already shown that the transformation started on the surface and a continuous layer of the product was formed before any larger y-blocks appeared in the sample interior (Fig. 113). Conclusions drawn from the inspection of Figs. 119-1 to 119-4 are: - 67 -(1) The y-phase patches have external shapes which are quite different from that of a sphere. (Lens-shaped grains were more common in structures obtained at 160°C than in those obtained at 200°C.) (2) The interfaces moved with different velocities (i.e., a non-linear growth rate) . (3) The shape of some y-blocks drastically changed during growth. These observations show that the growth mechanism in the t, —^- y transformation (for alloy 24 at %) does not exhibit the characteristics of the thermally activated growth by random atom jumps across the interface. Nor is constant growth rate in a l l directions observed. Although the kinetic data conformed to the JMA equation, i.e., the average growth of the transformed volume fraction was linear with respect to time (Fig. 118 shows straight lines in the plot log In — — vs. log t ) , the metallography showed discrepancies between the assumed model and the real transformation. The discontinuity in growth could be explained i f cooperative atom movement (by shear or twinning) occurs during the transformation. Other evidence that the y-phase does not grow completely randomly in the ^-matrix was obtained by conducting the following experiment: The c-phase was transformed at 210°C for 3 minutes and the sample was mechanically polished. Mechanical grinding and polishing introduced deformation into the surface layer of the £-phase but did not affect the y-phase. The deformation marks in the t;-matrix appearing as bands of narrow lamellae having perfect crystallographic directionality were revealed after chemical etching. Thus i t was possible to detect a relationship between the y-phase facets and crystallographic directions - 68 -in the £-phase matrix. Results are shown in thefollowing sequence of figures: Fig. 119-5 - Deformation marks in mechanically polished (--phase (polarized l i g h t ) . Arrows indicate directions of mechanical grinding. Fig. 119-6 Partially transformed r,^ (210°C for 3 min.) and 119-7 mechanically ground and polished, chemically etched. 10. ISOTHERMAL TRANSFORMATIONS IN 25 AT % Al ALLOY The decomposit ion of the c,^ In the alloy 25 at % Al was examined using: (1) optical metallography (2) X-ray D.S. technique (polycrystalline rods) (3) X-ray Laue technique (single crystals). 10.1 Optical Metallography and X-ray D.S. Patterns In order to establish the morphology of the transformation product in the 25 at % Al alloy, only structures i n the bulk of the samples were accepted as representative. It was found that the kinetics of the £-phase decomposition and product structures are completely different from those observed in the 24 at % alloy. The morphology of the product and the progress of the decomposition of the £-phase are shown in the following figures: - 69 -Fig. 120 - SQ. tempered 5 hrs. at 200°C. Fig. 121 - SQ. tempered 50 hrs. at 200°C. Fig. 122 - As Fig. 121, higher magnification. Fig. 123 - SQ, annealed 100 hrs. at 200°C. X-Ray D.S. patterns representative of these structures were obtained using rods of 0.020" diameter produced by electropolishing bulk samples. The patterns are shown in Fig. 124. Pattern a - annealed 7 hrs. at 200°C pattern b - annealed 80 hrs. at 200°C. Both patterns are consistent with the metallography. The intensity maxima on the y-phase Debye rings implies less randomness than was observed in the y-phase of the 24 at % alloy. The product of isothermal transformation in the 25 at % alloy is comparable to that obtained by continuous slow cooling from the s~field to R.T. (Fig. 39). Both can be classified as Widmannstatten structures. 10.2 Crystallographic Relationship s//y (25 at % Al Alloy) The morphology of the y-phase suggests that a specific crystallo-graphic relationship exists between the parent and the product-phase. Experiments were conducted to determine: (a) The habit plane in the matrix; (b) the orientation relationship s//y» The following two procedures were employed: (1) A traces analysis from two perpendicular surfaces. (2) A fixed X-ray beam method. - 70 -Surface Traces Analysis Broad plates of the y-phase produced by annealing the £Q at 2 0 0 ° C for 30 hrs. intersected two perpendicular polished surfaces. From the angle of inclination on the two surfaces and the orientation of the £q crystal i t was found that the following habit planes existed in the Cq-matrix: ( 2 0 2 5 ) , ( 1 1 2 4 ) , and ( 1 2 3 4 ) . A l l were planes having a small h, k and a larger 1 indices. Thus the Wldmannstatten y-plates were parallel to one of the determined habit planes in the matrix. Fixed X-Ray Beam Method A sample partially transformed at 2 0 0 ° C for 30 hrs. was cleaved through a band of the y-plates. The brittleness of the y-phase made i t possible to cleave the crystal without significant distortion of the adjacent c-matrix. The cleavage path was essentially parallel to the surfaces of the plates within the band. The cleavage surface at 2 6 0 times magnification i s shown in the micrograph in Fig. 1 2 5 . The cleaved surface of the y-plates was aligned perpendicularly to the X-ray beam and a Laue pattern was obtained. A layer of the y was then removed by successive electropolishing. Polishing was carried out u n t i l a perfect pattern of the c-phase appeared on the Laue photograph. (In the scheme in Fig. 1 2 5 the distance A represents the amount of material removed). Following each polishing procedure the sample was examined on the X-ray machine. The orientation of the sample was identical for each test. Fig. 1 2 5 shows y and t, Laue patterns obtained for the fixed beam-condition. Both patterns can be superimposed on one stereographic - 71 -projection in order to obtain the required crystallographic relationship between the two planes. Results of Analysis: (1) Plates of the y-phase were parallel to the (2025) plane in the matrix. This i s one of the habit planes found by the trace analysis. (2) (111) plane of the y was parallel to the (0001) plane of the S-phase. (3) Directions of the type [Oil] in the y-phase were parallel to the directions [1120] in the s-phase. Thus the crystallographic relationship between the s a n (* the y-phase in the 25 at % alloy can be written: (0001) ? / / ( l l l ) y (1120) s//L011]y This i s a common crystallographic relationship between h.c.p. and f.c.c. structures when one phase grows as a Widmannstatten plate. Generally, the Widmannstatten plates are also parallel to the close packed planes. This was not found in this investigation. 10.3 Transformation Sequence for Single Crystals When rods cf alloy 25 at % Al (0.020" diameter) were transformed into the y-phase at 200°C and examined in the X-ray cylindrical camera without removing the surface layers, the Laue patterns showed more y-phase than was detected in the bulk of the samples after the same annealing periods. The CQ—«» u transformation of a Cg-single crystal (not electro- . polished after annealing), is shown in the following sequence: Fig. 126 - CQ CU(W) radiation, sample fixed., Fig. 127 - CQ tempered 2 hrs. at 200°C, Cu(w) rad., sample fixed. Fig. 128 - CQ tempered 4 hrs. at 200°C, Cu(w) rad., sample fixed. Fig. 129 - CQ tempered 32 hrs. at 200°C, CuKa, sample rotated. A comparison between Fig. 127 and Fig. 124a indicates that a significant amount of the y-phase has probably nucleated on the rod surface (Fig. 127). No y-lines are visible in Fig. 12^abecause the surface was removed by electropolishing. Fig. 129 demonstrates the coincidence of many z,-l±nes with y-lines. After 32 hours the c-phase has decomposed only partially, giving rise to the intense spotted arcs. The weaker continuous lines belong to the y-phase. 10.4 c~Phase Nucleation on Free Surfaces In a l l alloys within the composition range Ag-20,1 at 27 at % A l , the y-phase showed a strong tendency to nucleate on free surfaces. Generally a continuous layer of the y-phase developed on well-polished, non^deformed surfaces before any significant internal transformation product appeared. Samples of 25, 26, and 27 at % alloys were electropolished prior - 73 -to the S —»- y transformation and the surfaces were microscopically examined (without repolishing). The following figures show results: Fig. 130 - SQ after electropolishing. The structure is represent-ative of 25, 26, 27 at % alloys. Fig. 131 - SQ alloy 25 at %, tempered for 30 hrs. at 200°C, surface not polished. Fig. 132 - SQ alloy 26 at %, tempered for 50 hrs. at 200°C, surface not polished. Fig. 133 - Alloy 27 at %,-SQ tempered 100 hrs. at 200°C, surface not polished. Fig. 134 - Alloy 26 at %, structure of sample shown in Fig. 132 after removing approximately 0.010" from the surface. The morphology of the u-surface product in 25, 26, and 27 at % alloys i s different from the y-product found in the bulk of the samples (Fig. 133, Fig. 134). The y-phase lines seen in the D.S. patterns in Fig. 127 and Fig. 128 are probably reflections of the surface y-phase product having the morphology shown in Fig. 131. 10.5 Effect ofi Deformation on SQ~Phase Room temperature deformation of the Sq-phase has a pronounced effect on the kinetics of the transformation and the morphology of the product. The interdependence of the two processes, deformation and the S •—s>- y transformation was examined qualitatively. - 74 -Surface Deformation It was mentioned before that mechanical polishing introduced deformation into the £-phase. j n surface layers of the £-phase which were mechanically ground, a, large amount of the y-phase developed during tempering at 200°C. In the 25 at % alloy this product has a morphology entirely different from the product observed in the undeformed sample interior. The undeformed £-phase of the 25 at % alloy transformed into the Widmannstatten y-phase after isothermal annealing 80-100 hours at 200°C. Whereas the y-phase i n deformed surfaces transformed into irregular y-patches after 15 minutes at 200°C (Fig. 135). Bulk Deformation Bending and introducing a hardness indentor into the £ -phase using a load 5 kg was always accompanied by distinct audible clicks similar to "crying" known to be associated with twinning in t i n . Qualitative evidence suggesting that the £-phase deforms readily by a twinning mechanism was obtained by examining deformed structures metallographically and using X-ray Laue methods. No attempt was made to establish the crystallography of the deformation processes in the £-phase. A specimen of the £-phase [25 at % Al] consisting of three grains 5 mm in diameter was deformed by cold r o l l i n g , electropolished and electroetched prior to X-ray exposure and metallographic examination. - 75 -X-ray Laue patterns and microstructures corresponding to different degrees of deformation are listed in the following table: Structure Laue Metallography Undefbrmed-phase Pig. 136a S-Phase deformed 3% Fig. 136b Fig. 136d S-Phase deformed 17% Fig. 136c Fig. 136e S-Phase deformed 10% (different grain) Fig. 136f Fig. 136g The Laue photographs showed several effects characteristic of the S-phase room temperature deformation. Most typical effects were: (1) Decomposition of the original Laue spots into groups of faint and streaked spots resembling a polygonized but strained single crystal structure (Fig. 136b). (2) Formation of multiple Laue patterns by apparent coordinated displacement of the original pattern (Fig. 136f). (3) Vanishing of the original single crystal pattern and development of Debye arcs (Fig. 136c). It appears that with increasing deformation the single crystal of the S-phase breaks into a large number of very small blocks which are significantly misoriented as evidenced by Debye rings produced by Ka radiation. It was already established that these rings belong to the h.c.p. phase (Fig. 100) and no crystal structure change occurs during deformation up to approximately 20% at room temperature. The work hardening was very pronounced and deformation higher than 20% resulted - 76 -i in crackingfcf the specimen. The microstructure of the deformed £-phase consisted of bands of narrow lamellae arranged along distinct crystallographic directions. Generally only three sets of directions were observed i n the planar section through the deformed crystal. Figures 136d, 136e and 136g confirmed that similar straight lamellae present in the deformed surface layer (Figs. 119-5 to 119-7) have the same origin, i.e., mechanical deformation. Deformation in metals and alloys occurs either by s l i p or mechanical twinning. Unrestricted s l i p preserves i n i t i a l atomic arrangement within the deformed structure. Slip lines observed on the surface cannot be revealed again on the repolished surface by etching. Since the twins are regions of different orientation with respect to the matrix, they respond to etching and become visible in the structure* It was concluded that the narrow lamellae developing during deformation of the £-phase i n the 24 and 25 at % Al alloys are twin lamellae similar to Neumann bands observed i n f e r r i t e , some h.c.p. 22 metals and several other alloys. Mathewson described structures reproduced in Fig. 137-1 as twin lamellae (photograph on the l e f t side shows deformed bronze and on the right side deformed brass respectively). A similarity between the deformed structures observed in the present 22-24 work and those reported in the literature to be twins is obvious. - 77 -Comments on Laue Patterns of Deformed Structures 22 It was suggested that there is no plane that would permit twinning in h.c.p. metals without distortion even when the axial ratio is ideal. Localized twin lamellae thus give rise to a strain in the 24 adjacent parent l a t t i c e . In Cahn's review this region was called the accommodation region. X-ray diffraction of the part of the crystal containing one twin lamellae produced separate reflections from the twin and showed a s p l i t of the parent lattice reflection due to distortion in the accommodation region. It is believed that the Laue pattern i n Fig. 136b (3% deformed material) demonstrates i n part a similar effect. Though twins were visible in the micrograph C (Fig. 136d) the Laue pattern failed to show .reflections suggesting the presence of a twin. Fig. 136f of the sample deformed 10% indicates however, the presence of two patterns as expected from a twinned part of the crystal. A meaningful crystallographic analysis of the observed twinning would require a more careful choice of crystal orientation, deformation conditions and refined X-ray technique (small beam collimator). A breakdown of the s-phase single crystal into an apparent poly-crystal as deformation increases is also in agreement with reported 25 facts. Mathewson analyzed twinning processes in zinc and concluded that in h.c.p. structures various twinned parts may be considered as independent different crystals, i.e., misoriented grains. Twinning in zinc was described essentially as a process of "crystallograhic ; fragmentation". An X-ray diffraction pattern from the fracture surface - 78 -25 of the zinc single crystal exhibited solid Debye rings. Fig. 137-2 illustrates drastic lattice reorientation within twin lamellae i n zinc crystal accomplished by relatively small shuffle of atoms. (The [0001] plane is rotated almost 90 degrees). The explanation of the C-phase fragmentation in a similar manner seems feasible. The frag-mentation is presumably further enhanced by the mutual interaction of ^ ^ the twins as their density increases and they intersect eventually. The absence of twin reflections in Laue patterns of the £-phase deformed 3% can be explained assuming that twinned lamellae are thin and asverlsi sets of la t t i c e orientation within lamellae exist. Thus, the Laue diffraction from particular twins are weak and lost in the background. Local Deformation The appearance of a deformed region of the r;-phase i n the v i c i n i t y of Vickers indentations (load 2 kg) is shown in Fig. 137-3. The surface was not repolished. The similarity between an apparent apex formed by the combination of two sl i p or twinning systems and the surface re l i e f in the samples partially transformed to the y-phase should be noted (Fig. 90). When locally deformed sample was annealed at 200°G, the regions adjacent to the indentation transformed to the y-phase. The transformation product grew s t r i c t l y through areas ofthe high density cf twin bands. An experiment was conducted using a single a. crystal of <;the alloy 25 at % Al deformed by1,Rockwell indentor under a load of 60 kg. After deformation, the sample was annealed 75 minutes - 79 -at 200°C and sectioned through the indentation. The observed micro-structures are shown i n the following figures: Fig. 138a - area adjacent to indentation. Fig. 138b - area 1 mm from indentation. Fig. 138c - area 3 mm from indentation. The process of the transformation within heavily twinned regions of the (--crystal resembles a recrystallization process taking place in other materials exhibiting twin bands after deformation. Bands of narrow twin lamellae coalesce into separate grains upon annealing. An analogy between recrystallization in twinned materials and elimination of the twin bands by the y-phase formation i n the alloy 25 at % Al might exist. Fig. 138c supplements a previously made assumption regarding possible fragmentation occurring at twin intersections. It is assumed that the role of deformation in this transformation is more complex thanta simple enhancement of the nucleation process. Not a l l of the regions at the twin intersections in Fig. 138c transformed to Ihe y-phase. The necessary degree of deformation (fragmentation) was not attained at this distance from the hardness mark. The y-patches visible in Fig. 138a and b grew to a maximum size after 75 minutes of tempering and did not increase in size after an additional annealing-period of 3 hours. 30 Masson has shown that deformation markedly affects phase transformations i n Ag-Cd alloys. - s o -i l . y-PHASE CRYSTAL STRUCTURE ANALYSIS The y-phase in the Ag-Al system is reported as being isomorphic with the B-manganese crystal structure, i.e. , complex cubic and a ° 1 lattice parameters of: a = 6.92 - 6.93 A. The specification of the crystal structure of the alloy composed of atoms which have a great difference in their atomic scattering powers (ratio -yP^- = 3.4j implies A l a random atom distribution. A l l atomic sites in the alloy are qualitatively|equivalent, being occupied by an "average" atom at the same coordinates as the chosen standard. The y-phase was f i r s t analyzed by Westgren and Bradley.^ They observed a close similarity between the X-ray patterns of the y-phase in the A*-A1 system and the B_manganese structure. Because of this similarity they concluded that the y-phase had a g-manganese structure. An analysis of the y-structure in terms of relative intensities of vthe diffraction lines was not done in the present work. The y-structure determination was carried out by comparing X-ray D.S. patterns with the published data. 11.1 Experimental Details of X-Ray Analysis (24 at % Al) The quality of the X-ray patterns obtained from time powder material was excellent. However, powder could not be used for the analysis due to surface composition changes (Section 7). The b r i t t l e y-phase crushed into fine powder yielded blurred diffraction patterns when used for analysis without annealing. These patterns were also - 81 -useless for identification of the presence of the retained £-phase. Solid rods(0.020" diameter) electropolished after heat treatment gave D.S. patterns of good quality. D.S. patterns of the y-phase were obtained after the following heat treatments: (1) Equilibrium cooling from the £-field at 0.4°C per minute , e v to room temperature (y„„). (2) Tempering of the quenched 5-phase at 200°C for 30 minutes (3) Tempering of the quenched £-phase at 200°C for 3 hours. The D.S. patterns corresponding to the different heat treatments and radiation used are list e d in Table 2. Table 2. Fig. Pattern Alloy (at % Al) Structure Time at 200° C Radiation 139 68 24 e yRT - CuKa 72 25 e UET - CuKa 140 62 24 A 30 min. CuKa 63 24 4 3 hours CuKa 141 . 81 24 e PRT - FeKa 82 24 3 hours FeKa 142 91 24 e VRT - CrKa 92 24 4 3 hours CrKa - 82 -The scale superimposed on the patterns shows the 26 angles with an accuracy of + 1 deg. The exact 26 angles were measured on thee original patterns with an accuracy of + 0.05 deg. FeKa, CrKct, and CuKa radiation was used to permit a direct comparison with published results and ensure reproducibility of results under different experimental conditions. The analysis of patterns 68 and 72 are included in Appendix -Table 2 and that of patterns 81, 82, 91 and 92 in Appendix-Table 3. 11.2 Conclusions from X-Ray Analysis of u-Phase e c Crystal Structures y„_ and y.: J KRT p i Though both crystal structures are identical,t"he diffraction patterns differ in the following details: The doublets in the y,,--' Ri r patterns are sharp (Fig. 139), whereas those i n the y^-patterns of samples annealed for time 30 minutes are unresolved (pattern 62 -Fig. 140). The resolution of the doublets improved with extended heating times at 200°C (3 and more hours - pattern 63, Fig. 140). The gradual elimination of the line broadening after longer annealing periods can be attributed to either stress relaxation or an increase of the y-crystallite size. . Both effects may occur s imult aneously. Comparative y-Structure Analysis The structure factor for thejy-phase was not available in the - 83 -literature. However, i t was calculated for the B-Mn structure by "t 26 Preson. If the y and the B-Mn structures are identical and i f the Ag and Al atoms in the y-phase are randomly arranged on the atomic sites, the relative intensities of the specific reflections should be identical. The calculated relative intensities for (h, k, 1) reflections in the g-Mn structure are listed in Appendix - Table 2. No reflections should be present for: S h k 1 1 100 4 200 16 400 On the y-patterns obtained using CuK radiation, the 200 reflection was always present (the line at approximately 26 degrees (Fig. 139)). The 400 reflection was very weak and did not appear in a l l patterns (line at 54.2 deg. (Fig. 139)). Using FeKa and CrKa radiation the 200 and 400 reflections were frequently recorded in the patterns. When a l l of the vis i b l e reflections were taken into account, i n addition to the "forbidden" B-Mn lines, several other non-cubic reflections were found. The "forbidden" and non-cubic reflections are referred to as "non $-Mn reflections" in Table 3. It is not considered lik e l y that the non-6-Mn lines were caused by the experimental setup since they appeared at the same angles when different cameras and collimators were used. There are, however, several possible sources which would explain the presence of the non-B-Mn reflections. They are: (1) (2) (3) Unfiltered Kg radiation. Ordering in the y-phase. Retained C-phase. Description of Table 3. The observed non-g-Mn reflections are listed in columns II and 2 2 2 III respectively. The corresonding values of S = h +k +1 for possible Kg reflections are in column IV. Since both Ka and Kg wavelengths produce diffraction from the same plane, then: where 9 = diffraction angle and X = wavelength of the radiation. Using values of Ka, and Kg- then: 2 For Fe rad. 2 1.2060 For Cu rad. 1.2273 Ratios necessary for Kg identification were obtained from measured reflections and are li s t e d in columns V and VI (Table 3). An intensity consideration ruled out the possibility that the Kg radiation i s the source of the (200) reflection. The intensity of the Ka and Kg reflection from the same plane can be evaluated from the following Debye rings in Fig. 141: h k l 2e (degrees) Ka Kg 300 221 50 45 330 411 73 65.2 431 510 91 80.8 Providing the (200) non-g-Mn lines is due to Kg radiation, then the similar intensity difference should be also seen comparing the following lines: Line at 29 and 26 degrees - Fig. 139 Line at 365 and 32.5 degrees - Fig. 141 Line at 44 and 39 degrees - Fig. 142. An inspection of patterns shows that the difference in intensities is much less than would be anticipated in the case of Kg and Ka reflections. Some D.S. patterns revealed the presence of the S-phase (column VII Table 3). The possible sources of the non-g-Mn lines are summarized in column VIII (Table 3). TABLE 3 I II III IV V VI VII VIII Rad. Observed non-B-Mn Reflections Possible S for KB 2 Sin 0K«c Identical C Ref1ec-Possible Reason for Reflection Si n 2 0 ; Sin* 0 Kp, (observed) y RT yRT t i o n s 2 Sin 0 0.0789 0.0791 5 1 .247.1 1 .251 6 cubic ordering 0.1451 0.1460 9 1 . 2233 1.2151 K e 0.2253 14 1.2162 K 6 0. 2898 0.2903 18 1.2136 1.2167 K B 0.3220 . 0.3239 20 1.2152 1.2075 0.3238 cub. ord. KB C 0.4183 0.4190 26 1.2132 1.2126 K B 0.4343 27 1 .2107 . K 3 Fe 0.4834 30 1.2071 K ft 0,5929 37 1 .2139 ' K B 0.6092 0.6097 38 1 .2440 0.6072 Q 0.77 90 0.7787 0.7792 C 0.8370 0.8364 52 0.8398 C 0.8562 0.8553 0. 914 0 0.9127 57 0.94 53 0.9448 59 0.9923 62 \ ^ \ \ \ \ \ v \ x \ x x ' .. . v • TABLE 3 - continued. 7> "\ k v v '. >. v v \ '.. >. •, y v '.j v. \ \ W . W \ '\ k-. V v \ '. v v I II III - . I V V . . VI VII VIII \ V \ V \ \ \ s \ \ - \ s \ \ N \ \ \ W \ X \ S \ X ^ S \ \ ^ \ . . . . . s . . . . . . . . V S \ ' . \ V - \ N N N W . W \ ' v \ \ N N N N s " > ' • 0.1109 0.1097 5 1 .2498 1 .2543 cub. ord. 0.2050 0.2051 0.2385 9 1 .2054 1 .2048 U Cr 0.4385 cub. ord. 0.4532 20 1.2037 0.4547 KB C 0. 5876 26 1 .2032 KB 0.8350 37 0.8492 C 0. 04 9-6 5 1 .25 cub. ord. 0.0914 9 1 . 22 K& 0.0960 0.0978 0.2069. 20 1.19 0.2080 Cu 0.2885 0.2912 0.3424 34 1 .23 0.3432 KB , C 0.3483 - 0.6415 0.7746 0.7761 / - 88 -The X-ray analysis verified that the y-phase has cubic symmetry. o The la t t i c e parameter was found to be: a = 6.9366 A. ° 1 The reported values are 6.92-6.93 A. The presence of the (200) and (400) cubic reflections in the y-phase patterns implies ordering in this structure. The obtained results suggest that the specifica-tion of the y-phase structure as a disordered g-Mn structure may not be ju s t i f i e d . The original investigations of this structure were carried out on powder material which was shown to be very unsuitable for this analysis. 12. ELECTRON MICROSCOPY OF g Q AND jti-PHASE IN ALLOY 24 AT % Al Foils were examined using a 100 KV-HITACHI electron microscope (model HU 11A)• The conditions for f o i l preparation are li s t e d in Table 4. Heat treatment for the s^-phase (Foils A, B, C): Heating at 700°C for 1 hour, cooled to 500°C at 10 deg. per minute, held at temperature for 30 minutes and quenched into water. 12.1 gq-Phase Foils of the s^-phase showed structure differences dependent on the nature of the starting material and the time maintained at R.T. after quenching. The ^-structures are shown in the following, figures: Fig. 143 -Fig. 144 -Fig. 145 -Fig. 146 -Table 4. Preparation of Foils for E l . Transm. Microscopy. Group Cast Hot Cold E l . Polishing homogenized rolled rolled 10% KCN AC Window DC A + + B + + + C + + + +. Foils A were cut and electropolished after heat treatment of the bulk material. Foils B were prepared from the material heat treated as a sheet of thickness =0.040". Foils C were produced by electropolishing from cold rolled sheets heat treated at a thickness < 0.005". Fo i l A, Cg-phase 5 days after quenching. Diffraction effects resembling small G.P. zones are v i s i b l e . F o i l A, S^-phase less than 4 hours after quenching. F o i l C, Sq-phase two hours after quenching. Fine dark clusters are visible in the sharp contrast within and along the edges of the diffraction contours. Selected area diffraction pattern of Fig. 145 (zone [3301]). - 90 -12.2 Structural Changes of CQ in El-Microscope Stage In many cases f o i l s of the Cg-phase underwent significant structural changes during the examination in the e l . microscope. Changes occurred rapidly when the intensity of the beam was either high or suddenly varied. The same effect took place in f o i l s held in the beam tor a long time ( 5 hours). Dark loops and disc-shaped contrast areas were commonly observed growing homogeneously from the matrix. The Cq-foll after 2 hours in the microscope i s shown in Fig. 147. Diffraction contrast resembling that associated with the presence of coherent precipitates or G.P. zones i n the matrix appeared after 5 hour exposure to the beam (Fig. 148). Dense and highly mobile bands of parallel striations formed i n the f o i l s when the beam penetrated areas near the edge of the f o i l . These striations had the appearance of stacking faults (Fig. 149). Bands of stacking faults were less common in f o i l s prepared from heavily cold rolled sheets (foils C). 12. 3 Transformation c n '—» V^-X-ray techniques detected measurable changes after isothermal e periods of 2-3 minutes at 200°C. The optical microscopy gave meaningful information after a period of approximately 4 minutes at 200°C. The use of the transmission microscopy revealed changes in the CQ~nmtrix believed to be associated with an early stage of the CQ — V reaction. The loops and zones formed i n c f o i l s due to beam heating could be - 91 -G.P. zones or small precipitates of the product phase. However, the temperature at which these changes took place could not be determined. The structuralchanges in the c ^ - f o i l s resulting from heating by the electron beam were not investigated. An attempt to follow the decomposition of the f^-phase in the hot stage under a controlled temperature failed as f o i l s lost transparency due to contamination. It was possible to suppress these effects using a low intensity electron beam andjh^plding f o i l s i n the stage for a short time only. (Generally, the f o i l s from thin cold rolled sheets prepared by procedure C - Table 4 were less susceptible to these structural changes. The product of the £Q—»- y transformation was examined using f o i l s obtained after an isothermal heat treatment at 200°C. The product appearing in the transformed structure after 1-2 minutes at temperature i s shown in Fig. 150 and 151. Fig. 141 is an enlarged and underexposed print of Fig. 150. It shows overlapping of two plates in the centre. The structures of the partially transformed material is shown in Fig. 152. The f o i l was produced by B treatment and the annealing time at 200°C was 6 minutes. The parallel lamallae resemble two sets of microtwins. Foils from cold rolled sheets (procedure C) showed a presence of the y-phase which either extended throughout the f o i l thickness or appeared as layers not extending throughout the f o i l thickness. F o i l C transformed 210 seconds at 200°C i s shown i n Fig. 154. The "grains" at the y^-phase boundary may be thin layers of y^-phase growing over the f o i l surface. The possibility that these plates are - 92 -polygonized subgrains in the matrix which is deformed due to the specific volume difference between the parent and product phase seems improbable. Polygonized subgrains should appear only in structures having a high dislocation density. No distinct dislocation patterns were observed in the matrix adjacent to the y-phase boundary. Furthermore, i f the volume difference between the y and the i^ -phase would effectively distort the matrix so as to i n i t i a t e polygonization during tempering, then each y-phase grain should be surrounded by sub-grains. This was not observed. Transmission microscopy provided evidence that the y^-product i s composed of thin layers. The structures in Fig. 155 and 156 obtained from the y-patch in Fig. 154 are representative of the transformation product found in partially transformed f o i l s prepared from cold rolled sheets (foils C). The structure exhibits parallel Moire fringes. This effect arises generally when layers of two different structures are superimposed or when two layers of the same structure are mis-oriented. This structure was never found in the B f o i l s . Selected, area diffraction pattern of Fig. 156 is shown in Fig. 157. Streaks similar to those seen in the SAD pattern shown in Fig. 157 have been related ±i the literature to the presence of a duplex f.c.c.-h.c.p. 31 structure. Streaks were not perpendicular to the fringes present in Figures 155, 156. Moire fringes and SAD patterns imply that the structure shown in Figures 155 and 156 cannot be a single crystal. The structure probably consists of a few thin y-layers or i t may be a duplex £+y structure. - 93 -Within layers of the u^-product areas with darker contrast were also observed. These areas could be zones of a different composition -the Ag and Al atoms have different scattering factors. To establish that internal compositional inhomogeneities existed in the y-phase i t was necessary to examine the effect of electropolishing on the appearance of the structure. This was possible because the y-phase breaks readily into thin flakes suitable for direct examination in the electron microscope. The microstructure of such a fragment of the y-phase after 30 minutes at 200°C is shown in Fig. 158. Dark zones resembling clusters are clearly evident and appear to be the same as observed on the electropolished f o i l s . The SAD pattern of the structure in Fig. 158 shows two diffracting y-phase zones (Fig. 159). The strong spots belong to the [121] zone and the weak spots near the centre belong to the incomplete [001] zone. 12.4 Double Diffraction Effects Double diffraction effects observed i n partially transformed fo i l s were considered direct evidence that the single c r y s t a l l i t e of the y-phase i s a thin plate. Double diffraction from the overlapping layers of the y and the S-phase is illustrated in the series of photographs in Figures 160 to 163. The f o i l was made by procedure C and transformed 210 sec. at 200°C. A SAD,pattern from the bottom part of Fig. 160 showed that this - 94 -area was the ?-phase (Fig. 161). When two layers of the different phases were examined (upper part of the Fig. 160) the same ^-pattern (A-B-C-D-E-F-G—H.) acted as a source for 8 additional primary beams which are again diffracted by the structure of the second phase. The SAD pattern from the y-phase ([001] orientation) i s shown in Fig. 162 (a-b-c-d- being the closest reflections). This pattern i s repeated around a l l diffracted beams A, B, C, D, E, F, E, H, giving rise to the complex network of spots seen in Fig. 163. The SAD ' pattern shown in Fig. 162 was obtained from the v i c i n i t y of the straight interface (Fig. 160). The SAD pattern i n Fig. 163 shows another significant fact. Spots H and D are common for the c, and the y-phases. This means that they have the same d spacing and that these planes are parallel in the structure, i.e., (1011)?//(310)y. The zone giving rise to the ^ -pattern in Fig. 160 i s : £ll0lj. The zone of the y-phase superimposed on the £-pattern i s : (OOlj. This Fragments of the y-phase < 0.5 microns in diameter were pressed into an acetate f o i l and a thin carbon layer was evaporated on the surface. After dissolution of the acetate, the y-particles on the carbon layer were examined in the electron microscope. Fig. 164 shows the particles and a corresponding SAD pattern. The indices of the y-phase cubic structure for the f i r s t 6 reflections are: also implies that: (1101)?//(001)y. Diffraction from Separate Misoriented y-Crystallites - 95 -Ring h,k,l 1 211 2 220 3 300 4 310 5 320 6 321 The (100), (110), (111), (200), (210) reflections which would occur at lower angles are missing in this pattern. The reason for thistts not known. 12.5 u^-Phase After Long Isothermal Tempering The optical microscopy results suggested that the isothermal reaction 5^ —*~ y^ at 200°C i s completed after approximately 12-15 minutes No visible changes could be detected in the y-phase after this period. However, X-ray D.S. techniques revealed that the y-phase underwent some changes as the tempering time increased from 30 minutes to 3 hours. Transmission microscopy qualitatively confirmed the X-ray results. The electron diffraction patterns of the y-phase present in the partially transformed cold rolled f o i l s (after tempering at 200°C for 3-4 mintes) were streaked and did not show single-zone patterns (Fig. 156). (The diffraction patterns in Figures 162 and 163 are not representative of the y-phase in partially transformed f o i l s since they are obtained from a thin y-layer not extending throughout the - 96 -f o i l ) . As the time was increased to 30 minutes the diffraction patterns showed less streaking and yielded more complete diffracted zones. This is obvious from comparison of Figures 157 and 159. (There i s , however, a difference in the f o i l preparation: Fig. 157 is from an el-polished sample and Fig. 159 is from a fragment). Microstructures of the y-phase produced isothermally at longer tempering periods exhibited fewer Moire fringes and the dark cluster-like zones were more distinct. Fig. 165 shows the y-phase after tempering 2 hours at 200°C. The iicrostructure and the SAD pattern from the y ^ - f o i l after tempering 7 hours at 200°C i s shown in Figure 166. The zone axis of the pattern i s [212]. The observation that line broadening on X-ra^d i f f r a c t i o n patterns decreases and electron diffraction patterns become more perfect as the isothermal tempering time increases implies that the y^-crystallites are very thin i n the i n i t i a l stages of the transformation and grow in thickness with tempering time. However, the possibility that la t t i c e distortion i s also involved injthe line broadening effect i n the y^-structure could not be excluded. 12.6 y^-Phase Structure A transmission micrograph of the y^-phase obtained by cooling at 0.4°C per minute from 650°C to R.T. i s shown in Fig. 167. Apparent "subgrains" seem unaffected by the light and dark contours which extend across the boundaries. - 97 -The structure shown in Fig. 167 is comparable to that shown in Fig. 158 and Fig. 165 (i.e., the y-phase produced isothermally). 13. REPLICAS FROM FRACTURE SURFACES Carbon replicas taken from unetched fracture surfaces of the y-phase confirmed the conclusion made from observations on the f o i l s : The y-phase cannot be considered as a simple homogenous phase. The structures as revealed by replicas showed features comparable to those observed in the f o i l s . Typical structures from replicas are shown in Fig. 168, Fig. 169, Fig, 170 and Fig. 171. The following heat treatments were used: Figure Heat Treatment prior to Fracture 168 Quenched from 500°C, tempered 30 min. at 200°C. 169 Quenched from 500°C, tempered 15 min. at 300°C. 170 Same as 169, different sample. 171 Cooled from 650°C at 0.4°C/min to R.T. • F The agreement of the transmission results with replica observations can be demonstrated by comparing the structures obtained by identical heat treatments. - 98 -Heat Treatment Transmission Replica Quenched from 500°C, tempered 30 minutes at 200°C Cooled from 650°C at 0.4°C per minute to R.T. Fig. 158 Fig. 167 Fig. 168 Fig. 171 The circular features seen in the structures obtained after tempering at 300°C (Fig. 169, Fig. 170) are assumed to be of the same origin as those present in the structure after tempering at 200°C (Fig. 168). Some coarsening may have occurred at the higher temperature. X-ray analysis showed no difference between the crystal structure i n Fig. 169 (300?C) and that shown in Fig. 168 (200°C). Thus the circular particles are not the second phase. 14. DISCUSSION OF RESULTS General Properties of Alloys Ag 20.1-27 at % Al A l l alloys were ductile when quenched from the (C+a) and the C-phase f i e l d . The hardness of the quenched structures increased with increasing Al content and was also dependent on the extent of the transformation to the y-phase. The y-phase as a product of either continuous cooling or isothermal treatment was extremely b r i t t l e and exhibited a hardness of HV = 420-450. A hardness peak in structures transformed to the y-phase was found to occur at the composition 24 at % Al. - 99 -Phase Fields and Experimental Conditions The approximate experimental conditions for which the phase diagram is applicable were established during this investigation. In the studied composition range the most significant discrepancies between non-equilibrium and equilibrium structures were observed when the following transformations occurred at low temperatures: (orK) >• (a y) (?+a) ——> y Z 3* (V+S) — * u t. (u+s) The high temperature transformations occurred without obvious d i f f i c u l t i e s at cooling-heating rates of 7°C per minute. (The cooling rate was chosen as a convenient parameter for the specification of different experimental conditions). In the alloys 24, 25, 26, and 27 at % Al the low temperature transformations required cooling rate-s of 0.3-0.6°C per minute in order to obtaintphases predicted from the phase diagram. The alloys 20.1 and 23 at % Al showed disagreement with the phase diagram when cooled at 0.3 C per minute to room temperature. .The alloy 20.1% had much less y-phase than anticipated at low temperatures. The a-phase decomposition into the y-structure within the (?+y) f i e l d presumably did not take place during continuous cooling at 0.3°C per - 100 -minute. Similar s t a b i l i t y of the a-phase was found i n the alloy 23 at % Al. This implies that the diffusion rates are probably slower in the f.c.c. a-phase than in the h.c.p. £-phase. Structures i n Two-Phase Fields Alloy 24 at % A l : When the alloy was cooled from the £-field at 5-7°C per minute, the transformation t, (c+y) expected within the two-phase f i e l d was completely suppressed. However, on cooling to 300°C, a large proportion of the £-phase transformed rapidly to the y-phase. This implies that the transforma-tion process involving long-range diffusion in the two-phase f i e l d i s sluggish, but is rapid in the single-phase f i e l d where the composition invariant reaction is possible. From the phase diagram i t could be concluded that theyU.-phase grows from the £-matrix by diffusional growth when the alloy i s slowly cooled oiheated within the (c+y) f i e l d . No other possible processes associated with this reaction were proposed. The present investigation showed that long-range diffusion occurs but certainly i s not the only atomic process involved in this transformation. X-ray techniques provided evidence of a homogenous breakdown of the £;-matrix into a polygonized substructure. Transmission microscopy showed faulted regions in the structure giving rise to X-ray patterns indicat-0 ?n ing polygonization. No equiaxed subgrains common deformed and annealed material were observed. The observed substructure cannot be - 101 -explained by matrix distortion due to volume differences between the £ and u-phases. Breakdown of the single crystal into lightly mis-oriented snail blocks occurs during continuous cooling from the £-field to temperatures 400-435°C in the (s+y)-phase f i e l d . Under these conditions, any effects on the parent matrix should be detected in the advanced stage of the transformation only. Small nuclei of the y-phase growing by diffusion cannot disturb the matrix structure homogenously at the beginning of the transformation. During this work a large number of X-ray Laue patterns were obtained from either partially transformed slowly cooled or isothermally treated materials. In general, whenever the s-phase single crystal started to decompose into the y-phase the original Laue spots underwent splitting and smearing. The same change of the Laue spots was also observed in the course of the s-phase deformation.A previously discussed polygonized structure found in the (s+y) mixture at temperatures 400-435°C was presented in Fig. 48. This pattern has a perfect counterpart in Fig. 136b which shows diffraction pattern^of lightly deformed s _phase. Having assumed that the deformation mechanism is twinning then the similarity suggests that the y-phase formation within the two phase f i e l d involves a process identical with twinning in the S-phase. Strain imposed on matrix in order to accommmodate twins may cause t i l t i n g of small parts of the matrix crystal resulting in the s p l i t of Laue spots. Since the temperature is relatively high during this process the polygonized blocks are strain free as indicated-by the sharpness of the small spots at particular S-phase reflections. It must be admitted that the substructure is only indirect - 102 -evidence of the twinning process. Direct evidence, i.e., reflections crystallographically related were not recorded in Laue patterns of the equilibrium (£+y) structures. The twinning process occurs probably on a very fine scale and reflections from twins may be scattered in the background of the pattern as assumed i n the case of the deformed t;-phase. The y-phase formed under the described conditions resembles plates growing parallel to certain crystallographic directions. This ? fact supports the suggestion that a cooerative atomic process i s probably involved in this transformation. The redistribution of elements by long range diffusion was however, also established. Alloy 25 at % Al: At a cooling rate of 7°C per minute the f;-phase was retained to room temperature. The t, — » y transformation similar to that occurring in the alloy 24 at % was never observed. This means that at 25 at % Al the alloy i s in the two-phase f i e l d at the transformation temperatures examined. Upon cooling at 0.4°C per minute the y-phase developed in the form of Widmanstatten plates.A Similar structure was obtained after long annealing periods in the two-phase f i e l d . The alloys 26 and 27 at % Al also showed some amount of Widmannstatten structure when cooled slowly to room temperature. y-Phase as Product of Slow Cooling to Room Temperature In the alloy 24 at % Al cooled to room temperature at 0.6°C per minute, the resulting structure consisted of a layer-like, presumably - 103 -duplex (C+u) structure. During slower cooling the y-layers grew thicker and formed plates grouped in bands apparently parallel to a set of habit planes in the y-phase. A single grain of the y-structure is a plate. " The thickness of the plates was dependent ojk the cooling rate. The plate-shape in two dimensions as seen on the micrograph was irregular and seemed to be affected by the polishing and etching conditions (Figs. 36 and 37). X-ray diffraction patterns showed that some £-phase was retained in the slowly cooled structures of the alloy 24 at %, though i t was not detected metallographically. The "background" £-phase patterns visible i i the Laue photographs of structures obtained by slow cooling to room temperature always exhibited a substructure similar to that observed in samples after transformations in the two-phase f i e l d . Dilatometry of Phase Transitions Dilat^ometric measurements on the alloy 24 at % Al during anjL--sothermal and isothermal transformations showed.the volume changes related to the phase transitions and revealed the dependence of the transformation processes on the cooling or heating rate. The volume change associated with the transformation £—»-/U. estimated from dilatometry was + 0.13%. This is an approximate value as it was obtained from the linear extension only. The theoretical volume change calculated from the lattice parameters i s : . V - V - • 100 <= + 0.237% - 104 -volume per atom i n the y-structure. volume per atom i n the ^ -structure. Dilatometry established that the expected volume changes corresponding to the reaction £ —>» y occurred only when the sample was cooled continuously from the g-field. Metallography and X-ray observations also confirmed that reproducible results can be obtained only on samples normalized in the g-field. The isothermal dilatometry indicated a step-like discontinuous process when the quenched r;-phase transformed to y at 200°C. Decomposition of g^-Phase in 24 at % Alloy The g-phase could be retained to room temperature by quenching in brine at -4°C, providing the sample thickness was less than 0.005". The quenched g-phase was unstable at room temperature. Precipitation hardening was examined during the g _»•• y decomposition at 100°C. Metallography established that the pre-precipitation stage in which the hardness of the g-phase increased was followed by the stage of the true y-phase precipitation. X-ray measurements substantiated that the g-phase transformed into a transition phase during a pre-precipitation stage. Laue patterns indicated a similarity between the i n i t i a l ' g^ and the transition structure - thus the transition phase was designated g. Though i t s nature was not further investigated, an assumption was made that i t probably possesses compositional and dimensional modulations due to a G.P. zone formation. where V" = V V = - 105 -The jli-phase grown in the B^-matrix had a spherical shape. It is believed that the Bg-phase decomposition at temperatures within the y-phase f i e l d takes place in the following steps: y At 200°C the rate of the reaction did not allow the Bg 8 step to be examined separately. The growth kinetics data obtained at 200°C should therefore be related to the B — y step. The growth rate of the y-particles was constant and equal to 2.5 x 10 cm sec \ The spherical shape of the precipitates and the constant growth rate imply that the process is interface controlled. The morphology, age hardening and kinetic data were obtained from samples of thickness < 0.005". It would be desirable to conduct experiments on bulk samples in order to generalize the conclusions drawn from the present observations (massive samples would require special quenching techniques). g;-Phase This phase showed an extreme sensitivity to deformation. Upon light deformation sharp deformation lines appeared on the prepolished surfaces. Generally three sets of crystallographic directions have been identified in themjajority of the deformed samples. This means that the deformation process does not occur by sl i p on the basal planes only. Metallo-graphy X-ray techniques : \ - 106 -established that the predominant mode,of deformation is twinning. The deformed £-phase always contained narrow-twin lamellae resembling Neumann bands regardless of the nature of the deformation (bending, ro l l i n g , scratching or local deformation). X-ray patterns show that in the i n i t i a l stage of deformation single crystal break, into a strained substructure. In terms of published results this effect i s understood as a t i l t i n g of parts of the crystal in the accommodation region at the twin lamellae. Due to the large number of different twins exposed to an X-ray beam of white radiation, the Laue photographs of the deformed 5-phase did not yield quantitative data necessary to examine the crystallography of twinning. Twinning converted the £-phase single crystal into a polycrystal when deformation was extensive. It is a known fact that la t t i c e reorientation within twin lamellae in h.c.p. crystals can be significant. The degree of randomness developed during twinning increases with the number of possible twinning modes, being largest in h.c.p. structures of the axial ratio — near the theoretical 18 value 1.635. (Mg can for example form twins on the 9 twinning systems. ) The axial ratio of the £-phase (1.628) suggests that several twinning modes may be ..available. The s-phase of the 24 at % alloy contains probably small clusters of the solute atoms as indicated by X-ray scattering effects and cluster-like contrast in microstructures of the S-phase examined in the electron microscope. To substantiate this conclusion more precise experimental procedures would be required. - 107 -Anisothermal Transformation Cgr-»• y in 24 at % Alloy The tjgphase transformed into the y-phase either during continuous heating from room temperature or isothermally. Higher heating rates resulted in a shift of the cj-^y reaction to higher temperatures. Under certain experimental conditions the material undergoing the ?^ -»»- y transformation could be significantly overheated above the temperature of the surroundings due to the released latent heat of the reaction. Surface Relief Observations The surface r e l i e f associated with the phase transformation is necessary, though insufficient evidence of the cooperative character of (he transformation. Various surface rumplings may arise also from volume differences between the parent and the product phases. In the case of surface distortions observed on samples partially transformed into the y-phase the following facts were taken into account: (1) The volume change accompanying the £-#ry-phase transfer i s relatively small. Using the theoretical value of the volume change per atom (0.23%);then in order to form a protrusion of 0.7 micron height above the surface a block of the y-phase should extend at least 1 mm below the surface. It was found that the product nucleated on the surface grew mostly in two dimensions over the surface. Growth into the sample interior was must slower. The penetration of the surface crystals did not exceed 0.1 mm below the surface. It is thus unlikely that the crystal of this thickness would produce optically observed r e l i e f . - 108 -Growth of plates in reverse direction, i.e., from the interior to the surface was impossible as a continuous layer of the product formed rapidly on the surface. If the strain arising from relative volume' differences produced stresses exceeding the yield points of the involved phases then either the s or the y-phase would be deformed. Since the hardness of the s-phase is HV = 90 and of the s-phase HV = 450, i t is clear the s_phase must be deformed. This was not observed. The surface r e l i e f was s t r i c t l y confined to the regions that had transformed to the y-phase. This rules out the possibility that a volume difference can account for the observed surface markings. (2) A comparison of the internal structure of the y-phase blocks and the surface r e l i e f details show that close similarities do exist. A conclusion was drawn that the surface r e l i e f i s directly associated with a cooperative atomic process involved in the y-phase formation. Isothermal Transformation s — ^ y Alloy 24 at % Al: X-ray methods substantiated that single crystals of the s_phase always transformed into a polycrystalline structure of the y-phase. However, the degree of randomness as estimated from X-ray D.S. patterns of fine rods rotated during exposure, was much higher i n comparison with Laue patterns from stationary bulk samples. Since the sample rotation i t s e l f could not explain the observed differences i t was concluded that either surface nucleation or size of the sample - 109 -may affect the transformation process. X-ray Laue patterns from bulk £-phase single crystals which transformed isothermally into the y-phase exhibited effects inconsistent with random, non-cooperative growth of the product. A distinct parallelism exists between the process of deformation and the isothermal £ V transformation. This suggests that , cooperative atom trans-lations occur during this transformation. It i s important to note that the y-phase change taking place during continuous cooling from high temperatures showed the same characteristics. Common effects were summed up in the following Table. Single Crystals Deformed £ • Sq—». u Isothermal £ —^- y Anisothermal Amount of deform. X-ray Laue effects Structure Treatment X-ray Laue effects Structure Temper-ature X-ray Laue effects Structure 3-10% strained substracture spl i t t i n g of Laue spots twinned s 20 to 40 sec. at 240°C splitting of Laue spots u-plates and s 415°C strain free substruc-ture 17% i £-Debye rings (fragment-ation to polycrystal) twinned X,' 100 sec. at 240°C y-Debye rings y-phase 300°C y-Debye rings y-phase - I l l -I t i s not assumed t h a t the tj — y t r a n s i t i o n i s a c c o m p l i s h e d by t w i n n i n g o r any c o o r d i n a t e d atom s h u f f l i n g o n l y . The r e a c t i o n s h o u l d be viewed as a combined p r o c e s s s i m i l a r t o m a r t e n s i t e f o r m a t i o n s i n which two atomic mechanisms are i n v o l v e d : (1) The change of the c r y s t a l s t r u c t u r e ( f o r example, c u b i c l a t t i c e changes to t e t r a g o n a l ) . (2) The shape d e f o r m a t i o n by s h e a r o r t w i n n i n g n e c e s s a r y to s a t i s f y the c o n d i t i o n of an u n d i s t o r t e d h a b i t p l a n e . A n e c e s s a r y c o n d i t i o n f o r the t r a n s f o r m a t i o n t o o c c u r m a r t e n s i t i -c a l l y i s an e x i s t e n c e o f a c l o s e g e o m e t r i c a l r e l a t i o n s h i p between ^ £ p a r e n t and product l a t t i c e s . The h.c.p. s t r u c t u r e o f the i d e a l a x i a l r a t i o can be c o n v e r t e d i n t o f a c e c e n t e r e d c u b i c s i m p l y by a s h e a r d i s p l a c e -ment o f d e n s e l y packed b a s a l p l a n e s o v e r a d i s t a n c e l e s s than the c l o s e s t i n t e r a t o m i c approach. The t,-phase has the a x i a l r a t i o — c l o s e t o the i d e a l v a l u e and the p r o d u c t y-phase has a c u b i c symmetry. Thus the p r i n c i p a l c o n d i t i o n f o r phase change by a c o o p e r a t i v e atom movement i s f u l f i l l e d . A g r e a t number o f the £-lines i n the X-ray D.S. p a t t e r n s e i t h e r c o i n c i d e w i t h the l i n e s o f the y-phase o r a r e v e r y c l o s e . T h i s i m p l i e s a s i m i l a r i t y of t h e i r p l a n a r s p a c i n g s . E l e c t r o n d i f f r a c t i o n p a t t e r n s e s t a b l i s h e d t h a t some of these p l a n e s a r e a l s o p a r a l l e l f o r b o t h m a t r i x and p r o d u c t phase. An X-ray e x a m i n a t i o n o f the C JQ—3»- y r e a c t i o n i n the C ^ - s i n g l e c r y s t a l s showed t h a t p a r t i c u l a r tj-phase r e f l e c t i o n s p r o d u c e d by Ka r a d i a t i o n d i d not i n d i c a t e a c r y s t a l s t r u c t u r e change. They m e r e l y decomposed i n t o a number o f s p o t s l o c a t e d on the y-phase Debye r i n g s d e m o n s t r a t i n g the f r a g m e n t a t i o n o f the s i n g l e c r y s t a l i n t o a p o l y c r y s t a l . - 112 -Many phsical features of martensitic transformations are displayed by deformation twinning. Hence the qualitative correspondence between features observed on the deformed tj-phase surface and surface markings arising from the u-formation furnish further evidence in favor of the martensitic-like cooperative mechanism. The cracks frequently observed in the centre of the y-phase blocks were considered a consequence of a stress concentration at twin intersections (Fig. 90C). This possibility has been reported in the l i t e r a t u r e . ^ The complexity of X-ray D.S. patterns of the y-phase indicates, however, that in addition to a cooperative atom translations other processes must take place in the transformation event. The kinetic measurements established that the overall rate of the CJQ — i» y reaction is controlled by a thermally activated rearrange-ment of atoms over a short distance. An average activation energy of 25 .6 kcal per mole was obtained. Non-constant growth rate and the shape change during growth confirmed the assumption that more than one simple atomic process must operate in the tj—«•» y transformation. Many y-phase nuclei exhibited a lens-like shape common to a growing twin or martensite plate. Deformation of 15% at R.T. increased the rate of the £ —*». y transformation by almost 100%. Alloy 25 % Al: The resulting structure of the isothermal transformation £Q—*- y consisted of Widmannstatten y-plates in the £-matrix. The reaction - 113 -rate was much slower than in the alloy 24 at %. (A complete transformation required 100 hours at 200°C.) The u-plates and the ^-matrix exhibited the following crystallo-graphic relationship: (0001)e//(111)y [1120]c/7 [011]y The habit planes of the £-phase were: (2025) (1124) (1234) The deformation of the j^-phase and the subsequent isothermal treatment resulted in the nucleation and growth of a product similar to the y-phase in the 24 at %. The amount of this product was directly related to the amount of deformation. Thin layers of the product were also formed on electropolished undeformed surfaces. y-Phase Structure Analysis X-ray analysis and electron transmission microscopy provided evidence that the y-phase i s probably an ordered cubic structure and contains clusters rich in Ag or Al respectively. This i s not in agreement with published results. The original y-phase structure specification was based on analysis of the X-ray powder patterns. This - 114 -method was shown to be u n r e l i a b l e for t h i s structure. The correct y-phase c r y s t a l structure analysis from X-ray patterns only would be very d i f f i c u l t . (There are 80 d i f f r a c t i o n l i n e s i n the CuKa d i f f r a c t i o n pattern of the y-phase.) Clu s t e r i n g e f f e c t s observed i n the e l e c t r o p o l i s h e d t h i n f o i l s were i n agreement with observations made on t h i n fragments and r e p l i c a s . The s i n g l e c r y s t a l of the y-phase i s a t h i n p l a t e . Though the thickness of plates was not measured i t i s estimated to be l e s s than the toil thickness used for the electron microscope. Evidence was presented showing a d e f i n i t e c r y s t a l l o g r a p h i c r e l a t i o n s h i p between the tj-matrix and the y-plates. Both the p l a t e - l i k e shape and the c r y s t a l l o g r a p h i c correspondence between phases are i n d i c a t i v e of the suggested coordinated atomic growth mechanism. Martensitic products are known to consist of stacks of f i n e twin-rel a t e d lamellae resolvable only by electron microscopy. I t i s thus possible that the y-phase i s b u i l t up of p a r a l l e l twin p l a t e s . The twin components i n some martensitic products grow d i r e c t l y drom the 24 parent and no deformation twinning of the product occurs. This d e s c r i p t i o n seems to f i t to the ej —o- y transformation. A twin plate developed i n the matrix becomes apparently a y-phase p l a t e . For the i n i t i a l transformation stage the product appears to form a duplex l a y e r - l i k e (Jj+y) structure. Thin y-phase layers grew i n t o thicker plates during longer annealing periods. - 115 -Comments on Transformation Mechanism The SQ »- y reaction is a composition invariant transformation. There are two possible mechanisms for this type of transformation: Martensitic or interface controlled mechanism. The phase structure change can occur eventually by a combination of both. The SQ —*»• y reaction is considered to be a complex process involving a thermally activated mechanism and coordinated atom transla-tions which are not thermally activated. In the h.c.p. metals an almost unlimited volume of the structure can be converted into a twin. An atomic configuration in thin twin lamellae on particular crystallographic systems may be identical with 29 stacking faults and the term twin fault i s used. Thus the term twin used in this work to describe the coordinated atom process in the S —=»- u transformation includes a l l shear displacements occurring within the original lattice capable of producing either a structure change or a random orientation of the product. There is a close association between the stacking fault energy and twinning. The lower the stacking fault energy the lower is the 2 8 stress to generate twins. Hendrickson and Fine found that the stacking • fault energy of Ag is low and rapidly decreases with increasing Al content, at composition 25 at % being zero. This implies a great susceptibility of the S-phase to twinning and stacking faults formation. Results obtained during the present investigation provided a complete - 116 -29 agreement with this statement. According to Bell and Cahn twins nucleate by the action of locally enhanced shear stress. Existing twins help to create others due to an autocatalytic action at the edge of the twin. This offers an explanation why the y-phase grew mostly as a bulky conglomerate of thin c r y s t a l l i t e s . The Widmannstatten y-phase of the 25 at % alloy also developed by a successive addition of new lamellae sidewise to existing bands. In order to explain the extreme brittleness of the y-phase, the following mechanism is suggested: During the isothermal transformation solute atoms cluster together. The difference in size of Ag and Al atoms is not great (the ratio of ladii i s 1.03) and small clusters would not produce significant strain. However, separation of the one element into large layer-like clusters can give rise to a remarkable misfit and strain at the cluster-matrix interface. The dependence of the strain 2. magnitude on the cluster size is shown schematically in Fig. 17$. The ratio of atomic radii was chosen as 1.08. For simplicity only a one-dimensional, arrangement was considered. Each set (A, B, and C) in r Fig. 172. represents a composition of the formula W^ D (W = white, D = ,dark). Arrangement A: Due to the random distribution in two adjacent rows, there is negligible short-range disregistry between two adjacent atoms located at the same atomic sites and no long-range disregistry (over the length X = 20 atomic sites). - 117 -Arrangement B: Clustering of the five D-atoms introduces a short-range disregistry = X' over the five atomic sites. These small clusters may not produce any long-range disregistry (the length X of the adjacent rows is the same). The distortions could be accommodated by elastic lattice strain. The hardness would be expected to be higher than in case A. Arrangement C: One element (D) is completely separated from the larger volume of the structure. In this case the disregistry between the adjacent atoms is equal to the closest interatomic distance "a" at the 13th atomic site. The disregistry over the 20 atomic sites (distance Y) is Z = 1.6 a. Arrangement C is considered most lik e l y to be j u s t i f i e d by experimental results. When the y-phase broke upon impact, the flake-like fine particles f e l l off the fracture surface. No specific mechanism for the £ ^ y transformation is proposed. However, from experimental results i t became evident that a short-range rearrangement of the Ag or Al atoms into energetically more favorable positions takes place. This rearrangement gives rise to the long range superlattice indicated by X-ray. The nature of - 118 -the ordering may be i d e n t i c a l with a separation of one element i n t o r e g u l a r l y stacked layers (described previously as two-dimensional c l u s t e r s ) . This would be consistent with b r i t t l e n e s s of the y-phase and "domains" showing dark contrast i n the y-structure revealed by electron microscopy. K i n e t i c s analysis also established that an atomic process equivalent to ordering i s the rate c o n t r o l l i n g step of the z, —=»- y reaction. The cooperative atomic displacements evidenced by the surface r e l i e f s may be a t t r i b u t e d e i t h e r to an i n v a r i a n t plane . s t r a i n or anisotropic dimensional changes r e s u l t i n g from the ordering process. The z, — _ u transformation exhibited the most s t r i k i n g s i m i l a r i t y with the 8 (or g ' ) — > . z, reaction occurring i n the 27 32 33 Ag-Zn a l l o y . , J Z , J J In the present study no difference was made between the z, —*»» y isothermal transformation occurring at 160°C, 200°C and 300°C. Having accepted the ordering as a part of the reaction then the transformation temperature may,however,affeet the transformation mechan-ism (the attained degree of order w i l l be greater at lower temperatures). The following observations seem to confirm t h i s assumption: The surface r e l i e f was more d i s t i n c t when the gj-phase transformed at 160°C then at 200°C. The y-phase produced isothermally at 300°C contained large c i r c u l a r c l u s t e r s not observed i n the y-phase obtained at 200°C. - 119 -CONCLUSIONS The contribution of this work to an understanding of the phase transformations in the alloys Ag-20.1 to 27 at % Al can be summarized in the following conclusions: (1) The morphology of the structures that develop during continuous cooling and long periods of annealing were established. (2) The designation of the y-phase as a disordered homogenous phase which is isomorphic with a B-Mn structure was found to be incorrect. A re-investigation of this structure is desirable. The electron transmission microscopy would be the most effective method for a complete analysis of this phase. (3) The g-phase quenched to room temperature exhibits age harden-ing at low temperatures and a transition structure (g) is formed. (4) The growth of the y-phase during the isothermal gq—*- g —*- y transformation is an interface controlled process. (5) The rate of the isothermal transformation SQ =»• y is interface controlled, but a cooperative atomic process is also incorporated in this reaction. - 120 -(6) The y-phase in the 25 at % alloy precipitates as Widmannstatten plates either during continuous cooling or under isothermal conditions. There are several habit planes for the y-phase plates in the t;-matrix. (7) The quenched £-phase deforms easily by twinning. Twins are always narrow lamellae comparable to Neumann bands. A single crystal of the s-phase fragments gradually into a polycrystal during deformation. (8) There is an analogy between the s-phase single crystal deformation and the s — u single crystal transformation in the sense that the fi n a l result of both processes i s a polycrystal. - 121 -A P P E N D I X X-RAY CRYSTAL STRUCTURE ANALYSIS Content: TABLE 1 OL - Phase calculated ( 20 at % <tlloy ) fo - Phase measured ( 24 at % alloy ) £ - Phase measured ( 24 at % alloy ) TABLE 2 n| - Phase produced by slow cooling to R T room temperature. Cu Kpi. radiation, 24 and 25 at % alloys. TABLE 3 M e and Misoth.- Phases. Reflections calculated R T for Fe and Cr radiation. Observed reflections of the <4 at % alloy. - 122-APPENDIX - TABLE 1 cL- PHASE Calculated 20 h k 1 38.4 111 44.6 200 65 220 78,10 311 82.30 222 98.86 400 111,2 331 116.6 420 137.04 422 161.6 333 (h- PHASE Measured 20 h k 1 39 110 56.9 200 71,5 211 8J.25 220 98,1 310 112 222 127.1 321 146.5 400 £ - PHASE Measured 20 h k 1 36.5 100 38.85 002 41.45 101 54.26 102 65.32 110 71.75 103 75.87 200 78.47 112 80il3 201 82,83 004 89.48 202 93.74 104 105.14 203 110.25 210 113.55 211 116,45 1 U 123,51 105 124.. 11 212 129.71 204. 136,82 300 146.47 213 160.93 302 162.88 006 - 123 -LEGEND TO TABLE 2 Column I Column I I S - h2+ k2+ l 2 sin © calculated for cubic l a t t i c e of the a =» 6.92 A and Cu K a radiation - 1.^ 1+05 A ) Estimated relative intensity of reflection on D.S patterns: Column VII Column VIII Column IX VW » very weak W = weak M = medium S «* strong VS " very strong Rel.intensity, y- phase (present study) " y " (Westgren and Bradley) " B-Mn (Westgren and Phragmen) Values l i s t e d i n column X have quantity : N ( A2+ B 2) where N =» multiplicity factor A - cos 2 TT ( hx + ky + l z ) B - £J sin 2 TT ( hx + ky + l z ) x,y,z = atom coordinates APPENDIX - TABLE 2 I II III IV V VI VII VIII IX X s h k 1 Calculated Sin 2 0 (Cu K<*- ) Observed X-Ray D. S. Patterns of h \ - Phase /i-Mn Patterns Present work W & B W & P Calculated Relative Intensity (Preston) S i n 2 6 2 0 25 at % Intensity 25 at % Intensity Intensity 1 24 at % 25 at % 1 100 0.012 - - - - - - 0 2 110 0.024 - - - - - - 2,4 3 111 0,036 0,0374 0.0375 22,34 W - W 2,7 4 200 0,048 0,0502 0.0496 25.74 W - - 0 5 210 0.060 0.0627 0.0620 28,48 M M If 9,1 6 211 0,072 0,0751 0.0746 31.70 M W - 1,6 0.0918 0.0914 35.20 W - - -- 0,0960 36.10 v.w. - - -8 220 0.096 0.1029 0,1015 37,15 v.w. - - 2,4 9 300 221 0.108 0.1126 0.1123 39,15 v.s. s s 272.0 10 310 0.120 0.1253 0.1250 41.40 v.s. s M 163,0 I II III IV V VI VII VIII IX X 11 311 0,132 0.1337 0.1371 43,46 5 M M 89.6 12 222 0*144 0.1484 0.1487 45.36 W - - 0,3 13 320 0.156 0.1623 0yl6l0^ 47.31 W - - 2.0 14 321 0.168 0.1739 0.1740 49.31 M . W 28.8 16 400 0.192 - 0.2069 54-12 W - - 0 17 410 322 0*204 0.2109 0.2116 54-77 W v.w. w 4.0 18 330 411 0*216 0-2235 0.2235 56-42 s s w 30.5 19 331 0.228 0.2357 0.2363 58.17 w - - 0.4 20 420 0.240 0.2484 0.2479 59.72 s s M 28.6 21 421 0.252 0.2602 0,2608 61.42 w w - 1,8 22 332 0.264 0.2733 0.2728 62.97 M M M 5,1 - 0.2885 64.97. W - -24 422 0.288 0.2974 0.2973 66.077 w - • - 9,1 25 430 500 0.300 0.3098 0.3101 67.68 M M S 9.7 I II III IV V VI VII VIII IX X 26 431 510 0,312 0,3228 0,3219 69,13 V.S. s s 142,2 27 333 511 0,324 0.3347 0,3342 70,63 M M V.S. 25,3 0,3424 71,63 W - -0.3483 72.33 W - -29 432 520 0,348 0,3605 0,3595 73,68 V.S. s V.S. 127,2 30 521 0.360 0,3720 0,3717 75,13 s M 38,9 32 440 0,384 0.3969 0,3989 78.33 M - 2.8 33 522 441 0.396 0,4132 0,4132 80,00 V.W. - - 3.6 34 433 530 0,408 0,4205 0,4213 80,93 M - 9.3 35 531 0,420 0,4331 0.4334 82,34 s S 36,2 36 442 600 0.432 0,4448 0.4443 83.74 s s 51,4 37 610 0.444 0,4582 0,4572 85,10 s M 33.6 38 532 611 0,456 0,4709 0.4703 86,60 s S 32.3 40 620 0.480 0,4949 0.4947 89.39 w V.W. 7.3 41 540 621 0.492 0,5067 0.5065 90,74 M M 19.1 I I I I I I IV V V I V I I V I I I IX X 42 541 0 .504 0 .5189 C 5 1 9 2 9 2 . 2 0 v .w . V.W. 43 533 0 .516 0.5320 0.5314 9 3 . 6 0 w -4 4 622 0 .528 0.5420 0,5418 9 4 . 8 0 w -45 630 542 0 . 5 4 0 0 .5572 0 .5558 9 6 , 4 0 M s 4 6 631 0-552 0 ,5681 0 ,5683 97-85 w v .w . 48 444 0*576 0 ,5935 0*5937 100 .80 v .w . -49 700 632 0 ,588 0 .6046 0»6040 1 0 2 , 0 0 s s 50 710 550 o»6ob 0 .6165 0 .6172 103-55 M s 51 711 551 0 .612 0»6288 0 .6299 105 .05 W 0-6415 0 .6415 106 .49 V.W. 52 640 0 .624 0 , 6 5 2 7 0.6538 107-91 M 53 720 641 0 . 6 3 6 - - - -54 721 633 0.648 0 ,6660 0 .6662 109 .41 M 56 642 0 ,672 0 ,6892 0 ,6911 112 ,46 M I I I I I I IV V VI V I I V I I I IX X 57 722 544 0 ,684 0 , 7 0 1 3 0 ,7031 1 1 3 , 9 6 M 58 730 0 . 6 9 6 0 , 7 1 4 4 0 ,7154 115 ,51 W 59 731 553 0 ,708 0 . 7 2 6 9 0 .7275 117 .07 S 61 650 0 .732 0 , 7 5 1 9 0 .7524 120 ,32 s 62 723 651 0 ,744 0 ,7 6 3 8 0 ,7644 121 ,92 s 0 .7746 123,72 w 64 800 0 ,768 0 . 7 8 8 6 0 ,7909 1 2 5 . 5 7 w 65 810 744 0 ,780 0 , 8 0 0 6 0 .8014 1 2 7 . 0 7 M 66 811 741 0 .792 0 .8123 0 ,8138 128 ,88 s 67 733 0 ,804 - 0 . 8 2 6 6 130 .78 V.W. 68 820 644 0 . 8 1 6 0 .8378 0 .8383 132 .58 M 69 742 0*828 0 . 8 4 9 3 0-8516 134»68 M 70 653 0 .840 0 . 8 6 1 9 0-8632 136>58 S 72 822 660 0 .864 0 , 8 8 7 0 0 ,8874 140.78 M I II III IV V VI VII VIII IX X 73 830 0.876 0,8989 0,9000 143,13 M 74 831 743 0.888 0,9109 0,9123 145,54 S 75 751 0.900 0.9237 0.9247 148,14 w 76 662 0,912 0.9364 0,9371 150,94 w 77 832 654 0.924 0,9482 0,9498 154.10 M 78 752 0.936 0.9605 0.9618 157,45 s 80 840 0.960 0,9855 0,9863 166,56 M 81 900 841 0.972 82 910 833 0.986 • APPENDIX - TABLE 3 I I I I I I IV V VI VII VIII IX X XI Fe - Koc radi a t i o n Cr - KpC radiation Calculated f o r cubic l a t t i c e Observed Calculated f o r cubic l a t t i c e Observed 2 0 S i n 2 e 2 G non cubic Sin 2 0 2 9 S i n 2 9 2 9 non cubic S i n 2 9 /•R.T. j%soth M R.T. Jit Isoth. JIL R.T. J Isotl Jli \ R.T. P- Isoth. 1 16*08 0,0196 19.16 0.0274 2 22.80 0.0391 - - 27.05 0.0547 - -3 28.04 0,0587 0.0586 0.0599 33.30 0,0821 - 0.0832 4 32,50 0.0783 0.0789 0»0791 38.66 0.1095 0,1109 0.1097 5 36.44 0.0978 0.0984 0.0990 73.44 0.1369 0.1386 0,1376 6 40.08 0.1174 0.1174 0.1185 47.84 0.1644 0.1655 0,1654 44.8 44.9 0.1451 0.1460 53,8 53,8 0,2050 0.2051 8 46.60 0,1565 0,1611 0.1614 55.84 0,2192 58,5 0,2282 . 0,2385 9 49.64 0.1761 0,1775 0,1774 59,56 0,2466 0,2471 0,2471 10 53.52 0.1957 0,1962 0.1968 63.12 0,2740 0,2739 0,2747 I II III IV V VI VII VIII IX X XI 11 55.28 0,2152 0,2156 0.2163 66.60 0.3014 0,3024 0*3012 57.1 0,2253 12 57,06 0,2348 0.2331 0.2360 69.96 0.3288 0,3288 13 60.58 0.2540 0,2544 0.2555 73.28 0.3562 0.3557 0,3562 U 63.12 0.2739 0,2745 0.2740 76.54 0.3836 0.3831 0,3832 65,1 65.2 0.2898 0,2903 16 68.08 0.3131 0,3220 0.3239 82.92 0.4384 0.4385 -84.6 0.4532 17 70.44 0,3326 0.3318 0.3341 86,08 0.4658 0,4650 0,4652 18 72.81 0.3522 Of3517 0-3532 89.22 0.4932 0.4912 0.4927 19 75-14 0.3718 0.3740 0.3729 92.36 0.5206 0.5202 20 77.56 0.3913 0.3913 0.3911 95-50 0.5480 0.5457 0,5455 21 79-74 0,4109 0,4106 0.4104 98.68 0.5754 0.5734 0.5732 80.6 80.7 0-4183 0.4190 100.1 0.5876 22 81.00 0.4305 0.4295 0.4302 101.86 0.6028 0.6003 82.4 0.4343 24 86.52 0.4696 0.4677 0.4675 108.36 0,6576 0.6553 0.6552 88.1 0.4834 25 88.76 0.4892 0.4887 0.4889 111.70 0.6850 0.6800 0.6812 I II III IV V VI VII , VIII IX X XI 26 91.00 0.5087 0.5075 0.5081 115.14 0.7124 0.7070 0,7085 27 92,26 0,5283 0,5258 0,5268 118,66 0,7398 0.7352 0,7351 *29 97.77 0.5674 0,5645 0,5651 126,10 0,7946 0,7894 0,7915 30 100,02 0.5870 100,7 102,6 102, •; 0.5835 0,5929 0.6092 0.5845 0.6097 130,08 0,8220 132 0,8165 0,8350 0,8182 32 104.6 0.6261 0,6232 0.6237 138,92 0,8768 0,8711 0,8723 33 106.94 0,6457 0.6446 0.6447 143,84 0,9042 34 109.30 0.6653 0.6617 0,6634 149,66 0,9316 0,9253 0,9252 35 111.68 0,6848 0.6805 0.6814 156,62 0,9590 0.9524 0,9532 36 1U.12 0,7044 0.7007 0.7012 166,60 0,9864 0,9801 0,9804 37 116.60 0,7240 0.7197 0.7198 38 109.04 0,7435 123.9 123.4 0.7395 0.7790 0.7395 0.7787 40 124.44 0,7827 0,7859 41 127,08 0.8022 0,7983 0.7976 I II III IV V VI VII VIII IX X XI 42 130.06 0,8218 132,4 132,3 0,8179 0,8370 0,8172 0,8364 43 133,06 0,8414 135,4 135,3 0,8491 0.8562 0,8491 0,8553 44 136.2 0,8609 0.8656 0,8649 45 139,54 0,8805 0,8764 0,8755 46 143,16 0,9001 145.9 145.6 0.8947 0.9140 0.8952 0.9127 48 151.44 0,9392 152,9 152.8 0.9345 0.9453 0.9346 0,9448 49 156,6 0.9588 0.9537 0,9527 50 163.08 0,9784 169.9 0.9733 0.9923 0.9727 PHASE TRANSFORMATIONS IN THE AG - 20.1 TO 27 AT % AL ALLOYS by LUDVIK CAMBAL B.A.Sc., Technical University, Brno, Czechoslovakia VOLUME 2 ( F I G U R E S ) FIG.i - 135 -FIG. 2 FIG. 3 800 2.860 20 40 60 80 Atomic Per Cent Aluminium FIG. h 400 mm ^ . w 1 Q U A R T Z T U B E 2 G R A P H I T E MOLD 3 INDUCTION C O I L A M U L L I T E ROD 5 C O P P E R M O L D 6 S T A N D F I G . 5" Q U E N C H I N G C Y C L E VALVE 7 VALVE 10 HEATING UNDER NEUTRAL GAS JOUT FURNACE--iFURNACE GAS PUMP EVACUATION BEFORE QUENCHING QUENCHING WATER iFIG. 7 - 139 -FIG. 8 - il*o -FIG. 9 \hN\N\ FIG. FIG. IJ - 1U1 -A CAST,HOMOGENIZED, SLOWLY COOLED O QUENCHED FROM 5 0 0 ° C a QUENCHED AND TEMPERED 30 MIN. AT 2 0 0 ° C o QUENCHED AND TEMPERED 40 HRS.AT 2 0 0 ° C ! , , ( ! 20.I 23 2 4 25 2 6 27 At %AI FIG.12 - 1)42 -20.1 % 5 0 0 ° C 400 ° C 3 0 0 ° C 200° c FIG. 13 - 1U3 -23 % FIG - 110. -24 % 500° C 435 ° C 400 ° C 11 H V H H H H H H H H H H H H H H i 300° C 200° C FIG. 15 - 11£ -25 °/o FIG. 1 c - 146 -FIG.1? FIG.^B FIG. 19 111111111^^ uij||iH FIG.?o - 1U7 -- me -- 1U9 -FIG. 28 - i5o -- 151 -FIG. 33 - 152 -FIG. 35 - 153 -FIG. 36 FIG. V - 15U -- 155 -FIG 1*6 -FIG. , P FIG.! 9 FIG. 50 FIG.51 - 1# -- 1 6 0 FIG. Sh - 161 -FIG31 FIG."0 - 163 -- 16U -- 165 -- 166 -30J FIG. 69 440-420-400-380-360-340-320 300-280-260-240-220 200-180-160-140-120-100 p PHASE AFTER TEMPERING AT 2 0 0 ° C p QUENCHED ANDTEMPERED AT 1 0 0 ° C Ofi HR. /•PHASE AFTER TEMPER. AT 100 °C 43 HRS. J3 AGED AT ROOM TEMP O 90 HRS. — , 1 10.000 100.000 100 1000 TIME SEC. FIG. 71 - 169 FIG. 73 171 -- 173 -- 17U -FIG. R? = 175 -- 1 7 6 -- 177 -- 178 -- 179 -8 9 FIG.92 ... j. --—T - 1 1 - | . ! .- i .| ; 1 — — 7 ISOT HF me y i FVFL 201 °C 206»e t ..:._! .' L_._ I • 6 1 • • j " ., ! ••--j- — -—f— UJ 5 1 ! \ \ . • i 1 r - • -ALUM — TEMR DF THE i — i SAMPLI DURING rRANSF % — M _i ui i — — . ASTRA MSFORM ED TO / - — _— u „>... £ ? i - - -\ - - -O: -••5--P 1 j j ..... , 1 24°C i | j ! | • i i j • j 4 5 0 100 i 1^0 i 2^ )0 290 SEC •o. 3&0 4< 500 5&0 j 600 TIME i : :|: F I G 93 • — [ — -Q - ! « U » C j ISOTHE RMAL 1 .EV :L !• • i • 1 i u . i • i -i - •• _:J 7 i ' i i —•• I i • " j ' . — L — 6 • | i i 1 .• . 2 — s : 5 u 4 .... i i . ].-.•: I ''' a I C — - - .... i . ! •; > E TEH DUR I P OFT ING TR HE SAM MtSF. $ PLE i i I AS rRANSI ORMEC 1 '. "i , ; : OL ? TO f ..... ' S a. i. — - T - -----•j | — -I I 1 1 i • ! • • c < K 10 a 4 )0 ;;:;( 2( 4 » S Tl IE ,C.: i l l: • 1FIG. 9k [ _ _.. . . . . . i — i i i i ..  -- --;••]•--._......! 32 4°C -—--j—-u| 10 K mm >MAL, L [VEL L.JP4J f d 8 I ai 4 TEMP 0 -THE S AMPLE DURING VNSf f AS TRA HSFORH IED;TO J1 I • . . . . . \ / 243 ""j i | 1 : ( ) | 5 ! 0 K KD C •0 2< K> 2! TO 0 IE < SE 5 0 K e 10 a )0 2! O K) ; " ""j "" j ---— . . . . . . -— . . . . . . . — . I :— j . . . . . . — • j - . -.: |:-:  1  FIG. 9 5 FIG. 10 L ' 9 | J - 2.9 L -- 188 -FIG.106 FIG.1 — - 190 -- 192 -- 193 -o < IOO 200 300 400 TIME SEC, FIG.116 500 600 i r 120 90 TIME MIN " i — r 150 FIG. 117 .: 8ti'9IJ 03S 3Vi\l 00001 0001 001 01 I " I I I I — I — I Il 1 1 I I I I I 111 I I I I I I - 195 -1000. I00 I rj J 2.0 2.I 2.2 , 2.3 2.4 -fxlO 3 FIG. 119 - 196 -- 200 -a - 2 201. -FIG. FIG.131 FIG.13? FIG. 133 FIG.135 - 206 -136 a 136 b 13 6 c 208 -138a - 210 -13 8c - 211 -212 -- 213 -- 21 It -FIG. - 216 -FIG.1£? FIG.' - 2 2 0 -- 221 -FIG. 1 223 -REFERENCES 1. . . . M. Hansen, Constitution of Binary Alloys (1958) 2. . . . M. H i l l e r t , B.L. Averbach and M. Cohen, Acta Met., Jan., 1956. 3. . . . C S . Barret, Trans. AIME 1931, p. 88 4. . . . A . Guinier, Acta Met. 1953, p. 568 5. . . . V . Gerold, Advances i n X-Ray Analysis 1963, Vol. 7. 6. . . . R.B. Nicholson and J . Nutting, Acta Met. 1961, p. 333. 7. . . . T.B. Massalski and B. Cockayne, Acta Met. 1959, p. 762. 8. J.P. Neumann, Acta Met. 1966, p. 505. 9. T.B. Massalski and H.W. King, Progress i n Materials Science 1961, Vol. 10, No. 1. 10. . . . E.B. Hawbolt and L.C. Brown, Trans. AIME 1968, Vol. 242, p. 1182. 11. . . . D. Arias and J.E. K i t t l , Trans. AIME 1969, p. 182. 12. . . . E.B. Hawbolt and T.B. Massalski, Trans. AIME 1971, p. 1771. 13. . . . CW. Spencer and R.J. Knight, Journal of Metals 1957, p. 689. 14. . . . F. Foote and R. J e t t e , Trans. AIME 1941, p. 151. 15. . . . A. Westgren and A.J. Bradley, P h i l . Magazine 1928, p. 280. 16. . . . V . Daniel and H. Lipson, Proc. Roy. Soc. A. 17. J . Burke, The Ki n e t c i s of Phase Transformation i n Metals 1965. 18. . . J . W. C h r i s t i a n , The Theory of Transformations i n Metals 1965. 19. . . . E.B. Hawbolt, unpublished work. 20. . . . D i f f u s i o n data No. 1, 1971. 21. . . . E.B. Hawbolt, personal communication. 22. . . . C.H. Mathewson, Proceedings AIME Vol. 80 1928, p. 311 and Vol. 78, p.7. 23. . . . T.A. Wilson and S.L. Hoyt, Proceedings AIME, Vol. 78, 1928, p. 241 . 24. . . . Deformation Twinning, M e t a l l u r g i c a l Society Conference, Vol 25, 1963. 25. C.H. Mathewson and A. P h i l l i p s , Proceedings AIME, 1927, p. 186. 26. . . . CD. Preston, P h i l . Magarine 1928, 5, p. 1207. 27. . . . J.E. K i t t l and A. Cabo, Symposium on Phase Transformation, I n s t i t u t e of Metals, London 1969 . 28. . . . A.A. Hendrickson adn M.E. Fine, Trans. AIME, 1961, p. 103. 29. . . . R.L. B e l l and R.W. Cahn, Proc. Roy. Soc. 1957 A, p. 494. 30. . . . Masson, Acta Met. 1960, Vol. 8, p. 71. 31. . . . Electron Microscopy, International Congress Sept. 1966, Vol. 1, p. 314. 225. 32. W.J. Kitchingman, Acta Met., 1962, p.799. 33. H.M. Clark and CM. Wayman, Seminar on Phase Transformations ASM, October 1968. 


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