FINE GRAINED STRUCTURE FORMATION IN ALUMINUM-COPPER BASED ALLOYS DURING SEVERE PLASTIC DEFORMATION

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(1)FINE GRAINED STRUCTURE FORMATION IN ALUMINUM-COPPER BASED ALLOYS DURING SEVERE PLASTIC DEFORMATION. Inna Mazurina. UEC Tokyo (The University of Electro-Communications) March 2008.

(2) FINE GRAINED STRUCTURE FORMATION IN ALUMINUM-COPPER BASED ALLOYS DURING SEVERE PLASTIC DEFORMATION. by Inna Mazurina. A Thesis Submitted in Partial Fulfillment of the Requirements for the Degree of Doctor of Engineering. Department of Mechanical Engineering and Intelligent Systems UEC Tokyo, (The University of Electro-Communications) Tokyo, Japan March 2008.

(3) FINE GRAINED STRUCTURE FORMATION IN ALUMINUM-COPPER BASED ALLOYS DURING SEVERE PLASTIC DEFORMATION. APPROVED BY SUPERVISORY COMMITTEE:. CHAIRPERSON:. TAKU SAKAI. MEMBER:. MAKOTO MURATA. MEMBER:. YASUO OCHI. MEMBER:. HIROMI MIURA. MEMBER:. TAKASHI MATSUMURA.

(4) Copyright by Inna Mazurina 2008.

(5) ABSTRACT. The microstructure evolution taking place during severe plastic deformation was studied in a commercial based coarse-grained Al alloy 2219 and a dilute Al-3%Cu by using equal channel angular pressing (ECAP). The effect of deformation temperature on ultrafine grain (UFG) structure development was investigated in both the alloys in detail. The process of new grain formation was analyzed and the operating structural mechanism was discussed. The deformed microstructures developed during ECAP at various temperatures are mixed ones and categorized into the two types, i.e. strain-induced fine grains with high angle boundaries (HABs) and coarse grains containing subgrains with low angle boundaries (LABs). The misorientation and microstructure characteristics in the fine grained regions are hardly affected by temperature, but the volume fraction of the fine grains is largely decreased with increase in temperature due to less frequent formation of large-scale deformation bands (DBs). The process of fine grained structure formation during ECAP can be classified into the three stages in a wide temperature interval. Namely, the formation of conventional dislocation substructures with LABs accompanied with embryos of DBs in stage 1; then development of large-scale DBs with HABs followed by grain fragmentation, leading to new grain formation occurring along DBs in stage 2, and finally development of fine grained structure taking place at high strains in stage 3. Microstructural development is hardly influenced by pressing temperature in stages 1 and 2, while the formation of the new fine grained structure at high strains is significantly affected by it in stage 3. It is concluded, therefore, that the strain-induced grain formation results from dynamic formation of HABs in stages 1 and 2, subsequently followed by frequent operation of dynamic recovery in the interiors of HABs of DBs in stage 3. The main process for formation of deformation-induced grains can be controlled by grain fragmentation due to development of DBs with moderate angles at low to medium strains, followed by transformation of the boundaries into high angle ones assisted by dynamic recovery with further straining, that is similar to in-situ or continuous dynamic recrystallization (cDRX)..

(6) LIST OF ABBREVIATION. SPD. - severe plastic deformation. ECAP. - equal channel angular pressing. HPT. - high pressure torsion. ARB. - accumulative roll bonding. MDF. - multi directional forging. HABs. - high angle boundaries. LABs. - low angle boundaries. DRV. - dynamic recovery. cDRX. - continuous dynamic recrystallization. gDRX. - geometric dynamic recrystallization. dDRX. - discontinuous dynamic recrystallization. EBSP. - electron back scattering diffraction patterns. GBS. - grain boundary sliding. GNBs. - geometrical necessary boundaries. OIM. - orientation imaging microscopy. Ref.. - reference. SAED. - selected area electron diffraction. SEM. - scanning electron microscopy. UFG. - ultrafine grain. TEM. - transmission electron microscopy. DBs. - deformation bands. MSBs. -microshear bands. deg. - degree. min. - minute. ks. - kilo seconds. SFE. - stacking fault energy. Al. - aluminum. Cu. - copper. AA 2219 - aluminum alloy 2219.

(7) T. - temperature. Tx. - testing temperature. Tm. -melting point. K. - Kelvin. GPa. - Giga Pascal. PD. - pressing direction.

(8) LIST OF SYMBOLS. φ. - channel angle. ψ. - arc of curvature. f. - frequency. VUFG - fraction of ultrafine grains N. - number. ε. - true strain. εcr. - critical strain. Θav. - average misorientation angle. ΣΔΘ - cumulative misorientation ΔΘ. - point-to-point misorientation. θ. - theta phase (Al2Cu). dfine. - size of fine grains. d. - crystallite size. FHAB. - fraction of high angle boundaries. σ. - true stress. φ. - diameter. t. - thickness.

(9) TABLE OF CONTENTS. CHAPTER 1. SEVERE PLASTIC DEFROMATION AND. Page. 1. GRAIN REFINEMENT 1.1 PHYSICAL PROPERTIES OF ULTRA FINE GRAINED. 2. MATERIALS 1.2 METHODS OF LARGE STRAIN DEFORMATION 1.2.1 Equal channel angular pressing. 6 6. 1.2.1.1 Fundamental parameters in ECAP. 7. 1.2.1.2 Experimental parameters influencing ECAP. 9. 1.2.2 Multidirectional forging. 10. 1.2.3 High pressure torsion. 12. 1.2.4 Accumulative roll bonding. 14. 1.3 MICROSTRUCTURE EVOLUTION AND MECHANISMS OF. 16. GRAIN REFINEMENT 1.3.1 Mechanisms of fine grain formation. 16. 1.4 SUMMARIES AND UNRESOLVED PROBLEMS. 19. 1.5 MAIN AIMS OF THE PRESENT STUDY. 21. 1.6 REFERENCES. 24. CHAPTER 2. EXPERIMENTAL TECHNIQUE 2.1 INVESTIGATED MATERIALS. 29 30. 2.1.1 Composition. 30. 2.1.2 Specimen Preparation and Heat Treatment. 30. 2.1.3 Initial microstructures. 32. 2.2 EQUAL CHANNEL ANGULAR PRESSING. 35. 2.3 MICROSTRUCTURAL ANALYSIS. 39. 2.3.1 Optical microscopy. 39. 2.3.2 Transmission Electron Microscopy. 39. 2.3.3 Scanning Electron Microscopy/ Electron Back Scattering Diffraction. 40. 2.3.4 Quantitative Analysis of Microstructures. 41.

(10) 2.4 REFERENCES. CHAPTER 3. MICROSTRUCTURE EVOLUTION IN Al-3%Cu ALLOY. 43. 44. DURING HOT ECAP 3.1 MICROSTRUCTURAL CHANGES IN Al-3% Cu AT 523 K. 45. 3.1.1 Optical Metallography. 45. 3.1.2 Orientation Imaging Microscopy. 45. 3.1.3 Misorientation characteristics of microstructure. 49. 3.1.4 TEM microstructures developed at 523 K. 53. 3.2 DISCUSSION 3.2.1 Strain-induced grain refinement process 3.3 TEMPERATURE EFFECT ON MICROSTRUCTURE. 55 55 58. DEVELOPMENT 3.3.1 Microstructures developed at high strain. 58. 3.3.2 Quantitative analysis of deformed microstructures. 61. 3.4 CONCLUSIONS. 63. 3.5 REFERECES. 66. CHAPTER 4. GRAIN REFIENEMENT IN ALUMINUM ALLOY 2219. 68. DURING EQUAL CHANNEL ANGULAR PRESSING AT 523 K 4.1. MICROSTRUCTURE EVOLUTION IN AA 2219 DURING ECAP. 69. AT 523 K 4.1.1 Deformed microstructures. 69. 4.1.2 OIM microstructures. 69. 4.1.3 TEM microstructures. 76. 4.1.4 Misorientation distribution. 88. 4.2 DISCUSSION. 85. 4.2.1. New grain formation process during ECAP. 85. 4.2.2. Mechanism of strain-induced grain refinement. 87. 4.3 CONCLUSIONS. 89. 4.4 REFERENCES. 91.

(11) CHAPTER 5. EFFECT OF DEFORMATION TEMPERATURE ON. 93. MICROSTRUCTURE DEVELOPMENT IN AA 2219 DURING HOT ECAP 5.1 MICROSTRUCTURAL EVOLUTION AT VARIOUS. 94. TEMPERATURES 5.1.1 Optical microstructures in temperature range 523 – 748 K. 94. 5.1.2 OIM microstructures at 573 K. 94. 5.1.3 OIM microstructures at 673 K. 97. 5.1.4 OIM microstructures at 748 K. 101. 5.1.5 OIM microstructures at ε = 12 from 573 K to 748 K. 104. 5.1.6 Quantitative analysis of strain-induced grained structures. 104. 5.2. DISCUSSION. 114. 5.2.1 Strain-induced grain formation process. 114. 5.2.2. Temperature effect on new grain formation. 115. 5.3. CONCLUSIONS. 120. 5.4. REFERENCES. 122. CHAPTER 6. SUMMARY. 124. ACKNOWLEDGEMENTS. 130. LIST OF PUBLICATIONS. 131.

(12) LIST OF FIGURES. Page. Fig.1-1. A comparison of yield strength and ductility for an AA 3004 processed by cold rolling and ECAP (5).. 3. Fig. 1-2. Relationship between Vickers hardness and crystallite grain size in iron (4) .. 3. Fig.1-3. Samples of the 1570 Al alloy after ECAP and pulling to failure at different temperatures and (b) the variation of strain rate sensitivity coefficient, m(6).. 5. Fig. 1-4. Schematic illustration for equal channel angular pressing (ECAP). The microstructure developments in (a) and (b) depending on parameters of die: φ is channel angle, ψ is arc of curvature (13), (c) the variation of routes in ECAP (16).. 8. Fig. 1-5. Schematic illustration for multi directional forging (MDF). Changes in the dimension sample under one cycle MDF with pass strain of 0.4(27).. 11. Fig. 1-6. Schematic illustration for high pressure torsion (HPT) (37).. 13. Fig. 1-7. Schematic illustration for accumulative roll bonding (ARB) (42).. 13. Fig. 2-1. Schematically drowing of solution heat treatment for Al alloys tested, i.e. Al-3%Cu alloy and AA 2219; (b) A part of the equilibrium diagram for Al-3%Cu alloys(1).. 31. Fig. 2-2. Initial microstructure of Al-3%Cu alloy after homogenization at 793 K for 4 hours revealed by (a) optical metallography, (b) TEM.. 33. Fig. 2-3. Initial microstructure of 2219 Al alloy after homogenization: (a) optical micrograph; (b) TEM microstructure.. 34. Fig. 2-4. Schematic illustration of ECAP facility used at high temperature. 36. Fig. 2-5. Rod dimension (in mm) for ECAP (a) and position of plane for microstructural investigation after ith pass of ECAP (b). PD is pressing direction.. 37. Fig. 2-6. Schematic representation for cyclic ECAP at temperature of Tx: (a) heating in a separate furnace, (b) exposure of sample in a die at temperature Tx before ECAP for ~45 min and ECAP followed by rapid water quenching (WQ).. 38. Fig. 3-1. Optical micrographs of Al-3%Cu deformed by ECAP at 523 K up to strains: (a) ε =1 (b) ε = 2, (c) ε = 4, (d) ε = 8 (e) ε = 12.. 46. Fig. 3-2. Typical OIM micropgraphs of Al-3%Cu alloy deformed at 523 K to. 47.

(13) various strains: (a) ε = 3, (b) ε = 8, (c) ε = 12. Fig. 3-3. Point-to-point and cumulative misorientations of strain-induced boundaries developed along lines marked in Fig. 3-2. (a) and (b) ε = 3, (c) ε = 8, (d) ε = 12. 48. Fig. 3-4. Changes in crystallite size with strain for Al-3%Cu. Transverse and longitudinal sizes were measured in perpendicular and parallel directions of developed microstructures. 50. Fig. 3-5. Changes in the misorientation distribution of deformation-induced boundaries developed in Al-3%Cu alloy deformed by ECAP at 523 K.. 51. Fig. 3-6. Changes (a) in the average misorientation angle Θav and (b) the fraction of strain-induced HABs, FHAB as well as new grains VUFG with repeated ECAP at 523 K. 52. Fig. 3-7. Transmission electron micrographs of Al-3%Cu deformed by ECAP at 523 K up to strains: (a) ε = 2 (b) ε = 8, (c) ε = 12.. 54. Fig. 3-8. Optical micrographs of Al-3%Cu deformed by ECAP to strain of 12 at (a) 573 K, (b) 673 K, (c) 748 K.. 57. Fig. 3-9. Typical optical micrographs of Al-3%Cu alloy deformed up to strain 12 at various temperatures: (a) 573 K, (b) 673 K, (c) 728 K.. 59. Fig. 3-10. Point-to-point misorientations of strain-induced boundaries developed in Al-3%Cu alloy at ε =12 along lines marked in Fig. 3-9. (a) 573 K, (b) 673 K, (c) 748 K.. 60. Fig. 3-11. Misorientation distributions for deformation-induced boundaries developed in Al-3%Cu at ε =12 after ECAP at various temperatures.. 62. Fig. 3-12. Changes in the characteristics dfine and Θav of deformation-induced grain structure developed in Al-3%Cu at various temperatures after ECAP at ε =12.. 64. Fig. 3-13. Temperature dependence of the volume fraction of new fine grains VUFG developed at ε =12 in Al-3%Cu. The data for AA 2219 deformed at similar conditions are shown for reference (10).. 65. Fig. 4-1. Optical micrographs of 2219 Al alloy deformed by ECAP at 523 K up to strains: (a) ε = 0; (b) ε = 2; (c) ε = 4; (d) ε = 12. PD is pressing direction.. 70. Fig. 4-2. OIM micrographs and corresponding misorientation profiles for 2219 Al alloy deformed by ECAP at 523 K. Thin white lines correspond to the boundaries with misorientation Θ >2o, thin black lines Θ >5o and bold lines Θ ≥15o, respectively. (a) ε = 1; (b) ε = 2 .. 71. Fig. 4-3. OIM micrographs and corresponding misorientation profiles for 2219 Al. 72.

(14) alloy deformed by ECAP at 523 K and ε = 3. Fig. 4-4. OIM micrograph, the misorientation changes along line T3 (b), and the inverse pole figure (c) of 2219 Al alloy deformed by ECAP to ε =3 at 523 K.. 73. Fig. 4-5. OIM micrograph of the 2219 Al alloy deformed by ECAP to ε = 6 at 523 K.. 74. Fig. 4-6. The misorientation changes were measured along lines T4 (a) and T5 (b) and the inverse pole figure (c) derived from OIM map in Fig.4-5.. 75. Fig. 4-7. TEM microstructures with diffraction patterns of 2219 Al alloy deformed by ECAP at 523 K at: (a), (a’) ε = 2 and (b) ε = 4. Numbers indicate misorientation angle of boundaries.. 77. Fig. 4-8. TEM microstructures with misorientation maps of 2219 Al alloy deformed by ECAP at 523 K. The numbers indicate misorientation angle in degrees. (a) ε = 8 and (b) ε = 12.. 79. Fig. 4-9. Changes of the average crystallite size of new (sub)grains and spacing of deformation bands with repeated ECAP for 2219 Al alloy.. 80. Fig. 4-10. Changes in the misorientation distribution of deformation-induced boundaries with repeated ECAP for 2219 Al alloy at 523 K. The broken line indicates the random misorientation distribution for annealed materials.. 83. Fig. 4-11. Changes in (a) average misorientation angle Θav, and (b) the fraction of strain-induced HABs F as well as new grains VUFG with repeated ECAP at 523 K for AA 2219.. 84. Fig. 4-12. Schematic drawing of the relationship between the average misorientation angle Θav of strain-induced low and high angle boundaries (LABs and HABs) and strain during severe plastic deformation in Al alloys at 523 K.. 88. Fig. 5-1. Optical micrographs of 2219 Al alloy deformed by ECAP to a strain of 3 and 12 at various temperatures: (a)-(a’) 523 K; (b)-(b’) 573 K; (c)-(c’) 673 K; (d)-(d’) 748 K. PD is pressing direction.. 95. Fig. 5-2. Typical OIM micrographs of 2219 Al alloy deformed by ECAP at 573 K. Thin white lines correspond to the boundaries with misorientation Θ >2o, thin black lines Θ >5o and bold lines Θ ≥15o, respectively. (a) ε =3; (b) ε = 6.. 96. Fig. 5-3. Typical OIM micrographs of 2219 Al alloy deformed by ECAP at 673 K. Thin white lines correspond to the boundaries with misorientation Θ >2o, thin black lines Θ >5o and bold lines Θ ≥15o, respectively. (a) ε =2; (b) ε = 4, (c) ε = 8.. 98.

(15) Fig. 5-4. Misorientation distribution of strain-induced boundaries in 2219 Al alloy developed along the lines (a) T1, (b) T2 , (c) T3 indicated in Fig. 53 (a), (c) and (d), respectively.. 100. Fig. 5-5. Typical OIM micrographs of 2219 Al alloy deformed by ECAP at 748 K. Thin white lines correspond to the boundaries with misorientation Θ >2o, thin black lines Θ >5o and bold lines Θ ≥15o, respectively. (a) ε =3; (b) ε = 6.. 102. Fig. 5-6. Strain dependence of the average crystallite size and for AA 2219 deformed by ECAP at various temperatures. The transverse grain size in (b) and (c) was measured perpendicular to the deformed microstructure direction.. 103. Fig. 5-7. OIM micrographs of 2219 Al alloy deformed by ECAP to a strain of 12 at various temperatures: (a) 523 K; (b) 573 K; (c) 673 K; (d) 748 K. PD is pressing direction.. 105. Fig. 5-8. Changes in the misorientation distribution of strain-induced boundaries developed in 2219 Al alloy deformed by ECAP to ε = 3 and ε = 12 at various temperatures. FHAB and solid arrows indicate fraction of HABs and the average misorientation angle measured in the whole volume Broken arrows show the average misorientation angle measured in fine grained regions.. 110. Fig. 5-9. Changes in average misorientation angle of strain-induced boundaries developed in 2219 alloy with straining by ECAP at various temperatures (a) in the whole volume and (b) in regions consisting of new fine grains only.. 111. Fig. 5-10. Temperature dependence of (a) average misorientation angle of straininduced boundaries and (b) new grain and subgrain size developed in 2219 alloy at ε = 12.. 113. Fig. 5-11. Dependence of volume fraction of ultrafine grained structure developed in 2219 alloy during ECAP with (a) repeated strain, (b) temperature.. 116. Fig. 5-12. The OIM microstructures of 2219 Al alloy deformed by ECAP at 673 K to ε = 8 (a) after deformation followed by quenching in a water (b) after static annealing in a furnace at 673 K during 50 min and cooling in a water.. 118. Fig. 5-13. Relationship between volume fractions of new fine grains VUFG and high angle boundaries FHAB developed in 2219 alloy during ECAP.. 119.

(16) CHAPTER 1 SEVERE PLASTIC DEFROMATION AND GRAIN REFINEMENT. This chapter gives an overview of the methods of large strain deformation, commonly called as severe plastic deformation (SPD), which is considered an attractive technique for producing fine grained structure in numerous metallic materials. Basic principles of widely investigated SPD methods and especially characteristics of equal channel angular pressing (ECAP) procedure used in this study are briefly described. The proposed mechanisms of ultra fine grain (UFG) structure formation and some factors promoting microstructural changes are discussed in detail. Finally, the main aim of the present study will be described based on the reviews of the above mentioned.. 1.

(17) 1.1 PROPERTIES OF ULTRA FINE GRAINED MATERIALS. Many metal-forming methods, such as rolling, extrusion, forging or drawing, introduce large plastic strains and as a result very fine grained microstructures can be formed in metals and alloys(1-3). During such processing, one or more dimensions of the work pieces are continuously reduced and eventually, foil or wire is produced which has limited use for industrial application. On the other hand, there is a number of large strain processing methods in which a sample can be deformed without any change in its dimensions and so there is no limit to the strain that can be achieved. Generally, the methods of ultra high strain deformation for ultra fine grain structure formation in bulky materials are usually called severe plastic deformation (SPD)(1-10). There are many SPD techniques available for grain refinement at present time: equal channel angular pressing (ECAP)(3,11-26), multi-directional forging (MDF)(27-34), high pressure torsion (HPT)(35-41), accumulative roll bonding (ARB)(42-45) and will be considered in present work. Based on the Hall-Petch relationship, the strength of a material increases with the inverse square root of the grain size, produced through SPD materials having ultrafine grain structures (UFG) demonstrate generally higher strength at ambient temperature (Figs. 1-1 and 1-2) and excellent superplastic behavior at elevated temperatures (Fig. 1-3). It is well known that superplastic forming is a highly efficient method of processing of shape products and forming of hard deformable materials(1,3,6,23,27,39). In connection to this, the superplasticity of UFG materials at relatively low temperatures and also high strain rates is a very attractive property leading to increasing production by forming (3,6). It has been recently demonstrated for various metals and alloys that grain refinement may lead to a very high hardness in exchange for invariably low ductility under tensile testing (1,3,5). A similar tendency is well known for metals subjected to heavy straining by other processes such as rolling (1,20), extrusion or drawing. The strength and ductility are key mechanical properties of any structural material but they usually have opposing characteristics (3,5,22,23,27,35-37,43). It is illustrated in Fig. 1-1 that the yield strength increased monotonically with the increasing equivalent strain during ECAP or rolling. The ductility decreases rapidly after one pass ECAP and then remains roughly constant with further ECAP. By contrast, ductility after cold-rolling continues to. 2.

(18) Fig. 1-1. A comparison of yield strength and ductility for an AA 3004 processed by cold rolling and ECAP (5).. Fig. 1-2. Relationship between Vickers hardness and crystallite grain size in iron. 3. (4)..

(19) decrease with increasing rolling strain of ε > 1 . Accordingly, the processing by ECAP may lead to a greater retention of ductility than that of conventional cold-rolling. In this connection, recent finding of very high strength and good ductility in several bulk UFG materials produced by SPD are of special research and commercial interests. It has been established experimentally for a number of materials that ultrafine structures usually exhibit higher fatigue resistance and improvement of the fatigue limit than materials with a conventional grain size under stress-controlled loading(7). For example, in the case of Cu, which is probably the most closely-studied material, an improvement in the fatigue life time was observed in UFG copper when compared to Cu with a conventional grain size (8). The processing by SPD can have strong effects not only on mechanical properties (3,9) but also on several functional properties of materials. Magnetic measurements demonstrate that there is advancement in magnetic properties with respect to coercitivity and specific magnetization that results in significant change of the hysteretic properties and the modification of the corrosion behavior of some ferromagnetic materials after SPD(3). The reduction in a grain size to nanometer scale results in greater applicability in the microelectronic industry, where the component sizes are required to be small, and this further makes some techniques of SPD, particularly HPT(35-40), very attractive for fabrication of nanocomposite materials (4,5,40). It should be noted that methods of SPD allow producing materials not only with grain sizes ranging from submicro- to nanometer but also with fairly homogeneous and equiaxed microstructures having high fraction of high angle boundaries (HABs). The latter important characteristic can be used in the grain boundary design or in the concept of grain boundary engineering, where it has been proposed that the properties of materials may be effectively changed by intention and careful modeling of the distributions of boundary angles(9,10). The tendency for the average boundary misorientation to increase with increasing strain during SPD, provides a unique opportunity to make use of SPD methods for the production of materials having different, but controlled, boundary misorientation distributions. This approach has been used successfully for several studies including, for example, improving the sensitivity to intercrystalline stress and corrosion cracking(10).. 4.

(20) (a). (b). Fig. 1-3. (a) Samples of the 1570 Al alloy after ECAP and pulling to failure at different temperatures and (b) the variation of strain rate sensitivity coefficient, m(6).. 5.

(21) 1.2 METHODS OF SEVERE PLASTIC DEFORMATION. The processing by SPD may be defined as a metal forming procedure by which a very large strain is imposed on a bulky material leading to significant grain refinement without any significant changes in the overall dimensions (1-3). The most important principles of application of SPD is to avoid size reducing and any mechanical damages during deformation and to achieve significant grain refinement effect in materials, resulting into extremely good mechanical properties(4-10,20-23,27,39,43). Moreover, the application of SPD to bulky alloys is one of requirements to engineering technologies with consequent economic and environmental benefits. The SPD has a great potential to be widely used in industrial process. The most well known techniques for providing large plastic deformation and formation of UFGed structure are equal channel angular extrusion or pressing (ECAE or ECAP)(11-26), multidirectional forging (MDF)(27-34), high pressure torsion (HPT)(35-41) and accumulative roll bonding (ARB)(42-45). Let us consider these methods in detail comparing them to each other.. 1.2.1 Equal channel angular pressing In the Equal Channel Angular Pressing (ECAP), samples are deformed repeatedly without changing of the cross-section(3,11-26). In the case of a hard-to-deform material, ECAP can be conducted at elevated temperatures(15,17,22,24,26). Thus, ECAP is advantageous to production of UFG structures in bulky materials from a position that it can be applied in a wide temperature range, including room temperature. This is very important feature in contrast to the other methods, such as HPT and ARB. In this connection, ECAP is seen as the most suitable and preferred method for producing UFG materials among other SPD techniques. As a result, the present work will be focused on ECAP at elevated temperatures and so role of ECAP parameters on workability and UFG structure formation at high temperatures will be considered in detail. When a material is pressed by ECAP, the material behavior and its microstructure formation can be influenced by several factors, which can be categorized into 2 groups (3,15,24): (i). ECAP facility configuration including die geometry (Fig. 1-4 (a) and (b)), i.e. the inner angle inside the die between the two parts of the channel φ and the outer arc 6.

(22) of curvature where the channels intersect ψ, (ii). Experimental factors, which are controlled during deformation and associate with the processing regimes, i.e. speed of pressing, temperature, absence or presence of back pressure, processing routes (Fig. 1-4I), and total number of passes.. 1.2.1.1 Fundamental parameters in ECAP The principles of ECAP is that the rod or bar-type samples are multiple number of times extruded through a special L-shaped configuration die consisting of the two channels of equal cross-section which intersect generally at an angle, φ, of 90o or higher ( Fig. 1-4(a)and (b)). Intense plastic deformation is introduced by a simple shear in a thin layer at the crossing plane of channel(12-16,25). Assuming there is a sharp die corner and no friction (Fig. 1-4(a)), i.e. the outer angle ψ =0o, the value of total strain. εN after N passes can be expressed through the strain. increment, Δε, at each pass by equation (1.6)(12,13,16,25):. ε N = Δε ⋅ N =. 2N φ cot 2 3. (1.1). In case when ψ ≠ 0o (Fig. 1-4), a relationship allowing one to calculate the strain after N passes of ECAP, εN, can be given in the following form:. εN =. N ⎡ ⎛φ Ψ ⎞ ⎛ φ ψ ⎞⎤ ⎢2 cot⎜ 2 + 2 ⎟ + ψ cos ec⎜ 2 + 2 ⎟⎥ 3⎣ ⎝ ⎠ ⎝ ⎠⎦. (1.2). From this relationship, it follows that at φ =90o and ψ =20o, which are the frequently used angles, each pass strain approximately equals to 1(12,13,16,25). The angle at the arc of curvature has a minor effect on the equivalent strain except only for channel angles less than 90o. High strains may be achieved at a single pass through the die with low values of φ and Ψ. The processing parameters exert a strong influence on the development of UFG structure(3,1216,18-25). . At present time it has been investigated the effect of ECAP on the structure formation and. influence of channel angle on the development of ultra fine grains(13,17,25). Thus, it was revealed. 7.

(23) (a). (b). (c). Fig. 1-4.. Schematic illustration for equal channel angular pressing (ECAP). The. microstructure developments in (a) and (b) depending on parameters of die: φ is channel angle, ψ is arc of curvature (13), (c) the variation of routes in ECAP (16). 8.

(24) that, from a practical point of view, it is reasonable to construct an optimum die with a channel angle of 90o (see Fig. 1-4(a)) which results in development of a fine-grained structure with large proportion of high angle boundaries (HABs)(3,13). An increase of the internal angle φ up to 157.5o leads to the formation of microstructure with high fraction of low angle boundaries (LABs) (Fig. 1-4(b)), by contrast, HABs are introduced into sample when the value of φ is in the range of 90115o at the same strain(13).. 1.2.1.2 Experimental parameters influencing ECAP Such conditions of ECAP as pressing speed, temperature of deformation, route of pressing and number of passes are usually controlled during pressing and play an important role in microstructure refinement (3). ECAP is usually conducted using hydraulic press that can operate with the pressing speeds over a very wide range. In general the pressing speed from 10-2 to 10 mm/s has no significant influence on UFG structure formation, while more equilibrium structure can form when pressing at the slower speed takes place(3). Thus, microstructure differences with increasing pressing speed tend to be minor. In contrast, with increasing pressing temperature at which deformation is conducted, number of passes and route of ECAP. The pressing temperature is a key factor in any use of ECAP because it can be controlled relatively easily. Although it is generally experimentally easier to press material at high temperatures, optimum ultrafine-grained microstructures will be attained when the pressing is performed at the lowest possible temperature where the pressing operation can be reasonably conducted without the introduction of any significant cracking in the billets. An increase in the temperature of deformation can lead to an increase in the grain size in material and grains became more equiaxed(3,17,20,22,26). It is found an increase in the fraction of low angle boundaries with rising temperature, due to activation of recovery with higher rate(17). Process of transformation of low into high angles takes place hardly, however depends on nature of material. The influence of temperature factor on microstructure formation is scarcely investigated at present time.. 9.

(25) In order to achieve very significant and uniform grain refinement in a wide range of materials, the detailed characteristics of the material prior ECAP and other various processing factors, including the processing route (whether A, Ba,c or C), the total number of passes imposed on the sample and temperature have been examined in numerous works(3,11-26). Analysis of the results shows that repeated ECAP provides an opportunity to develop different microstructures by rotating the sample between subsequent passes so that there are changes in the slip systems activated on each consecutive pressing. (11-13,15,16,25). . Between four possible ways to deform a. material with an intent to get homogeneous UFG microstructure, routes Bc and A and resulting microstructures were under intensive examinations Fig. 1-4 I. Route A is a simplest method for industrial conditions to deform a sample without any rotation of the billet at each pass, gives larger fraction of high angle boundaries (HABs) in fine grained microstructure, which, however, can be inhomogeneously formed. Route Bc includes a 90o rotation of the billet around its longitudinal axis with a constant sense of rotation after each pass. It was suggested that route Bc can be a preferable procedure for most expeditiously achieving a homogeneous fine grained microstructure (11-16). The total number of passes also influence the microstructure refinement. (16,18-23). . With. increasing strain a higher fraction of HABs is introduced into material and fully homogeneous fine grained microstructure is formed. It depends, however, on nature of materials tested, its initial grain size, texture, phase composition, etc.(3,25). Moreover, the process of such fine grained microstructure evolution during ECAP is not clear and is main subject for study at present time and will be considered in the following sections.. 1.2.2. Multidirectional forging. This method is one of the easiest SPD method without any specific device and has a great potentiality for producing of relatively large workpiece that can be used in a mass production industry (27-34). The principle of MDF is compression process which is performed change in the direction of the applied strain (i.e. x—y—z—x…) at each step of compression. A schematic illustration of the changes in the sample dimension during such multiple compressions is shown. 10.

(26) Fig. 1-5. Schematic illustration for multi directional forging (MDF). Changes in the dimension sample under one cycle MDF with pass strain of 0.4 (27).. 11.

(27) in Fig. 1-5. Here the loading direction between compressions is consequently changed in 90o in each pass. The dimension ratio of a rectangular sample is fixed during repeated compression, when a pass strain, Δε, is constant as shown in Fig. 1-5. The total strain value after repeated number N of compression is given as (27,34),. ε =NΔε. (1.3). Since the sample tested does not change its shape under MDF conditions, large plastic strain can be introduced into material during repeated compression at low to elevated temperatures. Moreover, the resistance of material to deformation, i.e. flow stress can be measured during MDF (27-34). . In contrast it is rather difficult for the other methods of SPD mentioned above, such as. ECAP, HPT and ARB, may not give any information of flow stress developed during deformation. As can be seen, the grain refinement during MDF can be controlled not only by a total accumulated strain and strain per each pass, but also by strain rate and temperature(27,28,33,34). When the mechanisms of plastic deformation and microstructural development are under discussion, it is necessary to evaluate and analyze the interrelationship between them developed in high strain range deformation(30). From this point, MDF has a superior advantage not only for new grain formation itself, but also for discussion of the mechanisms operating during SPD. The method of MDF has been used for microstructure refinement in a number pure materials and alloys, e.g., pure Ti and its alloys(27), stainless steel(28), pure Cu(29), Al and its alloys(30-31,.33), Mg alloy(32), high strength Ni based alloys(34), etc. The formation of fine grained microstructures can be obtained in rather brittle materials by using this method, when the process starts at elevated temperatures and then deformation temperature decreases (32), which make MDF is much more attractive method of SPD.. 1.2.3. High pressure torsion. Among the various methods of SPD, torsion under high pressure (HPT) produces much finer grain size than the other processes, although the sample size is small compared with the others. It uses a thin disc sample with a thickness less than 1 mm. The microstructure observation shows that there is significant difference in grain refinement between the center and the outer regions(35-. 12.

(28) 1. Upper anvil 2. Low anvil 3. Sample t - thickness φ – diameter Fig. 1-6. Schematic illustration for high pressure torsion (HPT). (37).. 1. Cutting 2. Degreasing and wire brushing 3. Stacking 4. Heating and roll bonding Fig. 1-7. Schematic illustration for accumulative roll bonding (ARB) (42).. 13.

(29) 41). . However, the further development in design of HPT tools takes place at present time that. makes possible the formation of homogeneous nanostructures with high angle boundaries under HPT(35-37). In processing, by HPT shown in Fig. 1-6, the sample is in the form of thin disc and placed in a HPT facility between upper and lower anvil, subjected to a very high pressure and then strained in torsion where the straining is usually achieved by rotating lower anvil. A method of HPT can be used for fabrication of disc type samples with a diameter between 10 and 20 mm and a thickness between 0.1-1 mm(35). An experiment can be conducted at ambient temperature nder applied pressure P of about 5-8 Gpa. When the lower holder rotates, surface friction forces deform the sample by shear. A true logarithmic strain ε is calculated by equation (1.2)(37): ⎛α ⋅r ⎞ ⎟ ⎝t 3⎠. ε = ln⎜. (1.4). Where α is rotation angle in radian, r and t are the radius and thickness of the disc, respectively. HPT can deform metallic materials to large strains above 10 to 20 and more at room temperature. Moreover, recently HPT can be used for consolidation of powders and can provide a rather high density in a disc type nanostructured samples processed due to high pressure(36,41). Thus, HPT have a great potential for industrial application to fabricate nanocomposite materials, compared to conventional method of fabrication, HPT have significant advantages: high performance, low cost, although small sizes of pressed samples.. 1.2.4. Accumulative roll bonding. Another promising technique for potential industrial application is the accumulative roll bonding (ARB) (42-45), a process able to impose severe plastic deformation of large sample, for example bulk materials. This process is schematically represented in Fig. 1-7. In ARB a strip is placed on the top of another strip, stacked together and rolled. The resulting length of the material is then sectioned into two halves that are again stacked together and rolled. The whole process can be repeated again and again to accumulate a strain. In ARB, rolling is not only a deformation process but also a bonding process carried out to obtain one-body solid final material. There. 14.

(30) exists a minimum limit of reduction of thickness, i.e. the threshold deformation to attain a sufficient bonding. It is well known that the threshold deformation decreases with temperature(1,3,10). Before rolling often the strips are reheated to facilitate bonding and to reduce the rolling force. In such case the heating temperature must be below the temperature for occurrence of recrystallizitaion, which can cancels out the accumulated strains. In order to achieve a good bonding, the surfaces of the strips are cleaned both mechanically and chemically, for example degreasing and wire brushing the surfaces to remove the oxide scale. When the 2 strips are overlapped and roll bonded by reduction 50%, the thickness t of the strip after N cycles is t=. to 2N. (1.5). Where to is the initial thickness of strips. Then the total reduction rt after N cycle is. rt = 1 −. 1 t = 1− N to 2. (1.6). An equivalent plastic strain ε can be expressed by von Misses yield criterion and plane strain condition(13): ⎡ 2 ⎛ 1 ⎞⎤ ln⎜ ⎟⎥ N=0.8N ⎣ 3 ⎝ 2 ⎠⎦. ε =⎢. (1.7). Thus, the total equivalent strain ε after N cycles of the 50% roll bonding ARB is ~0.8. The experiments show that ARB gives a possibility to obtain microstructures with improved combinations of strength, ductility/toughness through adequate grain refinement(43).The method of ARB was used for development of fine-grained microstructures in a number of materials, including of aluminum and its alloys (44), steels (45,46) and etc.. 15.

(31) 1.3 MICROSTRUCTURE EVOLUTION AND MECHANISMS OF GRAIN REFINEMENT. It is well established that methods of SPD have been used for the formation of new fine grain structures in bulky metals and alloys. However, most of works are carried out to study the microstructural developments at high strains and the effect of such microstructures on the mechanical properties of the products (1-10,20,22-24,27-29,35-37,39,43-45). Only a limited number of works were concentrated on the process of microstructure evolution(3) and the mechanisms of fine grain formation during described above methods of HPT(37,40), ARB(44), ECAP(16,21) and MDF(29-31). It should be noted, that there are few papers discussed in detail the mechanisms of new grain formation under different ECAP conditions, which will be summarized below.. 1.3.1 Mechanisms of fine grain formation. The evolution of new grain structures occurring under plastic working is frequently discussed as dynamic recrystallization (DRX)(3,47-63). A considerable refinement of the microstructure due to operation of DRX can be obtained under cold to warm deformation, i.e. T ≤ 0.5 Tm where Tm is the meeting point. The critical strain for the initiation of DRX can be incredible large at relatively low temperature and so makes the evolution of ultra-fine grains during cold to warm deformation difficult. There have been, many studies discussed the structural mechanisms for the evolution of new fine grains even under cold and warm deformation conditions(11-14,18,22,29,37,40,44). Many of them is considered that in materials with low-to-medium and high SFE, the microstructure evolution results to the operation of a kind of continious dynamic recrystallization (cDRX), containing the fragmentation of initial grain by development of geometrical necessary dislocation boundaries (GNBs)(47,49,51) and/or several deformation bands(50,52,53,) These boundary misorientations increase with deformation to high strains, resulting in the development of new grains with medium to high angle boundaries. The effect of temperature on the dynamic restoration process operating and the phenomenological terminologies, such as discontinuous dynamic recrystallization (dDRX) and continuous dynamic. 16.

(32) recrystallization (cDRX) in high and low-to-medium SFE materials are compared in the works (47-55, 57-59,61). .. The main characteristics of DRX during warm and hot deformation, i.e. T ≥ 0.5 Tm, have been clarified in the following works(47,48,54-56,58,62,63). Two types of DRX based on operating structural mechanisms that result in new grain development are commonly discussed in literature(54,58). The new grain evolution that takes place in materials with low-to-medium stacking fault energy (SFE), e.g. Cu, Ni, gamma-Fe, etc., is associated with discontinuous DRX (dDRX); i.e. the formation of a new grain structure that results from nucleation due to grain boundary bulging occurring on grain boundary serration and subsequent long-distance migration of high angle boundaries consuming the strain hardened substructures. (47). . The other type is the. continuous DRX (cDRX), which is discussed as the mechanism that is responsible for the formation of new grains in materials with high SFE, e.g. Al, Mg, alpha-Fe and their alloys. The cDRX consists of the formation of arrays of low angle boundaries in grain interiors and a gradual increase in the boundary misorientations during deformation, finally leading to new grain development in high strain (48,54-56). The formation of new grains at elevated temperatures and at high strains can be resulted from another mechanism of DRX, which commonly named geometrical DRX (gDRX). (54-56,58). .. Original grains are flattened and pancaked by large strain deformation and their boundaries become progressively serrated at high temperatures. When the original grain thickness is reduced to about two subgrain sizes, the grain boundaries begin locally to come into contact with each other, causing the grains to pinch-off(48,58). At the same time, it was clearly shown in works(61-63) that some low-to-medium angle subboundaries, sometimes called as deformation / microshear bands or geometrically necessary boundaries. (51). , may be generated in the original grain interiors at relatively low strains due to. high strain heterogeneity and result in grain subdivision(62,63). They are progressively transformed into high-angle boundaries with further deformation, leading to a full development of new submicrocrystalline grain structures at large strains. At the same time, a limited number of studies (17,26,48,62,63). dealt with the influence of pressing temperatures on the mechanisms of fine grain. formation and microstructural changes during ECAP. There are several points of view about the. 17.

(33) mechanisms of grain formation at high temperatures. Increasing pressing temperature leads to the formation of arrays of low angle boundaries in aluminum and Al-3%Mg alloy (17). The inability to achieve high angle boundaries at higher pressing temperatures is attributed to the higher rate of dynamic recovery which lead to dislocation annihilation rather than the absorption of dislocations into the subgrain walls. It is considered that, the microstructural development in Al alloys under rising temperature of ECAP conditions can be mainly affected by an increased rate of dynamic recovery, which enhances the rearrangement of the dislocation substructures and assists the formation of new coarser and more equiaxed (sub)grains (3,17,26,62). In contrast, several works(55,58) concluded that formation of fine-grained structures occurs due to gDRX in Al alloys during ECAP at high temperature of deformation. On the other hand, it should be noted that deformation at very high temperatures can affect the grain refinement process. The as-deformed samples can be exposured for a long time in the equipment at elevated temperatures and thus, may be statically annealed. Sitdikov et al. observed bimodal grain structures developed in Al−6%Mg 0.3%Sc alloy during ECAP at 573 K. (62,63). .. They concluded that two structural components can be resulted from simultaneous operation of cDRX occurring during deformation, and static recrystallization during each ECAP pass(63). As a result, the mechanisms of grain refinement during warm-to-hot ECAP are currently a matter of some debate and not clear. Unfortunately, major factors controlling such microstructural development remain unknown due to the lack of the related experimental data, especially for Al alloys.. 18.

(34) 1.4 SUMMARIES AND UNRESOLVED PROBLEMS. Most promising approach in producing ultra fine grain (UFG) structure with grain sizes in the submicro- to nanometer range is a use techniques of severe plastic deformation (SPD), such as equal channel angular pressing (ECAP), multidirectional forging, torsion under high pressure, accumulative roll bonding, etc. are popular for laboratory investigations in present time. SPD processing carried out on many pure metals, including Al, Ti, Cu and their respective alloys and some intermetallic compounds. UFG materials exhibit improved mechanical and physical properties based on the Hall-Petch relationship, e.g. a higher strength, low temperature and high strain rate superplasticity, enhanced corrosion and magnetic behavior, and may be considered as advanced engineering materials for commercial applications. Because of many advantages of mechanical and physical properties of UFG materials are determined by their microstructure, grain refinement process taking place during SPD is very interesting from the scientific and academic point of view. Despite numerous experimental investigations conducted to date, very little information is currently available on the mechanism responsible for the microstructure refinement in high rate recovery materials. As it was described in the previous sections, there have been several approaches to explain the microstructure changes taking place during SPD. Most of investigations agreed, that in the range of low and elevated temperatures a mechanism of new grain formation is connected with formation of various kinds of banded structures and an increase in their boundary misorientations with further deformation results in new grain formation that is continuous dynamic recrystallization. An increasing the temperature can effectively increase the rate of dynamic recovery that can suppress the formation of new fine grains surrounded by HABs. In the pure Al and its alloys this effect can have considerable result, because of high stack fault energy. There are only few experimental data on grain formation in such materials during warm to hot ECAP, in which the temperature effect was clarified briefly. Thus, the mechanism which is respecting for ultra fine grain formation is main important phenomena and it is at present time scarcely investigated especially at high temperature. 19.

(35) deformation of coarse grained materials with high rate of recovery. Therefore, the problems which have an academic interest can be summarized as following: 1. To study the process of grain refinement during SPD in a wide temperature interval, especially at hot SPD. 2. To study the mechanism of grain refinement and effect of rising the deformation temperature on UFG formation.. 20.

(36) 1.5 MAIN AIMS OF THE PRESENT STUDY. The main goal of present work is to study and analyze the microstructural changes in Al-Cu basis materials during SPD by ECAP in order to establish the mechanisms of grain refinement with the formation of (sub)micron scale grains in a wide temperature interval. The aluminum alloys, Al-3%Cu and AA 2219 have initial coarse grained structures and characterized by high stacking fault energy and associated with the high recovery rate. For this purposes the alloys were subjected by ECAP via route A to a total strains of 12 in a temperature interval ranging from 523 to 748 K. A considerable full grain refinement can be obtained in both alloys during deformation. A dilute alloy Al-3%Cu contains many coarse precipitates of θ - phase according to equilibrium diagram. The microstructure in this alloy after ECAP in a wide temperature interval can be characterized by a low thermal stability. On the other hand, the commercial base Al alloy, AA 2219, was selected for use in this investigation because its chemical composition and combination of second phases provide rather high thermal stability of deformed microstructures that is favorable for investigation of mechanisms of new grains formation during deformation. At room and elevated temperatures AA 2219 contains precipitates of θ-phase (Al2Cu), that will be dissoluble with temperature increasing, precipitates of Al6Mn and fine dispersoids of Al3Cr and Zr phases, that is thermally-stable in the wide temperature interval. In the present work, on the both alloys, therefore, the role of dissoluble secondary θ-phase and role of fine particles in grain refinement is considered. Thus, a great attention was paid to the investigation of microstructure development during repeated deformation depending on alloys at investigated temperature interval of above 0.5 Tm. Particular emphasis was given to the effects of deformation temperature on the microstructure evolution at different stages of ECAP. The microstructures of these materials were characterized including analysis of their boundary misorientations obtained by EBSP technique. A specific attention was paid to reveal main factors promoting grain refinement in the aluminum alloys and to discuss the operated microstructural mechanisms in detail. The following aims in the present study were formulated to achieve the main goal:. 21.

(37) 1.. To investigate the microstructure changes during ECAP in a temperature range from 523 K to 748 K taking place in a single phase dilute Al-3%Cu alloy.. 2.. To examine the microstructure development in a temperature range from 523 K to 748 K taking place in the aluminum alloy, AA 2219.. 3.. To examine the temperature effect on fine grained structure formation during ECAP and to discuss the mechanisms of strain-induced grain refinement and temperature effect on occurrence of this process.. 22.

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(44) CHAPTER 2 EXPERIMENTAL TECHNIQUE AND CONDITIONS. In this chapter, the materials tested and a method of severe plastic deformation by ECAP are described in detail accompanied with the experimental conditions used. Main important techniques for microstructure investigations and analysis by using optical microscopy (OM), transmission electron microscopy (TEM) and electron back-scattering diffraction pattern (EBSD) are explained.. 29.

(45) 2.1 INVESTIGATED MATERIALS. 2.1.1. Composition. The two kind of aluminium alloys were used as materials tested in the present study. The one of them was binary aluminium alloy, Al-3%Cu, produced by direct chill casting. The other was commercial base aluminium alloy, denoted as AA 2219, manufactured by semicontinuous casting method at the Kaiser Aluminum Center for Technology. The both alloys have following chemical compositions presented in Table 2-1.. Table 2-1. Chemical composition of Al-Cu base alloys investigated (in mass pct). Al alloys. Cu. Mn. Cr. Zr. Fe. Ti, Si. Al-3%Cu. 3. -. -. -. -. -. AA 2219. 6.4. 0.3. 0.18. 0.19. 0.06. 0.01. 2.1.2. Specimen Preparation and Heat Treatment. The alloys in as-cast state were subjected to the following solid solution treatment, as shown in schematically for each alloy in Fig. 2-1(a). A homogenization with subsequent slow cooling to room temperature was first conducted in air in a furnace for both alloys. The initial stage of homogenization for duralumin alloys is to raise the temperature sufficiently for the copper go into solution, and to hold it long enough for all copper to be dissolved (Fig. 2-1(b)) (1,2). Thus, the Al3%Cu alloy was solution treated at 793K for 4 hours (see Fig. 2-1 (a)), followed by cooling with a rate of about 50 K/h leading to formation of relatively coarse particles of Al2Cu and Al-0.1% Cu matrix at room temperature (2).. 30.

(46) Fig. 2-1. (a) Schematically drawing of solution heat treatment for Al alloys tested, i.e. 3%Cu alloy and AA 2219. (b) a part of the equilibrium diagram for Al-Cu alloys (1) .. 31. Al -.

(47) The aluminium alloy 2219 was homogenized at a higher temperature of 803 K for Cu and other alloying elements, such as Mn, hardly soluble Cr, Zr to be dissolved in Al for 6 hours, followed by over-ageing at 683 K for 8 hours and then air cooling. This second stage accompanied with slow rate cooling was applied in order to obtain hardening in the alloy due to full precipitation of second phase particles and its minimization of concentration in a solid solution (2,3).. 2.1.3. Initial microstructures. The microstructures after homogenization of the both alloys are shown in Figs. 2-2 and 2-3. An initial microstructure of Al-3%Cu alloy was composed of roughly equiaxed coarse grains with an average size varied from 100 to 350 μm (Fig. 2-2 (a)). At room and elevated temperatures Al-3% Cu contains precipitates of one phase, which has the composition Al2Cu. The θ-phase is dissoluble with increasing temperature, as can be seen in Fig. 2-1 (b). The precipitates of Al2Cu (θ-phase) shows round platelet shape in Fig. 2-2(b) and the size ranging from 0.5 to 1μm. An increase in the temperature to 673 K leads to raising the limit of solubility and decreasing the size and density of θ-phase precipitates in the Al-3%Cu alloy according to the phase diagram in Fig. 2-1 (b). On the other hand, an initial structure of the 2219 Al alloy consisted of essentially equiaxed grains with an average grain size of ~140 μm (Fig. 2-3 (a)). The alloy AA 2219 contains several types of secondary phases at room temperature (4), i.e. the precipitates of θ-phase (Al2Cu), and dispersion particles of Al6Mn – phase having a size of about 200 nm and the dispersoids of Al3(Cr, Zr) phases. The latters were identified in AA 2219 by transmission electron microscopy analysis, as shown in Fig. 2-3 (b). Coarser precipitates of Al2Cu (θ) phase are uniformly distributed in the grain interiors and at the boundaries An equiaxed spherical Al3Cr and Al3Zr dispersoids have an average size of about 100 nm and 20 nm, respectively, and precipitates of intermetallic Al6Mn phase are elongated and has “pencil” like shape in the AA 2219 alloy (Fig. 2-3 (b)). 32.

(48) (a). 100 μm (b). Al2Cu. 500 nm. Fig. 2-2. Initial microstructure of Al-3%Cu alloy after homogenization at 793 K for 4 hours. (a) optical metallography and (b) TEM micrograph.. 33.

(49) (a). (b). Al6Mn. Fig. 2-3. Initial microstructure of 2219 Al alloy after homogenization: (a) optical micrograph; (b) TEM microstructure.. 34.

(50) 2.2 EQUAL CHANNEL ANGULAR PRESSING. The equipment used for ECAP at elevated temperatures is shown schematically in Fig. 2-4. The die for ECAP fabricated from a tool steel had L-shape configuration with an angle φ of 90o between the two circular in the cross section channels 3 and 4 and an angle ψ of 90o at the outer arc of curvature at the point of intersection(5). According to the geometry of the die, each passage leads to a strain of ~ 1. (5). . All pressings were conducted using hydraulic press 1 operating at a. speed about 1 mm/s. The die was placed inside an electric furnace 8 (heating jacket) and the pressing temperature was controlled by a thermocouple 9, which was inserted into the channel to a distance of about 10 mm from the shearing corner. The temperature was controlled within ±3K before and during each pressing. The samples for ECAP were machined parallel to the ingot axis into rods with a diameter of 20 mm and a length of around 100 mm (Fig. 2-5 (a)), as described in the section 1.2.1. The procedure of each ECAP technique is schematically represented in Fig. 2-6. ECAP was carried out on the sample which was heated to the testing temperature in a separate furnace and prior deformation was inserted into the vertical channel and then pressed by a plunger after the stabilization of ECAP temperature Tx. After pressing, specimen was removed from the horizontal channel 4 and subsequently inserted into the vertical channel 3, waited for temperature stabilization and pressed again. The specimen was quenched in the water (WQ) after each pressing. The period time for intermediate holding in a furnace (a) and temperature stabilization after insertion of rod in the die (b) was in general 40-45 min (about 2.7 ks). The specimen was held at deformation temperature Tx in the die during 1 pass of ECAP for approximately 15 minutes. The samples were pressed repeatedly at temperatures ranging from 523 to 748 K up to total strains of 12 by using rout A, i.e. without any rotation of specimen’s orientation in channels during deformation at each pass. The strains tested are summarized in Table 2-3. The route A was selected in the present work because the recent experiments on Al alloys have shown that it is the most efficient processing route for forming a fine grain structure with high fraction of high angle. 35.

(51) Fig. 2-4. Schematic illustration of ECAP facility used at high temperature.. 36.

(52) Fig. 2-5. Rod dimension (in mm) for ECAP (a) and position of plane for microstructural investigation after ith pass of ECAP (b). PD is the pressing direction.. 37.

(53) 38. die at temperature Tx before ECAP for ~45 min and ECAP followed by rapid water quenching (WQ).. Fig. 2-6. Schematic representation for cyclic ECAP at temperature of Tx: (a) heating in a separate furnace, (b) exposure of sample in a.

(54) boundaries (HABs) (6). Moreover, this is the simplest method of ECAP with a constant strain path and, hence, it can be most suitable for grain refinement under industrial conditions.. Table 2-3. Strains produced by ECAP for investigated materials Alloy. Temperature, K 523. 573. 623. 748. AA 2219. 1, 2, 3, 4, 6, 8, 12. 1, 2, 3, 4, 6, 8, 12. 1, 2, 3, 4, 6, 8, 12. 1, 2, 3, 4, 6, 8, 12. Al-3%Cu. 1, 2, 3, 4, 8, 12. 1, 2, 4, 8, 12. 1, 2, 4, 8, 12. 1, 2, 4, 8, 12. 2.3 MICROSTRUCTURAL ANALYSIS. 2.3.1 Optical microscopy. Following ECAP, samples for metallographic observations were cut from the central part of pressed rods parallel to PD, as shown in Fig. 2-5 (b). The surface of each sample was mechanically ground using 400, 800 and 1200 emery papers in sequence, then polished using 1 μm diamond pastes on a polishing cloth by using a polishing machine Struers RotoPol-22. Then all samples were etched by immersion in a standard Dicks-Keller etchant with following chemical composition: 1% HF + 1.5% HCl + 2.5% HNO3 + 95%H2O. The metallographic analysis was carried out by using an Olimpus PME3 optical microscope with a digital camera Olimpus D11. The deformed microstructures were analyzed in conventional and polarized lights.. 2.3.2 Transmission Electron Microscopy. The samples for TEM analysis with a thickness of about 0.5 mm were cut from the central places of deformed samples by using a cutting machine, and mechanically ground to about 250. 39.

(55) μm. Final polishing and thinning to perforation was performed by using a Tenupol – 3 twin jet polishing unit and cooled by nitrogen solution of 30% HNO3 + 70% CH3OH at 243 K and 20 V. After polishing the thin foils were look like discs with a diameter 3-4 mm and examined using a JEOL transmission electron microscope JEM-2000EX with a double tilt stage at an accelerating potential 160 kV. The misorientation of (sub)grain boundaries were studied using a conventional Kikuchi-line technique(7). The accuracy of the measurements was about 1o. The misorientations of (sub)grain boundaries were measured at least from three selected typical areas for each specimen. A total number of boundaries analyzed were about 80-100 for each sample. SAED patterns were taken from the regions of ~6 μm in diameter. The crystallites surrounded partially by LABs and HABs were identified as (sub)grains. The average grain size with HABs and (sub)grains were measured by linear intercept method on the TEM photos of deformed microstructure.. 2.3.3 Scanning Electron Microscopy/ Electron Back Scattering Diffraction. The samples for EBSD analysis was similar in shape and in way of preparation for the metallographical analysis described above. Before EBSD analysis, deformed samples of 2219 Al alloy were electropolished in a solution of 30% HNO3 and 70% CH3OH which was cooled by nitrogen for the samples deformed at 523 and 573 K or in solution of 80% C2H5OH, 12% 2nButoxyethanol, 8% HclO4 at room temperature for the samples deformed at 673 and 748 K. The orientation imaging microscopy (OIM) with automated indexing of electron back scattering diffraction (EBSD) patterns was performed in Hitachi S-4300H scanning electron microscope (SEM) with OIM analysis software provided by TexSem Lab Inc. The examining areas were automatically scanned with steps of 0.3, 0.4, 0.5 and 1.5 μm which were chosen for samples deformed at 523, 573, 673 and 748 K, respectively. ECAPed samples of Al-3%Cu alloy were electropolished in a solution of 80% C2H5OH, 12% 2n-Butoxyethanol, 8% HclO4 at room temperature at accelerating potential voltage ~17.5 V. Orientation imaging microscopy (OIM) and misorientation analysis was carried out by EBSD technique. OIM maps were acquired using a LEO-1530 SEM fitted with automated HKL-EBSD pattern collection system provided by HKL Technology, Inc. 40.

(56) 2.3.4 Quantitative Analysis of Microstructures. The microstructure analysis was based mainly on the data of EBSD method as well as TEM. The (sub)grain boundary misorientation distributions were obtained from EBSD data in the whole scanned area and also in fine grained regions developed. In the data presented, HABs were defined as being greater than or equal to 15o in misorientation and low angle boundaries as having a misorientation less than 15o. The boundaries with misorientation less than 2o were not taken into account. HABs and LABs are depicted in OIM maps as different in color lines. Thin white lines correspond to the boundaries with misorientation from 2 to 5o, thin black lines to misorientation from 6 to 14o and bold black lines to the boundaries with misorientation more than 15o. The crystallites bounded partially by LABs and HABs were identified as subgrains and grains, respectively. A significant advantage of EBSD technique over all the microscopic methods is that EBSD determines various kinds of boundaries, whereas conventional optical metallography reveals some of the boundaries after etching. (7). . On the other hand, TEM method allow to make. observations of grains as well as (sub)grains with LABs even less than 2o, which are omitted by EBSD(8). In addition, it should be noted that TEM observations are possible to carry out in relatively small and limited regions. In present work all the data of TEM, EBSD and OM are analyzed together. The volume fraction of fine grains VUFG surrounded by HABs with misorientations above 15o was calculated on the OIM maps using the conventional point counting technique. (9,10). . The. volume fraction measured by grid method was estimated for early deformation of ECAP as:. VUFG =. (2-1). nr N. nr is the number of points in the fine grained regions N is the total number of grid points This point counting method is the most efficient one for partially recrystallized materials (11). If nr is the number of the points which are recrystallized in a sample whose recrystallized fraction is. VUFG, the confidence limits (σ) are given as:. 41.