No description available.

<p>This thesis deals with the concept of ECAP applied to a commercial Al Mg-Si alloy (AA6082.50). A detailed analysis of the strains introduced by ECAP in a single passage through the die has been made, based on direct measurements on partially pressed samples. </p><p>Further, the mechanical properties of ECAP’ed samples have been investigated. The effect of initial material temper and post-ECAP artificial aging was investigated in addition to the effect of strain accumulation and processing route.</p><p> Finally, a detailed study of the textural and microstructural development was made. The continuous evolution of texture and microstructure was followed through the ECAP deformation zone up to an accumulated strain of 2 (2 passes) by route A, and linked to strain measurements from the same zone. </p><p>Strain measurements on sectioned samples have validated the plane strain assumption for ECAP. The shear angle has been measured and some typical features of ECAP have been corroborated, i.e. friction and material temper affect the strain distribution, the strain homogeneity and the workpiece corner angle, friction being the most significant here. Also, new conclusions have been drawn. The analysis of material element deformation histories along path lines reveals that ECAP can be interpreted as the combination of shearing and stretching (i.e. tension and/or compression). Furthermore, the final shear strain angle obtained in ECAP appears to be friction and material temper independent in the zone of homogeneous deformation. </p><p>The 6082 alloy has been successfully processed by ECAP at room temperature to strains ε =6 to ε =8. The alloy has been pressed in the solutionized, T4, homogenized and soft annealed states. The highest tensile strength was obtained from the solutionized material, followed by T4, homogenized and soft annealed. This behaviour is linked to the solid solution content prior to ECAP and the potential for dynamic precipitation during ECAP processing.</p><p> The tensile elongation to failure drops to a constant level between 4% and 9% after ECAP and is highest for the soft annealed and lowest for the solutionized material. The ductility in the solutionized material can, however, recover to ~18% elongation to failure (i.e. an increase by a factor 2-3) by low temperature heat treatment with only a small drop in tensile strength. </p><p>Soft annealed and ECAP’ed material has been compared to cold rolling to similar strains. The tensile strength response to accumulated strain is similar, but the ductility and post uniform deformation is different. However, the ECAP’ed material can be processed to higher strains and, thus, achieving higher strength. </p><p>The tensile yield strength behaviour of the homogenized and ECAP’ed material can be described by a simple relation to the grain size and the fraction high and low angle boundaries.</p><p> The typical texture components related to ECAP of aluminium, pressed by route A, have been identified. The typical ECAP texture starts to develop already at ~25% strain and increases in intensity during the first pass. In the second pass, two of the stable texture components continue to increase in intensity, while the other texture components decrease.</p><p> The microstructural development during the first pass is dominated by deformation banding leading to grain-subdivision. The average linear intercept distance (grain size) is reduced from ~80μm to ~4μm for high angle boundaries and from ~10μm to ~0.7μm for low angle boundaries.</p><p>During the second pass, the linear intercept distance is further reduced to ~1.8μm for high angle and ~0.3μm for low angle boundaries. Deformation twins are observed during the second pass and are believed to play an important role in severe plastic deformation when the grains reach the sub-micron or nano-metre scale.</p><p>The deformation banding have been explained in terms of the low energy dislocation structure (LEDS) theory, and has been shown to be an important mechanism in the early stages of grain subdivision, and is further believed to be the main source of high angle grain boundary formation by grain subdivision down to a grain size of approximately ~0 6μm, when other deformation mechanisms may be energetically more favourable.</p>

No description available.

This thesis deals with the concept of ECAP applied to a commercial Al Mg-Si alloy (AA6082.50). A detailed analysis of the strains introduced by ECAP in a single passage through the die has been made, based on direct measurements on partially pressed samples. Further, the mechanical properties of ECAP’ed samples have been investigated. The effect of initial material temper and post-ECAP artificial aging was investigated in addition to the effect of strain accumulation and processing route. Finally, a detailed study of the textural and microstructural development was made. The continuous evolution of texture and microstructure was followed through the ECAP deformation zone up to an accumulated strain of 2 (2 passes) by route A, and linked to strain measurements from the same zone. Strain measurements on sectioned samples have validated the plane strain assumption for ECAP. The shear angle has been measured and some typical features of ECAP have been corroborated, i.e. friction and material temper affect the strain distribution, the strain homogeneity and the workpiece corner angle, friction being the most significant here. Also, new conclusions have been drawn. The analysis of material element deformation histories along path lines reveals that ECAP can be interpreted as the combination of shearing and stretching (i.e. tension and/or compression). Furthermore, the final shear strain angle obtained in ECAP appears to be friction and material temper independent in the zone of homogeneous deformation. The 6082 alloy has been successfully processed by ECAP at room temperature to strains ε =6 to ε =8. The alloy has been pressed in the solutionized, T4, homogenized and soft annealed states. The highest tensile strength was obtained from the solutionized material, followed by T4, homogenized and soft annealed. This behaviour is linked to the solid solution content prior to ECAP and the potential for dynamic precipitation during ECAP processing. The tensile elongation to failure drops to a constant level between 4% and 9% after ECAP and is highest for the soft annealed and lowest for the solutionized material. The ductility in the solutionized material can, however, recover to ~18% elongation to failure (i.e. an increase by a factor 2-3) by low temperature heat treatment with only a small drop in tensile strength. Soft annealed and ECAP’ed material has been compared to cold rolling to similar strains. The tensile strength response to accumulated strain is similar, but the ductility and post uniform deformation is different. However, the ECAP’ed material can be processed to higher strains and, thus, achieving higher strength. The tensile yield strength behaviour of the homogenized and ECAP’ed material can be described by a simple relation to the grain size and the fraction high and low angle boundaries. The typical texture components related to ECAP of aluminium, pressed by route A, have been identified. The typical ECAP texture starts to develop already at ~25% strain and increases in intensity during the first pass. In the second pass, two of the stable texture components continue to increase in intensity, while the other texture components decrease. The microstructural development during the first pass is dominated by deformation banding leading to grain-subdivision. The average linear intercept distance (grain size) is reduced from ~80μm to ~4μm for high angle boundaries and from ~10μm to ~0.7μm for low angle boundaries. During the second pass, the linear intercept distance is further reduced to ~1.8μm for high angle and ~0.3μm for low angle boundaries. Deformation twins are observed during the second pass and are believed to play an important role in severe plastic deformation when the grains reach the sub-micron or nano-metre scale. The deformation banding have been explained in terms of the low energy dislocation structure (LEDS) theory, and has been shown to be an important mechanism in the early stages of grain subdivision, and is further believed to be the main source of high angle grain boundary formation by grain subdivision down to a grain size of approximately ~0 6μm, when other deformation mechanisms may be energetically more favourable.

<p>Two different solution heat treatments (2hours at 570ºC and 10minutes at 520ºC) have been used to study precipitation in two 6xxx (Al-Mg-si) dispersoid-freealloys with composition: 0.721 at % Si, 0.577 at % Mg (alloy A3) and 0.57 at %Si 0.72 at % Mg (alloy A12). The relation between their microstructure and macroscopical properties such as hardness and ductility has been investigated. Tensile tests, hardness measurements, electrical conductivity sigma tests, grain size measurements in optical microscope and microstructure characterization in Transmission Electron Microscope (TEM) have all been done. The effect of alloy composition and solution heat treatment temperature and time on the microstructure and the resulting macroscopical properties (hardness, yield stress, tensile strength and ductility) was investigated. The results indicate that when alloy A3 is solution treated at 520ºC for 10 minutes and then annealed for 3 hours at 175ºC, its hardness, yield stress and tensile strength as well as ductility is optimised i.e.A3 has better mechanical properties and low cost of production at these conditions. It has been proved that the strengthening was solely due to precipitation particles and not grain size.</p>

Two different solution heat treatments (2hours at 570ºC and 10minutes at 520ºC) have been used to study precipitation in two 6xxx (Al-Mg-si) dispersoid-freealloys with composition: 0.721 at % Si, 0.577 at % Mg (alloy A3) and 0.57 at %Si 0.72 at % Mg (alloy A12). The relation between their microstructure and macroscopical properties such as hardness and ductility has been investigated. Tensile tests, hardness measurements, electrical conductivity sigma tests, grain size measurements in optical microscope and microstructure characterization in Transmission Electron Microscope (TEM) have all been done. The effect of alloy composition and solution heat treatment temperature and time on the microstructure and the resulting macroscopical properties (hardness, yield stress, tensile strength and ductility) was investigated. The results indicate that when alloy A3 is solution treated at 520ºC for 10 minutes and then annealed for 3 hours at 175ºC, its hardness, yield stress and tensile strength as well as ductility is optimised i.e.A3 has better mechanical properties and low cost of production at these conditions. It has been proved that the strengthening was solely due to precipitation particles and not grain size.

<p>The overall objective of this work has been to develop aluminium alloys, which after hot and cold deformation are able to withstand high temperatures without recrystallising. This has been done by investigating aluminium alloys with various additions of hafnium, scandium and zirconium, with a main focus on Hf and to which extent it may partly substitute or replace Zr and/or Sc as a dispersoid forming elements in these alloys.</p><p><i>What is the effect of hafnium, alone and in combination with Zr and/or Sc and how do hafnium containing alloys perform?</i></p><p>It is shown that hafnium may alter or modify the casting structure, though, not to the better as it can form TCGs in combinations with Zr and/or Sc. This is not advantageous neither as far as it concerns grain refining nor precipitation of dispersoids.</p><p>When precipitation of binary Al-Hf is compared to Al-(Hf)-(Zr) alloys, hafnium shows even slower precipitation than in Al-(Hf)-(Zr) alloys and also much slower and with a poorer spatial distribution of dispersoids than in Al-Sc or Al-Hf-Sc alloys. As a consequence, it may be concluded that binary aluminium-hafnium alloys are of limited interest as they display a poor recrystallisation resistance when no other alloying elements are added.</p><p>However, when hafnium is added together with scandium and/or zirconium, precipitation may actually improve both in Al-Hf-Zr alloys and in scandium containing alloys like Al-Hf-Sc and Al-Hf-Sc-Zr. Hafnium can still not completely replace neither Zr nor Sc due to its poor precipitation properties. However, Al-Hf-Sc-Zr alloys show a better dispersoid distribution (number density and volume fraction) than what was observed in Al-Sc-Zr alloys after extrusion. The Al-Hf-Sc-Zr alloys also show extreme high temperature properties, by withstanding recrystallisation at high temperatures and long annealing times. After severe cold deformation, the extruded profiles of both the Al-Sc-Zr alloy and the Al-Hf-Sc-Zr alloy displayed a remarkable recrystallisation resistance.</p><p>The reason why the Al-Zr-Sc- and the Al-Hf-Zr-Sc alloys behave so well has been investigated by detailed 3D Atom Probe investigations. Due to the homogeneous precipitation of Al3Sc dispersoids and the retarding effect from Hf and/or Zr containing shells, thus limiting the coarsening of these dispersoids, these combinations have been shown successful.</p><p>As we have seen when Hf and Zr are added in combination equal or improved recrystallisation properties can be obtained. This is specially the case when both these elements are added together with scandium. Since Hf and Zr are extremely difficult (and thus costly) to separate, Al-Zr master alloys used in industry today can in the future probably contain more Hf, lowering the cost of master alloys.</p><p>The work presented in this thesis have hopefully added some new insight and a better understanding of the effects of adding various dispersoid forming elements to aluminium, alone and in various combinations, which may be useful for industry today and a basis for further alloy development.</p>

The overall objective of this work has been to develop aluminium alloys, which after hot and cold deformation are able to withstand high temperatures without recrystallising. This has been done by investigating aluminium alloys with various additions of hafnium, scandium and zirconium, with a main focus on Hf and to which extent it may partly substitute or replace Zr and/or Sc as a dispersoid forming elements in these alloys. What is the effect of hafnium, alone and in combination with Zr and/or Sc and how do hafnium containing alloys perform? It is shown that hafnium may alter or modify the casting structure, though, not to the better as it can form TCGs in combinations with Zr and/or Sc. This is not advantageous neither as far as it concerns grain refining nor precipitation of dispersoids. When precipitation of binary Al-Hf is compared to Al-(Hf)-(Zr) alloys, hafnium shows even slower precipitation than in Al-(Hf)-(Zr) alloys and also much slower and with a poorer spatial distribution of dispersoids than in Al-Sc or Al-Hf-Sc alloys. As a consequence, it may be concluded that binary aluminium-hafnium alloys are of limited interest as they display a poor recrystallisation resistance when no other alloying elements are added. However, when hafnium is added together with scandium and/or zirconium, precipitation may actually improve both in Al-Hf-Zr alloys and in scandium containing alloys like Al-Hf-Sc and Al-Hf-Sc-Zr. Hafnium can still not completely replace neither Zr nor Sc due to its poor precipitation properties. However, Al-Hf-Sc-Zr alloys show a better dispersoid distribution (number density and volume fraction) than what was observed in Al-Sc-Zr alloys after extrusion. The Al-Hf-Sc-Zr alloys also show extreme high temperature properties, by withstanding recrystallisation at high temperatures and long annealing times. After severe cold deformation, the extruded profiles of both the Al-Sc-Zr alloy and the Al-Hf-Sc-Zr alloy displayed a remarkable recrystallisation resistance. The reason why the Al-Zr-Sc- and the Al-Hf-Zr-Sc alloys behave so well has been investigated by detailed 3D Atom Probe investigations. Due to the homogeneous precipitation of Al3Sc dispersoids and the retarding effect from Hf and/or Zr containing shells, thus limiting the coarsening of these dispersoids, these combinations have been shown successful. As we have seen when Hf and Zr are added in combination equal or improved recrystallisation properties can be obtained. This is specially the case when both these elements are added together with scandium. Since Hf and Zr are extremely difficult (and thus costly) to separate, Al-Zr master alloys used in industry today can in the future probably contain more Hf, lowering the cost of master alloys. The work presented in this thesis have hopefully added some new insight and a better understanding of the effects of adding various dispersoid forming elements to aluminium, alone and in various combinations, which may be useful for industry today and a basis for further alloy development.

<p>The processing of aluminium alloys from casting to end product is associated with a large number of metallurgical phenomena. In order to further improve and optimise process routes and alloys, a thorough understanding of the thermomechanical treatments by experimental observations and physically based modelling is necessary.</p><p>In part A of this thesis the work hardening behaviour of non-heat treatable alloys is followed up to large strains. The evolution in strength, microstructure and texture during cold rolling is analysed in commercially pure aluminium, AlMg alloys and AlMn alloys, covering both solid solution strengthening, particle strengthening, grain size effects and effects of impurities. A recently developed model for work hardening is applied to these alloys and discussed in relation to the experimental observations. The model successfully predicts the effects of Mg in solid solution, and is used to explain the effects of nonshearable particles and grain size on the strength and work hardening in stage II-III. Mn, Fe and Si in solid solution are suggested to create solute clusters that increase the strength significantly. At higher strains, stage IV, an unexpectedly low work hardening rate of high-Mg alloys is ascribed to a different storage pattern of dislocation and an increased amount of shear bands. High-resolution EBSD in FE-SEM is demonstrated to be a promising tool for substructure characterisation of cold rolled alloys.</p><p>In part B the mechanical anisotropy of flat extruded profiles of heat-treatable alloys is investigated. Two recrystallised alloys, AA6063 and 7030, and two non-recrystallised alloys, AA6082 and 7108, are tested in uniaxial tension in the solution heat treated condition, exhibiting strong directionality of yield stress, work hardening rate, elongation and r-value. The initial textures are very strong, and the anisotropy is analysed in terms of the Taylor model. The texture is found to be responsible for most of the strength anisotropy in the non-recrystallised alloys, while other sources of anisotropy must be present in the recrystallised alloys. Variations in the deformation structure development, indicating different slip activity in different directions, are believed to be partly responsible. The Taylor RC model predicts the r-values very well in all alloys, and a number of active slip systems in each position between two and three is assumed to be realistic for all alloys and directions.</p>

The processing of aluminium alloys from casting to end product is associated with a large number of metallurgical phenomena. In order to further improve and optimise process routes and alloys, a thorough understanding of the thermomechanical treatments by experimental observations and physically based modelling is necessary. In part A of this thesis the work hardening behaviour of non-heat treatable alloys is followed up to large strains. The evolution in strength, microstructure and texture during cold rolling is analysed in commercially pure aluminium, AlMg alloys and AlMn alloys, covering both solid solution strengthening, particle strengthening, grain size effects and effects of impurities. A recently developed model for work hardening is applied to these alloys and discussed in relation to the experimental observations. The model successfully predicts the effects of Mg in solid solution, and is used to explain the effects of nonshearable particles and grain size on the strength and work hardening in stage II-III. Mn, Fe and Si in solid solution are suggested to create solute clusters that increase the strength significantly. At higher strains, stage IV, an unexpectedly low work hardening rate of high-Mg alloys is ascribed to a different storage pattern of dislocation and an increased amount of shear bands. High-resolution EBSD in FE-SEM is demonstrated to be a promising tool for substructure characterisation of cold rolled alloys. In part B the mechanical anisotropy of flat extruded profiles of heat-treatable alloys is investigated. Two recrystallised alloys, AA6063 and 7030, and two non-recrystallised alloys, AA6082 and 7108, are tested in uniaxial tension in the solution heat treated condition, exhibiting strong directionality of yield stress, work hardening rate, elongation and r-value. The initial textures are very strong, and the anisotropy is analysed in terms of the Taylor model. The texture is found to be responsible for most of the strength anisotropy in the non-recrystallised alloys, while other sources of anisotropy must be present in the recrystallised alloys. Variations in the deformation structure development, indicating different slip activity in different directions, are believed to be partly responsible. The Taylor RC model predicts the r-values very well in all alloys, and a number of active slip systems in each position between two and three is assumed to be realistic for all alloys and directions.