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Effect of homogenization on the microstructural development in a d.c. cast aa3104 aluminum alloy used… Gandhi, Chetak 1999

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EFFECT OF HOMOGENIZATION ON THE MICROSTRUCTURAL DEVELOPMENT IN A D.C. CAST AA3104 ALUMINUM ALLOY USED FOR CANBODY STOCK  by CHETAK GANDHI B . Tech., Indian Institute o f Technology, Kharagpur, India, 1994  A THESIS S U B M I T T E D IN P A R T I A L F U L F I L L M E N T OF T H E R E Q U I R E M E N T S FOR THE DEGREE OF M A S T E R OF APPLIED SCIENCE in T H E F A C U L T Y OF G R A D U A T E STUDIES (Department o f Metals and Materials Engineering)  '  W e accept the thesis as comforming to the required standard  T H E UNIVERSITY OF BRITISH C O L U M B I A A p r i l , 1999  © Chetak Gandhi, 1999  in  presenting  degree  this  at the  thesis  in  partial  fulfilment  University of  British  Columbia,  of  the  requirements  for  that permission  copying  granted  department  this or  thesis by  publication of this  for scholarly  his thesis  or  her  may be  representatives.  It  is  for extensive  by the head of  understood  that  for financial gain shall not be allowed without  permission.  Department of The University of British Columbia Vancouver, Canada  DE-6 (2/88)  purposes  advanced  I agree that the Library shall make it  freely available for reference and study. I further agree of  an  copying  my or  my written  Abstract A s customer demands become more stringent for canbody stock, it becomes  essential to understand  the complex interaction between the  processing  conditions and resulting product properties. This research focused on investigating the influence o f homogenization process parameters (heat-up rate, soak temperature and time) on the microstructural evolution o f an A A 3 1 0 4 aluminum alloy used for canbody stock. Experiments were conducted on samples taken from an industrial D . C . cast ingot and homogenized i n a programmable temperature controlled laboratory furnace under various thermal profiles (i.e. homogenization temperatures 5 5 0 ° C , 5 8 0 ° C and 6 1 0 ° C at various heating rates and homogenization soak times o f up to ten hours). The samples were then characterized i n terms o f their microstructure (retained manganese i n solid solution, percentage a-phase, and size distribution and density o f dispersoids). The homogenization process parameters were found to affect the evolving microstructure profoundly with: • A n increase i n heat-up rate favoring a reduction i n the number o f evolving dispersoids. • A n increase i n soak temperature increasing the M n i n solid solution, and decreasing the number o f dispersoids that form. • A n increase i n soak time up to 3 hrs increasing the volume percent  of a-  Ali (Fe,Mn) Si. 2  3  Based o n this work, a homogenization profile for optimum microstructure and texture development w o u l d include a fast heat-up rate to a high soak temperature (610°C) w i t h moderate soak times (up to 3 hrs). ii  Table of Contents Abstract  ii  Table of Contents  iii  List of Figures  vi  List of Tables  ix  Acknowledgements  x  1  Introduction  1  2  Literature Review  5  2.1  Industrial C a n M a k i n g Operation  5  2.2  Solidification - As-cast Microstructure  6  2.3  Microstructural Changes During Homogenization  10  2.3.1  Equilibration o f A l l o y i n g Elements and Removal o f Microsegregation  11  2.3.2  Precipitation o f Dispersoids  14  2.3.3  Modification o f As-Cast Constituent Phases  15  2.3.4  M n i n S o l i d Solution  17  2.4  Effect o f Composition on the As-Cast and Homogenized Microstructure  19  2.4.1  Silicon  19  2.4.2  Iron  20  2.4.3  Copper  20  2.4.4  Magnesium  21  2.4.5  Manganese  22 iii  2.5  Effect o f Homogenization Process Parameters on Microstructure  2.5.1  Heat-up Rate  24  2.5.2  Soak Temperature  24  2.5.3  Soak time  29  2.5.4  Second Stage Soak  30  2.6  Models for Microstructural Changes during Homogenization  2.7 Summary  3.  4.  23  32 34  Scope and Objectives  35  3.1  Scope  35  3.2  Objectives  36  Experimental  37  4.1  Start Material  37  4.2  Homogenization Experiments  38  4.2.1  Laboratory furnace  38  4.2.2  Homogenization profiles  39  4.3  Manganese Retained i n Solid Solution  42  4.4  Microstructure Characterization / Metallography  43  4.4.1  Grinding and Polishing Procedure  43  4.4.2  Etching  44  4.4.3  Electron Probe M i c r o Analysis ( E P M A ) Line Scans  47  4.4.4  Image Analysis  47  4.5  As-Cast Characterization...  48  iv  5.  Results and Discussion 5.1  50  Microstructure Evolution during Homogenization  5.1.1  50  Effect o f Heat-up Rate  50  5.1.1.1 Effect on M n i n S o l i d Solution  51  5.1.1.2Constituent Particles and percent a - A l i ( F e , M n ) S i  53  5.1.1.3Evolution o f Dispersoids  54  5.1.1.4Removal o f Microsegregation  58  2  5.1.2  3  Effect o f Soak Temperature and Soak Time  59  5.1.2.1 Effect on M n i n Solid Solution  59  5.1.2.2Constituent Particles and Percent a - A l i ( F e , M n ) S i  62  5.1.2.3 Evolution o f Dispersoids  64  5.1.2.4Cool D o w n to R o l l i n g Temperature  66  2  5.2  Influence  o f Homogenization Microstructure on  after H o t - R o l l i n g  6.  Texture/Microstructure 70  Summary and Conclusions 6.1  3  73  Summary  73  6.2  Conclusions  75  6.3  Future W o r k  77  Bibliography  78  Appendix A  .....83  Appendix B  86  Appendix C  90 V  List of Figures Figure 1.1- A l u m i n u m beverage can demand [1 ].  1  Figure 1.2- Typical processing sequence for the production o f final gauge sheet [3].  3  Figure 2.1 - Photomicrographs showing the size and shape o f the constituents i n ascast ingot at: (a) edge and (b) center locations [12].  7  Figure 2.2 - Diffusion coeffients for various elements i n aluminum at different temperatures [18].  12  Figure 2.3 - Effect o f 0.06% Fe and 0.02% C u on the loss o f supersaturation i n A l - 1 M n A l l o y s at 5 5 0 ° C [30].  21  Figure 2.4 - M n precipitation vs. time and temperature for an as-cast 3104 alloy [34].  25  Figure 2.5 - M n content i n a-phase constituents as a function o f homogenization temperature [26].  27  Figure 2.6 - Examples o f secondary precipitates i n T E M thin foil specimens o f A A 3 0 0 4 alloy at (a) 6 0 0 ° C and (b) 5 6 0 ° C [27].  28  Figure 2.7 - T E M photomicrographs showing the a-dispersoids i n ingot homogenized at 577°C for (a) 14 hrs and (b) 26 hrs, followed b y water quench [12].  30  Figure 4.1 - Laboratory furnace set-up: a ) overview and b ) cross-sectional view.  38  Figure 4.2 - Experimental homogenization profiles with varying heat-up rates (dots mdicating where samples were taken for microstructure characterization).  40  Figure 4.3 - Experimental homogenization profiles w i t h varying soak temperatures (550°C, 5 8 0 ° C and 610°C).  40  Figure 4.4 - Experimental thermal profiles used for samples homogenized at different soak temperatures.  41  Figure 4.6 - Resistivity o f aluminum alloy 3004 as a function o f the percentage o f M n in solid solution [43]. 43 Figure 4.7 - As-cast microstructure (1000X) showing the various second phase particles.  45  Figure 4.8 - E D X / S E M results confirming the etching difference between a A l i ( F e , M n ) S i and Al6(Fe,Mn) particle.  46  2  3  vi  Figure 4.9 - Variation o f the cumulative average per field with the measured quantity per field for volume fraction second phase.  48  Figure 5.1 - Effect o f heat-up rate on % M n in solid solution.  51  Figure 5.2 - Effect o f heat-up rate on % a-Ali2(Fe,Mn)3Si after a 3 hour hold at 580°C.  54  Figure 5.3 - Microstructures o f samples for 580°C soak temperature after 3 hour hold for: a) 279°C/hr and b) 47°C/hr.  55  Figure 5.4 - Microstructures o f samples at 4 5 0 ° C during heat-up for: a) 2 7 9 ° C / h r and b) 47°C/hr.  57  Figure 5.5 - M n i n solid solution at different soak temperatures at a) 4 7 ° C / h r and b) 279°C/hr.  61  Figure 5.6 - Effect o f soak temperature on %a-Ali2(Fe,Mn)3Si after a soak period o f three hours.  62  Figure 5.7 - Effect o f soak time on %a-Ali (Fe,Mn)3Si at various soak temperatures at a heat-up rate o f 47°C/hr.  63  Figure 5.8 - Microstructures (1000X) o f samples after three hour hold w i t h a heat-up rate o f 47°C/hr for soak temperatures of: a) 550°C, b) 5 8 0 ° C and c) 610°C.  66  Figure 5.9 - M n i n solid solution for homogenization profiles o f 5 5 0 ° C , 5 8 0 ° C and 6 1 0 ° C for three hour hold followed by cool down to 5 0 0 ° C .  67  Figure 5.10 - Microstructures (1000X) o f samples after a cool down to 5 0 0 ° C for soak temperatures of: a) 5 5 0 ° C , b) 5 8 0 ° C and c) 610°C.  69  Figure A l - Temperature variations between furnace (580°C) and sample o f original furnace at a) two sample positions i n the furnace, and b) on a raised platform near the furnace thermocouple.  84  2  Figure A 2 - Temperature variations between furnace (580°C) and sample temperature in the retrofitted furnace. 85 Figure B l - Microstructures o f samples at 580°C soak temperature after 3 hr hold for a) 37°C/hr, b) 47°C/hr, c) 70°C/hr, d ) l 12°C/hr, and e) 279°C/hr. Figure C l - E P M A line scans ( M g and M n ) for the as-cast microstructure.  89 91  Figure C 2 - E P M A line scans ( M g and M n ) for 47°C/hr profile at a) 4 0 0 ° C , b) 500°C, 580°C (0 hrs), and d) 580°C (3hrs). 93 vii  Figure C 3 - E P M A line scans ( M g and M n ) for 279°C/hr profile at a) 4 0 0 ° C , b) 500°C, 5 8 0 ° C (0 hrs), and d) 580°C (3hrs).  viii  List of Tables Table 1.1 - Mechanical properties o f aluminum sheet i n H I 9 condition.  2  Table 2 . 1 - N o m i n a l weight percent composition o f selected 3xxx alloys.  5  Table 2.2 - Size variation o f cells and stringers with various casting methods [13].  8  Table 2.3 - Relative solid solubility and rate o f diffusion o f major alloying elements in 3xxx series alloys [10].  12  Table 2.4 - Effects o f homogenization treatment and S i content on oc-phase constituents (i.e. a-Ali2(Fe,Mn) Si) [4]. 3  16  Table 2.5 - Change i n a-Ali2(Fe,Mn) Si dispersoid mean size during preheat heat-up and soak [16].  29  Table 4.1 - Composition o f A A 3 1 0 4 alloy used.  37  Table 4.2 - Range o f homogenization process parameters used i n the study.  39  Table 4.3 - Composition o f A A 3 1 0 4 and effect o f elements on resistivity [18].  42  Table 4.4 - Grinding and polishing procedure followed.  44  Table 4.5 - As-cast characterization between 5 - 2 5 c m from surface.  49  Table 5.1 - M i n i m u m M n i n solid solution and corresponding temperature for different heat-up rates.  52  Table 5.2 - M n and M g segregation results for 4 7 ° C / h r and 2 7 9 ° C / h r profile at various temperatures during heat-up.  58  3  ix  Acknowledgements  I would like to thank, firstly, m y advisor, D r . Mary A . Wells, for her constant support and encouragement throughout this work. I would also like to thank N S E R C , Canada and A l c a n International L t d . for the material and financial help, and D r . D . J. L l o y d ( K R D C , Alcan Int. Ltd.) for many useful discussions and suggestions. Special thanks to M r . Ross M c l e o d and M r . Carl N g for preparing all m y samples, to M r . R u d y Cardeno for the image analysis work, to M s . M a r y Mager for her instructions and advice with the S E M , to M r . Serge Millare for the help with the furnace and other electrical issues, and to Peter M u s i l for help with the furnace repair and fittings. I would like to convey many thanks to M r . Gary Lockhart for his help with experimental set-up and valuable advice and Rob Stevenson for his help i n the project. Finally, this work would not have been completed without the encouragement and support o f m y family and friends.  x  1  Introduction  Over the last decade there has been considerable growth in the use o f aluminum alloys for rigid packaging o f food products, with the most significant increase being in the beverage container market. The world market for beverage containers has seen a steady rise over the last ten years [1] (Figure 1.1). The continuing strong demand for aluminum alloys i n this very competitive market is forcing aluminum sheet producers to critically examine their production processes in terms o f both cost and final material properties.  Mideast & Africa  Hi Europe M U S . & Canada  Latin America & M e x i c o  Pacific  Figure 1.1 - A l u m i n u m beverage can demand [1].  1  A l u m i n u m beverage cans are deep drawn from the final cold rolled sheets o f A A 3 0 0 4 or A A 3 1 0 4 aluminum alloys [2, 3]. These alloys are essentially A l - M n - M g alloys also containing C u , S i and Fe. The current demand is for A A 3 1 0 4 final canbody sheet that is < 0.3 m m gauge and with strength corresponding to the H I 9 condition (Table 1.1)  Table 1.1 - Mechanical properties o f aluminum sheet in H I 9 condition. Tensile strength (UTS)  295 M P a  Y i e l d Strength  285 M P a 2%  Elongation  Besides strength, elongation and gauge, another important requirement for the final sheet is its crystallographic texture. Undesirable textures i n the sheet leads to "earing" which may cause trouble during can making as well as yield losses as they must be trimmed off. The phenomenon o f "earing" refers to undulations on the r i m o f cups formed during cup drawing or w a l l ironing, the high points named as "ears". During processing, earing is controlled b y balancing two types o f textures; i.e., b y developing a strong cube texture (0-90° earing) upon recrystallization after hot rolling followed b y the superposition o f a rolling texture (45°-45° earing) during subsequent heavy cold rolling to final gauge [4, 5 , 6 ] . The resulting product w i l l not show zero earing, but rather small mixed ears o f similar size.  2  Typically processing o f Direct C h i l l (D.C.) cast aluminum ingots (generally 300 760 m m thick) to final gauge sheet (-0.3 m m thick) for the beverage can market involves a number o f processing steps (Figure 1.2).  Process time Figure 1.2 - Typical processing sequence for the production o f final gauge sheet [3].  The process steps include: i) homogenization (temperatures above ~ 5 5 0 ° C and soaked for a controlled period). ii) breakdown rolling (the temperature falls from ~ 5 0 0 ° C to ~ 3 0 0 ° C and the ingot is reduced to a slab o f ~25 m m thickness). iii) tandem rolling (multiple stand continuous m i l l w i t h 3-4 stands produce coiled metal ~2.5 m m thick that recrystallizes without furnace annealing).  3  iv) cold rolling (-87% reduction in thickness to final sheet with < 0.3 mm thick). v) final heat treatment. One area in sheet production that plays a crucial role in the final sheet gauge properties of canbody stock is homogenization. Homogenization, also known as preheating, is typically used to: (i) remove microsegregation of alloying elements in the ingot and (ii) raise the ingot temperature for hot rolling. The most economical way to run this process would be to have a cycle as short as possible and at as low a temperature as possible. Balanced with this is the need to produce an optimum microstructure capable of producing the desired properties in the final gauge sheet. In addition, the microstructure of the homogenized ingot can influence the strength, formability and the texture of the final gauge sheet for the canbody stock. Some of the important microstractural features in the as-homogenized ingot include relatively large constituent particles, fine dispersoids that are precipitated during processing and the solute levels in solid solution [7]. Each of these can affect the recrystallization process and in particular, the strength of the cube texture after hot deformation, which is of paramount importance in controlling anisotropy in the final product [5, 8]. In an effort to understand and quantify the microstructural changes that occur during homogenization, this study examined the effect of the homogenization process on key microstructural features.  4  2  Literature Review  2.1  Industrial Can Making Operation  The manufacture  o f beverage  cans is one o f the most demanding sheet  applications for aluminum [9]. The primary alloys used for canbody stock are A A 3 0 0 4 and A A 3 1 0 4 , which are non-heat-treatable aluminum alloys. The nominal specification for these alloys is given i n Table 2.1 but typically beverage manufacturers operate on much tighter specifications than those given. A A 3 0 0 4 and A A 3 1 0 4 are essentially the same alloys with the exception o f the tighter M n range i n A A 3 0 0 4 .  Table 2 . 1 - N o m i n a l weight percent composition o f selected 3xxx alloys. Alloy  Mg  Mn  Fe  Si  Cu  AA3004  0.8-1.3  1.0-1.5  0.7 max.  0.3 max.  0.25 max.  AA3104  0.8-1.3  0.8-1.4  0.8 max.  0.4 max.  0.05-0.25  Conventional body stock for beverage cans is produced from 600 to 750 m m diameter ingots b y semi-continuous direct-chill (D.C.) casting or electromagnetic casting ( E M C ) processes. Other methods for casting canbody stock that have been tried experimentally but are currently not widely used are: (a) continuous casting o f 13 to 25  5  m m thick slabs by introducing liquid metal between parallel mold surfaces (e.g., the Hazelett water-cooled steel belt casters, Alusuisse Caster II and Lauener block casters which employ large metal chill blocks) (b) continuous casting o f 6 to 12 m m thick strips produced by introducing liquid metal between water cooled rolls, followed by cold rolling to final gauge without any preheating or hot rolling (e.g., Pechiney Jumbo 3 C caster and the Hunter process).  2.2  Solidification - As-cast Microstructure  The as-cast microstructure  defines  the starting structure for the  study o f  microstructure evolution during homogenization. The cast structure o f A A 3 x x x consists o f a cellular dendritic aluminum matrix with stringers containing intermetallic phases along the cell boundaries, remnants o f the last highly segregated melt to freeze (Figure 2.1). The main features defining the cast structure are: (a) the cell size and intermetallic particle size, (b) the degree o f solute supersaturation, and (c) the morphology o f the intermetallics. Due to the thickness o f the ingot, the freezing rate varies from the edge to the center (Figure 2.1) [10, 11, 12]. The freezing rate can also vary greatly based on the casting practices. A s the freezing rate during casting increases, the dendrite cell size decreases while the stringers become thinner and envelop the cell boundaries more continuously (Table 2.2). The cell size plays an important role i n the equilibration process o f the alloying elements and the spatial arrangement o f the precipitating dispersoids during homogenization. Typically, the finer the cell size the smaller the distance which the elements must travel within the cell and therefore homogenization  6  occurs more quickly and the precipitate stringers tend to be finer yet engulf the cells more continuously [10, 11, 12].  I  I  50 Lim  a  I  I  50 nm  b  Figure 2.1 - Photomicrographs showing the size and shape of the constituents in as-cast ingot at: (a) edge and (b) center locations [12].  7  Table 2.2 - Size variation of cells and stringers with various casting methods [13]. Cell Size  (mm)  Freezing Rate- edge (°C/s)  (urn)  Stringer thickness (um)  Thickness Process  Shape Ingot  D.C*  660  1  60  1.0  Billet  D.C*  200  10  45  0.7  Slab  Belt/Block  19  20  30  0.5  Stationary Strip  Mold  13  100  10  0.2  Strip  Twin Mold  6  700  6  0.1  *NOTE: The freezing rate at the center of a D . C cast ingot is ~0.25°C/sec whereas at the edge is ~l°C/sec) [13].  The intermetallic phases that form in AA3104 consist of two major types: orthorhombic Al6(Fe,Mn) and cubic a-Ali (Fe,Mn)3Si. The relative amounts of these 2  phases are determined by the alloy composition as well as the solidification rate [11, 12, 14,15]. In AA3104 under commercial D.C-cast ingot cooling rates of about l°C/s (Table 2), about 85% of the primary constituent particles in the as-cast structure correspond to the Al (Fe,Mn) phase and the remainder to oc-Ali2(Fe,Mn) Si [14]. At these cooling rates, 6  3  about 25-30% of the Mn precipitates out of solid solution in the form of intermetallic particles during solidification while the rest of the Mn (-70-75%) remains in solid solution, thereby producing an as-cast structure which is in a supersaturated metastable  8  solid solution condition [14]. Moreover, the as-cast structure exhibits microsegregation o f M g , S i , and to a lesser extent M n , across the dendritic cellular cast structure [16]. The formation o f phases and the temperatures at which they nucleate during the solidification process o f A A 3 0 0 4 have been summarized by Backerud et al [17]. Initially as the molten aluminum begins to cool and its temperature reaches ~ 6 4 8 - 6 5 2 ° C the liquid w i l l begin to solidify and form an aluminum dendritic matrix according to : 1.  L i q u i d (L) -> A l , dendritic network  @ 648-652°C  The first constituent particles A l ( F e , M n ) start to nucleate at the periphery o f the growing 6  dendrites at ~ 6 4 3 ° C according the eutectic phase transformation 2; 2.  L i q u i d (L) -> A l + A l ( F e , M n )  @ 643°C  6  A t lower temperatures (~638°C), a peritectic reaction occurs between the (Mn,Fe)Al6 particles and the remaining liquid which has been enriched i n S i according to 3 a: 3a.  L i q u i d (L) + A l ( F e , M n ) -> a - A l ( F e , M n ) S i 6  12  @ 638°C  3  A s the liquid continues to be enriched i n solute the a-Ali2(Fe,Mn)3Si phase can form directly from it according to 3b: 3b.  L i q u i d (L) -> A l i ( F e , M n ) S i 2  @ 638°C  3  Finally, at slightly lower temperatures, eutectic P-Mg Si particles precipitate from the 2  melt given the high S i and M g i n the alloy according to 4; 4.  L i q u i d (L) -> A1 + A l ( F e , M n ) S i + M g S i 12  3  2  @ 630°C  Homogenization, which occurs after casting o f the ingot, modifies the as-cast microstructure significantly and has a direct influence on both the hot-rolled and cold-  9  rolled sheet properties. A s a result it is important to understand the microstructural changes which occur during homogenization, the effect o f homogenization parameters on the microstructure evolution and how this behaviour can be represented mathematically.  2.3  Microstructural Changes During Homogenization  During homogenization, the D . C . cast ingot is heated and soaked at a temperature in the range o f about 10-100°C below the solidus for the alloy and 2 0 - 1 0 0 ° C above the desired initial hot working temperature.  For A A 3 x x x alloys, the solidus ranges from  630-645°C, homogenization temperatures are typically 530-620°C and the initial hot rolling temperatures are 480-520°C [10]. The basic purpose o f the homogenization process is to reduce microsegregation that exists i n the as-cast ingot as well as preheat the ingot to the desired hot rolling temperature.  Concurrent with the  constitutional equilibration and reduction i n  microsegregation o f some elements i n solid solution during homogenization, several other rmcrostructural changes occur i n the matrix, such as: i)  precipitation o f supersaturated elements as dispersoids,  ii) dissolution o f unstable phases or precipitates, iii) transformation o f Alg(Fe,Mn) to a-Ali2(Fe,Mn)3Si, and iv) coarsening o f the stable intermetallic particles.  10  2.3.1  Equilibration of Alloying Elements and Removal of Microsegregation  Microsegregation i n the as-cast ingot and equilibration o f alloying elements during homogenization depend upon a number o f factors: their solid solubility i n the aluminum matrix, their rate o f diffusion i n the aluminum matrix for the element and the cell size o f the cast microstructure.  In A A 3 x x x aluminum alloys both M n and Fe have  relatively low solubility in the aluminum matrix and hence they tend to be the elements which w i l l be i n a supersaturated condition i n the as-cast aluminum matrix (Table 2.3). During homogenization, the supersaturation o f M n and Fe within the dendritic cells is largely relieved b y the precipitation o f fine (Fe,Mn)-rich dispersoids consisting primarily o f A l i 2 ( M n , F e ) S i phase [16], or Ale(Mn,Fe) i f the S i content is less than about 0.07 wt% 3  [4]. In the as-cast aluminum structure both M g and S i w i l l exhibit a significant amount o f microsegregation or coring across the dendritic cellular as-cast microstructure. The segregation pattern existing i n the as-cast microstructure w i l l be unstable as it is heated up to and held at the homogenization temperature and equilibration o f these elements w i l l begin to occur during heat-up to the homogenization soak temperature. In fact, b y the time the ingot has reached 500°C, the M g and S i levels have typically homogenized completely across the cell [16].  11  Table 2.3 - Relative solid solubility and rate o f diffusion o f major alloying elements i n 3xxx series alloys [10].  Moderate solubility, Mn, L o w diffusion rate Appreciable solubility, Cu, M g , Si H i g h diffusion rate Extremely low solubility, Fe V e r y low diffusion rate  cr i 8  1200 1000 1 <1 i 1 • 1 600 500  I  P T — i — i c — i —- i  800 I I | 400 1  600 F i 1  i  i  300 C 1—l  1.4  (000/T Figure 2.2 - Diffusion coefflents for various elements i n aluminum at different temperatures [18].  12  The mechanism behind the homogenization process is diffusion and M n and Fe have much lower diffusion coefficients (Figure 2.2). Since most F e stays out o f solid solution, the rate determining element is M n . Diffusion o f M n \ is influenced b y both temperature and time and can be described mathematically by Fick's law:  dt  dx  2  This equation states analytically how the concentration (C) w i l l change w i t h time (t) according to a temperature dependent constant D which is called the  diffusion  coefficient. D depends on the temperature according to the Arrhenius equation:  D = D exp-  Q / R T  0  (2)  Where: D = constant 0  Q = activation energy (kJ/mole) R = gas constant (8.314 J/mole°K) T = temperature (°K)  It is this exponential dependence on temperature that causes diffusional removal o f concentration gradients to take place more rapidly at higher temperatures. Since this 13  process is governed entirely by diffusion, three inain factors w i l l determine the time needed for this to occur: i) ii)  temperature, the actual values o f D and Q in the Arrhenius equation which w i l l differ for 0  different elements, and iii)  the cell size o f the cast microstructure since this w i l l determine the distance over  which diffusion must take place. 2.3.2  Precipitation of Dispersoids  Dispersoids are secondary precipitates that form during homogenization as a result o f supersaturation o f particular elements i n the matrix. In A A 3 x x x alloys, these dispersoids are mostly oc-Ali (Fe,Mn)3Si, or sometimes Al6(Fe,Mn) i f the S i is lower than 2  0.08% i n the ingot [4, 10, 16].  The formation o f these precipitates during the  homogenization process is a complex procedure involving the nucleation o f the new phase into the aluminum matrix followed by growth until all of the excess solute has been consumed. In A A 3 x x x aluminum alloys, as the ingot is heated above 300°C, needles o f and more sparsely distributed square platelets o f P-Mg2Si begin to precipitate i n the A l matrix [16]. A t a slightly higher temperature o f around 400°C, precipitates o f ocAli2(Fe,Mn)3Si  begin  to  nucleate  on  the  pre-existing  (3-Mg2Si  needles  which  subsequently dissolve as the temperature is increased. A s the temperature is further increased, a-Ali (Fe,Mn)3Si dispersoids continue to precipitate while the existing 2  precipitates begin to grow through coarsening mechanisms. B o t h the volume fraction and average size o f dispersoids increases until ~ 480°C when the M n precipitation rates peak  14  [4, 19]. Above that temperature, re-dissolution o f dispersoids take place due to increased solid solubility o f M n i n A l at higher temperatures [4, 16, 20]. The re-dissolution o f dispersoids causes their number to decrease, however, those dispersoids that remain do undergo growth [16]. This shows that while the dispersoids nucleation period stops at ~ 4 8 0 ° C , the growth period continues at higher temperatures. Subsequent cooling from the peak preheating temperature to the hot rolling temperature causes renewed precipitation and growth o f oc-Ali2(Fe,Mn)3Si dispersoids. The final average size o f dispersoids depends on the homogenization profile and can vary from 50nm to about 0.5 urn i n diameter [4,16, 19].  2.3.3 The  Modification of As-Cast Constituent Phases iron-rich  intermetallic  particles  which  form  during  solidification  (Ai^(Fe,Mn)) are not removed during homogenization since the alloy is never heated into a single phase field and the diffusion coefficient o f F e i n aluminum is so low. However, some changes, w h i c h are both time and temperature dependent, i n the as-cast constituent particles do begin to occur during the homogenization process: i)  Ale(Fe,Mn) starts to transform to a - A l ( F e , M n ) S i , and  ii)  Spherodization o f the particles. Between 4 8 0 ° C and 593°C, the normally dominant Ale(Fe,Mn) constituent phase  gradually transforms  to Ali2(Fe,Mn)3Si phase. This transformation occurs due to  localized enrichment o f the matrix with S i , originally present i n the form o f M g S i i n the 2  as-cast microstructure but which dissolves during the early stages o f the homogenization  15  process.  The transformation is initiated by nucleation o f a-Ali2(Fe,Mn)3Si at the  interface between Ai6(Fe,Mn) and the solute enriched matrix [21]. The a-Ali2(Fe,Mn)3Si phase then grows into and consumes the Ale(Fe,Mn) particles.  Depending on the  composition o f the alloy, this transformation can either go quickly to completion or be stopped before completion because o f a lack o f S i . In order for this reaction to proceed, M n and S i must diffuse through the matrix, and the rate o f the reaction is usually controlled by M n diffusion. If S i is tied up as M g S i , due to high M g or low S i , the 2  transformation is slowed down due a lack o f S i . In general, this transformation is promoted by a higher S i content and a higher homogenization temperature (Table 2.4) [4, 22, 23].  Table 2.4 - Effects o f homogenization treatment and S i content on a-phase constituents (i.e. a - A l i ( F e , M n ) S i ) [4]. 2  Simulated Homogenization  3  Percent o f a - A l i ( F e , M n ) S i 2  3  Treatment  0.21% S i Ingot  0.30% S i Ingot  As-cast  10  20  Heated to 427°C  10  20  Heated to 4 8 2 ° C  25  35  Heated to 538°C  50  75  Heated to 566°C  70  97  Heated to 593°C  85  100  Heated to 616°C  85  100  Heated to 632°C  50  100  Heated to 638°C  0  100  16  The exact significance o f the transformation o f Al6(Fe,Mn) to a - A l ( F e , M n ) S i is not completely understood although an increased level o f a - A l ( F e , M n ) S i has been reported to improve the drawing and ironing behaviour in canmaking; the a - A l ( F e , M n ) S i particles are harder than the Al6(Fe,Mn) particles and help i n cleaning the dies thereby prevent galling [4]. The  constituent  particles  can  also  change  their  morphology  during  homogenization. Initially, the particles have a plate-like elongated and irregular shape with a high surface area. During homogenization, spherodization o f particles can start to take place as the elongated particles break up into smaller fragments having a more globular appearance [16, 24, 25].  2.3.4  M n in Solid Solution  The level o f M n i n solid solution after the homogenization process for A A 3 x x x alloys can influence the properties o f the final gauge sheet. During the homogenization process, the amount o f M n i n solid solution w i l l initially drop as M n precipitation rates peak at temperatures between 4 7 5 ° C and 500°C depending on the alloy composition [4, 16, 20]. A s the temperature is further increased, the amount o f M n i n solid solution increases as the solubility o f M n i n the matrix increases. A t the end o f homogenization, however,  there is considerably less M n i n solid solution than  microstructure. M n depletion i n the solid solution takes place i n three ways: i)  Precipitation o f M n - r i c h dispersoids.  17  i n the  as-cast  ii)  Transformation o f A l ( F e , M n ) to a - A l i ( F e , M n ) S i : Total (Fe+Mn) i n A l ( F e , M n ) 6  2  3  6  is about 25 wt.% compared with 32 wt.% o f (Fe+Mn) i n a - A l i ( F e , M n ) S i phase 2  3  [4, 26, 27]. W i t h very little Fe available i n the matrix, it is mostly achieved through M n enrichment. iii)  Enrichment o f M n i n both dispersoids and eutectic constituents [10, 26, 27]. During homogenization, M n is known to diffuse into the particles thereby  increasing the % M n i n both A l ( F e , M n ) [10] and a - A l i ( F e , M n ) S i [26, 27] particles. 6  2  3  Hence, though the ( F e + M n ) / A l i ( F e , M n ) S i ratio o f the particles stays fairly constant, the 2  3  Mn/(Fe+Mn) increases with an increase i n homogenization temperature or decrease i n cell size o f the ingot.  18  2.4  Effect of Composition on the As-Cast and Homogenized Microstructure  The exact composition o f the alloy w i l l have a significant influence o n both the structure that forms  during solidification o f the ingot as w e l l as some o f the  microstructural changes which occur during the homogenization process  2.4.1  Silicon Commercial producers o f beverage canbody stock presently have between 0.12%  and 0.30% S i i n their alloys. Varying silicon i n the ingot affects the developing microstructure i n the following ways i)  Effect on constituent particles: an increase i n percentage o f Si i n the alloy gives a  higher percentage o f cc-Ali2(Fe,Mn)3Si i n the as-cast structure and also promotes the transformation  of  Al (Fe,Mn) 6  to  a-Ali (Fe,Mn)3Si 2  during  various  stages  of  homogenization treatment [4, 28] (Table 2.4). ii)  Effect on M n solubility: S i enhances the precipitation o f M n as w e l l as other  constituent elements such as F e and M g . Increasing S i reduces the amount o f M n i n solid solution i n the as-cast as well as homogenized ingot [4, 28]. S i reduces M n content i n solid solution by causing a higher proportion o f a-Ali2(Fe,Mn)3Si to form, thereby removing more M n from solid solution (Table 2.4). W h i l e A l ( F e , M n ) contains about 25 wt% (Fe + M n ) , o c - A l i ( F e , M n ) S i 6  2  contains about 32 wt% [26]. 19  3  iii)  Effect on Dispersoids: A s mentioned earlier, the dispersoids formed are mostly a -  Ali2(Fe,Mn)3Si, but i f S i is below 0.08%, the dispersoids w h i c h form are o f the type Al (Fe,Mn). 6  2.4.2  Iron  The effect o f iron is very pronounced i n an A l - M n - S i alloy. E v e n a very small addition o f Fe to the A l - M n alloy greatly enhances the number o f constituent particles which form facilitating the M n precipitation and preventing the formation o f the ternary metastable phases ( A l M n F e ) usually encountered i n the A l - M n - F e alloys [29]. N o x  y  z  ternary precipitation phase appears, but rather F e atoms substitute for M n atoms i n A l e M n . D u r i n g homogenization, M n enters Al6(Fe,Mn) until the phase has achieved its maximum stability at Al<s(Feo.5,Mno.5) [29]. Iron additions influence the number o f constituent particles w h i c h form and this w i l l have a direct impact o f the recrystallization kinetics and texture w h i c h forms i n hot rolled sheet as particle stimulated nucleation (PSN) becomes more dominant.  2.4.3  Copper  Copper has a very similar though milder effect than iron. B o t h copper and iron promote M n precipitation i n aluminum alloys. The effect o f 0.06 F e and 0.20 C u on the loss of M n supersaturation i n A l - 1 M n alloys at 550°C is shown i n Figure 2.3 [30].  20'  10  9  -A]  17  u  C3  Oi  >-  10 pi  »  >  10  Exposure Time (Hrs) Figure 2.3 - Effect of 0.06% Fe and 0.02% Cu on the loss of supersaturation in Al-1 M n Alloys at 550°C [30]. 2.4.4  Magnesium  Magnesium (Mg) contributes to the overall strength of the alloy through solid solution hardening and work hardening [31]. AA3004/AA3104 have often been considered as non-heat-treatable A l - M n alloys to which M g is added to improve its work hardening characteristics. Furthermore, a comparison of similar AA3xxx alloys indicates that the M g containing AA3104 requires substantially less homogenization time than the equivalent AA3003 without M g [32]. This has been explained by the role of Mg^Si matrix precipitation acting as nucleation sites for a-dispersoids during the early stages of ingot heating [4,16, 32].  21  2.4.5 Manganese M n i n canbody stock provides large hard particles effective i n die-cleaning during drawing operations. It also imparts some strength to the final gauge sheet through solution hardening and excellent corrosion resistance [33]. V a r y i n g the M n content o f the ingot from 0.6wt% to 0.9wt% w i l l give a higher M n in solid solution throughout the homogenization process [26]. However, the resulting microstructure, i n terms o f volume fraction and size distribution o f constituent particles and precipitation and dissolution o f dispersoids, is not significantly different.  22  2.5  Effect of Homogenization Process Parameters on Microstructure Although studies have been completed over the past decade [10, 24, 25, 34] in an  effort to understand the effect o f the homogenization profile on the microstructural evolution, the ability to predict an optimum homogenization profile for a given starting structure and alloy composition has not yet been developed. In fact, currently, industry uses a wide variety o f homogenization profiles to process the same alloy even though the starting structure and chemistry may be very similar. Industrial homogenization profiles can vary significantly i n terms o f the soak temperature and time as well as the overall homogenization profile (i.e, a second stage soak at a lower temperature after the primary soak is complete). The final microstructure o f the as-cast material after homogenization is chiefly dependent on the type o f thermal profile (temperature and time) it has undergone. The parameters that can be varied i n any thermal profile are: •  heat-up rate,  •  soak temperature,  •  soak time, and  •  cool down rate. Each o f these process parameters can affect the microstructural evolution during  the homogenization process..  23  2.5.1  Heat-up Rate  Currently, industrial heat-up rates during homogenization o f ingot material typically range from around 45°C/hr (soaking pit furnace) to around 70°C/hr (walking beam furnace)[35]. T o date, almost no research has been done on the effect o f heat-up rate on the microstructure evolution or properties o f homogenized canbody stock material. Sheppard [36] was the only researcher who studied the effect o f heat-up rate on the microstructure evolution o f a 7xxx series alloy. H i s findings showed that a slower heat-up rate to homogenization temperature had no influence on the constituent particles however that it did affect the dispersoid size and distribution, with the slower heating rate giving more dispersoids with a more homogeneous  distribution. H i s results were  explained using classical precipitate nucleation theory; a slower ramp provided a longer period within the temperature zone corresponding to maximum nucleation rates hence more dispersoids could precipitate. Later during the soak period at the homogenization temperature, precipitate growth and precipitate nucleation become competing processes whereby pre-existing nuclei grew at the expense o f further nucleation.  2.5.2  Soak Temperature  Currently, the homogenization soak temperature for A A 3 x x x alloys used in the industry lies between 560°C and 610°C [35]. The ingots are heated to the soak temperature and held at this temperature for a period o f time to achieve the desired microstructure. Soak temperature also known as the homogenization impacts most aspects o f the microstructure such as:  24  temperature,  i)  percentage M n i n solid solution,  ii)  constituent particles, and  iii)  dispersoids.  Goodrich  [34]  investigated  the  effect  o f soak  temperature on  the  final  homogenized structure and percent M n left i n solid solution. He estimated the M n in . solid solution using resistivity measurements (Figure 2.4).  6.Sr  4" 100  " 150  200  250  300  330  400  450  SOO  S50  600  Tenparatu-* QU)  Figure 2.4 - M n precipitation vs. time and temperature for an as-cast 3104 alloy [34].  25  A s shown, after around 550°C, the difference between resistivity values ( % M n i n solid solution) for 1 hr and 200 hrs is not very different indicating that M n equilibration rates are sufficiently faster at those temperatures. It can also be inferred from the graph that equilibrium M n i n solid solution rises with an increase i n temperature above 4 8 0 ° C . This temperature can be referred to as the peak M n precipitation temperature since at this temperature there is least amount o f M n i n solid solution.  This is due to the higher  solubility o f M n i n A l at higher temperatures [4, 16, 34]. Constituent particles undergo a variety o f changes as the homogenization temperature is raised over the peak M n precipitation temperature such as: (i)  Al6(Fe,Mn) to a-Ali2(Fe,Mn)3Si transformation.  (ii)  Particle coarsening and spherodizing.  (iii)  M n enrichment i n particles.  The effect o f temperature on Al<5(Fe,Mn) to a-Ali2(Fe,Mn)3Si transformation can be estimated from Table 2.4. It is clear that an increase i n temperature increases the %ccA l i ( F e , M n ) S i up to 5 9 3 ° C for both S i levels. A t higher temperatures o f 6 3 2 ° C and 2  3  638°C, where incipient melting takes place, transformation kinetics become greatly enhanced and a reduction i n the %a-Ali2(Fe,Mn) Si is seen i n the lower S i alloy. 3  The constituent particles have shown a rounding o f their edges (spherodization), coarsening and break up o f the three-dimensional dendritic network when held at higher homogenization temperatures [22, 37,27]. W i t h an increase i n homogenization temperatures, the M n levels i n particles is known to stabilize at different levels (Figure 2.5) [26]. Sun [26] suggested that it was  26  possible to determine the homogenization temperature from the composition of the eutectic constituents.  14.0  Sample B  Sample A  10.54 580.  — i  1  1  590 600 610 620 Preheating Temperature, C  630  Figure 2.5 - Mn content in cc-phase constituents as a function of homogenization temperature [26].  Quite a few researchers [10,12, 16, 26, 27] have found that dispersoids that form during ingot heating become fewer and coarser as homogenization temperature increase.  27  There are fewer and coarser dispersoids present i n a sample homogenized at 6 0 0 ° C than one homogenized at 560°C (Figure 2.6) [27].  4j  Figure 2.6 - Examples o f secondary precipitates i n T E M thin foil specimens o f A A 3 0 0 4 alloy at (a) 600°C and (b) 560°C [27].  Bolingbroke [16] has characterized the increase i n size o f the dispersoids with increase i n temperature during heat-up (Table 2.5). It is clear that dispersoid coarsening is a direct function o f temperature.  28  Table 2.5 - Change i n a - A l i ( F e , M n ) S i dispersoid mean size during preheat heat-up and 2  3  soak [16].  Preheat treatment  Average a-Ali2(Fe,Mn)3Si precipitate size (nm)  As-cast  -  Heat-up to 400°C  35  Heat-up to 4 5 0 ° C  47  Heat-up to 500°C  66  Heat-up to 600°C  108  Heat-up to 6 0 0 ° C , cooled to 500°C at 30°C/hr  136  Heat-up to 6 0 0 ° C , soaked for 8 hrs and cooled to 5 0 0 ° C at 30°C/hr  150  2.5.3  Soak time  The soak time at the high temperature allows the reaction kinetics to proceed. B o t h M g and S i have a high diffusion coefficient and w i l l equilibrate even as the temperature reaches the homogenization temperature. M n on the other hand, has a low diffusion coefficient and it is the rate determining step for the transformation o f Al6(Fe,Mn) to ot-Ali2(Fe,Mn)3Si and dispersoid precipitation during the homogenization treatment. A s the  homogenization temperature increases,  29  the  time  required  for  homogenization decreases [12, 38]. This is partly because the M n diffusion rates become faster and also partly because the equilibrium M n in solid solution is higher at higher temperature. Besides M n equilibration, time at homogenization temperature also allows for precipitate coarsening. For example, there is no change in M n level for an ingot homogenized at 5 7 7 ° C for 14 hrs and for 26 hrs but T E M microstructures shows that a dispersoids are almost twice the size after 26 hrs at 577°C [12](Figure. 2.7)  a  0.5 um  b  0,Sj,m  Figure 2.7 - T E M photomicrographs showing the cc-dispersoids i n ingot homogenized at 577°C for (a) 14 hrs and (b) 26 hrs, followed by water quench [12].  2.5.4  Second Stage Soak  Traditionally, industrial homogenization cycles consisted o f heating the ingot to a prescribed temperature and holding it there for a certain length o f time followed by air  30  cool to the hot rolling temperature. Recently, researchers have investigated the merit o f altering the traditional homogenization cycle to include a second soak period at a lower temperature on the microstructural evolution i n the ingot [26, 39]. Hence, a typical two stage homogenization profile consists o f heat-up, hold at high temperature (560°C 610°C) followed by controlled cooling to a lower temperature (450°C - 500°C) and a soak at this temperature prior to hot rolling. One major difference i n microstructure after a second stage soak is the lower M n i n solid solution compared to a traditional homogenization profile [39]. During the second stage cool, M n precipitation occurs i n a controlled manner such that the wt % M n i n solid solution is lower and at the same time the number o f dispersoids is not increased [12].  31  2.6  Models  for  Microstructural  Changes  during  Homogenization Very  few  mathematical  models  have  been  developed  to  predict  the  microstructural evolution i n an aluminum ingot as a function o f time and temperature during homogenization. Only two models were found in the literature that had been developed  to  model microstructural  changes during homogenization  of  AA3xxx  aluminum alloys. A mathematical model is an attractive way to describe microstructural evolution, because the response o f the microstructural features such as dispersoids and constituent particles to time, temperature and composition can be quantified. To date, two researchers have attempted to model dispersoid formation during homogenization o f A A 3 x x x alloys [34,40]. Goodrich [34] used an A v r a m i type formulation (equation 3) to model the amount o f M n i n solution for A A 3 1 0 4 . The precipitation or dissolution rate constants i n the model were dependent on temperature and the state o f the preheat cycle and these were determined experimentally from isothermal treatments o f an ingot. The fit was apparently quite good with A v r a m i exponents between 0.2-0.3.  X = l-exp(-kt )  (3)  n  where X is the fraction o f M n precipitated, k is the rate constant,  32  t is the t i m e i n hours, n is related to the type o f reaction.  X is expressed as:  (4)  X = ^ l c.  - c,  W h e r e : C(t) is the d i s s o l v e d M n at a n y t i m e t, Cj is the initial concentration o f d i s s o l v e d M n , C  e  is the " p r a c t i c a l " M n s o l u b i l i t y for this alloy.  G o o d r i c h ' s m o d e l was  never validated u s i n g  industrial m e a s u r e m e n t s  so it is  d i f f i c u l t to c o m m e n t o n the a c c u r a c y o f the predictions.  S u n i [40, 41] d e v e l o p e d a p h y s i c a l l y  b a s e d m o d e l to p r e d i c t the  dispersoid  evolution during homogenization o f A A 3 0 0 3 and A A 3 0 0 4 a l u m i n u m alloys. H i s m o d e l c o m b i n e d the processes o f nucleation, g r o w t h and coarsening so that p a r t i c l e size, density and  volume  fraction  of  dispersoids  could  be  predicted.  Model  parameters  were  determined u s i n g industrial A l c o a data a n d the m o d e l predictions w e r e t h e n v a l i d a t e d u s i n g industrial data a n d i n general p r e d i c t e d sizes c o m p a r e d w e l l to measurements.  33  2.7 Summary Industrial can making is a complex process comprising various steps with many variable process parameters. Homogenization is the starting step i n the processing o f the ingot to the final sheet. It is important as the microstructure o f the homogenized ingot can influence the strength, formability and the texture o f the final gauge sheet for the canbody stock. Although studies have been completed over the past decade [10, 24, 25, 34] in an effort to understand the effect o f the homogenization on the microstructural evolution, the ability to predict an optimum homogenization profile for a given starting structure and alloy composition has not yet been developed. Although some studies have been done investigating the effect o f temperature and time o f homogenization on the microstructures o f A A 3 x x x aluminum alloys [21-26], there is almost no work studying the effect o f heatup rate. A s such, there is no comprehensive study on the combined effect o f the various homogenization parameters (heat-up rate, temperature, time, cool-down rate) on the microstructure evolution.  34  3.  Scope and Objectives  3.1  Scope  The goal o f homogenization process is not only to preheat the material to the desired hot rolling temperature but also allow segregated areas i n the material to equilibrate. Currently i n industrial practice, a wide variety o f homogenization practices exist i n terms o f soak temperature, soak times and heat-up rate. In addition, newer cycles involving a second stage cool followed by a soak at a lower temperature have started to be investigated. To date, a comprehensive study quantifying the effect o f homogenization practice on the microstructural evolution has not been done. Such a study w i l l be o f benefit to aluminum sheet metal producers i n that it w i l l address the balance required i n terms o f developing the most economic practice i n conjunction with one that obtains an optimum rmcrostructure for hot rolling. Although some studies have been done investigating the effect o f temperature and time o f homogenization on the microstructures o f A A 3 x x x aluminum alloys [21, 26], there is almost no work studying the effect o f heat-up rate. A s such, there is no comprehensive study o n the combined effect o f the various homogenization parameters (heat-up rate, temperature, time). A s a result, the purpose o f this study was to examine various homogenization profiles and their effect o f on the microstructural development during homogenization o f an A A 3 1 0 4 aluminum alloy.  35  3.2  Objectives  The objectives o f the present study are as follows: to simulate typical industrial homogenization profiles i n a laboratory furnace, to  characterize  the  effect  o f homogenization  process parameters  microstructural evolution o f aluminum alloy A A 3 1 0 4 .  36  on  4.  Experimental  4.1  Start Material  The experimental work was undertaken on an as-cast commercial A A 3 1 0 4 aluminum alloy. The material was received from A l c a n International and was taken from an ingot that had been previously D . C . cast. The chemical composition o f the alloy as supplied by A l c a n and verified through low magnification E D X / S E M is shown i n Table 4.1.  Table 4.1 - Composition o f A A 3 1 0 4 alloy used.  Alloy  M g (wt%)  M n (wt%)  Fe (wt%)  S i (wt%)  C u (wt%)  AA3104  1.20  0.87  0.40  0.20  0.18  Microstructure and composition can vary dramatically across an ingot and the starting  microstructure  can greatly  influence  the  microstructure  evolution  during  homogenization [10, 11, 12]. To ensure homogeneity i n the starting specimens, samples were cut 7-15cm from the ingot edge and were tested to ensure similar initial conductivity values. This ensured that conductivity (a measure o f the amount o f solute i n solid solution) and cell size differences were minimized.  37  4.2  Homogenization Experiments  4.2.1  L a b o r a t o r y furnace  The samples were homogenized i n a programmable controlled laboratory furnace. To  ensure accurate temperature control and homogeneity within the furnace,  an  aluminum tube was w e l l fitted into the furnace with the controlling thermocouple inside the aluminum tube (Figure 4.1). Measurements o f thermal profiles inside the modified furnace using instrumented samples indicated that the temperature variations were ± 2°C between the samples (Appendix A ) .  K-thermocouple  Aluminum Tube A l u m i n u m Tube Samples Furnace  \ Samples  Programmable temperature controller  b ) Cross-sectional view  a ) Overview  Figure 4.1 - Laboratory furnace set-up: a ) overview and b ) cross-sectional view.  38  4.2.2 Homogenization profiles A wide range o f homogenization profiles were simulated i n the laboratory furnace so that the effect o f heat-up rate, soak temperature, and soak time on the microstructural evolution could be quantified (Table 4.2).  Table 4.2 - Range o f homogenization process parameters used i n the study. Process Parameter  Range  Heat-up rate  37- 279°C/hr  Soak temperature  550-610 °C  Soak time  0 - 1 0 hrs  In addition, the effect o f cool down to the rolling temperature after  the  homogenization treatment on the microstructure evolution was also studied. Samples w h i c h were taken for microstructural characterization were quenched immediately representative  from  the  homogenization cycle  to  ensure that  the  structure  o f what is seen during homogenization. Figures 4.2-4.4 show  was the  experimental homogenization cycles used as w e l l as where samples were taken for microstructural characterization.  39  0  5  10  15  20  25  30  Time (hrs)  Figure 4.2 - Experimental homogenization profiles with varying heat-up rates (dots indicating where samples were taken for microstructure characterization). 700  0  5  10  15  20  25  Time (hrs) Figure 4.3 - Experimental homogenization profiles with varying soak temperatures (550°C,580 Cand610°C). o  40  After an ingot is homogenized i n industry, it cools down to the temperature at which hot rolling starts, also known as the lay-on temperature. Due to cool down, microstructural changes, especially i n terms o f M n in solid solution are expected. Therefore, experiments that were conducted at different homogenization temperatures were cooled down to 5 0 0 ° C from the soak temperature in about half an hour (Figure 4.4). In order to cool the samples down to 500°C, the furnace was turned off and the door was partially left open.  700  Time (hrs) Figure 4 . 4 - Experimental thermal profiles used for samples homogenized at different soak temperatures.  41  4.3  Manganese Retained in Solid Solution  Electrical conductivity measurements were used to determine the amount o f solute i n solid solution. The sample conductivity was measured at 2 0 ° C as a percentage o f International Annealed Copper Standard ( % I A C S ) [18] using a Verimet 4900C eddy current probe. The readings (in % I A C S ) could then be directly used to estimate the % M n in and out o f solid solution. This is because M g , a major alloying element, has a much lower effect on conductivity than M n (Table 4.3) [18], almost all the Fe stays out o f solution during the treatment, and S i and C u levels i n the alloy are too l o w to significantly affect the conductivity readings. Resistivity values could be obtained from conductivity measurements by the simple equation [42]: p =172.41/(%IACS)  (5)  where p is resistivity i n p.Qcm. Conversion o f resistivity values to wt.% M n i n solid solution can then be done using the graph shown i n Figure 4.6 [43].  Table 4.3 - Composition o f A A 3 1 0 4 and effect o f elements on resistivity [18]. A A 3 1 0 4 composition  Average increase i n resistivity per  Element  (wt%)  wt% i n solution (p.Qcm)  Mn  0.87  2.94  Mg  1.20  0.54  Fe  0.4  2.56  Si  0.2  1.02  42  1.0 12 1.4 MoC%) in SS  Figure 4.6 - Resistivity o f aluminum alloy 3004 as a function o f the percentage o f M n i n solid solution [43].  4.4  Microstructure Characterization / Metallography  4.4.1  Grinding and Polishing Procedure  Samples were cold mounted i n an epoxy resin. A l l samples were polished using the Buehler Ecomet I V Polisher/Grinder which is a ten mount automatic polisher. The following schedule (Table 4.4) was standardised after trial and error and consultation with Tech-Met C a n a d a :  43  Table 4.4 - Grinding and polishing procedure followed.  Abrasive Abrasive type  Lubricant  Size  Pressure  Time  (psi)  (min)  Additional comment  Repeated step till surface S i C paper  180 grit  Water  20  2  S i C paper  320 grit  Water  40  2  40  3  40  3  30  2  was flat Changed paper after 1 m i n  Metadi  Diamond suspension on  6um  Texmet 1000 cloth Diamond suspension on  lum  Texmet 1000 cloth Colloidal S i 0  2  fluid Metadi fluid  0.05um  suspension  4.4.2  i f murkiness not noticed  Washed w i t h water for last 10 sees  Etching  The chemical etchant selected had to adequately distinguish between the two predominant  constituent phases namely  a-Ali (Fe,Mn)3Si and 2  literature, two etchants were found to do this, 10% H P 0 3  4  AJ.6(Fe,Mn).  From  at 2 0 ° C [16], and H S 0 2  4  at  80°C [12]. H3PO4 was tried and found to distinguish the two phases adequately for image analysis. Quantification o f dispersoids was also attempted, and H3PO4 was effective i n a qualitative viewing o f dispersoids. 0.5 % H F , an etchant used i n literature [18] was also tried but was unsuccessful.  44  Samples were etched with 10% H3PO4 for 8 minutes at room temperature to distinguish between the second phases (a-Ali2(Fe,Mn) Si and A l ( F e , M n ) ) . The a3  6  Ali2(Fe,Mn)3Si etched considerably darker while Al6(Fe,Mn) was grey when viewed optically (Figure 4.7). The differentiation was confirmed by analysis using E D X / S E M (Figure 4.8) where only the darker particle showed a S i peak confirming it to be ocAli (Fe,Mn) Si. 2  3  oc-Ali (Fe,Mn) Si (etched black) 2  3  &Si (black)  M  Al (Fe,Mn) (etched grey) 6  Figure 4.7 - As-cast microstructure (1000X) showing the various second phase particles.  45  17-Sep-19S8 Dark  14:41:37 200 200  Preset-  Particle  Vert. 5000 counts Quantex> fi 01  Dlsp=  Elapsed^  1  sees sees  SI  J n  4-  rg  0.000  B Range-  n  ns  -  1 0 . 2 3 0 keV  IB  Ta  TT  Integral  8  1 0 . 1 1 0 -> 306306  -  Particle etched black showing A l , S i , M n and Fe peaks.  17-Sep-1998  14:46:41  Vert5000 c o u n t s Quantex>  Disp-  286 20B  PresetE l apsedc  Greu P a r t i c l e 1  sees sees  • •  01  n 4-  0.000  re  IT  Range-  nm  TT  rs~  10.H38 keV  Integral  TO 0  °  V3  1 0 . 1 1 0 -t 338423  Particle etched grey showing A l , M n and Fe peaks. Figure 4.8 - E D X / S E M results confirming the etching difference between cxA l i ( F e , M n ) S i and Alo<Fe,Mn) particle. 2  3  46  4.4.3  Electron Probe Micro Analysis ( E P M A ) Line Scans  It is known that the as-cast microstructure has microsegregation o f eutectic elements such as M g and M n , which is eliminated during homogenization. To find out the temperature at w h i c h M g and/or M n diffuses so that microsegregation is reduced across the cell during homogenization, E P M A line scans were taken using the Scanning Electron Microscopy coupled with W D X .  4.4.4  Image Analysis  Quantitative metallography was done on a C-Imaging Systems optical image analyzer. The coarse eutectic intermetallic particles were analyzed at (32X) objective magnification  whereas  the  dispersoids  were  distinguished  at  (50X)  objective  magnification. A s there was a lot o f microsegregation seen i n the as-cast and partially homogenized samples, data was collected over a large number o f fields till  the  cumulative averages o f the data collected stabilized as shown i n Figure 4.9. For the coarse constituent particles, data was collected from 40 to 80 fields depending on how quickly the average stabilized, with the stabilized average being used for subsequent analysis. E v e n so, due to the high variations i n the readings between individual fields, the standard error was quite large  47  0%  1  1  1  1  1  1  11  1  21  1  1  31  '  N o . o f fields  Figure 4.9 - Variation o f the cumulative average per field with the measured quantity per field for volume fraction second phase.  4.5  As-Cast Characterization  Initial material characterization involved quantifying the as-cast microstructure and manganese ( M n ) retained i n solid solution. Due to the large variances i n as-cast microstructure (cell size and microsegregation) reported i n the literature, samples were chosen from 5 to 25 c m from the surface o f the ingot to minimize differences i n the starting composition and microstructure. Table 4.5 gives the as-cast characterization i n terms o f wt% M n i n solid solution converted from conductivity measurements and second phase quantification.  48  Table 4.5 - As-cast characterization between 5 - 25 cm from surface.  Wt.% M n in  Second phase V o l .  a-Ali2(Fe,Mn)3Si  Mg Si  solid solution  Fraction (%)  (% second phase)  (% second phase)  0.66  2.75  -16  -11  2  Total second phase volume fraction included a-Ali (Fe,Mn)3Si, Ale(Fe,Mn) and 2  M g S i (Figure 4.7). The majority o f the second phase comprised o f Al6(Fe,Mn) which 2  amounted to about 73% o f the total second phase. The M g S i was seen as a dark 2  agglomeration o f circular particles. The a-Ali (Fe,Mn)3Si also etched dark and was 2  usually present as part o f the eutectic interdendritic structure. The Ale(Fe,Mn) phase and was grey when viewed optically.  49  5.  Results and Discussion The experimental results o f this study are outlined i n this section, describing the  qualitative and quantitative effect o f process parameters on the microstructural evolution during homogenization.  The effect o f the homogenization cycle on the final sheet  properties i n terms o f microstructure and texture are also described.  5.1  Microstructure Evolution during Homogenization  The effect o f homogenization cycle on the rm^rostructural evolution was assessed varying process parameters such as heat-up rate, soak temperature and soak time, and characterizing the microstructure o f the samples at various points i n the homogenization cycle. Microstructure evolution during the homogenization cycle was characterized by conductivity measurements (to assess the % M n i n solid solution) as w e l l as image analysis to quantify the percentage o f different constituent phases and examine the dispersoid distribution.  5.1.1 Effect of Heat-up Rate The effect o f heat-up rate to homogenization temperature on the microstructural evolution is described i n this section i n terms of: •  the % M n i n solid solution (Figure 5.1),  •  the transformation o f Ale(Fe,Mn) to cc-Ali2(Fe,Mn) Si (Figure 5.2),  •  evolution o f dispersoids(Figure 5.3) as w e l l as,  •  microsegregation.  3  50  5.1.1.1  Effect on M n in Solid Solution  M n being supersaturated i n solid solution i n the as-cast condition starts to come out as the temperature rises during heat-up. Depending on the heat-up rate, M n precipitation is affected i n terms o f its initiation temperature and the amount that comes out (Figure 5.1). W i t h increasing heat-up rate, M n precipitation starts at a later temperature and less M n precipitates out and dissolves back i n during heat-up. In addition, the temperature at which the minimum level o f % M n i n solid solution is found to increase as the heat-up rate is increased (Table 5.1).  -±-279°C/hr  -«-lll C/hr 0  -*-70°C/hr  -»-47 C/hr 0  Time (Hrs) Figure 5.1 - Effect o f heat-up rate on % M n i n solid solution.  51  -*-37°C/hr  Table 5.1 - M i n i m u m M n i n solid solution and corresponding temperature for different heat-up rates. Heat-up Rate  M i n i m u m % M n i n solid  Temperature o f minimum  (°C/hr)  solution (wt %)  % M n i n solid solution (°C)  37  0.32  475  47  0.33  500  70  0.35  500  111  0.38  525  279  0.43  550  The following observations can be made from the results: i)  The temperature which attains the lowest M n i n solid solution (maximum  precipitation rates) increases as heat-up rate increases ranging between 4 7 5 ° C and 5 5 0 ° C . This point corresponds to the rate o f maximum precipitation rates [4, 16, 20]. A s seen i n the literature, the temperature o f maximum precipitation rates for M n is ~ 4 7 5 ° C . This corresponds to the temperature measured i n our slowest heat-up rate o f 37°C/hr. ii)  The amount o f M n that precipitates out o f solid solution at the temperature o f  maximum precipitation rates decreases as the heat-up rate increases. This is expected based on the classical precipitation theory as a faster heat-up rate means there is less time at temperature to allow M n precipitation. iii)  The total amount o f M n that precipitates out o f solid solution at the end o f the 3 hr  hold at 580°C is almost identical for each heat-up rate. This was again expected since there was enough time during the soak period, to allow the precipitation o f M n into the  52  matrix until the equilibrium amount of Mn at the soak temperature was left in solid solution. 5.1.1.2  Constituent Particles and percent a-Alj2(Fe,Mn)3Si  The shape and size of the constituent particles was qualitatively observed after a three hour hold at the homogenization temperature of 580°C for each of the heat-up rate (Appendix B) and was not seen to be affected by the heat-up rate. The Al6(Fe,Mn) -> cc-Ali2(Fe,Mn)3Si transformation was measured as percent aAli2(Fe,Mn) Si of total second phase and plotted for a three hour hold at 580°C for the 3  various heat-up rates (Figure 5.2). Although heat-up rates do not have a significant effect on percentage transformation, there is a trend showing a higher a-phase at faster heat-up rates. Due to the high standard error during measurement, 95% confidence interval error bars suggest that no conclusive statement can be made about the increasing trend indicated on the graph.  53  60%  50%  CD co cd  40%  30%  20% 37°C/hr  47°C/hr  70°C/hr  112°C/hr  279°C/hr  Figure 5.2 - Effect of heat-up rate on % a-Ali (Fe,Mn) Si after a 3 hour hold at 580°C. 2  5.1.1.3  3  Evolution of Dispersoids  Microstructures were taken from samples at a soak temperature of 580°C after a three hour hold for each of the heat-up rates (see Appendix B for a list of pictures for each profile). Dispersoids (seen as the fine dots in the microstructure) formed from a faster heat-up rate were larger and fewer than ones formed from a slower heat-up rate profile (Figure 5.3).  54  20um b) Figure 5.3 - Microstructures o f samples for 580°C soak temperature after 3 hour hold for: a) 279°C/hr and b) 47°C/hr.  55  M n precipitates during heat-up i n the form o f a-Ali2(Fe,Mn) Si dispersoids with 3  the M n precipitation kinetics bearing a direct relation to the dispersoid evolution [4, 19]. Since there has been no study on heat-up rate for A A 3xxx alloys, no results or theory was found i n literature to confirm or explain these findings. Evaluating % M n i n solid solution at specific temperatures during heat-up revealed that there was more precipitation i n the 47°C/hr heat-up profile than the 279°C/hr profile (Figure 5.4). It was reasoned that due to the small time range during heat-up for the faster heat-up rate, the M n precipitation rates were lower causing less dispersoids to nucleate. This was confirmed by the smaller minima o f M n i n solid solution, as illustrated i n Figure 5.1. The growth o f the dispersoids, however, continued during the soak period and the fewer dispersoids grew larger with the available M n i n solid solution.  56  Figure 5.4 - Micro-structures o f samples at 450°C during heat-up for: a) 279°C/hr and b) 47°C/hr.  57  5.1.1.4  Removal of Microsegregation  Electron Probe M i c r o Analysis ( E P M A ) scans carried out on samples at various temperatures during heat-up showed that M g and M n microsegregation was removed from the samples at different temperatures (Table 5.2) (for micrographs see Appendix C ) . M g microsegregation also referred to as M g coring [33] is high i n the as cast structure and is considerably reduced by the time the temperature reaches 500°C for the 47°C/hr profile and by 580°C for the 279°C/hr profile.  Table 5.2 - M n and M g segregation results for 47°C/hr and 279°C/hr profile at various temperatures during heat-up.  Condition  Heat-up profile 279°C/hr  Heat-up profile 47°C/hr Mg  Mn  Mg  Mn  As-cast  High  Medium  High  Medium  400°C  High  Medium  High  Medium  500°C  Medium  Medium  High  Medium  580°C  Low  Low  Low  Medium  Low  Low  Low  Low  580°C  3hr  58  5.1.2  Effect of Soak Temperature and Soak Time  5.1.2.1  Effect on M n in Solid Solution  A s the soak temperature increases, the amount o f M n i n solid solution at the soak temperature increases for both a fast (279°C/hr) and slow (47°C/hr) heat-up rate (Figure 5.5). This is i n tune with the literature since M n solubility increases with an increase i n temperature [4, 26]. The effect o f soak time at 550°C, 580°C and 610°C can also be estimated from Figure 5.5 for both heat-up rates. The results can be analyzed for two distinct regions: i)  Between 0 and 0.5 hrs - most profiles show an increase i n the M n i n solid solution  especially for the higher temperature profiles (580°C and 610°C). This could be attributed to the fact that at 0 hr hold even though the samples have reached the soak temperature the M n re-dissolution kinetics are still catching up to reach the equilibrium amount o f M n i n solid solution for that temperature. ii)  Between 0.5 hrs and 10 hrs: The results for this range show a slight drop i n % M n  i n solid solution with increasing soak time at higher soak temperatures (~600°C) whereas a steady amount or even a slight rise can be seen at lower soak temperatures (~550°C). The decrease i n % M n i n solid solution during the soak period at higher homogenization temperatures have also been noted by other researchers [12, 19, 21]. The reason for this could be that at higher soak temperatures other processes competing for M n become more favorable such as:  59  Al6(Fe,Mn) -> a-Ali2(Fe,Mn) Si: A higher % a - A l i ( F e , M n ) S i is reported at longer 3  2  3  soak periods [4] (Figure 5.7) and there is a higher % M n i n a-Ali2(Fe,Mn) Si as 3  compared to Al6(Fe,Mn), and % M n i n constituent particles: A higher % M n is seen i n both a - A l i ( F e , M n ) S i and 2  Al6(Fe,Mn) at higher soak temperatures [12, 26, 27].  60  3  • 550°C  ^580°C  H610°C  0.6  0  0.5  3  10  Soak time at temperature (hrs) b) Figure 5.5 - M n i n solid solution at different soak temperatures at a) 47°C/hr and b) 279°C/hr. 61  Constituent Particles and Percent a-Ali (Fe,Mn) Si  5.1.2.2  2  3  Particle coarsening occurs at high temperature hold during homogenization [10, 12, 24] with more coarsening at higher temperatures. However, in the present study, a significant coarsening was not observed in the particles. The particles at 610°C, however, were seen to be more rounded than at 550°C (Figure 5.8). With  higher  soak  temperatures,  the  Al6(Fe,Mn)  to  a-Ali2(Fe,Mn) Si 3  transformation is also enhanced such that the percentage of a-Ali2(Fe,Mn) Si is higher 3  for higher soak temperatures (Figure 5.6). This trend is seen at both slow and fast heat-up rates with the total %a-Ali (Fe,Mn) Si being higher at the faster heat-up rates. 2  3  147°C/hr 3 hr hold  22  • 279°C/hr 3 hr hold  550  • As cast  610  580  Temperature (°C) Figure 5.6 - Effect of soak temperature on %a-Ali (Fe,Mn) Si after a soak period of 2  three hours.  62  3  A s the soak time is increased the percentage a-phase increases and eventually levels off after a soak time o f 3 hours is reached. The leveling was more pronounced for higher soak temperatures (Figure 5.7). This could be attributed to the unavailability o f M n due to higher M n i n solid solution at the end o f 10 hrs for the higher soak temperature (Figure 5.5).  °550°C  ^580°C  S  610°C  60%  0  3  10  Soak time at temperature (hrs)  Figure 5.7 - Effect o f soak time on %oc-Ali2(Fe,Mn)3Si at various soak temperatures at a heat-up rate o f 47°C/hr.  63  5.1.2.3  E v o l u t i o n o f Dispersoids  The effect o f soak temperature on the evolution o f dispersoids can be estimated qualitatively from the microstructure inspection o f samples from 550°C, 580°C and 610°C soak temperature after three hour hold with a heat-up rate o f 47°C/hr (Figure 5.8). The dispersoids at 550°C are smaller and more i n number, the number decreasing and the size increasing with an increase i n the soak temperature till at 610°C there are very few dispersoids observed, and the ones present are very coarse. There are two main mechanisms that contribute to the observed results: i) Growth o f dispersoids is a diffusion-based process and as temperature rises from 550° to 610°C, the rates o f this process increase rapidly. ii) Increased solubility o f M n at higher temperatures lead to a dissolution o f the smaller dispersoids into the matrix thereby explaining the low number o f dispersoids seen i n the sample w h i c h had been homogenized at 610°C.  64  *.  20um  a) 550°C  20u:  b)580°C  65  /  '  f  ' 20um  c)610°C Figure 5.8 - Microstructures (1000X) o f samples after three hour hold with a heat-up rate o f 47°C/hr for soak temperatures of: a) 550°C, b) 580°C and c) 610°C.  5.1.2.4  Cool Down to Rolling Temperature  The effect o f soak temperature is seen to translate to a difference i n % M n i n solid solution even when the samples were cooled down to a similar temperature o f 500°C (Figure 5.9), with the samples which had been soaked at a higher temperature exhibiting a higher level o f M n in solid solution.  66  e - 550°C (46.5°C/hr)  10  580°C (46.5°C/hr) - * ^ 6 1 0 ° C (46.5°C/hr)  12  14  16  Time (hrs) Figure 5.9 - M n in solid solution for homogenization profiles of 550°C, 580°C and 610°C for three hour hold followed by cool down to 500°C.  During cool down, very little change takes place in the morphology or size of the constituent particles, the precipitating M n comes out enlarging the pre-existing dispersoids (Figure 5.10). Hence, the higher number of dispersoids in the 550°C sample (Figure 5.8) are also carried forward at the end of the cool down to 500°C.  67  18  b) 580°  68  / *  \ 1 %  %  < 20um  c) 610°C Figure 5.10 - Microstructures (1000X) o f samples after a cool down to 500°C for soak temperatures of: a) 550°C, b) 580°C and c) 610°C.  69  5.2  Influence  of  Homogenization  Microstructure  on  Texture/Microstructure after Hot-Rolling The size and shape o f the constituent and dispersoid phases formed i n the ingot during homogenization influence the size distribution o f the second phases i n the final gauge sheet [4, 12, 37, 45, 46] as well as the recrystallization kinetics and texture formation after hot rolling. During hot rolling, the eutectic networks o f constituents break down and depending on the rolling temperature, recrystallization o f the grains initiates during/after rolling [11, 12, 48]. The morphology and size distribution o f the constituent particles (a-Ali2(Fe,Mn)3Si and Al6(Fe,Mn)) affects the cube texture evolution due to particle-stimulated nucleation (PSN) during recrystallization after hot-rolling, diecleaning during ironing i n the bodymaker operation, and tear-off tendency [4, 12, 45, 48, 47]. The dispersoids, on the other hand, affect the texture b y retarding recrystallization [45, 46,48]. The morphology o f the constituent particles i n terms o f the percent  a-  Ali2(Fe,Mn) Si phase determines the die-cleaning efficiency o f the alloy during drawing. 3  The constituent a-Ali2(Te,Mn)3Si phase is relatively harder (Vickers micro-hardness o f 900-950 H v ) than the A l ( F e , M n ) phase, which has hardness o f 700-750 H v [4, 16, 22]. 6  Hence a certain fraction (greater than 40%) o f a-Ali2(Fe,Mn)3Si constituent particles is desirable for prevention o f galling (the seizure o f metal on the working surfaces o f the dies observed when the constituent particles are not hard enough). The majority o f the  70  manufacturers  i n the industry control the percent a-Ali2(Fe,Mn)3Si phase fraction  between 50 and 60% [23]. To minimize earing during the can drawing operation, adequate amounts o f cube texture (0/90° ears) are needed after hot rolling. However, large constituent particles can cause deformation zones to form around them during hot rolling, and during subsequent annealing, these deformed zones become nucleation sites for recrystallization. P S N results i n random texture, which reduces the effective cube content, and hence decreases the cube texture i n the final gauge sheet [45, 5]. Hence, it is not favorable to have a large number o f constituent particles. M n retained i n solid solution retards recrystallization during annealing after hot rolling [12]. It is desired to achieve a totally recrystallized microstructure before the final cold roll step to obtain a sufficient amount o f cube texture [39]. Hence, it is desirable to have as l o w a % M n i n solid solution after hot rolling as possible. The  number  and  size  distribution  of  dispersoids  that  evolve  during  homogenization are important parameters i n determining the amount o f cube texture formed. These dispersoids exert a retarding force on the low/high angle grain boundary, which can affect the kinetics o f recovery, recrystallization and grain growth. This effect is known as Zener drag [45] and can be expressed as: P z = (3Fv*y)/2r  (6)  Where F v is the volume fraction o f randomly distributed spherical particles o f radius r, and y is the grain boundary energy reported to be 324 m J / m [49] for aluminium. 2  F r o m equation (6) it can be deduced that the Zener drag is a function o f volume fraction  71  (Fv) and size r (radius) o f the dispersoids such that a large volume fraction o f small dispersoids increases Zener drag.  72  6.  Summary and Conclusions  During the manufacturing o f aluminum beverage cans from the ingot through the cold rolled sheet, it is important to control the microstructure and texture evolution as they w i l l control the final mechanical properties and performance o f the alloy during the can  making operation. Since the homogenization profile the ingot undergoes,  will  influence the starting microstructure prior to hot rolling, it is critical to characterize and quantify the influence o f various homogenization parameters on the microstructure evolution and determine the "optimum" homogenized microstructure.  6.1  Summary  The present research was primarily concerned with characterizing and quantifying the effect o f homogenization process parameters on the microstructural evolution o f an A A 3 1 0 4 aluminum alloy. The following observations can be made based on the study: i)  Heat-up rate to the soak temperature affected the amount o f M n precipitating out  and re-dissolving i n solid solution during heat up. However, the total amount o f M n that precipitated out o f solid solution at the end o f three hours at the soak temperature was almost identical for each heat-up rate. The variation i n M n during heat-up affected the size and shape o f the evolving dispersoids; the dispersoids became fewer and appeared coarser as the heat-up rate increased possibly due to the lower amount o f M n precipitation during heat-up. The shape and size o f the constituent particles was not  73  affected by the heat-up rate, but there was a slight increase i n the percent  a-  Ali (Fe,Mn)3Si seen for the faster heat-up rates. This may have occurred because o f the 2  extra M n that was available due to fewer dispersoids. A n increase i n heat-up rate also caused a delay i n reduction o f M g and M n microsegregation across the cell. ii)  A s soak temperature was increased, the amount o f M n i n solid solution increased  for each heat-up rate. The rise in M n i n solid solution was accompanied in parallel by a drop  i n dispersoids. Higher soak  temperatures  also resulted  i n a higher %oc-  Ali (Fe,Mn)3Si particles. 2  iii)  Increasing soak times also increased the %cx-Ali2(Fe,Mn) Si initially but started 3  tapering o f f for longer soak times o f 3 hrs as the equilibrium amount o f %cxAl] (Fe,Mn)3Si was approached for the soak temperature. Higher soak 2  temperature  profiles tapered faster due to the lower amount o f M n i n solid solution. M n i n solid solution was affected inversely with increasing soak times. A drop was seen i n % M h i n solid solution with increasing soak times between 0.5 to 3 hrs. Between 0 and 0.5 hrs the M n level still increased i n all profiles due to lag i n M n dissolution kinetics to reach the equilibrium amount o f M n for that soak temperature. Dispersoids on the other hand, were qualitatively assessed to have coarsened slightly without a noticeable increase i n their number. Since the microstructure after the homogenization process w i l l influence the microstructure and sheet properties that develop during and after hot rolling, it is important to understand the influence o f key microstructure features on the both the microstructure evolution during hot rolling as well as the final material properties. A summary o f desired microstructure characteristics after homogenization include  74  •  a fair amount o f %a-Al]2(Fe,Mn)3Si preferably over 50% and definitely not below 40% (to minimize galling),  •  minimum coarsening o f the constituent particles so as to obtain a m i n i m u m number o f large constituent particles after hot rolling (particles greater than l - 2 u m promote P S N during recrystallization [12]),  •  a small volume fraction o f dispersoids which are as large as possible so as to minimize the Zener drag effect which can inhibit recrystallization and thereby lower the total amount o f cube texture after hot rolling, and  •  low % M n i n solid solution after hot rolling so that the temperature at w h i c h recrystallization occurs is not too high and the material can self-anneal during coiling.  6.2  Conclusions  F r o m the results, the  following conclusions can be made  for a  suitable  homogenization profile: •  A faster heat-up rate is more favorable i n terms o f the formation o f fewer dispersoids and a slightly higher volume fraction o f a-Ali2(Fe,Mn)3Si.  •  A soak temperature o f 6 1 0 ° C is prefered as it achieves the required volume fraction o f a-Ali2(Te,Mn)3Si and leads to a lower volume fraction o f dispersoids with a larger size. The disadvantage o f a higher soak temperature is that there w i l l be a higher level o f M n i n solid solution during hot rolling which could inhibit recrystallization during  75  annealing after hot rolling. However, this could be lowered before annealing by adjusting the rolling temperatures. •  Soak times at temperature allowed the % a - A l i ( F e , M n ) S i to increase ( - 1 5 % from 0 2  3  to 3 hours), however, given a high soak temperature, high soak times could increase the size o f the constituent particles promoting P S N . Beyond three hours there did not appear to be much change i n the microstructure even for the material heated using a fast heat-up rate Based on this study, it appears that for an A A 3 1 0 4 alloy used for canbody stock the  optimum homogenization profile  includes: a  fast  heat-up  rate  to  a  high  homogenization temperature (580°C - 610°C) with moderate soak times (0.5-3 hrs). Other conclusions that can be drawn from the study are: i)  There is a significant difference i n the microstructure obtained between the  industrial standard heat-up rate range (45°C/hr - 70°C/hr) profile and very fast heat-up rates. Hence, i n order to apply laboratory based homogenization results, it is critical that similar heat-up rates to industrial conditions be applied to the samples (i.e., samples which have been heated using salt baths w o u l d give very different microstructures than the ones industrially heated even i f they have been held at 580°C for 3 hrs.) ii)  E v e n though the heat-up rate had a significant impact on the microstructure  evolution, within the current industrial heat-up rate range (i.e., soaking pits (~50°C/hr) V s . walking beam furnaces (~70°C/hr)), there is not a significant difference i n the microstructure obtained  76  6.3  Future Work  Based on this study, the following are some o f the topics that may be considered for future work: i)  Quantification o f dispersoids was not possible with optical image analysis or  S E M . Hence, T E M analysis o f multiple fields within samples in each profile is recommended to complete quantification and numerical analysis o f microstructure evolution during homogenization. ii)  One o f the interesting observations i n this study was the lowering o f M n i n solid  solution with time at higher temperatures. To explain the results and validate the theory proposed extensive fundamental relationship  between  Mn  in  research is required studying the solid  solution,  Al<5(Fe,Mn)  ->  thermodynamic  a-Ali2(Fe,Mn) Si 3  transformation and M n i n Ai6(Fe,Mn) and a-Aln(Fe,Mn)3Si. There is very little literature available covering the thermodynamics involved i n the A A 3 1 0 4 alloys. T o the best o f the author's knowledge the A l - M g - M n - F e - S i phase diagram is not yet known; it is under investigation i n the European C O S T 507  (Co-  operation i n Science and Technology) action [50, 51, 52]. iii)  Finally, modeling the microstructure evolution would be a useful undertaking  especially i f the thermodynamic data is available. The model would predict the microstructure characteristics given homogenization process parameters. could be tested against the results obtained i n this study.  77  The model  Bibliography  1. A L C O A website; Sources: A l c o a , Canadean, C C I , K A A L , I M E S , " http://www.alcoaxorri7frarneset.asp ?page=%2Finvestor%2Ffinancial%2Findex%2Eas  u" 2.  Sanders E . R . , Lege D J . ; Hartman T . L . , " A l u m i n u m rigid container sheet for the packaging industry ", A l u m i n i u m . 65 (9), 1989, 941-947.  3. Nes E . , Hutchinson W . B . , Proc. 10th Riso Symposium. 1989, Eds. Bilde-Sorensen et al, 233, Roskilde, Denmark, 4. Westerman E . J . , "Silicon: a vital alloying element i n aluminum beverage can body stock", A l u m i n u m A l l o y s for Packaging. Eds. J . G . Morris, H . D . Merchant, E . J . Westerman and P . L . Morris, The Minerals, Metals & Materials Society, Warrendale, P A , U S A , 1993,1-16. 5. Bolingbroke R . K . , Creed E . , Marshall G J . , R i c k s R . A . , "The effect o f Composition and Microstructure on the Recrystallization o f A A 3 0 0 4 " , A l u m i n u m A l l o y s for Packaging, eds. J . G . Morris, H . D . Merchant, E . J . Westerman and P . L . Morris, (Warrendale, P A : T M S , 1993), 215-224. 6. Hirsh, J.R., "Texture Evolution During Hot R o l l i n g i n A l A l l o y s " , Hot Deformation o f A l u m i n u m A l l o y s . Eds. Langdon T . G . , Merchant H . D . , M o r r i s J . G . , and Zaidi M . A . , The Minerals Metals & Materials Society, 1991, 379-388. 7. K y a n g Y . and M o r r i s J . G . , "The effect o f structure on the mechanical behavior and stretch formability o f constitutionally dynamic 3000 series aluminum alloys", Materials Science and Engineering. 1977, 59-74. 8. Hutchinson W . B . , Oscarsson A and. Karlsson A , "Control o f microstructure and earing behaviour i n aluminum alloy A A 3 0 0 4 hot bands", Materials Science and Technology, v o l . 5, 1989, 1118. 9. Regan, P . C . , "Recent Advances i n A l u m i n u m Strip Casting and Continuous R o l l i n g Technology—Implications for A l u m i n u m C a n B o d y Sheet Production", Light M e t a l A g e . February 1992, 58-62.  78  10. Merchant, H . D . ; Kattamis, T . Z . ; Scharf: "Homogenization o f A l u m i n u m A l l o y s " , Homogenization and Annealing o f A l u m i n u m and Copper A l l o y s , eds. H . D . Merchant, J. Crane, and E . H . Chia, (Warrendale, P A : T M S , 1988), 1-52. 11. D i n g S . X . , R e n B . & Morris J.G., "Influence o f Initial Structure on H o t Rolling and Recrystallization Textures i n A A 3 0 0 4 A l u m i n u m A l l o y " , Advances i n Hot Deformation Textures and Microstructures. Eds. Jonas J.J. Bieler T . R . and B o w m a n K . J . , 1994, 281-288. 12. Kamat R . G . , " A A 3 1 0 4 Can-body stock ingot: Characterization and homogenization", J . O . M . v o l . 48, N o . 6, 1996, 34-38. 13. M i k i I., Kosgue H . & Nagahama K . , " Supersaturation and Decomposition o f A l - F e A l l o y s During Solidification", J. Japan Inst. Light Metals. 1975, v o l 25, 1-12. 14. Chen L . , Morris J . G . & Das S.K., "The effect o f cooling rate during casting on the structure and mechanical property behavior o f A A 3 0 0 4 aluminum alloy", Continuous Casting o f Non-Ferrous Metals and A l l o y s Eds. Mercant H . D . , et al, The Minerals, Metals & Materials Society, Warrendale, P A , U S A , 1989, 269-284. 15. Mondolfo L . F . , A l u m i n u m A l l o y s . Structure and Properties. Butterworth & C o . Ltd., London, 1976. 16. Bolingbroke R . K . , Marshall G . J . & Ricks R . A . , "Microstructural Development D u r i n g Preheating o f A A 3 0 0 4 " , The 3 International Conference on A l u m i n u m A l l o y s . Trondeim, Norway, 1992, 285-290. r d  17. Backerud L . , K r o l E . , & Tamminen J., Solidification Characteristics o f A l u m i n u m V o l u m e 1: Wrought A l l o y s . Skan A l u m i n u m , Sweden, 1986. 18. Hatch J . E . , A l u m i n u m : Properties and Physical Metallurgy. American Society for Metals, Ohio, U S A , 1984, 61-65, 6-9,140 & 204-206. 19. Rouns T . N . , "Composition and Preheating on the Dispersoid and Insoluble Constituent Particle Evolution i n 3xxx A l l o y s " , A l u m i n u m alloys for Packaging HI. Eds. Das S.K., (Warrendale, P A : T M S , 1993), 3-20. 20. Palmer, S . L . ; L i , Z . : "The Single-Step Preheating o f A A 3 1 0 4 Can-Body Stock Ingots," J O M (6) 1996, 30-33. 21. Tromberg E . , Dons A . L . and Arnberg L . , "Investigation o f the Al6(Fe,Mn) -> ccAli2(Fe,Mn)3Si phase during homogenization o f A A 3 0 0 3 and A A 3 0 0 4 aluminum alloys", 3 International Conference on A l u m i n u m alloys. Trondheim, Norway, 1992, 270-275. rd  79  22. Watanabe H . , Ohori K . & Taheuchi Y . , "Phase Change i n 3004 Base A l l o y s at Elevated Temperatures", A l u m i n u m 60.1984, 310-321. 23. Wang, X . ; Kamat, R . G . : " A Technique to Measure Intermetallic Size Distribution i n A l u m i n u m C a n B o d y Stock," A l u m i n u m A l l o y s for Packaging II. eds. J . G . Morris, S.K. Das and H . S . Goodrich, (Warrendale, P A : T M S , 1996), 209-222. 24. Kattamis, T . Z . ; Merchant, H . D . ; Skolianos, S.; Scharf, G . : "Homogenization and Coarsening i n Cast 3004 A l u m i n u m A l l o y , " A l u m i n u m 65 (4) (1989), 367-376 25. Kattamis, T . Z . ; Merchant, H . D . ; Scharf, G . : "Influence o f Homogenization on Intermetallic Coarsening in A l u m i n u m A l l o y 3004," Homogenization and Annealing o f A l u m i n u m and Copper A l l o y s , ed., H . D . Merchant, J . Crane and E . H . Chia, (Warrendale, P A : T M S , 1988), 117-135. 26. T . C . Sun, "The effect o f preheating on A A 3 1 0 4 aluminum alloy ingot structure and particulate composition", A l u m i n u m A l l o y s for Packaging. Eds. J . G . Morris, H . D . Merchant, E . J . Westerman and P . L . Morris, The Minerals, Metals & Materials Society, Warrendale, P A , U S A , 1993, 31-46. 27. Karlson A , Oscarsson A , Lehtinen B & Hutchinson W . B . , "Influence o f Homogenization on Structure and Earing i n A l u m i n u m alloy 3004", Homogenization and Annealing o f A l u m i n u m and Copper A l l o y s , ed., H . D . Merchant, J . Crane and E . H . Chia, (Warrendale, P A : T M S , 1988), 99-116. 28. Anyalebechi, P . N . ; Rouns, T . N . ; Sanders, Jr., R . E . , "Microstructural Analysis o f Second Phases Developed During Casting and Preheating o f 3 x x x A l u m i n u m alloys", R A S E L M ' 9 1 . Tokyo, Japan, 1991, 923-928. 29. Theler J.J., & Furrer P., "Influence o f F e on the Course o f Precipitation i n A l - M n A l l o y s " , A l u m i n i u m . 50, v o l 7,1974,467-472. 30. G o e l D . B . , Furrer P. & Warlimont H . , " Recrystallization and Precipitation i n A l - M n A l l o y s " , A l u m i n i u m 50. v o l 8, 1974, 511-521. 31. U s u i E . , Takashi I. & Noboru S., "Influence o f M n and M g Additions on Hot Deformation o f A l u m i n u m and A l u m i n u m alloys", Z . Metallkde.. 1986, v o l 77, 179187. 32. Marshall G.J., "Microstructural Control During Processing o f A l u m i n u m Canning A l l o y s " , Materials Science Forum. 1996, vols. 217-222, 19-30.  80  33. Sanders R . E . , Baufmann S.F. & Stumpf H . C . , "Wrought Non-Heat-Treatable A l u m i n u m A l l o y s " , A l u m i n u m A l l o y s - Contemporary Research and Applications. 1988, Eds Vasudevan A . K . & Doherty R . D . , 75-83. 34. Goodrich H . S . , " A M o d e l for the Precipitation/Dissolution o f M n D u r i n g Commercial Homogenization o f A l u m i n u m A l l o y 3104," A l u m i n u m A l l o y s for Packaging, eds J . G . Morris, H . D . Merchant, E . J . Westerman and P . L . Morris, (Warrendale, P A : T M S , 1993), 47-60. 35. Private communication with A l c a n International. 36. Jackson A . & Sheppard T., "Structural Modifications occurring during the Homogenization o f some 7xxx alloys". E T 96: Profiles o f Change. 6 International A l u m i n u m Extrusion Technology Seminar. 1996, 541-550. l  37. Kattamis, T . Z . ; Merchant, H . D . ; Scharf, G . : "Influence o f Homogenization on Intermetallic Coarsening i n A l u m i n u m A l l o y 3004," Homogenization and Annealing o f A l u m i n u m and Copper A l l o y s , ed., H . D . Merchant, J. Crane and E . H . Chia, (Warrendale, P A : T M S , 1988), 117-135. 38. E . Tromberg, A . L . Dons and L . Arnberg, "Investigation o f the Al6(Fe,Mn) -> ocAli2(Fe,Mn)3Si phase during homogenization o f A A 3 0 0 3 and A A 3 0 0 4 aluminum alloys", 3 International Conference on Aluminum alloys. Trondheim, Norway, 1992, 270-275. rd  39. L i , Z . ; L i , C . X . ; Morris, J.G.: "Precipitate Behavior o f A A 3004 A l u m i n u m A l l o y " , A l u m i n u m A l l o y s for Packaging, eds. J . G . Morris, H . D . Merchant, E . J . Westerman and P . L . Morris, (Warrendale, P A : T M S , 1993), 61-69. 40. Suni J.P. & Shuey R . T . , "Modeling Dispersoid and Constituent Particle Evolution i n 3xxx A l l o y s " , A l u m i n u m alloys for Packaging III. Eds. Das S . K . , (Warrendale, P A : T M S , 1993), 21-37. 41. Suni J.P., Shuey R . T . & Doherty R . D . , "Dispersoid Modeling i n 3xxx A l l o y s " , A l u m i n u m A l l o y s for Packaging JJ. eds. J . G . Morris, S . K . Das and H . S . Goodrich, (Warrendale, P A : T M S , 1996), 145-159. 42. Anyalebechi P . N . , Rouns T . N , Sanders, Jr. R . E . , "Effects o f C o o l i n g Rate and Grain Refining on Constituent Phase Particle Size In As-Cast 3004 A l l o y , " Light Metals 1991. ed. E . L . R o y (Warrendale, P A : T M S , 1990), 821-850. 43. Tilak R . V . , . Morris J . G , "Studies o f the effect o f thermomechanical treatments on the supersaturation content o f strip-cast aluminum alloy 3004," Materials Science and Engineering. 73 (1985), 139-150. 81  44. Tromberg E . , Dons A . D . , A m b e r L . , "Investigation o f the Al6(Fe,Mn) -> cxAli2(Fe,Mn)3Si phase transformation during homogenization o f A A 3 0 0 3 and A A 3 0 0 4 A l u m i n u m A l l o y s " , The 3 International Conference o f A l u m i n i u m A l l o y s . 1992, 270-275. r d  45. Humphreys F . J . , Hatherly. M , Recrystallization and Related Annealing Phenomena. Elseview Science Inc., N e w Y o r k 1995. 46. Es-Said O.S., Morris J . G . & Merchant H . D . , "The Influence o f Particles on the Recrystallization o f A l u m i n u m A l l o y s " , Homogenization and Annealing o f A l u m i n u m and Copper A l l o y s , eds. H . D . Merchant, J. Crane, and E . H . Chia, (Warrendale, P A : T M S , 1988), 183-208. 47. Sun, T . C . : "Surface and Metallurgy Effects i n W a l l Ironing o f A A 3104," A l u m i n u m A l l o y s for Packaging. Eds. J . G . Morris, H . D . Merchant, E . J . Westerman and P . L . Morris, (Warrendale, P A : T M S , 1993), 183-192. 48. L i Z . , L i C . X . , Morris J. G . , " A Comparison o f Recrystallization behavior o f D . C . Cast and Continous Strip Cast A A 3 0 0 4 Aluminum A l l o y " , A l u m i n u m A l l o y s for Packaging, eds. J . G . Morris, H . D . Merchant, E . J . Westerman and P . L . Morris, (Warrendale, P A : T M S , 1993), 299. 49. M u r r L . E . , Interfacial Phenomena i n metals and A l l o y s . Reading, P . A . AddisonWesley, 1975, 131. 50. Sigli C , Vichery H . & Grange B . , "Computer Assisted Metallurgy for Packaging A l l o y s " , ", A l u m i n u m alloys for Packaging ITI. Eds. Das S . K . , (Warrendale, P A : T M S , 1993), 189-197. 51. K o l b y P., Sigli C . & Simensen C . J . , "Solubility Limit o f M n and S i i n A l - M n - s i at 5 5 0 ° C " , A l u m i n u m alloys: Their Physical and Mechanical Properties ( I C A A 4 ) , E d . T . N . Sanders & E . A . Starke, 1994, 513-519. 52. Fries, S . G . , Jantzen, T., "Compilation o f C A L P H A D formation enthalpy data- Binary Intermetallic Compounds i n the C O S T 507 Gibbsian Database", Thermochimica Acta, v o l . 314, no. 1-2, (1998), 23-32.  82  Appendix A Temperature profile of aluminum samples in original and retrofitted (i.e. with aluminum tube) furnace.  83  b) Figure A l - Temperature variations between furnace (580°C) and sample o f original furnace at a) two sample positions i n the furnace, and b) on a raised platform near the furnace thermocouple.  84  Figure Al - Temperature variations between furnace (580°C) and sample temperature i n the retrofitted furnace.  85  Appendix B Microstructures o f samples heated to 580°C using different heat-up rates and held for three hours:  86  87  88  e) 279°C/hr Figure B l - Microstructures o f samples at 580°C soak temperature after 3 hr hold for a) 37°C/hr, b) 47°C/hr, c) 70°C/hr, d)112°C/hr, and e) 279°C/hr.  89  Appendix C E P M A line scans ( M g and M n ) at various temperatures during heat-up to 580°C using two different heat-up rates (47°C/hr and 279°C/hr).  90  Figure C l - E P M A line scans ( M g and M n ) for the as-cast microstructure.  91  c) 580°C 0 hr hold  d) 5 8 0 ° C 3 h r h o l d  .^^^^^^H^HBMl  Figure C 2 - E P M A line scans ( M g and M n ) for 47°C/hr profile at a) 4 0 0 ° C , b) 500°C, 580°C (0 hrs), and d) 580°C (3hrs). 93  a) 4 0 0 ° C  b)500°C  10 um i 1 Figure C 3 - E P M A line scans ( M g and M n ) for 279°C/hr profile at a) 4 0 0 ° C , b) 500°C, 580°C (0 hrs), and d) 580°C (3hrs). 95  

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