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The effect of pre-deformation on the ageing behavior of 7030.60 and 7108.72 alloys Huang, Jin 1999

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T H E E F F E C T O F PRE-DEFORMATION O N T H E A G E I N G BEHAVIOR O F 7030.60 A N D 7108.72 A L L O Y S by Jin Huang B. Eng., Shanghai Jiao Tong University, 1991  A THESIS SUBMITTED IN PARTIAL FULFILLMENT O F T H E REQUIREMENTS FOR T H E DEGREE O F MASTER 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 of Metals and Materials Engineering)  We accept this thesis as conforming to the required standard  T H E UNIVERSITY OF BRITISH COLUMBIA March 1999 ® Jin Huang, 1999  In  presenting  degree  this  thesis  in  at the University of  partial  fulfilment  British Columbia,  of  the  requirements  I agree  for  of  department  this or  thesis by  publication of this  for scholarly  his thesis  or  her  purposes  may be  representatives.  It  for financial gain shall not  is  granted  yfr{f>fV/l  a<*J  l „ : . ~l r»_:i:_L_ r~ _ I L * The IUniversity of British Columbia Vancouver, Canada Tkn  Date  DE-6 (2/88)  by  understood  for extensive  the head that  be allowed without  permission.  Department of  advanced  that the Library shall make it  freely available for reference and study. I further agree that permission copying  an  j/ft,4tf V  of  my  copying  or  my written  ABSTRACT  This study is part of a research project whose goal is to develop a process model for the production of aluminum alloy automotive bumpers. The problem originates from the industrial bumper manufacturing practice w h i c h uses aluminum alloys, AA7030.60 and AA7108.72 i n particular.  In the industrial production schedule, a forming (pre-  deformation) process occurs after the solution heat treatment and water quenching, followed b y a T 6 ageing process. D u e to the specific shape of the bumper product, the forming process introduces a heterogeneous distribution of deformation i n the bumper.  A systematic investigation of pre-deformation o n the ageing behaviour of these alloys has been carried out. A series of pre-deformations (pre-strain from 0 to 1.2) was imposed on the materials, whose subsequent  ageing responses were recorded i n terms of the  mechanical properties (yield stress, tensile strength, etc.) and electrical resistivity. Some transmission electron microscopy measurements were also conducted b y other members of the research group. W i t h a combination of these results the evolution of the material's microstructure and mechanical properties can be traced during the ageing process.  Some of the significant results from the present study include the observation that the natural ageing response is reduced and the artificial ageing kinetics are accelerated for the pre-deformed alloys. L o w e r peak ageing strength is also observed for the pre-deformed alloys.  These changes i n the material behaviour are related to the changes i n the  microstructure when pre-deformation is present i n the materials. A t a microscopic level, dislocations are observed to arrange themselves into a loose cell structure. Pre-deformed alloys were observed to have larger average precipitate sizes. In addition, they also have different precipitate size distribution.  F o r pre-deformed samples t w o peaks can be  11  recognized i n the precipitate size distribution curves, corresponding to the precipitates i n the bulk and the precipitates at dislocation sites. Based on this study, it has been observed that deformation prior to artificial ageing leads to inferior final properties. A modified bumper manufacturing sequence is proposed i n an attempt to eliminate this effect.  However, further studies are needed to test the  feasibility of the proposal.  111  TABLE OF CONTENTS  ABSTRACT  ii  LIST O F FIGURES  vi  LIST O F TABLES  xi  ACKNOWLEDGEMENTS  xii  CHAPTER 1 INTRODUCTION  1  CHAPTER 2 BACKGROUND  4  2.1 7XXX A L U M I N U M ALLOYS USED IN A U T O M O T I V E INDUSTRY 2.2 INDUSTRIAL BUMPER MAKING PROCESS 2.3 A G E H A R D E N I N G  4 6 9  2.3.1 The Heat Treatment 2.3.2 Precipitation Process  9 12  2.3.2.1 Formation of G-P Zones 2.3.2.2 Formation of r\' 2.3.2.3 Formation of r| 2.3.2.4 Precipitation Kinetics  13 14 16 16  2.3.3 Dislocation-Precipitate Interaction 2.4 T H E EFFECT O F PRE-DEFORMATION O N T H E PRECIPITATION BEHAVIOR 2.5 T H E STATEMENT O F OBJECTIVES  C H A P T E R 3 EXPERIMENTAL P R O C E D U R E 3.1 INTRODUCTION 3.2 MATERIAL 3.3 METALLOGRAPHIC PREPARATION 3.4 EXPERIMENTAL H E A T A N D M E C H A N I C A L T R E A T M E N T 3.5 PROPERTY MEASUREMENTS  C H A P T E R 4 RESULTS  17 21 23  24 24 24 25 26 27  30  4.1 GRAIN STRUCTURE O F T H E AS-EXTRUDED MATERIAL 4.2 AGEING BEHAVIOR  30 33  4.2.1 Evolution ofMechanical Properties During Ageing. 4.2.1.1 Typical Tensile Curves  33 34  IV  4.2.1.2 Ageing Curves in Yield Stress, UTS and Strain-to-necking  37  4.2.2 Evolution of Electrical Resistivity During Ageing.  4.3 SUMMARY OF SOME SIGNIFICANT OBSERVATONS  CHAPTER 5 DISCUSSION  44  47  53  5.1 EFFECT OF Zr ON THE GRAIN STRUCTURES 5.2 OVERVIEW' OF THE AGEING BEHAVIOUR  53 54  5.2.1 Yield Stress Evolution During Ageing 5.2.2 Electrical Resistivity Evolution During Ageing.  5.3 EFFECT OF PRE-DEFORMATION ON THE YIELD STRESS OF THE ASDEFORMED MATERIAL 5.4 EFFECT OF PRE-DEFORMATION ON THE NATURAL AGEING BEHAVIOUR 5.5 EFFECT OF PRE-DEFORMATION ON THE ARTIFICIAL AGEING BEHAVIOUR  55 58  59 62 64  5.5.1 Artificial Ageing Stage I): AgeingatlOO°C 5.5.2 Artificial Ageing Stage II): Ageing at 5.5.3 Kinetics of Precipitate Coarsening.  5.6 RECOVERY DURING THE ARTIFICIAL AGEING 5.7 A SIMPLE MODEL FOR PREDICTING STRESS-STRAIN CURVES 5.8 OPTIMUM HEAT TREATMENT  CHAPTER 6 CONCLUSIONS AND FUTURE WORK 6.1 CONCLUSIONS 6.2 FUTURE WORK  :  64 66 74  75 78 85  87 87 88  BIBLIOGRAPHY  91  APPENDIX  94  v  LIST O F F I G U R E S  Figure 2.1: Schematic illustration of a typical European bumper  ...7  Figure 2.2: Schematic illustration of the thermal-mechanical history of the 7030.60 and 7108.72 aluminum alloys before bumper forming process  7  Figure 2.3: Schematic illustration of the bumper forming and ageing schedule for the 7030.60 and 7108.72 aluminum alloys  9  Figure 2.4: Schematic illustration of the route of phase transformation during quench  11  Figure 2.5: Schematic illustration of dislocation-particle interactions. (1) - (3): moving dislocation shearing a weak particle; (i) - (iii): moving dislocation bypassing a strong particle  19  Figure 2.6: Schematic illustration of the breaking angle <p for a dislocation holding by an obstacle  21  Figure 2.7: Influence of the precipitate size distribution width A on the evolution of yield stress during ageing  23  Figure 3.1: Schematic illustration of the experiment schedule for the 7030.60 and 7108.72 aluminum alloys  26  c  Figure 3.2:  Schematic illustration of two types tensile sample used in this research, a) for M T S tensile machine and b) for Instron tensile machine  29  Figure 4.1: Tri-planar optical microscopic graphs showing the grain structure of the 7030.60 alloy in the as-received condition  31  Figure 4.2: Tri-planar optical microscopic graphs showing the grain structure of 7108.72 alloy in the as-received condition  32  Figure 4.3a: Engineering stress-strain curves for 7030.60 alloy with no pre-strain, in various stages of ageing till the peak  35  vi  Figure 4.3b: Engineering stress-strain curves for 7030.60 alloy w i t h no pre-strain, i n various stages of ageing after the peak  35  Figure 4.4a: Engineering stress-strain curves for 7030.60 alloy w i t h a pre-strain of 0.30, i n various stages of ageing till the peak  36  Figure 4.4b: Engineering stress-strain curves for 7030.60 alloy w i t h a pre-strain of 0.30, i n various stages of ageing after the peak  36  Figure 4.5a: E v o l u t i o n of yield stress for the 7030.60 and 7108.72 alloys w i t h no pre-strain  39  Figure 4.5b: E v o l u t i o n of U T S and strain-to-necking for the 7030.60 and 7108.72 alloys w i t h no pre-strain  39  Figure 4.6a: E v o l u t i o n of yield stress for the 7030.60 and 7108.72 alloys w i t h 0.03 pre-strain  40  Figure 4.6b: E v o l u t i o n of U T S and strain-to-necking for the 7030.60 and 7108.72 alloys w i t h 0.03 pre-strain  40  Figure 4.7a: E v o l u t i o n of yield stress for the 7030.60 and 7108.72 alloys w i t h 0.10 pre-strain  41  Figure 4.7b: Evolution of U T S and strain-to-necking for the 7030.60 and 7108.72 alloys w i t h 0.10 pre-strain  41  Figure 4.8a: E v o l u t i o n of yield stress for the 7030.60 and 7108.72 alloys w i t h 0.30 pre-strain  42  Figure 4.8b: E v o l u t i o n of U T S and strain-to-necking for the 7030.60 and 7108.72 alloys w i t h 0.30 pre-strain  42  Figure 4.9a: E v o l u t i o n of yield stress for the 7030.60 and 7108.72 alloys w i t h 1.20 pre-strain  43  Figure 4.9b: Evolution of U T S and strain-to-necking for the 7030.60 and 7108.72 alloys w i t h 1.20 pre-strain  43  Figure 4.10: E v o l u t i o n of electrical resistivity for the 7030.60 and 7108.72 alloys w i t h no pre-strain  44  Vll  Figure 4.11: E v o l u t i o n of electrical resistivity for the 7030.60 and 7108.72 alloys w i t h 0.03 pre-strain  45  Figure 4.12: E v o l u t i o n of electrical resistivity for the 7030.60 and 7108.72 alloys w i t h 0.10 pre-strain  45  Figure 4.13: E v o l u t i o n of electrical resistivity for the 7030.60 and 7108.72 alloys w i t h 0.30 pre-strain  46  Figure 4.14: E v o l u t i o n of electrical resistivity for the 7030.60 and 7108.72 alloys w i t h 1.20 pre-strain  46  Figure 4.15: Y i e l d stress evolution during the ageing process for the 7030.60 and 7108.72 alloys without pre-deformation. (The dashed line corresponds to the industrial T 6 schedule.)  48  Figure 4.16: Y i e l d stress evolution during the ageing process for the 7030.60 alloy w i t h different levels of pre-deformation  49  Figure 4.17: Y i e l d stress evolution during the ageing process for the 7108.72 alloy w i t h different levels of pre-deformation  49  Figure 4.18: Electrical resistivity evolution during the ageing process for the 7030.60 alloy w i t h different levels of pre-deformation  51  Figure 4.19: Electrical resistivity evolution during the ageing process for the 7108.72 alloy w i t h different levels of pre-deformation  51  Figure 5.1: Schematic illustration of the flow stress as a function of the obstacle radius r  57  Figure 5.2: Y i e l d stress of the 7030.60 and 7108.72 alloys i n the as-deformed condition as a function of pre-strain  60  Figure 5.3: Plots of [(a—CT )/ M a G b ] vs. strain, illustrating the increases of dislocation density after the pre-deformation for the 7030.60 and 7108.72 alloys  62  Figure 5.4: Changes i n yield stress during the 24-hour natural ageing process as a function of pre-strain  63  Figure 5.5: Change i n yield stress between the end of natural ageing and after 1 hour at 100°C  65  2  0  viii  Figure 5.6: Variation i n the peak strength as a function of pre-strain  67  Figure 5.7: TEM-micrographs from a) the non-deformed and b) pre-deformed (0.10 strain) conditions for the 7030.60 alloy aged for 70 hours at 150°C  68  Figure 5.8: The time to peak strength as a function of pre-strain  70  Figure 5.9: T E M micrographs of 7030 alloy w i t h 0.1 pre-strain, aged for 5 hours at 100°C plus 20 hours at 150°C. a) dislocations i n contrast and b) dislocations out of contrast, showing the precipitates  71  Figure 5.10: Precipitate size distributions for pre-deformed and non-deformed a) 7108.72 alloy and b) 7030.60 alloy, aged at 100°C for 5 hours followed b y 70 hours at 150°C (From Saeter, et. al. [33])  72  Figure 5.11: Schematic illustration of the size distributions for bulk precipitates and precipitates o n dislocations for the pre-deformed materials  73  Figure 5.12: Plot of logarithm of the precipitate size vs. the logarithm of ageing time  75  Figure 5.13: F l o w stress components for the 7030.60 alloy w i t h pre-strain of 1.20 i n various ageing stages  77  Figure 5.14: Plot of the w o r k hardening rate vs. stress for the 7030.60 alloy w i t h no pre-deformation i n the as water quenched condition  80  Figure 5.15: Plots of the experimental stress-strain curve and the flow stress contribution from the dislocation hardening for the 7030.60 alloy w i t h no pre-deformation i n the as water quenched condition  81  Figure 5.16: Plots of the experimental stress-strain curve and the calculated ones using different averaging laws for the non-deformed 7030.60 alloy naturally aged for 24 hours  83  Figure 5.17: Plots of the experimental stress-strain curve and the calculated ones using different averaging laws for the non-deformed 7030.60 alloy aged at 100°C for 1 hour  83  Figure 5.18: Plots of the experimental stress-strain curve and the calculated, ones using different averaging laws for the non-deformed 7030.60 alloy aged at 100°C for 5 hours ,  84  ix  Figure 5.19: Plots of the experimental stress-strain curve and the calculated ones using different averaging laws for the non-deformed 7030.60 alloy aged at 150°C for 10 hours  84  Figure 5.20: Plots of the experimental stress-strain curve and the calculated ones using different averaging laws for the non-deformed 7030.60 alloy aged at 150°C for 60 days  85  Figure 6.1: Schematic illustration of the extrusion and bumper manufacturing processes  89  Figure 6.2: Schematic illustration of the modified bumper manufacturing process  90  Figure A . l : Hardness distribution through the thickness of the as-received a) 7030.60 and b) 7108.72 plate  95  Figure A . 2 : Ageing curves (in hardness) for 7030.60 alloy w i t h different resolution time at 4 8 0 ° C  96  Figure A . 3 : Ageing curves (in resistivity) for 7030.60 alloy w i t h different resolution time at 4 8 0 ° C  97  Figure A . 4 : Ageing curves (in hardness) for 7108.72 alloy w i t h different resolution time at 4 8 0 ° C  97  Figure A . 5 : Ageing curves (in resistivity) for 7108.72 alloy w i t h different resolution time at 4 8 0 ° C  98  Figure A . 6 : Ageing curves (in hardness) for 7030.60 alloy w i t h quenching conditions  different 99  Figure A . 7 : Ageing curves (in resistivity) for 7030.60 alloy w i t h quenching conditions  different  Figure A . 8 : Ageing curves (in hardness) for 7108.72 alloy w i t h quenching conditions  different  Figure A . 9 : Ageing curves (in resistivity) for 7108.72 alloy w i t h quenching conditions  different  99  100  100  LIST O F T A B L E S  Table 2.1: Compositions of some typical 7xxx aluminum alloys (wt.%)  5  Table 2.2: Typical mechanical properties of some 7xxx aluminum alloys  5  Table 3.1: Chemical composition of the 7030.60 and 7108.72 alloys (wt.%)  25  Table 4.1: Y i e l d stress data for 7030.60 and 7108.72 alloys w i t h different levels of pre-deformation i n the as-quenched, as-deformed and overaged conditions  38  Table 5.1: Summary of precipitate sizes after ageing at 1 0 0 ° C / 5 h r . + 150°C /70hr (From Saeter, et. al. [33])  69  Table 5.2: Summary of precipitate sizes for the non-deformed alloys i n the overaged conditions  74  Table 5.3: Y i e l d stress data for 7030.60 alloy w i t h pre-strain of 1.20 i n the asquenched, as-deformed and overaged conditions  75  xi  ACKNOWLEDGMENTS  I wish to thank m y supervisor, D r . Warren J . Poole, for his instructions and guidance throughout this study.  H e has taught me so much, not only what to learn, but also,  most importantly, h o w to learn. It is on his chassis that this thesis can run to an end. I w o u l d also like to acknowledge D r . Jan A . Saster and D r . Alexis Deschamps for the fruitful discussion, and many thanks to the group members: Shahrzad Esmaeili, M a r i o n Charleux, Partha Ganguly, and Gilles Guiglionda for your friendship and inspiring free discussions. I ' l l remember those exciting group hikes and the good time I had w i t h you. Special thanks extend to G a r y Lockhart, P.Eng., w h o had made it easy for me to get accustomed to the surroundings at the beginning, and helped me to overcome many problems a newcomer may face.  Being very versatile and handy, he also gave me so  much hands-on help w i t h the experimental set-ups. In the t w o years of study i n the department of Metals and Materials Engineering, U B C , I find it really a pleasure to be i n such a friendly atmosphere.  Thanks to everybody's  efforts to make it such an enjoyable place. So many good things left i n memory. Finally, I w o u l d like to present m y gratefulness to m y wife, D u y u n L i u , though no words are enough for your love and support that make this happen. A l l the happiness and success as I have, I share them w i t h y o u .  xu  Chapter  1  INTRODUCTION  A l u m i n u m alloys have been flying high i n the sky for decades w i t h many applications found i n the aerospace industry. However, i n response to the growing concern about environmental issues, more and more aluminum alloys have landed d o w n on the ground, finding their applications i n the automotive industry. The most significant advantage of applying aluminum i n automotive applications is the weight saving, not o n l y i n the component itself, but also on the other surrounding systems. F o r example, a light car body made out of aluminum makes it unnecessary to have a very strong and heavy suspension system matching it. Less weight leads to lower fuel consumption and green house gas emissions, and culminates i n a better environment. This advantage can offset the higher cost of aluminum compared to steel, its chief competitor i n the automotive industry. A s a result of these considerations, there has been a growing utilization of aluminum alloys i n automobiles during recent years. Their applications are diverse and one can find their existence i n almost all the structures of an automobile, ranging from the cylinder head, block and pistons, to suspension components, wheels, body panels and bumpers. A s an example, A u d i A . G . has made an all-aluminum car, the A u d i A 8 , where its aluminum body-in-white weighs 40% less than the steel counterpart. aluminum as a bumper material is the subject of this research project.  1  T h e use of  In manufacturing of aluminum bumpers, a heat treatable aluminum alloy is usually selected. The manufacturing process can then be designed to take advantage of the alloy's age hardening potential.  Bumpers are usually formed i n the solid solution condition  (low yield stress and high formability) and then go through a heat treatment to age harden the alloy. In attempt to improve the production process for aluminum bumpers, Hydro-Raufoss (a Norwegian company w h o supplies aluminum bumpers to many European automotive manufacturers) is currently developing a comprehensive model to predict the final properties of their bumpers.  The process model w i l l include finite  element method calculations to examine the fabrication process and a microstructurally based mechanical property model for the heat treatment cycle. Preliminary calculations and experimental observations show that there is a large spatial variation i n the plastic deformation w i t h i n the bumper after forming. Furthermore, initial experiments by Hydro-Raufoss suggest that the final yield stress of the alloy actually decreases when the alloy is pre-deformed prior to ageing process. This result is not intuitive and requires further investigation. Therefore, the goal of this w o r k is to examine the relationship between deformation after solution heat treatment and the agehardening response. The driving force comes not only from the academic curiosity, but also from the industry's desire to improve the performance of the product. O n the basis of this w o r k , as well as results from other research w o r k , a process model that predicts the yield strength of the products after the manufacturing process w i l l ultimately be developed, although it does not fall under the focus of the present study.  In order to examine the effect of pre-deformation on the ageing behavior of t w o 7xxx series aluminum alloys (AA7030.60 and AA7108.72), a set of experiments simulating the industrial bumper fabrication process was conducted.  Samples were taken from as-  received hot extruded plate and given a solution heat treatment at 4 8 0 ° C for 20 minutes. Then, a pre-determined strain ranging from 0 to 1.2 was imposed o n the sample by either  2  uniaxial tension or cold rolling. After the pre-deformation, samples were naturally aged for 24 hours, then they went through a t w o step artificial ageing processes: 100°C for 5 hours, and 150°C for up to 60 days. The ageing response of these alloys was recorded i n terms of yield stress, U T S , strain-to-necking and electrical resistivity properties, by conducting tensile tests and electrical conductivity measurements during various ageing stages. The results from the present study confirm the preliminary studies b y Hydro-Raufoss. Deformation after solution heat treatment has a negative effect o n the final strength of the material. T h i s is attributed to the effects of dislocations on the ageing behaviour of both 7030.60 and 7108.72 alloys. It is found that the presence of a high dislocation density lowers the natural ageing response of alloys, probably due to the dislocation acting as sinks for quenched i n vacancies. The higher dislocation density also accelerates the ageing kinetics by providing short circuit diffusion path due to pipe diffusion along dislocation lines.  Moreover, the peak strength of the material is lowered since the  presence of dislocations changes the precipitate size distribution so that not all the precipitates can make their full contribution to the strength simultaneously.  There is  also evidence of static recovery that occurs i n the pre-deformed samples during the artificial ageing processes.  3  Chapter  2  BACKGROUND  This chapter provides a brief review of the alloys that are used for bumpers and their processing history for this product.  This is followed by a literature review of  precipitation processes and the strengthening mechanisms of relevance to these alloys.  2.1 7XXX ALUMINUM ALLOYS USED IN AUTOMOTIVE INDUSTRY  In Europe, 7xxx series aluminum alloys are widely used for making automotive bumpers. The use of these alloys offers substantial weight savings compared to traditional steel bumpers. The choice of 7xxx series alloys is made due to their potential for producing high final strengths through precipitation hardening. The alloying elements for 7xxx alloys are mainly zinc and magnesium, which form a fine distribution of MgZn  2  precipitate under the appropriate heat treatment conditions. Usually copper is also added to increase the resistance to stress corrosion cracking [1], and some alloys have trace elements content such as zirconium [2, 3] or silver which are added to refine the grain structure of the material. The highest strength commercial alloys (usually aerospace alloys) contain larger amounts of zinc and magnesium, with the Zn content in the range of 5-7 wt.%. For lower strength alloys the Zn and Mg contents may be reduced [4]. Table 2.1 and 2.2 show the compositions of some typical 7xxx series aluminum alloys  4  and their mechanical properties, respectively. T h e alloys for the present study, 7030.60 and 7108.72, have also been included for easy comparison. Table 2.1: Compositions of some typical 7xxx aluminum alloys (wt.%)  Alloy  Zn  Mg  Cu  Si  Fe  Mn  Cr  Ti  Zr  Al  7090  7.3-8.7  2.0-3.0  0.6-1.3  0.12  0.15  -  -  -  -  bal.  7475  5.2-6.2  1.9-2.6  1.2-1.9  0.10  0.12  0.06  0.18-0.25  0.06  -  bal.  7050  5.7-6.7  1.9-2.6  2.0-2.6  0.12  0.15  0.10  0.04  0.06  0.08-0.15  bal.  7075  5.1-6.1  2.1-2.9  1.2-2.0  0.40  0.50  0.30  0.18-0.35  0.20  -  bal.  7009  5.5-6.5  2.1-2.9  0.6-1.3  0.20  0.20  0.10  0.10-0.25  0.20  -  bal.  7108  5.45  1.20  0.27  -  -  -  -  -  0.16  bal.  7030  5.45  1.22  0.3  -  -  -  -  -  -  bal.  7004  3.8-4.6  1.0-2.0  0.05  0.25  0.35  0.20-0.7  0.05  0.05  0.10-0.20  bal.  7039  3.5-4.5  2.3-3.3  0.1  0.3  0.40  0.10-0.40  0.10  0.10  -  bal.  Table 2.2: Typical mechanical properties of some 7xxx aluminum alloys  Alloy  Temper  Yield Stress (MPa)  UTS (MPa)  Elongation (%)  7090  T7E71  580  620  9  7475  T651  560  590  12  7050  T736  510  550  11  7075  T6  500  570  11  7009  T6  470  535  12  7108  T6  431  482  15  7030  T6  388  427  19  7004  T6  340  400  12  7039  T61  345  415  13  5  The most significant strengthening mechanisms normally utilized by the 7xxx series aluminum alloys are precipitation hardening and solid solution hardening. However, effectiveness of precipitation strengthening is much more pronounced than that of the solid solution strengthening [5]. W h e n the material is age hardened, the contribution to the strength from the solid solution strengthening is further reduced, since o n l y a small fraction of the total available solid atoms is retained i n the matrix. In some cases, further strengthening from dislocation hardening may be important, particularly when the alloy is formed into the shape of the product by plastic deformation. While it is clear that the increase i n strength due to the precipitation hardening during ageing is much more than enough to offset the loss of strength due to the reduced contribution from solid solution, the interaction between precipitation hardening and dislocation hardening is more complicated.  2.2 INDUSTRIAL BUMPER MAKING PROCESS  T w o 7xxx series aluminum alloys, 7030.60 and 7108.72, are currently being employed i n industry b y the N o r w e g i a n company Hydro-Raufoss for bumper manufacturing. Figure 2.1 schematically illustrates the shape of a typical bumper.  It has an aluminum body,  typically covered by a matching plastic shell. These materials are A l - Z n - M g alloys w i t h moderately high peak ageing strength.  The processing history for these alloys prior to  bumper forming is shown schematically i n Figure 2.2. First the alloys are smelted and D C cast into billets.  T h e y are then homogenized and hot extruded into C shaped  extrusions (i.e. a precursor shape for the final bumper).  6  Figure 2.1: Schematic illustration of a typical European bumper.  Temperature Casting  Time  Figure 2.2: Schematic illustration of the thermal-mechanical history of the 7030.60 and 7108.72 aluminum alloys before bumper forming process.  In the bumper forming process, the extrusions are given a solution heat treatment and quench before they are fed directly into the stretch bending line, where the bumpers are  7  formed. The thermo-mechanical process is accomplished as follows: i)  the extrusions are heated to 480°C and held for 20 minutes when they are fed o n a conveyer belt into a tunnel furnace,  ii) a water spray quenches the material upon emerging from the furnace, iii) the extrusions are then immediately formed into the shape of bumper by a series of stretch bending operations. After the forming process, the material is given an ageing treatment consisting of three processes, 1) 24-hour natural ageing (this is the typical value, the actual natural ageing time varies depending o n the industrial production schedule, etc.), 2)  1 0 0 ° C / 5 hour ageing,  3)  1 5 0 ° C / 6 hour ageing.  This is the industrial T 6 heat treatment schedule, and it is conducted to utilize the age hardening potential of the material to improve the final strength of the product.  The  whole forming and ageing process is shown i n Figure 2.3. It is w o r t h noting that the strain imposed on the plates during stretch bending varies, depending o n the shape of the bumper.  The material i n the corner areas experiences a  larger strain, while the material i n the centre part of the bumper receives a very small amount of strain. A finite element method analysis predicts strains i n the range of 0 0.30 [6]. This difference i n the degree of deformation leads to a different dislocation density and amount of w o r k hardening i n the various sections of the bumper.  8  Temperature Solution @480°C/20min.  Ageing IWater quench @100°C/5hr. Ageing @150°C/6hr. Stretch bending ^  I  Natural ageing/Id.  Time Figure 2.3: Schematic illustration of the bumper forming and ageing schedule for the 7030.60 and 7108.72 aluminum alloys.  2.3 A G E H A R D E N I N G  The very incentive for the development of 7xxx series alloys is their ability to achieve high mechanical properties through age hardening. The age hardening is the result of the interaction of dislocations with precipitates, which are formed by decomposition of a supersaturated solid solution during ageing process.  2.3.1 T H E H E A T  TREATMENT  The 7xxx series aluminum alloys are heat-treated in the usual way as other age hardenable aluminum alloys. The first step is homogenization, which is followed by a  9  solution heat treatment, quench and an ageing treatment. D u r i n g these treatments, the microstructures of the materials undergo a variety of changes w i t h variations of mechanical properties associated w i t h the different stages of heat treatment. The homogenization process is employed to even out the distribution of alloying elements i n the casting materials by high temperature diffusion of the atoms.  It is a  necessary process for cast aluminum alloys, however, we w i l l not go further into the details of this process as it falls outside the scope of this w o r k . In the following we w i l l focus mainly o n the remaining three stages of heat treatment of the alloys. The goal of the solution heat treatment is to obtain a high level of alloying elements i n solid solution.  The solution temperature is chosen to fall i n the single phase (solid  solution) region i n the phase diagram, to obtain a solid solution.  T h i s temperature  depends o n the chemical composition of a particular alloy. A n important factor that is associated w i t h the solution treatment is the formation of an equilibrium vacancy concentration at the solution temperature.  D u r i n g quenching, i f the cooling rate is  sufficiently high, a non-equilibrium vacancy concentration may be produced w h i c h can significantly affect subsequent ageing steps.  Higher solution temperature favors the  higher degree of non-equilibrium vacancy concentration that can be obtained by quenching. However, undesired local melting of the alloy due to the segregation around second phase particles may occur if the solution temperature is too high (close to the solidus temperature line). The solution time is usually short, i n the order of minutes, compared w i t h homogenization times i n hours, since homogenization relies upon long range diffusion to even out the distribution of these alloying elements [7]. O n the other hand, the solution treatment only requires short range diffusion to dissolve precipitates formed during previous processes.  After solution heat treatment, the material w i l l undergo a quenching process, w h i c h brings it from solution temperature to a predetermined lower temperature.  10  A s shown  schematically i n Figure 2.4, the alloy i n solid solution region (point A ) goes into a multiphase region (point B) during the quenching process.  The solid solution at high  temperature is "frozen" i n position thus a supersaturated solid solution (SSS) is obtained. This is essential to the following ageing process. After quenching, the material is i n a non-equilibrium state, and there is a strong driving force for it to reach equilibrium conditions by the change from a single phase to a multiphase structure. This process occurs during ageing, either naturally (i.e. at r o o m temperature) or artificially (i.e. at elevated temperatures).  Often this process is very  complicated and involves the formation of Guinier-Preston (G-P) zones, clusters, and formation of a series or sequence of metastable phases such as intermetallic precipitates. Furthermore, the precipitate composition and crystal structure often go through a transition during ageing. These phenomena w i l l be discussed i n the following sections.  Temperature  Composition Figure 2.4: Schematic illustration of the route of phase transformation during quench.  11  2.3.2 PRECIPITATION PROCESS  After quenching, the system is frozen into a supersaturated solid solution condition that is i n a high free energy state and is therefore unstable. Thus, driven b y the nature of the material to stay on the lowest energy level, the material undergoes a series of changes i n microstructures and consequently, mechanical properties, during the various stages of heat treatment.  The generally accepted sequence of precipitation i n 7xxx series  aluminum alloys is [8, 9]:  SSS -> G - P zones -> n ' -> q( M g Z n ^  where r\' is the metastable transition phase, and n is the equilibrium M g Z n phase. The 2  structure of the G - P zones is not fully understood. However, their role i n determining the density of nuclei are thought to be extremely important. The precipitation kinetics usually display the classical behavior of nucleation and growth, followed by cubic coarsening, as G u y o t and Cottignies [10] stated i n their study using 7050 aluminum alloy as an example. The decomposition of a supersaturated solution can occur even during the quenching process, as soon as the temperature drops below the solute solvus temperature [11, 12]. Usually, this leads to a rather coarse distribution of precipitates (often o n grain boundaries or dispersoids), w h i c h adversely affects subsequent hardening. This is one of the reasons w h y the quenching rate, as well as the rate of heating up to ageing temperature (heating ramp), can affect the size and distribution of the precipitates, as we w i l l discuss later.  12  2.3.2.1 FORMATION OF G-P ZONES  The first result when ageing at l o w temperature (below the G - P zone solvus temperature, Tc) is the formation of the G - P zones. The G - P zones are very small i n their initial stage, and apparently, nucleate homogeneously.  They have a size i n the range of 0.5-1.0  nanometers, w h i c h is beyond the capability of resolution of the conventional T E M . However, their existence can be traced by the change i n the hardness/yield stress and resistivity of the alloy [13], since these material properties are very sensitive to the microstructure of the material. Direct observation of G - P zones is n o w possible using Field Ion Microscopy or H R T E M (High Resolution Transmission Electronic Microscope). F o r example, G - P zones have been directly observed by Mukhopadhyay, et al. [14]. These observations have shown that the shape of G - P zones i n A l - Z n - M g is spherical i n their initial stages. A s the Z n atom is smaller than A l atom, and M g is larger, the atomic size misfit can be compensated b y the formation of clusters containing both Z n and M g atoms resulting i n stress-free spherical shaped coherent clusters, to reduce the free energy of the system. Most researchers agree that there are t w o types of precipitate nuclei: solute rich clusters that form during l o w temperature ageing [15], and vacancy rich clusters that form during quenching [16]. The solute rich clusters dissolve partially or completely during artificial ageing, depending o n temperature.  The vacancy rich clusters are thought to be more  stable and may act as nucleation sites for the metastable phase r|' at appropriate ageing temperatures [17].  It had once been considered that only the solute supersaturation should control the nucleation of the precipitates.  However, the distribution of precipitates i n alloys of  similar solute supersaturations is frequently very different, and i n many alloys the  13  distribution is sensitive to the quench rate, the time of natural ageing, the rate of heating to the ageing temperature and the presence of trace elements, none of w h i c h appreciably affect the solute supersaturation.  A s a result, other factors must be taken into  consideration. A s the research into precipitation hardening systems proceeded, the role of quenched-in vacancies o n the nucleation of precipitates was recognized [18]. In fact, the precipitate distribution i n A l - Z n - M g alloys is largely controlled b y the distribution and concentration of lattice defects, particularly vacancies, as well as the dislocations and grain boundaries. A l t h o u g h traditional T E M is not applicable for direct observation of the nucleation process since point defects i n the crystal lattice are not visible i n this technique, information o n the role of vacancies can be still obtained b y a detailed study of the microstructure of the alloy after various heat treatments.  After quenching, a  typical aluminum alloy has a vacancy supersaturation of the order of 10 times higher 10  than equilibrium, i.e. much greater than the solute supersaturation.  T h e y tend to  precipitate as loops, helices and possibly small clusters or they may be absorbed into vacancy sinks such as grain boundaries or dislocations.  The depletion of vacancies  around these defects can lead to the formation of precipitate free zones. B y gathering information on the form and number of precipitates i n these regions, the role of vacancies i n the precipitation process can be revealed [18, 19].  2.3.2.2 FORMATION OF r)'  Further ageing causes the precipitation of particles i n the shape of t h i n plates. This is a transition phase, r\', and can be identified as having a hexagonal structure w i t h lattice parameters a=0.496 n m and c= 1.403 n m , as Park and A r d e l l [20] reported i n their study of 7075 aluminum alloy i n the peak aged T651 and overaged T 7 tempers.  14  T h e y also  observed that the peak strength mainly comes from the contribution of these small particles, w h i c h are finely dispersed. Thomas and N u t t i n g [21] also confirmed that the n ' precipitates are responsible for the m a x i m u m hardness when ageing at 130°C.  Some  researchers have suggested that the n ' transition phases w o u l d form o n the basis of initial G-P zones [8], while others reported that on increasing the ageing temperature to an intermediate range (100-130°C), the coherent n ' transition phase was found to form without evidence of prior G - P zone formation [22]. Details of heat treatment conditions should therefore be considered to determine the actual precipitation behaviour.  The role of G - P zones on the precipitation of the n ' precipitates has received considerable attention.  A useful approach to this question is to consider the metastable solvus  boundary for G - P zones. A critical temperature is reached where G - P zone formation is no longer favorable. However, at lower temperatures, G - P zones may form w i t h smaller zone sizes being stable at lower temperatures.  The next critical question is what w i l l  happen to these zones when the temperature is raised.  G - P zones formed at a lower  temperature w i l l dissolve on heating to a higher temperature if they are smaller than a critical size but w i l l transform to the intermediate precipitate at higher temperature if they are larger than a critical size. Whichever being the case, the nucleation stage is of most importance i n controlling the precipitate distribution, and hence, the final strength of the material.  The heat treatment after the formation of G - P zones can affect the final properties very significantly. Staley [23] found i n his research of copper bearing 7050 alloy that the G - P zones formed during natural ageing dissolve o n subsequent  exposure  at  ageing  temperature above a critical temperature (the G - P zone solvus) if the rate of heating to the elevated temperature is high.  The crystalline precipitate initiates heterogeneously  under these conditions (on grain boundaries and dislocations), and the strength is low. W i t h lower heating rates many of the G - P zones nucleated during natural ageing do not  15  dissolve at the elevated temperature.  They i n fact can grow to larger sizes w h i c h are  stable at higher temperatures, and can therefore serve as nuclei for the crystalline precipitate. This provides a large number of very fine particles and increases the strength attainable by a treatment at that ageing temperature.  2.3.2.3 FORMATION OF r\  The last stage of ageing for the 7xxx aluminum alloys is the formation of the equilibrium crystalline r| phase, M g Z n , and the subsequent coarsening of the n precipitates. 2  The  crystal structure of the n phase is hexagonal (C14) w i t h a = 0.520nm and c = 0.862 n m [18]. It has a lath like morphology. The n precipitates can be transformed from the metastable plate shaped TI' phase.  In some cases, it can also directly nucleate at  dislocations [24]. The formation and coarsening of r\ phase occur i n the overageing stage, i n w h i c h the strength of the material is decreased from the peak strength.  Therefore, it is of less  interest to the industry, where high strength is usually desirable.  2.3.2.4 PRECIPITATION KINETICS  G u y o t et al. [10] suggested that the precipitation kinetics for an A l - Z n - M g - C u alloy follows the classical behavior of nucleation, growth and coarsening. The nucleation and growth periods are difficult to quantify, due to the existence of transients occurring  16  before the stabilization of the ageing temperature.  However, the volume fraction of the  precipitates seems to follow a Johnson-Mehl-Avrami law w i t h a time exponent n of ~1.3, w h i c h is close to the theoretical value of 1.5 expected for the growth by long range diffusion of a constant number of precipitates.  The coarsening, or Ostwald ripening  occurs at a nearly constant precipitate volume fraction.  2.3.3 DISLOCATION-PRECIPITATE INTERACTION  The plastic deformation of metallic materials occurs by the m o t i o n of dislocations. In the absence of dislocations, the strength of a material w o u l d be very high (i.e. G/20). Practically, it is not possible to remove all dislocations and therefore  strengthening  requires that obstacles be added w h i c h make the dislocation m o t i o n more difficult. It is well established that the age hardening of the material results from the interaction between precipitates and mobile dislocations. The theory was developed i n the 1950's and 1960's, and a wealth of literature on this subject is available.  Several excellent  reviews may be found including those by K e l l y and N i c h o l s o n [25] and by E m b u r y [26]. Although we have a solid theoretical basis for the origin of the age hardening behaviour, the details of the precipitate-dislocation interaction is still of considerable  interest.  However, i n most cases, it is difficult to quantify the exact nature of the precipitatedislocation interaction. Generally, the presence of pre-precipitates (G-P zones or clusters) and precipitates produced by the ageing process described i n the previous section w o u l d increase the stress required for dislocations to pass through them, thus increasing the flow stress of  17  the plastic deformation. The resistance for dislocations to pass the obstacles may arise from several factors. First, when the particles are small, such as the clusters formed at the beginning of the ageing process, they can be sheared by dislocations. The resistance to shearing derives from a number of possible sources including chemical strengthening, modulus strengthening, coherency strain strengthening, and stacking fault energy strengthening [25].  The  resulting flow stress contribution generally has the form:  shearable  = q f V  (2.1)  where f is the volume fraction of the shearable particles, r the particle radius, c  u  m and n  are constants, and for most dislocation-particle interactions, both m and n have the value of 0.5 [27]. The dislocation interaction w i t h shearable particles is shown schematically i n Figure 2.5 (1) - (3). W h e n the particles grow larger as the ageing process advances, they become stronger obstacles and can no longer be sheared by dislocations. Instead, it becomes easier for the dislocations to bypass the precipitate, leaving a dislocation loop around the unbreakable particle (Figure 2.5 (i) - (iii)). This also gives rise to the flow stress of the dislocation slip. The flow stress contribution from this mechanism has the form:  CT  non-snearable  ~ c  z 2 L  (2.2)  where q is a constant, G the shear modulus, b the Burgers vector, and L the particle spacing. L can be expressed i n the following way:  18  3)  iii)  Figure 2.5: Schematic illustration of dislocation-particle interactions.  (1) - (3): moving  dislocation shearing a weak particle; (i) - (iii): moving dislocation bypassing a strong particle.  19  L  =  c 3  ^  (2-3)  where c is a constant, r is the radius of the particle, f the volume fraction of the particle. 3  Substituting Equation 2.3 into Equation 2.2 gives:  Gbf non-shearable  1/2  (2.4)  r  where c contains all the constants i n Equation 2.2 and 2.3. 4  Foreman  and M a k i n  [25] studied the  dislocation-particle interaction i n a  more  complicated situation b y using computer simulations, where a random array of particles of varying strength is presented. Each particle is characterized by its breaking angle cp , c  illustrated schematically i n Figure 2.6, w h i c h is the angle of the t w o arms of dislocation line when it overcomes the particle and advances to new obstacles. Thus the strength of a particle can be expressed i n its cp value. A weak particle has a (p « 180°, while a strong c  c  particle has a cp « 0°. The results from the simulations showed that the flow stress can be c  expressed as follows:  (9 < 100°)  (2.5)  (cp > 100°)  (2.6)  C  c  where G , b and L have the same meaning as i n Equation 2.2. F o r strong particles whose critical breaking angle cp « 0°, Equation 2.5 reduces to c  Equation 2.2 (with a factor of 0.8, w h i c h is related to the difference between the random  20  distribution and square arrangement of the particles), w h i c h gives the flow stress for nonshearable particles.  Figure 2.6: Schematic illustration of the breaking angle cp for a dislocation holding by an c  obstacle.  2.4 T H E E F F E C T O F PRE-DEFORMATION O N T H E PRECIPITATION B E H A V I O R  In industrial production of automobile bumpers using 7xxx series aluminum alloys, the material undergo a stretching process, where a strain of up to 0.3 is applied to the material at r o o m temperature i n the as-quenched condition. It is therefore necessary to consider the effect of this deformation o n the age hardening response of the 7xxx series aluminum alloys.  Pre-deformation prior to ageing can affect the precipitation behaviour i n several ways. C o l d deformation increases the dislocation density of the material significantly. This increased dislocation density can have t w o main effects:  21  i) dislocations can act as a preferential nucleation sites for precipitates ii) enhanced growth and coarsening of precipitates can occur due to the high rates of solute diffusion along dislocations. The role of the increased dislocation density introduced by the pre-deformation on the ageing behaviour varies for different aluminum alloys series [5]. L o w levels of pre-strain (less than strains of 0.1) have only a minimal effect o n the yield strength i n T 6 tempers and negative effect o n that i n T 7 tempers for 7xxx series alloys, while the reverse is observed i n the behavior of 2xxx alloys. In research on the influence of dislocations o n precipitation i n an Al-6.1%Zn-2.35%Mg (wt.%) alloy, Deschamps et al. [28] found that the pre-deformation decreases the overall mechanical properties attained b y precipitation strengthening, particularly at peak hardness as well as during the coarsening stage. Deschamps et. al. also observed changes i n the precipitation sequence.  T h e change i n  precipitation sequences was attributed to the decrease of the activation energy for the formation of the stable n phase o n dislocations, due to the relaxation of elastic misfit i n the presence of the dislocation. It also introduces spatial heterogeneity of precipitates and wide precipitate size distribution, w h i c h were proposed to be related to t w o effects: first, the dislocations introduced after  quenching act as vacancy sinks, decreasing the  precipitation rate of G - P zones at r o o m temperature.  The l o w level of G - P zones near  dislocations then makes subsequent formation of n ' and r\ phases more difficult (A similar effect as i n classical precipitate free zone near grain boundaries).  T h e authors  attributed the lowering of the peak hardness to a smaller solute content available for homogeneous finely dispersed precipitation and to a broader precipitate size distribution. Deschamps and Brechet [29] have theoretically studied the effect of precipitate size distribution and have shown that as the size distribution widens, the peak strength decreases, as shown i n Figure 2.7.  22  0.01  0.1  1  10  100  1000  Time (h) Figure 2.7: Influence of the precipitate size distribution w i d t h A o n the evolution of yield stress during ageing. (From Deschamps and Brechet [29].)  2.5 THE STATEMENT OF OBJECTIVES  The  objective of this study is to obtain better understanding of the effect of pre-  deformation o n the ageing behaviour and mechanical properties  (yield stress, i n  particular) of 7030.60 and 7108.72 aluminum alloys. W e hope that this effort w i l l help the industry to improve the performance of the aluminum bumper product, and ultimately, help to build a better environment.  H a v i n g reviewed the basic mechanisms involved i n precipitation hardening alloys, the following chapters w i l l describe the experimental techniques and results and then discuss the results.  23  Chapter  3  EXPERIMENTAL PROCEDURE  3.1 INTRODUCTION  The  general approach that has been taken to study this problem is to simulate the  industrial process i n the laboratory. Samples were taken from the as-extruded plate and then processed i n a similar manner to the industrial process described i n Section 2.2. However, the samples were pre-deformed by tensile loading or rolling i n the laboratory, instead of stretch bending w h i c h is used i n the industrial process. Different levels of predeformation w i l l be imposed on the samples to study the response of the material behavior.  3.2 MATERIAL  The  aluminum alloys involved i n this study were received from the Hydro-Raufoss  Automotive Research Centre i n the form of hot extruded plate of dimension 150x750x5 mm.  T h e y were taken from the production line and cut to length. T w o alloys (7030.60  and 7108.72) were examined and their chemical composition is shown i n Table 3.1. N o t e that the major difference between these two alloys is the intentional addition of zirconium to the 7108.72 alloy. The effect of Z r o n grain structure w i l l be discussed later  24  i n Chapter 5.  Table 3.1. Chemical composition of the 7030.60 and 7108.72 alloys (wt.%) Alloy  Zn  Mg  Cu  Zr  Al  7030.60  5.45  1.22  0.3  -  bal.  7108.72  5.45  1.20  0.27  0.16  bal.  3.3 METALLOGRAPHIC PREPARATION  The metallographic examinations were carried out to observe the grain structure of 7030.60 and 7108.72 alloys.  F o r each alloy, three samples representing the normal,  extrusion and transverse directions of the extruded plate were taken by cutting from the as-received plates using a S i C cutoff wheel w i t h liquid cooling.  The samples were  mounted i n the L E C O S E T 100 polymer resin, and then they went through a mechanical grinding and polishing process to obtain smooth surfaces. done using a 0.06 p m colloidal A l 0 2  3  The last polishing step was  powder on a polishing cloth. T h e etchant used to  reveal the grain structure was Graff and Sargent's etchant, w h i c h has the following composition:  HN0  15.5 m l  3  HF Cr0  0.5 m l 3g  3  H 0  84 m l  2  The samples were immersed i n the solution of the etching reagents for 80 seconds w i t h  25  m i l d agitation, then rinsed and dried for the optical microscopic observation.  The  U n i m e t U N I T R O N 8919 microscope and the Polaroid 55 black and white instant sheet film were employed to take the pictures of the grain structure.  3.4 EXPERIMENTAL HEAT AND MECHANICAL TREATMENT  F o l l o w i n g the industrial route described i n Section 2.2, the experiments were carried out so that the materials w o u l d experience a similar heat and mechanical treatment.  The  experimental schedule is shown schematically i n Figure 3.1. Temperature Solutionizing @480°C/20min.  IWater quench  t  - - - Ageing @100°C/5hr.  t  .' Ageing @150°C/up to 60days Pre-deformation  V  f  ^, Natural ageing/24hr./ J Time  Figure 3.1: Schematic illustration of the experiment schedule for the 7030.60 and 7108.72 aluminum alloys. The samples from both alloys were solution heat treated at 480°C for 20 minutes, followed by water quench (water temperature was ~20°C). The samples were then pre-  26  deformed to a pre-determined strain. A l l the pre-deformations were finished w i t h i n 15 minutes after water quench.  After the pre-deformation, the samples were aged for 24  hours at r o o m temperature (20°C) followed by a two step artificial ageing procedure which consisted of 100°C for 5 hours (in a boiling water bath) and 150°C for up to 60 days i n an air furnace. Various levels of pre-deformation (0, 0.03, 0.10, 0.30 and 1.20) were imposed o n the samples of both the 7030.60 and 7108.72 alloys to study its effect o n the ageing behavior. F o r the pre-deformation levels of 0.03 and 0.10 strains, the samples were taken along the extrusion direction of the as-received plates.  The tensile and resistivity samples were  prepared. These samples were elongated using a M T S (100 K N frame) servo-hydraulic testing machine, at a strain rate of 0.02 s" , to give the designated pre-strain. 1  The  deformation was measured using an extensometer on the sample. F o r the higher pre-deformation levels of 0.30 and 1.20 strain, small pieces of specimen (30x100x5 mm) were sheared from the as-received plate. T h e y were rolled i n a smallscale laboratory rolling m i l l , w i t h the rolling direction parallel to the extrusion direction of the original plate, at an estimated strain rate of 7.7 s' . 1  It is w o r t h noting that the  samples experienced a m i l d temperature rise (i.e. the sample temperature rose to 60-70°C) when they were rolled through the m i l l . After rolling, tensile samples were punched out from the rolled material using a die.  3.5 PROPERTY MEASUREMENTS  The ageing behavior of the alloys studied was measured by using hardness, resistivity,  27  and tensile tests, w h i c h were conducted on the samples at various stages i n the ageing process. The hardness property of the alloys was measured using the M I C R O M E T 3 micro hardness tester w i t h the pyramidal diamond indentor and a test load of 1 kgf.  The  electrical resistivity of the alloys for various ageing times was also measured i n order to study the microstructure change.  This was done by conducting measurements of the  electrical conductivity of the samples, which had gone through the thermal-mechanical treatments described i n Section 3.3, using a Verimet 4900C eddy current probe, and converting the conductivity data into resistivity values. The tensile tests were conducted either on an Instron screw driven tensile test machine or on the above mentioned M T S testing machine at a nominal strain rate of 1.5x1c s' . 4  1  T w o different sizes of tensile sample were used as shown i n Figure 3.2. The w i d t h and thickness of the gauge section area of each tensile sample were measured using a micrometer before the test, so that the value of the initial gauge section area could be obtained. A n extensometer was attached to the specimen during each tensile test, and the displacement-load data were acquired and stored i n a data file, from w h i c h  the  engineering stress and strain, true stress and strain data, and, upon o n further data analysis, the yield stress data, can be obtained. The engineering strain, engineering stress, true strain, and true stress were calculated using the following equations.  51 •eng.  (3-1)  0  __F  (3.2)  C7, o ^true  l 0  —  tftrue  n  =  +  (3.3)  Seng.)  (3-4)  CT ng.*ln(l + S g.) e  en  28  where 81 is the elongation, 1 the gauge length, F the tensile load, and AQ the initial gauge 0  section area. The yield stress value was determined using the 0.2% offset method.  12.5 mm 55 mm  155 mm 4.5 mm  - ' J 5.3  « A  20 mm  mm  •  70 mm  •  b  Figure 3.2 Schematic illustration of t w o types tensile sample used i n this research, a) for M T S tensile machine and b) for Instron tensile machine.  In summary, the industrial process has been simulated i n the laboratory. T h e evolution of properties was measured w i t h tensile tests and electrical resistivity. chapter, the main results from these experiments w i l l be summarized.  29  In the next  Chapter  4  RESULTS  The experimental results, including the grain structure examinations and ageing curves (in terms of yield stress, U T S , strain-to-necking and electrical resistivity) for both 7030.60 and 7108.72 alloys w i t h different levels of pre-deformation, are given i n this section. Several important observations are also summarized.  Some preliminary experimental  w o r k was conducted o n the as-received, hot-extruded plates of the 7030.60 and 7108.72 alloys to obtain the basic characteristics of these materials, such as the homogeneity of the strength through the plate thickness, the sensitivity to solution time and the quench sensitivity. These results, w h i c h helped to design the schedule for experiments thereafter, are given i n the Appendix.  4.1 GRAIN STRUCTURE OF THE AS-EXTRUDED MATERIAL  The grain structure of these two alloys (7030.60 and 7108.72) i n the as-extruded condition is shown i n Figures 4.1 and 4.2. T h e y show very different grain morphologies and grain sizes, although these alloys have been processed w i t h the same thermal-mechanical treatment (see Figure 2.2 i n Section 2.2).  The 7030.60 alloy has an equiaxed grain  structure (with an average grain size of 60 pm), indicating that the material recrystallized after hot extrusion. The 7108.72 alloy has a elongated, fibrous grain structure (with an average size of 5 p m i n the transverse direction and 20.5 p m i n the longitudinal direction), a remnant of the deformation structure, w i t h no evidence of recrystallization.  30  Figure 4.1: Tri-planar optical microscopic graphs showing the grain structure of the 7030.60 alloy i n the as-received condition.  31  Figure 4.2: Tri-planar optical microscopic graphs showing the grain structure of 7108.72 alloy i n the as-received condition.  32  Obviously the grain size for the 7108.72 alloy is much smaller than that of the 7030.60 alloy.  The role of Z r i n the 7108.72 alloy is undoubtedly critical i n making this  difference, as w i l l be discussed later.  4.2 AGEING BEHAVIOUR  D u r i n g the ageing process, the evolution of the microstructure i n the alloys is reflected by changes i n the microstructure sensitive properties, such as mechanical properties (yield stress and U T S ) and electrical resistivity. The ageing behaviour of the 7030.60 and 7108.72 alloys was studied by experimentally measuring the ageing curves i n terms of mechanical properties (yield stress, U T S and strain-to-necking) and electrical resistivity. The results are described i n the following sections.  4.2.1 EVOLUTION OF MECHANICAL PROPERTIES DURING AGEING  The mechanical properties are among the most important material characteristics as that they affect both the production process and the final service properties for the bumpers. The ageing behaviour of these alloys was examined by studying the evolution of mechanical properties (yield stress, U T S and strain-to-necking) of these alloys w i t h different levels of pre-deformation after the solution heat treatment. Tensile tests were conducted o n samples i n various stages of ageing:  1) as-quenched, 2) after the pre-deformation (as deformed),  33  3) after natural ageing for 24 hours, 4) during ageing at 100°C, 5) during ageing at 150°C.  4.2.1.1  TYPICAL TENSILE CURVES  Figure 4.3 shows some typical engineering stress-strain curves for the 7030.60 alloy w i t h no pre-strain. A number of important observations can be obtained from this Figure. 1) the yield stress increases during the natural ageing and artificial ageing up to 10 hours at 150°C (Figure 4.3a), after w h i c h it decreases (Figure 4.3b). 2) serrations are observed i n the engineering stress-strain curves for the samples i n as-quenched, naturally aged and artificially aged ( 1 0 0 ° C / 1 hr.) conditions. However, no serrations are observed i n the samples artificially aged for longer times. 3) qualitative changes i n w o r k hardening behaviour as a function of ageing time can be observed.  In particular, the w o r k hardening rate decreases as the  strength increases. 4) i n the overaged sample (Figure 4.3b curve c), an unusual stress-strain curve was observed. The initial w o r k hardening rate is l o w and then there is an increase i n the w o r k hardening rate after a strain of around 1%.  Figure 4.4 gives some typical engineering stress-strain curves for the 7030.60 alloy w i t h a pre-strain of 0.30. W h i l e observations are similar to those i n the non-deformed cases, there are a number of differences between these two pre-deformation conditions.  34  - as waterk quenched — 24 hr. natural c —  9 9 1 hr. at100°C  a  e i n  -5hr. at100°C •10 hr. at150°C (peak)  0.05  0.1  0.15  0.2  Strain  Figure 4.3a: Engineering stress-strain curves for 7030.60 alloy w i t h no pre-strain, i n various stages of ageing till the peak.  0  0.05  0.1  0.15  0.2  Strain  Figure 4.3b: Engineering stress-strain curves for 7030.60 alloy w i t h no pre-strain, i n various stages of ageing after the peak.  35  • as waterquenched - as rolled  450 400 350 (0 CL  300  • 24 hr. natural ageing  250 U) CO  <D \ W  •1 hr. at100°C  200 150  •5hr. at150°C (peak)  100 50  '  0  -  Figure 4.4a: Engineering stress-strain curves for 7030.60 alloy w i t h a pre-strain of 0.30, i n various stages of ageing till the peak.  -5hr. at150°C (peak) -70 hr. at150°C -60 days at 150°C  Figure 4.4b: Engineering stress-strain curves for 7030.60 alloy w i t h a pre-strain of 0.30, i n various stages of ageing after the peak.  36  1) stepped serrations occur i n the engineering stress-strain curves for the samples i n 24 hour natural ageing condition and 1 h o u r / 1 0 0 ° C artificial ageing condition (Figure 4.4a). 2) a substantial change i n w o r k hardening behaviour from the  as-quenched  condition to the as-deformed condition. 3) relatively large post-necking elongation i n the overaged samples (Figure 4.4b).  4.2.1.2 AGEING CURVES IN YIELD STRESS, UTS AND STRAIN-TO-NECKING  In this section, the results of mechanical properties evolution for these alloys w i l l be displayed. The ageing curves i n yield stress, U T S and strain-to-necking are plotted for each pre-stain level. Figure 4.5a shows the yield stress evolution for these alloys without pre-deformation during t w o step artificial ageing. Generally, the yield stress of the alloys first increases and reaches a peak, then it begins to decrease.  The same trend is observed for the  evolution of U T S , as shown i n Figure 4.5b. It is also shown i n Figure 4.5b that the strain-to-necking decreases during the second step of artificial ageing. It is interesting to note that while the yield stresses for the non-deformed alloys run parallel, they converge at 70 hours of artificial ageing at 150°C (see Figure 4.5a). This is observed only i n the non-deformed alloys.  Figure 4.6a shows the yield stress evolution for the alloys w i t h a pre-strain of 0.03. Similar to the curves shown i n Figure 4.5a, single-peak ageing curves are observed. However, the yield stresses for these pre-deformed alloys run at a nearly constant offset, and the yield stress for the 7108.72 alloy is consistently higher that the 7030.60 alloy  37  throughout the ageing process.  Figure 4.6b gives the evolution of U T S and strain-to-  necking for the pre-deformed alloys (0.03 pre-strain). The U T S follows the same trend as the yield stress i n Figure 4.6a. The strain-to-necking increases slightly i n the first stage of artificial ageing, then it decreases i n the rest of the ageing period. The mechanical properties for the alloys w i t h pre-strain of 0.10, 0.30 and 1.20 are given i n Figure 4.7, Figure 4.8 and Figure 4.9, respectively. Similar behaviour is observed for these alloys pre-deformed to larger levels, although the magnitudes of each curve are different.  A s the pre-strain level increases, the magnitude of the changes i n the yield  stress and U T S becomes less, and the strain-to-necking also decreases.  A detailed  comparison between the behaviour of materials w i t h different levels of pre-strain w i l l be summarized later. Table 4.1 gives the yield stress data for 7030.60 and 7108.72 alloys w i t h different levels of pre-deformation i n the as-quenched, as-deformed and overaged (60 days at 150°C) conditions. Table 4.1: Y i e l d stress data for 7030.60 and 7108.72 alloys w i t h different levels of predeformation i n the as-quenched, as-deformed and overaged conditions. Alloy 7030.60  7108.72  Pre-strain  Yield stress (MPa) As-quenched  Yield stress (MPa) As-deformed  Yield stress (MPa) Overaged (60 d.)  0  60  /  182  0.03  60  127  194  0.10  60  184  154  0.30  60  223  164  1.20  60  300  191  0  106  /  193  0.03  106  165  217  0.10  106  235  188  0.30  106  255  188  1.20  106  346  211  38  500 -•-7108.72, yield stress  450 400  -•-7030.60, yield stress  350 300 250 200 150 100 50 • 0  ageing at 100°C  ageing at 150°C  10  no pre-strain  100  1000  10000  Artificial Ageing Time / Hr.  Figure 4.5a: Evolution of yield stress for the 7030.60 and 7108.72 alloys with no prestrain.  1  10  100  1000  10000  Artificial Ageing Time / Hr.  Figure 4.5b: Evolution of UTS and strain-to-necking for the 7030.60 and 7108.72 alloys with no pre-strain.  39  500  -•-7108.72, yield stress  450 400  ro  —•—7030.60, yield stress  350  ^ CL 300 « 250 £ 200  CO  i——  150 100 50  ageing at 100°C  ageing at 150°C  0.03 pre-strain  0 10  100  1000  10000  Artificial Ageing Time / Hr.  Figure 4.6a: Evolution of yield stress for the 7030.60 and 7108.72 alloys with 0.03 prestrain.  ro w (A <D  co  10  100  1000  10000  Artificial Ageing Time / Hr.  Figure 4.6b: Evolution of UTS and strain-to-necking for the 7030.60 and 7108.72 alloys with 0.03 pre-strain.  40  500 -•-7108.72, yield stress  450 400  —•—7030.60, yield stress  350 300 250 200 150 100 50  "  ageing at  ageing at  100°C  150°C  0.10 pre-strain  0 10  100  1000  10000  Artificial Ageing Time / Hr. Figure 4.7a: E v o l u t i o n of yield stress for the 7030.60 and 7108.72 alloys w i t h 0.10 prestrain.  ro  CL CO CO  CO  10  100  1000  10000  Artificial AgeingTime / Hr. Figure 4.7b: E v o l u t i o n of U T S and strain-to-necking for the 7030.60 and 7108.72 alloys w i t h 0.10 pre-strain.  41  500 -•-7108.72, yield stress  450 400  —•—7030.60, yield stress  350 300 250 200 150 100 50 0  ageing at 100°C  ageing at 150°C 10  0.30 p r e - s t r a i n 100  1000  10000  Artificial A g e i n g T i m e / Hr.  Figure 4.8a: Evolution of yield stress for the 7030.60 and 7108.72 alloys with 0.30 prestrain.  7108.72, UTS 7030.60, UTS 7108.72, strain-to-necking 7030.60, strain-to-necking  0.4 0.35 0.3 0.25 c  0.2 0.15  '5 CO  0.1 0.30 p r e - s t r a i n 100  1000  0.05  10000  Artificial A g e i n g T i m e / Hr.  Figure 4.8b: Evolution of UTS and strain-to-necking for the 7030.60 and 7108.72 alloys with 0.30 pre-strain.  42  500  -•-7108.72, yield stress  450 400  -•-7030.60, yield stress  350 300 250 200 150 100 50 " 0  ageing at 100°C  ageing at 150°C 10  1.20 pre-strain 100  1000  10000  Artificial Ageing Time / Hr.  Figure 4.9a: Evolution of yield stress for the 7030.60 and 7108.72 alloys with 1.20 prestrain.  1  10  100  1000  10000  Artificial Ageing Time / Hr.  Figure 4.9b: Evolution of UTS and strain-to-necking for the 7030.60 and 7108.72 alloys with 1.20 pre-strain.  43  4.2.2 EVOLUTION OF ELECTRICAL RESISTIVITY DURING AGEING  T h e e v o l u t i o n o f electrical resistivity was also m e a s u r e d as a n o t h e r m e t h o d t o s t u d y t h e ageing b e h a v i o u r o f t h e materials, since it is also a m i c r o s t r u c t u r e d e p e n d e n t parameter f o r the materials. T h e results are c o m p i l e d a n d p l o t t e d i n F i g u r e s 4.10 - 4.14.  In F i g u r e  4.10, it is o b s e r v e d that t h e electrical resistivity f o r t h e n o n - d e f o r m e d alloys decreases slightly i n t h e 1 0 0 ° C ageing stage. I n the f o l l o w i n g 1 5 0 ° C ageing stage, it decreases m o r e rapidly.  T h e electrical resistivity o f 7108.72 a l l o y runs above t h e 7030 a l l o y i n t h e  b e g i n n i n g , t h e n t h e y c o m e close at 20-70 h o u r s o f artificial ageing at 1 5 0 ° C , a n d finally, t h e y separate again u p o n f u r t h e r ageing.  48 - » - 7108.72  46 E  d  o> O  X  J >  1  44 • 42  >—  — • »  1  -•-7030.60  40 38  CO  'co  CD  36 34 32  ageing at 100°C  ageing at 150°C 1  i ii i  10  n o pre-strain i  ,  i  100  1000  10000  Artificial Ageing Time / Hr. F i g u r e 4.10: E v o l u t i o n o f electrical resistivity f o r the 7030.60 a n d 7108.72 alloys w i t h n o pre-strain.  F o r t h e alloys w i t h a pre-strain o f 0.03, t h e electrical resistivity r u n s almost parallel w h i l e decreasing i n a s i m i l a r m a n n e r as t h e n o n - d e f o r m e d alloys, as s h o w n i n F i g u r e 4.11. T h e same t r e n d is also o b s e r v e d f o r the alloys w i t h 0.10, 0.30 a n d 1.20 pre-strain, as s h o w n i n  44  Figure 4.12, 4.13 and 4.14, respectively. It is worth noting that the rate at which the electrical resistivity decreases is much higher for the alloys with 1.20 pre-strain.  10  100  1000  10000  Artificial Ageing Time / Hr.  Figure 4.11: Evolution of electrical resistivity for the 7030.60 and 7108.72 alloys with 0.03 pre-strain.  E  d  > cn 'c/> CD  CC  10  100  1000  10000  Artificial Ageing Time / Hr.  Figure 4.12: Evolution of electrical resistivity for the 7030.60 and 7108.72 alloys with 0.10 pre-strain.  45  E  d  > to  'tn  or  10  1000  100  10000  Artificial Ageing Time / Hr. Figure 4.13: Evolution of electrical resistivity for the 7030.60 and 7108.72 alloys with 0.30 pre-strain.  48 46 E  d  0>  O  II 44 it  -•-7108.72  •  1  *  1  -•-7030.60  42 40  > or  38 36 34 32  ageing at 100°C  ageing at 150°C 10  ^  1  _^  ~rZr ,—  100  1.20 pre-strain i 1000  1  1  —  10000  Artificial Ageing Time / Hr. Figure 4.14: Evolution of electrical resistivity for the 7030.60 and 7108.72 alloys with 1.20 pre-strain.  46  4.3 SUMMARY OF SOME SIGNIFICANT OBSERVATONS  •  The grain structure of the as-received (hot-extruded) 7030.60 alloy is a recrystallized structure, while it is a unrecrystallized, fibrous structure for the 7108.72 alloy.  •  Generally, the yield stress and U T S for both alloys increase during ageing until they reach a peak, after w h i c h decrease upon further ageing.  •  F o r the alloys w i t h pre-deformation, both show a similar ageing behaviour i n the evolution of mechanical properties. T h e i r ageing curves run parallel to each other.  •  F o r the alloys without pre-deformation, a difference i n yield stress evolution was observed for the t w o alloys during the overageing region.  T h e yield stress first  converges at 70 hours of artificial ageing at 150°C and then at longer ageing times separates again.  Similar behaviour can be observed i n the evolution of electrical  resistivity for these alloys. Figure 4.15 summarizes the ageing curves (for ageing times of up to 70 hours at 150°C) for the solution heat treated and water quenched 7030.60 and 7108.72 alloys without predeformation, showing the evolution of yield stress.  The industrial standard T 6 heat  treatment schedule is also located o n the plot. This is the baseline w i t h w h i c h the other results w i l l be compared.  The first data point i n the plot stands for the yield stress of the alloy i n the as-quenched condition. The 7030.60 alloy i n such a condition has a yield stress of 60 M P a , while the 7108.72 alloy starts off at a higher value of 106 M P a . In the next stages, these t w o alloys behave similarly. D u r i n g the 24 hours of natural ageing, the yield stress of both alloys increases; and i n the 5-hour artificial ageing at 100°C, the yield stress increases at a higher rate.  In the following artificial ageing at 150°C, the yield stress keeps rising until it  47  reaches a peak, after w h i c h it decreases. It is w o r t h noting that although these t w o curves run  at an almost constant offset at the beginning, they meet at a time of 70 hours of  artificial ageing at 150°C.  T i m e after W a t e r Q u e n c h / Hr.  Figure 4.15: Y i e l d stress evolution during the ageing process for the 7030.60 and 7108.72 alloys without pre-deformation.  (The dashed line corresponds to the industrial T 6  schedule.)  The  ageing curves (for ageing times of up to 70 hours at 150°C) for these alloys with  different levels of pre-deformation are summarized i n Figures 4.16 - 4.17. Generally, the pre-deformed 7030.60 alloy behaves i n a similar manner, i.e., the yield stress rises to a peak and then decreases. However, important differences can be observed between these ageing curves.  48  CO  CO CO  S>  CO  co >-  12  24  36  48  60  72  84  96  T i m e after W a t e r Q u e n c h / Hr.  Figure 4.16: Y i e l d stress evolution during the ageing process for the 7030.60 alloy w i t h different levels of pre-deformation.  12  24  36  48  60  72  84  96  T i m e after W a t e r Q u e n c h / Hr.  Figure 4.17: Y i e l d stress evolution during the ageing process for the 7108.72 alloy w i t h different levels of pre-deformation.  49  First, the yield stresses at the 'starting point'.  A l l the ageing curves start at the as-  quenched condition, w h i c h shows a yield stress value of 60 M P a . However, according to the experiment schedule (Figure 3.1), the materials were pre-deformed immediately after water quench, then naturally aged at r o o m temperature.  The data points w h i c h fall  along the y-axis (i.e. at time « 0) i n Figure 4.16 correspond to the yield stress of the alloy i n the as-deformed condition, and they are designated as the 'starting point' of the ageing process. A s one can see, as the pre-deformation level increases (from the strain of 0 to 1.2), the yield stress of the as-deformed material increases significantly.  Second, there is a yield stress drop for the samples pre-deformed for a strain of 0.1 and up at the first hour of artificial ageing at 100°C. Moreover, the drop i n yield stress becomes larger as the pre-strain level increases. T h i r d , the peak yield stress for those samples w i t h pre-deformation are always lower than the peak yield stress of the non-deformed sample. Furthermore, w i t h careful scrutiny of these curves, it can be observed that the time to reach peak yield stress is decreasing as the level of pre-deformation imposed on the material increases. F r o m Figure 4.17, the similar observations can be made on the ageing behaviour of the 7108.72 alloy w i t h pre-deformation. These phenomena i n each ageing stage w i l l be discussed i n details i n Chapter 5. The resistivity evolution for both alloys during ageing is summarized i n Figure 4.18 and 4.19, respectively.  F o r the non-deformed materials, the electrical resistivity increases  during the natural ageing, then drops to a lower step during the artificial ageing at 100°C. D u r i n g the artificial ageing at 150°C, the electrical resistivity keeps decreasing, firstly at a  50  higher rate, then the rate of decrease is reduced.  0  12  24  36  48  60  72  84  T i m e after W a t e r Q u e n c h / Hr.  Figure 4.18: Electrical resistivity evolution during the ageing process for the 7030.60 alloy w i t h different levels of pre-deformation.  0  12  24  36  48  60  72  84  96  T i m e after W a t e r Q u e n c h / Hr.  Figure 4.19: Electrical resistivity evolution during the ageing process for the 7108.72 alloy w i t h different levels of pre-deformation.  51  F o r the pre-deformed materials, the electrical resistivity is increased i n the as-deformed condition. The more the pre-deformation is imposed on the sample, the larger the value of electrical resistivity increases. D u r i n g the artificial ageing period, the electrical resistivity of the pre-deformed samples follows the same trend as the non-deformed material. However, it is w o r t h noting that the rates at w h i c h the electrical resistivity decreases are different for the materials predeformed to different level. The rate is higher as more pre-deformation is imposed on the material.  52  Chapter  5  DISCUSSION  There is a wealth of information to be explored i n the experimental results we obtained. In this chapter, we w i l l give a detailed analysis o n the results i n order to unveil the ageing characteristics of these materials.  The ageing behaviour of the materials i n each  individual ageing stages w i l l be discussed.  5.1 EFFECT OF Zr ON THE GRAIN STRUCTURES  The grain structure of an alloy is the result of its chemical composition and the thermalmechanical history it experienced. The grain structures of the 7030.60 and 7108.72 alloys after the hot extrusion appear to be very different, as illustrated i n Figures 4.1 and 4.2. Considering the fact that these t w o alloys have the same thermal-mechanical history (see Figure 2.2), and almost the same chemical composition i n the major alloying elements of Z n , M g and C u (see Table 3.1), it is suggested that the difference i n the trace elements (Zr addition) plays a role on the microstructure of the material that leads to the present result. The 7030.60 alloy has no Z r content. D u r i n g or after the hot extrusion, recrystallization may occur w i t h the help of the thermal energy, leading to the equiaxed coarse grain structure.  53  The 7108.72 alloy has a Z r addition of 0.16 wt.%. It is well documented that Z r acts as a grain refiner  and recrystallization inhibitor i n aluminum alloys [3, 4].  It  also  simultaneously improves the stress corrosion cracking resistance of the 7xxx series aluminum alloys [30]. Z i r c o n i u m has l o w solubility i n the 7xxx series aluminum alloys. In the as-cast ingot, Z r usually stays i n supersaturated solid solution.  D u r i n g the preheating i n the extrusion  process, Z r precipitates as A l Z r intermetallic particles to form dispersoids i n the 3  aluminum matrix. Due to the very l o w diffusion coefficient of z i r c o n i u m i n aluminum, the A l Z r particles are resistant to coarsening. 3  Moreover, since they are formed by a  solid-solid reaction, their size is small, ranging from 0.02 to 0.5 p m [5]. So the A l Z r 3  dispersoids act effectively as grain refiner and recrystallization inhibitor because of the Zener drag phenomenon [31], resulting i n the fibrous, unrecrystallized grain structure for the Z r bearing 7108.72 alloy after hot extrusion. The difference  i n the grain structure between these two alloys may affect  their  mechanical properties, as we w i l l see later.  5.2 OVERVIEW O F T H E A G E I N G B E H A V I O U R  Referring to the T E M works done by Stiller, et. al. [32] and Saeter, et. al. [33] on the same alloys, the general sequence of precipitation is given by: SSS -> G - P zones -> n ' - » n  where n ' is the metastable precipitate and n the equilibrium precipitate.  M a n y variants  of the precipitates, such as r\ n and r| , etc., have been reported [33]. G - P zones nucleate u  2  4  54  and grow during the natural ageing period, and they are also found to survive even at the beginning of the 100°C artificial ageing process.  The r\' phase is formed during the  artificial ageing process, and may transform into rj phase at some time i n the 150°C ageing process [32]. F o l l o w i n g the series of microstructure evolution, the mechanical properties (yield stress, U T S , etc.) and electrical resistivity of these alloys changes accordingly.  These microstructure-sensitive properties provide us a w i n d o w through  w h i c h the changes i n the microstructure can be traced.  The two-step ageing process is designed to take the full advantage of the precipitation hardening.  The natural ageing and the first step of artificial ageing are employed to  generate a high density of G - P zones, so that they can act as nucleation sites for the formation of the metastable n ' precipitates.  In such a heat treatment schedule, very  finely dispersed T I ' precipitates can be obtained, w h i c h provide high ageing strength, according to the precipitation hardening mechanism discussed later.  5.2.1 YIELD STRESS EVOLUTION DURING AGEING  The yield stress is an important mechanical property for the material that is to be adopted for making bumper product. study.  Therefore, it receives the most attention i n this  The yield stress is determined by many factors, such as chemical composition,  thermo-mechanical history, temperature, etc. It is a microstructure sensitive parameter. In this study, the most significant contributors to the yield stress property are precipitation hardening, dislocation hardening and solid solution hardening.  These  factors either make their o w n contributions independently, or they can interact w i t h each other and make the result no longer accountable by a simple linear addition law.  55  According to the classic theories of age hardening [34], the precipitates contribute to the strength of the materials i n t w o ways. W h e n they are small and weak, the flow stress for dislocation shearing is less than that for dislocation bypassing, so they are sheared by dislocations.  The overall strength is dominated by the shearing mechanism.  Since a  dislocation has some flexibility, the number of precipitates it touches per unit length increases as r increases and the precipitates grow stronger.  So the flow stress increases.  Their flow stress contributions can be expressed i n Equation 2.1, namely,  ^shearablf = q f V  (2.1)  W h e n the precipitates grow larger i n the coarsening stage and become strong obstacles to the dislocations, the dislocation-particle interaction changes to bypassing. T h e flow stress is then dominated by the bypassing mechanism and the flow stress is determined by Equation 2.4, namely,  Gbf  = c  CT non-shearable  4  , (2.4)  1 / 2  x  j-  v  '  In the coarsening stage, the increases i n the size of the precipitates usually leads to the increasing of the precipitate spacing L , since the volume fraction of the precipitates remains nearly constant.  Thus, under the dislocation bypassing mechanism, the flow  stress decreases as the precipitates grow larger. A t a certain point where the radius of the precipitates is r , the flow stresses for dislocation shearing and bypassing are equal, and it 0  reaches its m a x i m u m .  This is where the transition from dislocation shearing to  bypassing occurs, and also the peak strength occurs during the ageing process. Figure 5.1 illustrates the evolution of yield stress as a function of the precipitate size r, schematically. Bearing i n m i n d that the precipitate size r is a function of ageing time, this qualitatively explains the shape of the ageing curves i n Figure 4.15.  56  • *  %  Precipitation strength due to \ dislocations by*» passing particles * • * * * * *  _. *  Precipitation strength due to dislocations shearing particles  \ ^  >— Overall precipitation strength  Figure 5.1: Schematic illustration of the flow stress as a function of the obstacle radius r.  It is widely accepted that the age hardening of the 7xxx aluminum alloys is i n most cases associated w i t h the disperse precipitation of the metastable n ' phase, however, the details of the strengthening mechanism remain unclear. complicated process.  The dislocation-n interaction is a 1  A t present, the information on the dislocation-n' interaction is  very limited, due to the experimental difficulties. It is even u n k n o w n that whether the n ' phase is shearable or non-shearable, so the transition point can not be determined.  A n alternate approach to examine the age hardening response is to consider the precipitates as a series of obstacles lying on the glide plane. D u r i n g ageing, the obstacle strength (which may be characterized b y the critical dislocation breaking angle (pj increases.  Simultaneously the averaging, L , is also changing, first due to nucleation of  new precipitates and later due to coarsening. F r o m this point of view, the problem can be examined using the results of Foreman and M a k i n (i.e., Equation 2.5 and 2.6):  57  (9 < 100°)  (2.5)  (  (2.6)  C  9 e  > 100°)  where L is the precipitate spacing and 9 the critical breaking angle w h i c h is related to C  the precipitate strength.  A t the initial stage of precipitation, the increasing cos(9 /2) c  term, due to the precipitates growing stronger, is more than enough to compensate for the decreasing term (1/L), due to the increase i n precipitate size r. So the overall flow stress increases. However, the increase i n the value of cos(9 /2) term is limited. A t a c  certain point where the decreasing of the term (1/L) can not be compensated b y the increasing i n the cos(9 /2) term, the overall flow stress due to precipitation hardening c  then begin to decrease.  This is an alternative explanation for the occurrence of the peak  i n the age hardening curve. complementary.  It is w o r t h noting that these t w o views are actually  F o r example, Equation 2.1 can be derived from Equation 2.6 if it is  assumed that the obstacle strength of the precipitate is proportional to the size of the precipitate [29]. It is also w o r t h noting that all the arguments above assume that all precipitates have the same size at the same time. W e shall see later that the precipitate size distribution may have an interesting influence on the precipitation hardening response.  5.2.2 ELECTRICAL RESISTIVITY EVOLUTION DURING AGEING  The electrical resistivity measurement has been adopted by many authors to study the  58  precipitation process [10, 35, 36].  In fact, the electrical resistivity can be used as an  indirect size measurement of the precipitates [35].  In this w o r k quantitative analysis of  the electrical resistivity data proved to be more complicated than expected, however, useful information, for example, o n the precipitate coarsening kinetics, can be obtained by a qualitative analysis. V e r y small G - P zones can scatter free electrons and give a rise to an increase i n electrical resistivity.  This happens during the formation of the G - P zones, and explains the  resistivity increase during the natural ageing period, shown i n Figures 4.18 and 4.19. The m a x i m u m resistivity is reached when the size of these clusters is of the same order as the mean free path of conduction electrons [35]. W h e n the precipitates grow larger, they are less  effective  in  scattering  the  electrons,  and  their  population  also  decreases.  Furthermore, the matrix is depleted as well, therefore the electrical resistivity decreases. B y this means the electrical resistivity provides a statistical measurement characteristic of the precipitate size, particularly i n the artificial ageing stages.  F r o m Figure 4.18, it can be observed that i n the 150°C ageing stage, it takes less time for the 7030.60 alloy w i t h higher level of pre-deformation to reach the same resistivity value. This implies that the precipitates grow faster as the pre-strain level increases. The same is also true for the 7108.72 alloy, as seen from Figure 4.19.  5.3 EFFECT OF PRE-DEFORMATION ON THE YIELD STRESS OF THE AS-DEFORMED MATERIAL  F r o m Figures 4.16 - 4.17, it is evident that pre-deformation has several effects o n ageing behaviour of both alloys. Starting w i t h the pre-deformation process, we w i l l examine the effect of pre-deformation o n the as-deformed materials i n this section.  59  The pre-deformation process was carried out immediately after solution and water quench. Figure 5.2 illustrates the yield stresses of the 7030.60 and 7108.72 alloys i n the as-deformed condition as a function of pre-strain. The curve quite resembles the plastic region i n a tensile curve i n shape (although to much higher strain that is achievable i n a tensile test). The yield stress increases as the pre-strain increases, and it does not seem to have saturated, even at the pre-strain of 1.2.  450  j  400 -  Strain  Figure 5.2: Y i e l d stress of the 7030.60 and 7108.72 alloys i n the as-deformed condition as a function of pre-strain.  It can be observed that there is a nearly constant difference i n the yield stress for the two alloys, as shown i n Figure 5.2, and this offset remains more or less constant during ageing process. This may be attributed to the different grain structures (i.e., grain size and subgrain size effects) and textures [37] between these t w o alloy extrusions.  However, this  phenomenon is not comprehensively understood since the yield stress is a complicated function of many material parameters.  60  After solution and water quench, the material is brought into supersaturated solution condition.  solid  The solute atoms make their contribution to the flow stress.  Meanwhile, the pre-deformation process introduces new dislocations into the material, and makes the material w o r k harden due to the increased dislocation density. A s the level of pre-strain increases, the dislocation density becomes higher. A n approximation for the increases i n dislocation density can be made from the following equation [38], a = a  0  + MaGbp  (5.1)  1 / 2  where M is the T a y l o r factor (3.06 for random texture; for 7030.60 and 7108.72 alloy, texture effects are ignored i n this approximation), a is a constant w i t h a value of 0.2 for F . C . C . metals, G the shear modulus, b the magnitude of the Burgers vector and p the dislocation density. The dislocation density data calculated from Equation 5.1 is shown i n Figure 5.3. The higher dislocation density i n the alloy w i t h higher level of pre-deformation results i n the increased flow stress. The presence of a high dislocation density i n the pre-deformed materials can have many effects on the behaviour of the material. It is proposed that the dislocations can act as vacancy sinks [39].  D u e to this effect, the excess vacancy concentration obtained by  quenching from solution temperature is lower i n the pre-deformed materials where the dislocation density is higher than the non-deformed material. The dislocations can also act as heterogeneous nucleation sites for precipitates. So the distribution of dislocations i n the material has a strong, if not dominant, influence on the distribution of precipitates. The dislocation lines can also provide a high diffusivity short circuit diffusion path for solute atoms. This enhances the overall diffusion process, w h i c h dominates the kinetics of precipitate growth and coarsening. The dislocations also deplete solute atoms from  61  the matrix, due to a solute flux to dislocations caused by the interaction between elastic stress field around the dislocations and solute atoms.  The result is the lower solute  fraction available for bulk precipitation.  30  0 0  0.2  0.4  0.6  0.8  1  1.2  Strain  Figure 5.3: Plots of [ ( a - a ) / M a G b ] 0  2  vs. strain, illustrating the increases of dislocation  density after the pre-deformation for the 7030.60 and 7108.72 alloys.  5.4 EFFECT OF PRE-DEFORMATION ON THE NATURAL AGEING BEHAVIOUR  D u r i n g the natural ageing period, the yield stresses of the two alloys rise, as seen from Figures 4.16 - 4.17.  Figure 5.4 summarizes the yield stress gain for these alloys w i t h  different levels of pre-strain during the 24-hour natural ageing period.  62  It is clear that the gain i n yield stress becomes less as the pre-strain level increases. This indicates that the natural ageing response is lowered for the pre-deformed material. It is k n o w n from the literature that the precipitation reaction during natural ageing is the formation of G - P zones. These G - P zones contribute to the strength of the alloys by acting as weak obstacles to dislocations. The nucleation of G - P zones is strongly affected b y the excess vacancy concentration i n the material. In the non-deformed alloys, the excess vacancy concentration introduced by water quenching is available to enhance the nucleation of G - P zone zones, resulting i n a high density of G - P zones. O n the other hand, for the materials w i t h pre-deformation, the excess vacancy concentration is lowered by the increased number of dislocations acting as vacancy sinks. Therefore the nucleation of the G - P zones is suppressed, leading to a lower density of G - P zones. In this case, the contribution to the strength of the alloy is less, resulting i n the lowered natural ageing response for the pre-deformed materials.  0.2  0.4  0.6  0.8  1.2  Strain  Figure 5.4: Changes i n yield stress during the 24-hour natural ageing process as a function of pre-strain.  63  5.5 EFFECT OF PRE-DEFORMATION ON THE ARTIFICIAL AGEING BEHAVIOUR  F o l l o w i n g the natural ageing process, the materials were then artificially aged i n a two step ageing process. The ageing temperature for the first step is 100°C and the second is 150°C. The ageing behaviour of the material i n the artificial ageing process is discussed i n this section.  5.5.1 ARTIFICIAL AGEING STAGE I): AGEING AT 100°C  The yield stress of the non-deformed material rises rapidly i n this stage, as seen i n Figure 4.15.  The increasing yield stress is attributed to the increasing number of the r|'  precipitates w h i c h are formed i n this stage [32]. A n interesting material behaviour was observed i n this ageing stage. A yield stress drop occurs for some of the pre-deformed alloys i n the first hour of artificial ageing at 100°C, as shown i n Figures 4.16 and 4.17. The changes i n yield stress during this period are summarized i n Figure 5.5. F o r the 7030.60 alloy w i t h 0 and 0.03 pre-strain, the yield stress increases.  F o r all other cases, the yield stress drops w i t h the amount of drop  becoming larger as the pre-strain level increases  In the stage of artificial ageing at 100°C, there are t w o main strength contributors for the pre-deformed materials (the intrinsic strength of the aluminum matrix is always default), one is the precipitation strengthening and the other is the dislocation strengthening. Since no yield stress drop is observed i n the non-deformed materials (which w o u l d be an indication  of no  G - P zone dissolution), the  64  other  contributor, the dislocation  strengthening, seems to be responsible for the observed yield stress drop.  Strain  Figure 5.5: Change i n yield stress between the end of natural ageing and after 1 hour at 100°C.  W e suggest that static recovery occurs i n the pre-deformed materials during the artificial ageing process. After cold deformation, dislocation density becomes higher, and energy is stored i n the material.  This provides the driving force for the recovery to occur.  W h e n the material is artificially aged at a higher temperature, the thermal energy makes the recovery process kinetically possible. A s a result of static recovery, the dislocation density decreases, therefore its contribution to the strength becomes less.  The higher  degree of pre-strain imposed o n the material, the higher the driving force for recovery, so the dislocation density recovers more quickly, showing a larger drop i n the yield stress.  Further evidence for the static recovery w i l l be provided by analyzing the data from the overageing stage.  65  5.5.2 ARTIFICIAL AGEING STAGE II): AGEING AT 150°C  It can be clearly observed from Figures 4.16 and 4.17 that the peak strength and T 6 strength for the materials vary as the pre-deformation is imposed on the materials. It is w o r t h emphasizing again that the T 6 strength and the peak strength for the pre-deformed materials are lower than the non-deformed material. Referring to Figure 2.1, the corner area of the bumper receives a larger deformation (approximately 0.3 of strain) than the center during the bumper forming process. Therefore, the material i n this area should have a relatively higher level of w o r k hardening, due to dislocation hardening. Superimposed w i t h the precipitation hardening during the heat treatment, it is natural to think that the strength of the material at the corner area should be higher than the material at other area of the bumper. However, the experiments simulating the industrial process give the reverse result (see Figures 4.16 and 4.17 for the T 6 strength). In fact, the T 6 strength for the material w i t h 0.3 pre-strain is significantly lower than the non-deformed material.  This leads to the spatial  heterogeneity i n the strength throughout the bumper product.  A similar effect is observed for the peak strength of the materials. It is unusual that the sum of the dislocation hardening and precipitation hardening w o u l d be less than if we only had the precipitation hardening. Figure 5.6 summarizes the peak strength for both alloys as a function of pre-strain. Generally, the peak strengths for the materials w i t h pre-deformation (all levels for the 7030.60 alloy, 0.03 and up for the 7108.72 alloy) are lower than that without pre-deformation. Each of these t w o strengthening mechanisms, the dislocation hardening and the precipitation hardening, contributes to the strength of the material positively, when they are acting individually.  However, the total amount of strengthening obtained from a  66  combination of these t w o mechanisms, is less than the contribution from o n l y one of them w o r k i n g alone for this alloy system.  This scenario leads us to the following  suggestion: some sort of negative interaction between these t w o mechanisms exists under these circumstances. In the following, we w i l l discuss the negative effects of dislocations on the precipitation strengthening.  450  Strain  Figure 5.6: Variation i n the peak strength as a function of pre-strain.  A s it is k n o w n , the precipitation hardening is dominated b y the evolution of precipitate size and spacing. The precipitate size measurements were conducted i n the T E M works on the same alloys used i n this study by Salter, et. al. [33].  Figure 5.7 shows T E M  micrographs for the 7030.60 alloy, demonstrating influence of pre-deformation o n the precipitation behaviour. In the non-deformed alloy, the precipitates appear to be small and have a more or less rounded shape.  W h e n dislocations are introduced into the  material, the precipitates are larger, they vary more i n size and tend to differ more i n shape than i n the non-deformed condition.  67  b) Figure 5.7: TEM-micrographs from a) the non-deformed and b) pre-deformed (0.10 strain) conditions for the 7030.60 alloy aged for 70 hours at 150°C. (From Saeter, [40].)  The results of the precipitate size measurement are summarized i n Table 5.1.  68  Table 5.1: Summary of precipitate sizes after ageing at 1 0 0 ° C / 5 h r . + 1 5 0 ° C / 7 0 h r . (From Salter, et. al. [33]) Condition  A v g . size (nm)  S T D (nm)  M a x size (nm)  Avg. size/STD  7030.60 0 pre-strain  6.3  3.3  19.3  1.9  7030.60 0.1 pre-strain  11.3  8.5  53.2  1.3  7108.72 0 pre-strain  9.4  4.3  24.7  2.2  7107.72 0.1 pre-strain  13.4  8.6  56.5  1.6  It is clear that the precipitates are larger i n the pre-deformed materials than i n the nondeformed ones.  (Our electrical resistivity measurements also indirectly confirm this  qualitatively, see Figures 4.18 and 4.19).  Since the precipitate coarsening process is  controlled by diffusion, it is apparent that diffusion is enhanced b y additional short circuit diffusion paths, such as along dislocation cores, w h i c h are added to the normal bulk diffusion mechanism when more dislocations are introduced into the materials by cold w o r k . If this occurs, the precipitates are able to grow more quickly and rapidly reach the critical size where peak strength is developed. Therefore the shortening of time to reach peak strength can also be expected. In fact, this has been observed i n our results, as summarized i n Figure 5.8. The time to peak strength is shortened from 10 hours for the non-deformed materials to approximately 4 hours for the materials w i t h a pre-strain of greater than 0.3.  The T E M studies by Salter, et. al. [33] o n the same alloys used i n this study also give some other important observations.  First, they found that the dislocations organize  themselves into a loose cell structure, consisting of very broad cell walls and dislocationfree cell interiors i n the 7030.60 alloy w i t h 0.1 pre-strain, shown i n Figure 5.9a.  69  This  feature appears i n the 7108.72 alloy as well.  In Figure 5.9b, it is shown that the  precipitates i n the cell walls are much larger where the dislocation density is higher, indicating that the dislocations play an important role i n contributing to the nucleation and growth of precipitates.  A n i n situ small angle X-ray scattering (SAXS) study of an A l - Z n - M g - C u alloy by Gomiero, et. al. [41] shows that there is a significant difference i n the size of precipitates at the dislocation sites and i n the bulk. The precipitates at the dislocations are larger than the bulk precipitates.  T h e y also found that deformation prior to ageing accelerates  precipitate coarsening during ageing at 160°C (which is the second step i n a t w o step artificial ageing process), especially heterogeneous precipitation o n dislocations.  70  0.3  m j  b) Figure 5.9: T E M micrographs of 7030 alloy w i t h 0.1 pre-strain, aged for 5 hours at 100°C plus 20 hours at 150°C. a) dislocations i n contrast and b) dislocations out of contrast, showing the precipitates. (From Saeter, et. al. [33].)  71  Saeter, et. al. [33] also studied the precipitate size distributions o n our alloys. The results are shown i n Figure 5.10. The log-normal distribution for the precipitates are observed for these curves. The difference between the pre-deformed and non-deformed materials is that the tail of the size distributions is much longer for the former one. Furthermore, the existence of t w o peaks, one large and one small, i n these distributions indicates that there are t w o sets of precipitates i n the system.  W e suggest that t w o groups of  precipitates can be distinguished i n the pre-deformed materials: one for the bulk precipitates, and the other for precipitates on dislocation sites.  It is clear that the  precipitates o n dislocations are larger i n size than the bulk precipitates (see Figure 5.9). This is due to the enhanced diffusion provided along dislocation cores. Figure 5.11 gives a schematic illustration of the precipitate size distribution i n the pre-deformed materials, based on the measurement of Salter [33].  Size [nm]  Size [nm]  (a)  (b)  Figure 5.10: Precipitate size distributions for pre-deformed and non-deformed a) 7108.72 alloy and b) 7030.60 alloy, aged at 100°C for 5 hours followed by 70 hours at 150°C. (From Saeter, et. al. [33].)  72  A s the levels of pre-deformation increases, the dislocation density i n the material increases as well (a sample estimate was made i n Section 5.3). The leads to the following results: the heterogeneous nucleation of precipitates at dislocation sites is enhanced. O n the other hand, as more solute atoms are depleted from the matrix, the bulk precipitation becomes more difficult. These t w o process culminate i n the variation i n the distributions of precipitates between the bulk and dislocations. H i g h e r levels of pre-deformation w i l l result i n the rising of the peak for precipitation o n dislocations and the lowering of the peak for the bulk precipitation i n Figure 5.11.  Bulk precipitates  -  Precipitates on dislocations  t\  - /  I  1  1  1  ,  1  1  1  1  1  1  0 Size  Figure 5.11: Schematic illustration of the size distributions for bulk precipitates and precipitates on dislocations for the pre-deformed materials.  A s a brief summary, the presence of pre-deformation leads to higher dislocation density, larger precipitate size, and a t w o peak precipitate size distributions consisting bulk precipitation and precipitation on dislocations. A s the pre-deformation level increases, the precipitation on dislocations is enhanced. F o l l o w i n g the arguments i n Section 5.2.1, we n o w consider the evolution of strength under the condition that there is a precipitate size distribution as shown i n Figure 5.10.  73  Assuming that the peak strength occurs at a precipitate size r , regardless of the c  mechanisms whether it is a shearing-bypassing transition or the precipitate spacing (L) term offsetting the precipitate strength term (cpj.  U n d e r the condition that all the  precipitates have the same size at the same time, all the precipitates make their full contributions to the flow stress. However, when there is a precipitate size distribution, only a fraction of precipitates has the critical value of r . The remaining fraction contains c  precipitates either larger or smaller i n size than r , and make a smaller contribution to the c  flow stress than those w i t h a size of r (i.e., there are weak links i n the chain). A s a result, c  the whole population of precipitates do not make same contribution to the flow stress simultaneously.  Thus the lower peak strength is expected.  This explains w h y the  presence of pre-deformation leads to a lowered peak ageing strength of the alloys. The w o r k b y Deschamps [42] shows a similar results (see Section 2.4).  5.5.3 K I N E T I C S O F P R E C I P I T A T E  COARSENING  The precipitate coarsening occurs i n the overageing stage. Table 5.2 summarized the data of the precipitate size for the non-deformed alloys i n the overaged conditions. Figure 5.12 illustrates the average precipitate size versus ageing time at 150°C. F o r the 7108.72 alloy, a straight line can be fitted to the log-linear data w i t h a slope of 0.33, giving a coarsening law of:  r oc t°-  33  It is therefore apparent that the coarsening of the precipitates during the overageing stage follows the cubic coarsening law (i.e. exponent = 0.333).  74  Table 5.2: Summary of precipitate sizes for the non-deformed alloys i n the overaged conditions. Condition 7030.60 20 hr./150°C 7030.60 70 hr./150°C 7030.60 60 days/150°C 7108.72 20 hr./150°C 7107.72 70 hr./150°C 7107.72 60 days/150°C  Avg. size (nm)  STD (nm)  Max size (nm)  Avg. size/STD  5.4  -  -  -  6.3  3.3  19.3  1.9  23  10  72  2.3  6.3  2.4  16  2.7  4.3  24.7  2.2  12  80  2.0  9.4 24  .  5.6 RECOVERY DURING THE ARTIFICIAL AGEING  The overageing process (60 days/1440 hours at 150°C) greatly reduces the yield strength of the alloys. The evidence for static recovery can be revealed by the following analysis  75  o n the experimental data. Referring to the data i n Table 4.1 i n Section 4.2.1.2, we collect the yield stress data for the 7030.60 alloy w i t h pre-strain of 1.20 i n the as-quenched condition, o" ^, as-deformed a  condition, CT , and overaged (150°C/60 days) condition,CT^,as shown i n Table 5.3. AJ  Table 5.3. Y i e l d stress data for 7030.60 alloy w i t h pre-strain of 1.20 i n the as-quenched, as-deformed and overaged conditions. Alloy  Pre-strain  Yield stress (MPa) As-quenched  Yield stress (MPa) As-deformed  Yield stress (MPa) Overaged (60 d.)  1.20  60  300  191  7030.60  A t each ageing stage, the flow stress components are illustrated i n Figure 5.13. In the as-quenched condition, a . - o + o„ a  q  (5.2)  0  whereCT is the intrinsic strength (including grain size effect) for the aluminum matrix, 0  w h i c h can be estimated to a value of approximately 25 M P a [34], and  is the strength  contribution from the supersaturated solid solution. In the as-deformed condition,  a-d  CT  =  a-q  CT  +  a  di  s  (  t o  t [ i e  first order of approximation)  76  (5.3)  where the a  is the amount of w o r k hardening (dislocation hardening) obtained during  &  the pre-deformation. Therefore, ^d, = ^ - a ^ = 2 4 0 M P a . s  350 300  CD CL  Assuming no recovery  250 Experiment  200 CO  CO  °"dis.  °>pt  150  residual  °"di!  + CD 0"ss. residual  100 50  °"o  ° 0  0  As quenched  Ageing  Figure 5.13:  Overaged 150°C/60d.  As deformed  Overaged 150°C/60d.  Conditions  F l o w stress components for the 7030.60 alloy w i t h pre-strain of 1.20 i n  various ageing stages.  In the overaged condition, o n l y a fraction of solid solution strengthening precipitation strengthening  (o"  p p t  )  r e s  iduai  (aj  residual  and  remain. If there is no recovery of the dislocation  hardening at all, then,  77  o-a~  CT  °"0 +  (°"ss)residual +  Even taking the values of both (CTj  dis  CT  residual  +  (  a PP  (5.4)  t)residual"  and (a ) ppt  residual  to zero, we can see that  has a  value of at least (25 + 240) =265 M P a (see the data column to the right of the Figure 5.13), w h i c h is substantially larger than the actual value of 191 M P a i n Table 5.3. Therefore, the recovery of a  dis  has to be taken into account to accommodate this  discrepancy. The presence of recovery make the separation of dislocation hardening and precipitation hardening more complicated for the alloys w i t h pre-deformation, and this problem needs to be further explored.  5.7 A SIMPLE MODEL FOR PREDICTING STRESS-STRAIN CURVES  In this study, t w o principle strengthening mechanisms exist for the t w o alloys, i.e., the dislocation hardening and precipitation hardening.  They originate from t w o sets of  obstacles, forest dislocations and precipitates. In this section, we w i l l consider h o w the two mechanisms contribute to the strength of the alloys. Foreman and M a k i n [25] studied this problem of averaging over the various strength of obstacle i n their computer experiments. Obstacles of t w o different strengths distributed randomly i n the slip plane w i t h different volume fraction are considered. The flow stress if the first type of obstacle was acting alone is a  u  and the flow stress to overcome the  second type alone is a . T w o addition laws were proposed. The first is a simple linear 2  addition law:  78  0"™. = where a  tot  o-j  + c  (5.5)  2  is the total flow stress.  The second is a mean square law:  <*«. = ( ^  2  +  (5-6)  B y comparing the computer calculations by Foreman and M a k i n and the results from Equation 5.5 and 5.6, it is found that the linear addition law o n l y applies well to the circumstance where a few strong obstacles are introduced among many weak ones, while the mean square law is i n good agreement except for the special circumstance mentioned above [25]. In our case, t w o types of obstacles, forest dislocations and precipitates, are present. Their contributions to the total flow stress w i l l be discussed next. First, i n a simple situation of the stress-strain curve of the 7030.60 alloy i n the asquenched condition (with no pre-strain), the flow stress is the result of o n l y dislocation hardening  plus the intrinsic strength of the matrix CT (which also includes the solid 0  solution hardening factor for the consideration of simplicity here), i.e., tr The a  tot  tot  = CTQ + CT^.  can be modeled using Voce equation [43]:  CT-a I  -s ] = e x p [ - -s—  (5.7)  L  CT. - C Ts s  8 r  whereCTis the flow stress,CT;and e are the yield stress and yield strain, respectively,CT is ;  s  the saturation stress, w h i c h is the interception of the w o r k hardening rate 0 (=dCT/ds) w i t h the stress axis shown i n Figure 5.13, s is the relaxation strain, w h i c h equals to CT/0 0) r  (  79  where 0 is the interception of w o r k hardening curve linearly extrapolated to the 0 axis, O  also illustrated i n Figure 5.14.  3500  CO  e  0  -7030.60, 0 pre-strain, as water quenched.  3000 2500  ra 2000 c 'c d) T3 i  ra X  1500 1000 500 -+-  50  100  150 Stress /  200  250  300  MPa  Figure 5.14: Plot of the w o r k hardening rate vs. stress for the 7030.60 alloy w i t h no predeformation i n the as water quenched condition.  Using the Equation 5.7 and the parameters described above, the calculated total flow stress curve can be obtained.  Subtracting the a from the total flow stress curve, then we 0  obtain the flow stress contribution from the dislocation, w h i c h is shown i n Figure 5.15. Assuming that the w o r k hardening curve does not change during ageing, the stress-strain curves for different levels of precipitation hardening can be predicted using either the linear addition law, or the mean square law mentioned above.  80  300  j  a  250 -  7030.60, 0 prestrain, as water quenched.  CO D.  200  dislocation hardening  to to  150 --  CO  100 -  2  , k^i a lt  b  50 • 0 -•  0.02  0  0.04  0.06  0.08  0.1  Strain  Figure 5.15: Plots of the experimental stress-strain curve and the flow stress contribution from the dislocation hardening for the 7030.60 alloy w i t h no pre-deformation i n the as water quenched condition.  F o r the linear addition law, the flow stress is given by:  0\ot. = °"o + 0" pt.  (5.8)  +0-^  P  where cr is the intrinsic strength of the aluminum matrix, a 0  contribution from precipitation hardening, and  p p t  is the flow stress  is the flow stress contribution from  dislocation hardening. F o r the mean square law, the flow stress is given by:  tot. =  CT  where a , a 0  p p t  and a  0  CT  dis  +  ( ppt. + ^dis. ) CT  2  2  1/2  (5.9)  have the same meaning as i n Equation 5.8.  The results of the predicted stress-strain curves as well as the experimental curves are  81  plotted i n Figures 5.16 - 5.20. Figure 5.16 illustrates the predicted stress-strain curves for the non-deformed 7030.60 alloy i n the natural ageing condition. It is clear to observe that the linear addition law applies properly for this case, except for the stress at larger strains (0.07 and up), while the mean square law under-estimate the flow stress. Figure 5.17 illustrates the predicted stress-strain curves for the non-deformed 7030.60 alloy aged at 100°C for 1 hour. The same observation can be made as i n Figure 5.16. This implies that there are "a few strong obstacles introduced among many weak ones", i.e. forest dislocations among weak G - P zones i n the 7030.60 alloy i n such ageing conditions. F o r the 7030.60 alloy aged at 100°C for 5 hours, the linear addition law begins to overestimate the flow stress, while the mean square law still under-estimate the flow stress as the former cases, as shown i n Figure 5.18. A s the ageing process proceed, both the overestimate of flow stress for the linear addition law and the under-estimate for the mean square law become more significant, as clearly illustrated i n Figures 5.19 and 5.20. The deviation of the models and the experimental results indicates that the addition of the dislocation strengthening and precipitation strengthening is complicated. Since the mean square law actually averages the density of the t w o obstacles, the under-estimate of flow stress implies that the w o r k hardening behaviour of the 7030.60 alloy may be changed during ageing w i t h the presence of precipitates.  M o r e dislocations may be  generated due to the dislocation-precipitates interactions and extra flow stress may be added accordingly, as might be expected for non-shearable precipitates (i.e. dispersion hardened material) where geometrically necessary dislocations play an important role and thereby affect the dislocation w o r k hardening [25]. Nonetheless, further study is needed to fully understand this problem.  82  I  200 -linear addition law - experiment CO Q_  150  - mean square law  «! 100 D D  50  0.02  0.04  0.06  0.1  0.08  Strain Figure 5.16: Plots of the experimental stress-strain curve and the calculated ones using different averaging laws for the non-deformed 7030.60 alloy naturally aged for 24 hours.  200  a  linear addition law  ——experiment  150  a  — — mean square law  D_  <» 100 b »  c  50  0.02  0.04  0.06  0.08  0.1  Strain Figure 5.17: Plots of the experimental stress-strain curve and the calculated ones using different averaging laws for the non-deformed 7030.60 alloy aged at 100°C for 1 hour.  83  200 -linear addition law -experiment (0  150  - mean square law  «! 100 b D  50  o-l 0.02  0.04  0.06  0.08  0.1  Strain Figure 5.18: Plots of the experimental stress-strain curve and the calculated ones using different averaging laws for the non-deformed 7030.60 alloy aged at 100°C for 5 hours.  200 -linear addition law -experiment  ro 0_  150  -mean square law  «? 100 b 50  0.02  0.04  0.06  0.08  0.1  Strain Figure 5.19: Plots of the experimental stress-strain curve and the calculated ones using different averaging laws for the non-deformed 7030.60 alloy aged at 150°C for 10 hours (peak aged).  84  200  7  a b  co  150 +  C  linear addition law experiment mean square law  a b  c  0  0.02  0.04  0.06  0.08  0.1  Strain  Figure 5.20: Plots of the experimental stress-strain curve and the calculated ones using different averaging laws for the non-deformed 7030.60 alloy aged at 150°C for 60 days.  5.8 OPTIMUM HEAT TREATMENT  Conforming to criteria that the bumper should have a highest m i n i m u m strength, the optimum heat treatment can be determined w i t h the existing facility, based o n this study. Since the m a x i m u m level of strain imposed on the material during bumper forming is 0.3, and the material w i t h a pre-strain level of 0.3 has the lowest peak strength among the materials w i t h pre-strain levels of up-to 0.3 (referring to Figures 4.16 and 4.17), the ageing behaviour of the material at the local section of the bumper where the deformation strain is 0.3 controls the final mechanical properties of the whole bumper product.  Therefore  the heat treatment schedule should be selected such that the material w i t h the maximum strain (0.3) i n the bumper can develop a peak strength after the ageing process.  85  Referring to Figure 5.8, the time-to-the-peak for the 7030.60 and 7108.72 alloys w i t h prestrain of 0.3 are 5 hours and 4 hours at 150°C, respectively.  So the heat treatment  schedule can be chosen as:  480°C/ 20 m i n .  Water quench Bumper forming  natural ageing Vlday  100°C/ + 5 hrs.  ,7030.60 k 1 5 0 ° C / 5 hrs.  *<  150°C/ 7108.72  4 hrs.  Different from the above schedule, the actual industrial heat treatment has an ageing time of 6 hours at 150°C, i.e. the pre-deformed alloys are slightly overaged.  Some other  considerations, such as the susceptibility to stress corrosion cracking, may have been taken into account i n the selection of this schedule.  86  Chapter  6  CONCLUSIONS A N D FUTURE WORK  Experiments simulating the  industrial bumper  manufacturing process  have  been  conducted i n this study. Mechanical properties (yield stress, U T S and strain-to-necking) and electrical resistivity evolution during aging for the materials w i t h pre-determined levels of pre-deformation were obtained.  After analyzing these data, we come to the  following conclusions.  6.1 C O N C L U S I O N S  1. The presence of pre-deformation leads to heterogeneous distribution of mechanical property i n the bumper products. 2. The presence of pre-deformation lowers the natural ageing response of both 7030.60 and 7108.72 alloys. This is suggested to occur due to the loss of excess vacancies to dislocations sinks. 3. Static recovery is suggested to occur during artificial ageing, simultaneously w i t h the precipitation process. 4. The presence of pre-deformation lowers the peak strength that can be obtained during artificial ageing, possibly due to the dislocations affecting the precipitate size distribution.  87  5. The overageing kinetics are accelerated due to additional short circuit diffusion paths provided b y dislocations. 6. According to the criteria that the bumper should have the highest m i n i m u m strength, the o p t i m u m heat treatment schedule is determined as:  480°C/ 20 m i n .  Water quench Bumper forming  7030.60 natural ageing /lday  +  100°C/ 5 hrs.  150°C/ 5 hrs.  +'  150°C/ 7108.72  4 hrs.  6.2 FUTURE WORK  F r o m this study, it is k n o w n that the presence of pre-deformation between quenching and natural ageing culminates i n the lower peak ageing strength of the alloys, w h i c h is not appreciable for the industry to produce high performance bumper products.  We  w o u l d like to examine the possibilities where the negative effect of the pre-deformation on the final mechanical properties of the materials is reduced or eliminated.  The deformation (the shaping of the bumper) is obviously necessary i n the bumper manufacturing process. However, the sequence i n w h i c h the deformation is taking place may be changed to avoid the negative effects. So it is worthwhile to examine the manufacturing process again. Referring to Figures 2.2 and 2.3 shown again here i n Figure 6.1, the layout of the whole production line is i n such a way that the dislocations introduced by the deformation process can not be eliminated  88  before the following ageing treatment is carried out.  Thus the dislocations remain  through the ageing treatment and adversely affect the strength of the materials.  Temperature  Ageing lOCC/Shr.  Time  Time  Figure 6.1: Schematic illustration of the extrusion and bumper manufacturing processes.  According to this study, the dislocations introduced by the deformation process should be removed before the ageing process. So we suggest a modified production schedule i n w h i c h the  deformation is moved forward to the end of the hot extrusion process, as  shown i n Figure 6.2. In the modified process, it is expected that the dislocation introduced b y the deformation process w i l l be eliminated during the following solution heat treatment. Therefore the adverse effect of the dislocations w i l l be avoided i n this new procedure.  O n the other  hand, if the heterogeneous distribution of strength i n the bumper is acceptable, the solution heat treatment i n the modified production schedule may be removed, which means energy and resource savings.  89  Temperature feasting  Time Figure 6.2: Schematic illustration of the modified bumper manufacturing process.  The details of the proposed process need to be further investigated. F o r example: •  the hot extrusion temperature.  The temperature should be selected and  controlled more carefully so that a solid solution can be obtained during the stretch bending process. This is to reduce the resistance of the material to the deformation. •  water quench after the hot extrusion. Water quench, instead of air cooling i n the present process, should be employed, also to i n favor of the solid solution formation.  90  BIBLIOGRAPHY  1. M . O . Speidel, Metall. Trans. A , 6 A 631 (1975). 2. N . R y u m , Acta. M e t a l l , 17, 269 (1969). 3. E . Nes, Acta. Metall., 20, 499 (1972). 4. A . J . Bryant, J . Inst. Metals, 97, 311 (1969). 5. J . T . Staley, Proc. of the 3 Int. Conf. o n A l A l l o y s , T h e i r Physical and Mechanical Properties. E d . L . Arnberg, O . Lohne, E . Nes and N . R y u m , V o l . HI, N T H and S I N T E F , T r o n d h e i m , N o r w a y , 107 (1992). rd  6. Private communications, S. R . Skjervold, Hydro-Raufoss. 7. W . Rosenkranz, A l u m i n u m , 36, 250, 397 (1950). 8. L . F . Mondolfo, N . A . Gjostein and D . W . Levinson, Trans. A I M E . , 206, 1378 (1965). 9. J . Lendvai, Materials Science F o r u m , v o l . 217-222. 43-56 (1996). 10. P . G u y o t and L . Cottignies, Acta. M e t a l l , 44, 10, 4161 (1996). 11. P . N . T . U n w i n and R . B . Nicholson, Acta. M e t a l l , J7, 1379 (1969). 12. J . T . Staley, R . H . B r o w n and R . Schmidt, Metall. Trans., 3, 191 (1972). 13. M . M u r a k a m i , O . Kawano and Y . M u r a k a m i , Acta. M e t a l l , 17, 29 (1969). 14. A . K Mukhopadhyay, Q . R . Yang and S. R . Singh, A c t a M e t a l l , 42, 9, 3083 (1994). 15. G . W . L o r i m e r and R . B . N i c h o l s o n , The Mechanisms of Phase Transformations i n Crystalline Solids, Inst, of Metals, L o n d o n , 43 (1969). 16. A . K e l l y and R . B . N i c h o l s o n , Prog. Mater. Sci., 10, 245 (1963). 17. G . W . L o r i m e r and R . B . N i c h o l s o n , Acta. Met., 14,1009 (1966).  91  18. J . D . E m b u r y and R . B . N i c h o l s o n , Acta. Metall., 12, 403 (1965) 19. P . N . T . U n w i n , G . W . Lorimer and R . B . Nicholson, Acta. Metall., 17,1363 (1969). 20. J . K . Park and A . J . A r d e l l , Metall. Trans. A , 14, 1957 (1983). 21. G . Thomas and J . N u t t i n g , J . Inst. Metals, 88, 81 (1959-60). 22. D . W . Pashley, M . H . Jacobs and J . T . Vietz, P h i l . Mag., 16, 51 (1967). 23. J . T . Staley, Metall. Trans., 5, 929 (1974). 24. S. P . Ringer, B . C . Muddle and I. J . Polmear, Metall. Trans. A , 26A, 1659 (1995). 25. L . M . B r o w n and R . K . H a m , Strengthening Methods i n Crystals, E d . A . K e l l y and R . B . N i c h o l s o n , Halsted Press Division, J o h n W i l e y & Sons, Inc., N e w Y o r k (1971). 26. J . D . E m b u r y , Materials Sci. F o r u m , 217-222. 57 (1996). 27. E . A . Starke Jr., Mater. Sci. Eng., 29, 99 (1977). 28. A . Deschamps, Y . Brechet, P . G u y o t and F . Livet, Z . Metallkd., 88, 601 (1997). 29. A . Deschamps and Y . Brechet, Acta. Mater., 47, no. 1, 293 (1999). 30. I. J . Polmear, Light A l l o y s , Metallurgy and Materials Science Series, 2 A r n o l d , L o n d o n , 97 (1989).  n d  edn., Edward  31. J . Wert, Microstructural C o n t r o l i n A l u m i n u m A l l o y s : Deformation, Recovery and Recrystallization, E d . E . H . C h i a and H . J . M c Q u e e n , T M S , 67 (1986). 32. K . Stiller, V . Hasen, M . Knuston-Wedel, G . Waterloo and J . Gjonnes, Proc. of the 6 Int. Conf. o n A l and A l l o y s , Their Physical and Mechanical Properties, ed. T . Sato, S. K u m a i , T . Kobayashi and Y . M u r a k a m i , Toyohashi, Japan, 2, 615 (1998).  th  33. J . A . Saeter, G . Waterloo and W . J . Poole, i n the Proceedings of I C A A - 6 , ed. T . Sato, S. K u m a i , T . Kobayashi and Y . Murakami, Toyohashi, Japan, 2, 745 (1998). 34. H . R . Shercliff a n d M . F . A s h b y , Acta. Metall. Mater., 38,10, 1789 (1990).  92  35. V . Gerold, i n "Metallic Solid Solutions", ed. b y J . Friedel and A . Guinier, W . A . Benjamin Inc. N e w Y o r k , X L (1963). 36. Y . Quere, i n "Metallic Solid Solutions", ed. by J . Friedel and A . Guinier, W . A . Benjamin Inc. N e w Y o r k , X X X V (1963). 37. A . Fjeldly and H . J . Roven, Material Science and Engineering, A234-236, 606 (1997). 38. F . R . N . Nabarro, Z . S. Basinski, and D . B . H o l t , A d v . Phys., 13, 193 (1964). 39. M . Wintemberger, Acta. M e t a l l , 7. 549 (1959). 40. J . A . Saeter, unpublished w o r k , (1998). 41. P . G o m i e r o , A . Reeves, A . Pierre, F . Bley, F . Livet and H . Vichery, i n the Proceedings of I C C A - 4 , ed. T . H . Sanders, Jr. and E . A . Starke, Jr., Atlanta, Georgia, U S A , 644 (1994). 42. A . Deschamps and Y . Brechet, Acta. Mater., 47, n o . l , 281 (1999). 43. Y . Estrin and H . Mecking, Acta. M e t a l l , 32, no. 1, 57 (1984).  93  APPENDIX  Some preliminary experimental w o r k was conducted o n the as-received, hot-extruded plates of the 7030.60 and 7108.72 alloys to obtain the basic characteristics of these materials. In particular, the homogeneity of the strength through the plate thickness, the sensitivity to solution time and the quench sensitivity, were examined.  These results,  w h i c h helped to design the schedule for experiments thereafter, are given i n this section.  The Through Thickness Hardness Distribution in the As-received Plates  The hardness profile through the thickness of the 7030.60 and 7108.72 aluminum plates was measured b y taking a piece of material sheared from the as-received plate and grinding to different levels. The Vickers hardness at different levels i n the plate is shown i n Figure A . l .  The results indicate that there is only a small difference i n hardness  between the surface and center of the plates (6.8% for 7303.60 plate and 2.7% for 7108.72 plate). However, to ensure the consistency of the results, all the samples i n the following experiments were taken from the center i n transverse direction of the plates by grinding off the surface layer, where applicable.  94  a)  7030.60 Extruded Plate Thickness: 5mm average of 3 hardness measurements  Rolling s u r f a c e , Hv= 104.1 0 . 1 m m below, Hv=105.4 1 m m below, Hv=111.2 2 m m below, Hv=111.6 C e n t r e , Hv=111.7  b) 7108.72 Extruded Plate Thickness: 5mm average of 3 hardness measurements  s u r f a c e , Hv=110.4 0 . 1 m m below, Hv=111.0 1 m m below, Hv=111.3 2 m m below, Hv=112.2 C e n t r e , Hv=113.5  Figure A . l Hardness distribution through the thickness of the as-received a) 7030.60 and b) 7108.72 plate.  95  Sensitivity to Solution Time  The sensitivity of the materials to solution hold time was examined to determine the resolution heat treatment time for the experiment schedule.  Samples of 7030.60 and  7108.72 alloy were solution treated at 480°C for 10, 20 and 60 minutes i n a salt bath, followed by water quench. The samples were then naturally aged at r o o m temperature for 3 days, followed by an artificial ageing at 100°C (in boiling water bath) for 5 hours. The hardness and resistivity measurements were conducted for different ageing times to study the ageing response of the materials. The ageing curves are shown i n Figures A . 2 A.5.  F r o m these results, it can be concluded that there is no substantial difference i n ageing behavior of the 7030.60 alloy for solution heat treatment times of 10, 20 and 60 minutes. The same is true for the 7108.72 alloy. F o r convenience, a solution heat treatment time of 20 minutes was adopted for subsequent experiments.  140 120  > X  (/> co cu c  "E CO  X  ageing at room temperature  > x  100 80  tn  CD  60  c  TJ  L-  40 20 0  CO  —% 60 min. _ a _ 20 min. _o_10min. i  X  i  10  100  1000  60  10000  120 180 240 300 360  Ageing Time / min.  Ageing Time / min.  Figure A . 2 Ageing curves (in hardness) for 7030.60 alloy w i t h different re-solution time at 4 8 0 ° C .  96  48  E d  47 46  . ageing at room temperature  45  o>  cb 44  43  > tn  42 41  60 min. 20 min. 10 "in-  % _£  co 4 0 cu Cr: 39  — 0 —  38 10  100  1000  10000  60  120 1 8 0 2 4 0 3 0 0  360  Ageing Time/min.  Ageing Time / min.  Figure A . 3 Ageing curves (in resistivity) for 7030.60 alloy w i t h different re-solution time at 4 8 0 ° C .  140 120  ageing at room temperature  > x  100  > x  CO  80  CO  co  Cl)  c "S  60  X  20  CO  tn  cu c  40  "E ro x  % 60 rrin. _ £ _ 2 0 rrin. _o—10 min.  0 10  100  1000  60  10000  120 180 2 4 0 3 0 0 3 6 0  Ageing Time / min.  Ageing Time / min.  Figure A . 4 Ageing curves (in hardness) for 7108.72 alloy w i t h different re-solution time at 4 8 0 ° C .  97  48 47  a CD  o  ageing at room temperature  46 45 44  X  43  J  42  ;> 41 tn 40  .60 rrin. . 20 min. .10 min.  'to  <u 39 1 38 1  10  100  1000  10000  60  Ageing Time / min.  120  180  240  300  360  Ageing Time / min.  Figure A . 5 Ageing curves (in resistivity) for 7108.72 alloy w i t h different re-solution time at 4 8 0 ° C .  Sensitivity to Q u e n c h i n g M e d i u m  The quench sensitivity of the 7030.60 and 7108.72 materials was also studied.  Samples  for both alloys were solution treated at 480°C for 20 minutes, then quenched i n four different media: water at r o o m temperature, o i l , ice-water mixture, and air.  After  quenching, the samples were given a natural ageing at r o o m temperature for 3 days, followed by an artificial ageing at 100°C for 5 hours. Again hardness and resistivity were measured to observe the ageing behaviors of the materials, and the results are shown i n Figures A . 6 - A . 9 .  F o r 7030.60 alloy, o n l y a m i n o r difference i n ageing curves was observed against the differences i n quenching condition, although there are some deviations i n the initial stage of natural ageing.  98  For 7108.72 alloy, the ageing curve for air quenching condition is different from other curves i n that it showed lower hardness and resistivity values. This way be attributed to some loss of precipitation hardening potential due to the precipitate of course precipitate on the A l Z r dispersoids during the relatively slow cooling for air quenching. 3  140 120  > x co co OJ c  "S  CO  X  140  ageing at " room temperature  120  > x  100 80  co co <D c T3 i  60 40  03  _ • — w ater Q — A — oil Q . _ 0 — ice W.Q. —*— air cool  20 0 10  100  1000  X  ageing at100°C  100 80 60 40  -waterQ. . oil < eW.Q. . icel . air cool  20 0 60  10000  120 1 8 0 2 4 0 3 0 0  360  Ageing Time / min.  Ageing Time / min.  Figure A . 6 Ageing curves (in hardness) for 7030.60 alloy w i t h different quenching conditions.  48  E d CD  b  x  J '>  '-4—*  48  47  ageing at  47  46  room temperature  46  a  45  CD  44 43 42  CD  -  .  Cd 39 38  I  I Mill  10  co CO  1000  41  10000  Ageing Time / min.  . . . .  40  a> Cr: 39 38  I  100  44 42  —0—water Q. oil Q. _ o — ice W.Q. » air cool  co to 40 d>  45 43  ~+--4d^  41  ageing at100°C  H 60  1  120  water Q. oil Q. ice W.Q. air cool  h 180 2 4 0 300 360  Ageing Time / min.  Figure A . 7 Ageing curves (in resistivity) for 7030.60 alloy w i t h different quenching conditions.  99  140  140 ageing at room temperature  120  > X to tn CD c  "s ro X  100  > x  80  tn tn  100 80  60  CD  60  40  i  40  c %  water Q oil Q . _ 0 — ice W.Q. — air cool  CD  X  _ A _  20  ageing a t 1 0 0 ° C  120  .water Q. .oilQ. . ice W.Q. . air cool  20 0  0 10  100  1000  60  10000  120  180  240  300  360  Ageing Time / min.  Ageing Time / min.  Figure A . 8 Ageing curves (in hardness) for 7108.72 alloy w i t h different quenching conditions.  48 47  a O  46  48 47  ageing at room temperature  a  45  cn <3  44  X  43 ">  '-4—*  41  .water Q. -oil Q. . ice W.Q. . air cool  tn tn 40  CD 39  or  J >  42  '-4—" CO  100  1000  45 44 43 42  =8=  *  •  •  41  .water Q. .oilQ. . ice W.Q. . air cool  to 40 CD  a:  38 10  ageing a t 1 0 0 ° C  46  10000  39 38  —I  60  1 120  1  1  180  240  1 300  360  Ageing Time / min.  Ageing Time / min.  Figure A . 9 Ageing curves (in resistivity) for 7108.72 alloy w i t h different quenching conditions.  100  

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