University of Nairobi

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1 University of Nairobi School of Engineering A STUDY OF FATIGUE AND FRACTURE CHARACTERISTICS IN THE HAZ OF AA7075 ALLOY BY Rotich Kiplangat Sammy F56/75499/2014 A Thesis submitted in partial fulfillment of the requirements for the degree of Master of Science in Mechanical Engineering, in the Department of Mechanical and Manufacturing Engineering of the University of Nairobi February 2017

2 DECLARATION ii

3 DEDICATION This thesis is dedicated to my family for their support, patience and love during the entire period of this study. This thesis is also dedicated to Prof. Rading who has been a great source of inspiration and motivation. iii

4 ACKNOWLEDGEMENTS I would like to thank the following people who played a pivotal role to the success of this research. They contributed and extended their valuable assistance at one point or another during the course of this study. First, I would like to offer my profound gratitude to Prof. G.O. Rading for his unrelenting support, guidance and patience throughout the course of this study. I am also grateful to Prof. F.P.L. Kavishe for his guidance and offering valuable insights to the project. I also acknowledge with gratitude the guidance and effort of Dr. L. Chown who was instrumental in this project. I would also like to acknowledge the African Materials Science Engineering Network (AMSEN, A Carnegie-RISE Network) under the coordination of Prof. Rading for the financial support during the entire period of this course. I am also grateful to the support offered by the technical staff at the University of Nairobi, Mechanical engineering workshop in many ways during the experimentation and setting up stage of the various tests. My deep gratitude goes to Mr. Njue for his assistance during the polishing of the specimens. Finally, I thank my family for their love, patience and encouragement. iv

5 Anti-Plagiarism Statement v

6 ABSTRACT The fatigue and fracture characteristics of the heat affected zone (HAZ) of AA7075 have been studied comprehensively. To determine the effects of the thermal peak temperature on the fatigue and fracture characteristics of AA7075 in T651 temper condition, a number of tests were carried out. Both experimental and theoretical work was carried out in this study. The HAZ was simulated by subjecting the specimens to predetermined thermal cycles in a muffle furnace. This enabled an effective approach in studying the specimen as a whole and also gave a thorough understanding of the fatigue crack growth (FCG), fatigue strength and impact toughness properties in relation to the microstructure of the different parts of the HAZ. This enabled the weakest region of the HAZ to be identified. In this study, the HAZ was divided into five regions. The muffle furnace was heated up to the peak temperature and the specimens inserted for 60 seconds as earlier determined in actual welding. Since the specimens have thin cross-sections of 1.6 mm in addition to high conductivity of aluminium metal, the specimen were assumed to have reached the peak temperature after 60 seconds. The specimens were then removed and left to cool in open air. The emphasis was placed on the effects of the HAZ peak temperature on the fatigue and fracture characteristics. The optical micrographs of the base metal (BM) and the five regions of the HAZ were studied using a universal optical microscope (OPTIKA B-353 MET) and the micrographs were taken at a magnification of X200. The fatigue crack growth (FCG) characteristics of the BM and the five regions of the HAZ were obtained through cyclic loading at an R-ratio of 0.2. The fatigue crack growth threshold, test was obtained by successive reduction of the by no more than 10% for every three successive readings. High cycle fatigue test was carried out in a rotating fatigue machine (SM1090) to determine the fatigue strength across the HAZ. The fatigue vi

7 strength was approximated at 10 7 cycles. A linear interpolation scheme was used in approximating the stress level between a specimen that failed and one run-out specimen. The hardness profiles across the base metal and the HAZ regions were carried out using a digital Vickers macrohardness tester, LV 800. The impact toughness properties were determined using the Izod impact machine. The Hounsfield tensometer was used to carry out tensile tests. The tensile tests were carried out to determine tensile properties in the longitudinal direction of the BM and five parts of the HAZ. The impact toughness data and the (0.2%) yield strength from tensile test were used to calculate the approximate plane strain fracture toughness of the BM and five regions of the HAZ. The properties of region E were found to be consistently lower in the major tests that were carried out. Region E was 5 mm from the weld centerline and received a peak temperature of 590 C. It had the lowest FCG resistance and also exhibited low fatigue strength properties. These were consistent with the low hardness and coarse microstructure found in this region. It had the lowest hardness of 112 HV and ASTM Grain Size Number of (48µm). There was drastic drop in tensile properties from towards the weld centerline. The proof stress dropped from 380 MPa at a distance of 11 mm from the weld centerline to 185 MPa at a distance of 5 mm to the weld centerline. In this study, region E was confirmed to be the weakest link in the HAZ which is significantly affected by the thermal cycle. The weakness is caused by dissolution of strengthening phases. The degradation of the strengthening phases occurs severely at 590 C due to dissolution of the major strengthening phase,. There is evidence of high grain growth in this region as evidenced by the optical micrographs and the average grain size number. The hardness in the HAZ of AA7075-T651 depends on the interplay between precipitate dissolution and reprecipitation. vii

8 NOMENCLATURE a Half crack length for central crack d Distance from the crack tip da/dn Crack growth rate(length per cycle) K Stress intensity factor (MPa m) K IC Plane strain fracture toughness K C Fracture toughness Range of stress intensity factor Effective range of stress intensity factor η η Semi coherent N Number of cycles R Stress ratio T ( ( ) ) T6 Solution heat treated and then artificially aged viii

9 ABBREVIATIONS ASTM American Society for Testing and Materials BM Base Metal FCG Fatigue Crack Growth FL Fusion Line FSW Friction Stir welding GP Guinier Preston GMAW Gas Metal Arc Welding GTAW Gas Tungsten Arc Welding HAZ Heat Affected Zone LEFM Linear Elastic Fracture Mechanics SEM Scanning Electron Microscope SSS Supersaturated Solid Solution TMAZ Thermo-Mechanically Affected Zone WM Weld Metal WN Weld Nugget WRS Weld Residual Stresses ix

10 TABLE OF CONTENTS DECLARATION... ii DEDICATION... iii ACKNOWLEDGEMENTS... iv Anti-Plagiarism Statement...v ABSTRACT... vi NOMENCLATURE... viii ABBREVIATIONS... ix TABLE OF CONTENTS...x LIST OF TABLES...xiv LIST OF FIGURES... xv CHAPTER ONE: INTRODUCTION Background Statement of the Problem Objectives of the Study Limitations of the Study...5 CHAPTER TWO: LITERATURE REVIEW The Hardening and Precipitation Sequence of AA 7xxx Welding of 7xxx Aluminium Alloys Microstructural Evolution During Welding...9 x

11 2.4 HAZ Strength and Hardness Fatigue Crack Growth Characteristics Crack Closure Effect of Welding on Fatigue Crack Growth (FCG) Characteristics High Cycle Fatigue Effect of Welding on the Fatigue Strength Effect of Welding on Impact Properties Summary of Literature Review CHAPTER THREE: EXPERIMENTAL TECHNIQUES Introduction The Materials and Preparation of Specimens Fatigue Crack Growth Testing Fatigue Crack Growth Testing Method Fatigue Crack Growth Test Rig Fatigue Strength Testing Fatigue Strength Testing Method Fatigue Strength Testing Machine Microstructure Characterization Optical Microscopy Average ASTM Grain Size Number xi

12 3.6 Hardness Testing Tensile Testing Impact Testing CHAPTER FOUR: RESULTS Introduction Optical Microscopy Fatigue Crack Growth (FCG) Hardness Profile Tensile Properties Impact Test Izod Impact Energy Plane Strain Fracture Toughness K IC Fatigue Strength CHAPTER FIVE: DISCUSSION Introduction Thermal Cycle Simulation Microstructural Characterization Hardness Profile Fatigue Properties Fatigue Crack Growth xii

13 5.5.2 Fatigue Strength Impact Strength Tensile Properties CHAPTER SIX: CONCLUSIONS AND RECOMMENDATIONS Conclusions Recommendations REFERENCES APPENDICES Appendix A: ASTM Grain Size Number Data Appendix B: Hardness Values Appendix C: Impact Toughness Appendix D: Fatigue Crack Growth Data for the BM and the HAZ of AA 7075-T Appendix E: Published Paper From This Work xiii

14 LIST OF TABLES Table 3.1: Chemical composition of the tested material (wt %) Table 4.1: FCG Parameters K TH, C and m for the BM and the different regions of the HAZ Table 4.2: Longitudinal tensile properties of the base metal and five regions in the HAZ of AA7075-T Table 4.3: Izod impact test data, proof stress and calculated K IC for the base metal and the five regions in the HAZ of AA7075-T Table 4.4: Fatigue strength of the base metal and various parts of the HAZ of AA7075- T xiv

15 LIST OF FIGURES Figure 2.1: Typical fatigue crack propagation Curve [53] Figure 3.1: Thermal simulations cycles obtained by Sakwa [21] Figure 3. 2: Welding arrangement showing position of the thermocouples from the weld centerline Figure 3. 3: Peak temperature vs distance from the weld centerline Figure 3. 4: Center Cracked Tension specimen for fatigue crack growth tests (dimensions in mm.) Figure 3.5: Fatigue crack growth test rig Figure 3.6: Cantilever specimen for fatigue testing (Dimensions in mm) Figure 3.7: Rotating bending fatigue testing machine (SM1090) Figure 3.8: The Optical microscope connected to a computer [100] Figure 3.9: Vickers Macro hardness tester LV Figure 3.10: Tensile test specimen (Dimensions in mm) Figure 4.1: Microstructure of the base metal Figure 4.2: Microstructure for region B Figure 4.3: Microstructure for region C Figure 4.4: Microstructure for region D Figure 4.5: Microstructure for region E Figure 4.6: Microstructure of region F xv

16 Figure 4.7: Average ASTM grain size number of the base metal and the five regions of the HAZ Figure 4.8: Fatigue crack growth curve for the base metal Figure 4.9: Fatigue crack growth in region B Figure 4.10: Fatigue crack growth curve for region C Figure 4.11: Fatigue crack growth curve for region D Figure 4.12: Fatigue crack growth curve for region E Figure 4.13: Fatigue crack growth for region F Figure 4.14: Comparison of the FCG curve for the base metal and region F of the HAZ Figure 4.15: FCG curves for the base metal and various regions of the HAZ Figure 4.16: Graphical representation of Vickers Hardness Numbers across the HAZ of AA7075- T Figure 4.17: Proof stress of the base metal and five regions of the HAZ in the longitudinal direction Figure 4.18: Representative graph of impact energy across the HAZ of AA 7075-T Figure 4.19: The fatigue strength of the base metal and the five regions of the HAZ xvi

17 CHAPTER ONE INTRODUCTION 1.1 Background High strength lightweight aluminium alloys especially 7xxx series are used in construction of aircraft, automobiles, armoured fighting vehicles and high speed trains and boats due to their low density, high specific strength and excellent corrosion resistance [1]. Due to the lightweight nature of aluminium alloys, they offer economies in transport and increased mobility in light armoured vehicles [2]. The main alloying elements in 7xxx aluminium alloys are zinc and magnesium. A range of intermetallic compounds may form during solidification depending on alloy composition and solidification conditions. Mechanical properties, formability and fracture behaviour of polycrystalline materials such as aluminium alloys are influenced by the presence of intermetallic compounds and whether they are precipitated at the grain boundaries or within the grain [3, 4]. For some applications, 7xxx aluminium alloys are joined by welding [5]. The use of welding offers the additional benefit of mass saving and elimination of fastening techniques such as riveting. However during welding, a heat affected zone (HAZ) is created adjacent to the fusion zone in the weldment. This zone experiences high temperatures but does not melt [6]. Due to varying temperatures experienced in this zone, microstructures of varying properties are generated. The strength in the HAZ is reduced due to dissolution of the hardening precipitates [7]. Research has shown that heat treatable aluminium alloys such as 7xxx gain their strength by precipitation hardening which is sensitive to thermal influences [8, 9]. The precipitation of ( ) from solution heat treatment during cooling may lower the strength if a large number of coarse precipitate particles are formed. Two types of Guinier-Preston (GP) zones, considered 1

18 to be the precursor to the phase, form in different preferential planes of the Al matrix and at different rates of quenching and aging. In the early stages of solution treatment, and GP zones are formed [10]. The control of the size and volume fraction of these phases during early stages of treatment is critical in obtaining optimised properties in 7xxx aluminium alloys. Researchers have shown that damage during extrusion begins by fracture or decohesion of the inclusions. The resistance to damage and fracture of aluminium alloys may be strongly influenced by the volume fraction, shape and distribution of second phase particles [8-11]. The variation of physical and mechanical properties is attributed to various precipitation stages during multi-stage heat treatment. For AA 7xxx aluminium alloys, the precipitation sequence from supersaturated solid solution (SSS) may be represented as [2]: SSS GP zones (semi coherent MgZn 2 ) (incoherent MgZn 2 ) Rangantha et al. [10] investigating the microstructure of AA 7049 observed fine precipitates of GP zones distributed homogenously inside the grains of coarse grain boundary precipitates. The strengthening of the alloy was attributed to the presence of GP zones and metastable phase Quantitative analysis and thorough understanding of crack initiation and propagation behaviour is crucial in determining fatigue life of aluminium alloys. During cyclic loading, multiple crack nucleation sites occur in aluminium alloys [12, 13]. In 2xxx and 7xxx series, multiple fatigue cracks initiate in the coarse Fe-containing particles or pores due to stress/strain concentration in these particles. The Fe- and Si-containing particles are the coarse constituent particles in high strength aluminium alloys [14]. Fe- containing precipitates such as Al 7 Cu 2 Fe are relatively 2

19 brittle. During extrusion and rolling the precipitates fracture resulting in pre-cracks which act as crack initiation sites in the material [15-17]. The HAZ is a region adjacent to the weld metal zone composed of the base metal that did not melt but was exposed to high temperatures for a sufficient period so that grain growth and other microstructural changes occurred. The heat of welding affects the mechanical properties and microstructure of the HAZ [4, 18-23]. Formation of a broad HAZ during welding causes a coarse grain structure which has unfavourable mechanical properties. The hardness profile through the HAZ provides a simple and convenient method for studying the effect of heat during welding [24]. The weld bead geometry is considered critical in determining the mechanical properties of the weld [25]. Research has revealed that during welding, bead geometry, microstructure and chemistry are affected by metal transfer [26]. Several crack initiation processes occur within the welds and the HAZ with separately distinct implications on the fatigue performance of welded structures [27]. Researchers [27] investigating the effect of residual stress on fatigue performance of welded aircraft structures reported although crack initiation was found in the fusion zone, cracks forming at the peak of residual stresses was located in the HAZ. The low hardness in the HAZ and highest residual stress of 264 MPa in the longitudinal direction was found to be detrimental to fatigue performance in AA 2024-T351 aluminium alloy joints [28]. 1.2 Statement of the Problem It is well known that the welding process negatively affects the grain structure and precipitate formation and distribution in the microstructure of the heat affected zone in precipitation hardened aluminium alloys [18-20, 29-31]. Although several studies such as that by Sakwa [21] characterised the HAZ as a single zone, due to the different temperatures experienced in the 3

20 HAZ, the affected region consists of graded zones which have different microstructures and mechanical properties. There are three metallurgically distinct regions in the HAZ: the grain growth region, the grain refined region and the transition region [32]. This is due to the complex thermal cycle, since all temperatures from the melting range of the alloy down to much lower temperatures are experienced by the metal. Steep property gradients in the HAZ make it extremely difficult to study the fatigue properties from an actual weld. The behaviour of the HAZ with regard to fatigue crack growth characteristics and fatigue endurance in 7xxx aluminium alloy has not been investigated exhaustively since the HAZ has been treated as one region during the analysis of fatigue properties in most studies. This study proposes to address this problem by finding the endurance limit, the fracture properties and the fatigue crack growth characteristics of the throughout the HAZ regions of aluminium alloy AA T Objectives of the Study The aim of this study was to link the microstructural changes in the heat affected zone to the fatigue, fracture and hardness characteristics of welded AA 7075 based on the principle that the heat affected zone (HAZ) is composed of different microstructural regions. The following specific objectives were identified to achieve this aim. a) Simulation of five different regions of the HAZ by subjecting five different specimens to thermal histories as close as possible to thermal histories experienced in actual welding by the five different regions. b) Characterization of the microstructure of the base metal (BM) and different regions of the HAZ in the five simulated specimens using optical microscopy. 4

21 c) Determination of the fatigue characteristics (i.e fatigue strength and fatigue crack growth) of the BM and different regions of the HAZ in the five simulated specimens. d) Characterization of the hardness profiles of the HAZ in order to find the weakest region of the HAZ. e) Determination of the fracture energy from impact test results of various parts of the HAZ. 1.4 Limitations of the Study Fractography studies on the fractured specimen from the impact and fatigue tests were not conducted due to lack of functioning scanning electron microscope (SEM) locally. Comparison between simulation done in the muffle furnace and the gleeble simulator (the gleeble simulator purposely designed for process simulation such as weld HAZ cycles) would have greatly contributed to the quality of this work. The gleeble simulator could not be found locally. Moreover the gleeble simulator had the limitation of not being able to accommodate FCG specimens. 5

22 CHAPTER TWO LITERATURE REVIEW 2.1 The Hardening and Precipitation Sequence of AA 7xxx The 7xxx series aluminium alloys are age-hardenable. Precipitates form during age hardening of the supersaturated solid solution and as the number and size of precipitates increase, the strength increases. Overaging will result in reduction of mechanical properties such as toughness, tensile strength and hardness [3]. Afify et al. [29] reported that the GP zones, and (Mg 32 (Al, Zn) 49 ) phases are the most effective hardening phases in Al-Mg-Zn alloys. The increased density of precipitates was observed to be responsible for high hardness. The basal planes of the hexagonal precipitates are partially coherent with the {111} matrix planes but the interface between the matrix and the c-direction of the precipitate is incoherent [11]. This is responsible for the high hardness. It was further observed that increasing the aging temperature enhances the precipitates nucleation process. In 7xxx aluminium alloys, the GP zones and phase form early during the precipitation process. To obtain optimum properties in these alloys, control of the volume fraction and size of these phases is critical [21, 22]. Strengthening is attributed to the presence of GP zones and the metastable in AA7049 [10]. The hardening phase results from GP zones that form on vacancy-rich clusters from the supersaturated solution [9, 33]. It was observed that for AA 7049 in the T6 tempered condition, a high density of fine precipitates were homogenously distributed in the aluminium matrix together with small amounts of GP zones and [10]. The reversion of GP zones during rapid heating, e.g. during welding, makes precipitation of difficult [34]. Researchers have found that and phases have compositions similar to the 6

23 MgZn 2 and have a hexagonal plate-like structure which grows preferentially on the {111} planes of the Al matrix [35, 36]. Zinc is the strengthening component in 7xxx series alloys. If the alloy contains more than 3 wt% Zn with a Zn: Mg ratio > 2, the hardening phase is MgZn 2 [37]. Copper increases the quench sensitivity while a high Zn: Mg ratio allows for homogenous nucleation and decreases the quench sensitivity [38]. A high Zn: Mg ratio also decreases the size and increases the density of the precipitates. 2.2 Welding of 7xxx Aluminium Alloys Welding of 7xxx aluminium alloys such as AA 7075 is difficult due to solidification and liquation cracking, poor weld zone ductility and susceptibility of the joint to stress corrosion cracking [2]. Gas metal arc welding (GMAW) and Gas tungsten arc welding (GTAW) are commonly used. The loss of strength in the weld metal and HAZ is due to dissolution of strengthening precipitates during welding, while the high cooling rates after the weld cycles do not favour reprecipitation [37]. In 7xxx aluminium alloys, zinc rapidly forms an oxide during welding, which affects the surface tension of the weld pool and increases the risk of defects due to lack of fusion[2]. This can be solved by using a current which is ~10-15% higher than that used in welding 5xxx alloys [2]. The hot cracking susceptibility of high strength 7xxx series aluminium alloys is strongly affected by the chemical composition of the weld metal. The crack susceptibility increases with increasing copper content up to 3 wt % Cu and then decreases with increasing copper content for copper content higher than 3 wt%. The resistance to cracking is enhanced by grain refining agents such as Zr [38]. 7

24 As previously mentioned, aluminium alloys such as 7xxx are normally welded by GMAW or by GTAW. High deposition rates, high welding speeds and deeper penetration are achieved in GMAW due to the high heat input. However, a high heat input is detrimental to welding of thin aluminium sheets since it causes distortion and increases the width of the heat affected zone [15]. Alloying elements such as magnesium or zinc may be oxidised during welding and the loss of magnesium is worst during GMAW [38]. Grain refinement in GTAW increases the resistance to fatigue crack growth [18]. The heat input can be controlled precisely during the pulsed GTAW process. GTAW is therefore preferred over GMAW as it produces good quality welds [10, 18, 20]. The pulsing of welding current leads to grain refinement in the fusion zone, HAZ width reduction, reduction of residual stresses and hot cracking sensitivity [18, 24]. Solid-state welding processes such as friction stir welding (FSW) offers an alternative welding method for AA FSW is a fairly new technique developed by The Welding Institute (TWI) in Cambridge, UK in 1991 [40]. The technique is widely used in joining light metals since it has improved weld properties, reducing solidification cracking and porosity in the weld metal [19]. It achieves the solid phase joining by locally introducing frictional heat due to tool rotation and plastic flow which causes a change in the local microstructure [41]. The welding temperature does not cause melting but it is high enough to cause dissolution and overaging of strengthening precipitates. The FSW technique has been successfully used to weld dissimilar alloys such as AA 6061 and AA 7075 [42]. 8

25 2.3 Microstructural Evolution During Welding Reduction in mechanical properties such as hardness, tensile strength and fatigue strength in the HAZ of weldments is associated with microstructural evolution during welding [2, 15, 16, 30, 31]. The heat input in arc welding is proportional to the welding voltage and current and inversely proportional to welding speed. With high heat input, a wider HAZ is generated due to the high conductivity of the base metal. The increased number of coarse grains in the wider HAZ is detrimental to mechanical properties [5]. The temperature variation in pulse welding leads to a continuous change in the weld pool size and shape, which favours growth of new grains [20]. At a given frequency, long pulse duration produces a coarser grain structure than shorter pulse duration [22]. With increased heat input, the grains in both the fusion zone and HAZ were reported to coarsen, and the HAZ width increased [22]. In welding of heat treatable aluminium alloys such as 7xxx, there is a significant change in micro-hardness in the weld zone, fusion boundary and the HAZ compared to the base metal. Researchers [5, 16] have reported low hardness in the weld zone. The hardness increases in the fusion boundary due to finer grains and subsequently decreases in the HAZ due to coarser elongated grains, segregation of intermetallics to dendrite boundaries and dissolution of strengthening precipitates [43]. Mechanical properties in the weld zone depend on the grain size and the unfavourable mechanical properties of the weld zone are due to the coarse grain sizes. In FSW, the weld zone is stirred and therefore has finer grains but generally the hardness is lower than that of the base metal [23]. Therefore welding of Al alloys should be done by using minimum heat input which can be controlled by optimum process parameters. 9

26 A study of AA7039 heated in the C temperature range showed a drop in hardness which could not be restored by natural or artificial post-weld aging. The reduction of mechanical properties was due to loss of coherency between the precipitate and the matrix [5]. The weld nugget (WN), thermo-mechanically affected zone (TMAZ), heat affected zone and the flow arm zone have been observed in friction stir welds [39]. The flow arm zone at the top of the weld has fine grains. The WN consists of fine equiaxed grains which indicated that recrystallisation and limited grain growth had taken place [44]. Researchers investigating the stability of nugget zone in high strength aluminium friction stir welds reported that the weld nugget zone grain structures are inherently unstable and abnormal grain growth occurred in the entire weld nugget [45]. Although friction stir welded joints are free of solidification-related defects since no melting takes place, they are prone to defects such as pin hole tunnel, piping defect, zig-zag line cracks and kissing bond defects [46]. Kissing bond defect are defects in solid state bonding in which two solid materials are in contact but with little or no metallurgical bonding present. This is attributed to improper flow of the metal and insufficient consolidation in the weld nugget region [46]. 2.4 HAZ Strength and Hardness Metallurgical transformations in the HAZ are affected by welding temperature, the type of welding process and associated parameters which determine the final microstructure. High heat input and preheating leads to deterioration of strength and increase in HAZ width [9]. The HAZ loses its strength due to either precipitate growth or transformation when the alloy is at high temperature. The width of the HAZ was found to decrease with decrease in heat input [4]. The behaviour of Al-Mg-Zn alloys such as AA7075 is different, as age hardening occurs at room temperature and there is no over-ageing effect. Properties of the HAZ can almost be recovered 10

27 by natural ageing or artificial ageing of the joint. High strength alloys such as AA 7075, AA7079 and AA7178 suffer from liquation cracking and reduction of strength in the HAZ [16]. Low micro-hardness has been found in the HAZ of AA 7039 [5]. The hardness increases towards the fusion boundary due to finer grains and decreases in the HAZ due to coarse grains [37]. The degree of hardening depends on the number, size and distribution of the precipitates. The sudden high heat input is detrimental to the strengthening precipitates of the alloy, which results in softening in the HAZ [5, 15, 31]. In the base metal adjacent to the fusion zone line, coarse and intermediate sized second phase particles segregate to the grain boundaries. The slow dissipation of heat generated during welding can result in grain coarsening. The fusion zone of a hybrid laser/gma welded AA 7075 had the lowest microhardness in the weldment [4]. This was attributed to alloying elements segregating in the dendritic phase rather than forming precipitates during solidification. The alloying elements then concentrate along the dendrite grain boundaries, forming secondary phases, whereas depletion of the alloying elements occurs inside the dendrite branches. The micro-segregation cannot be modified by ageing treatments and therefore softening dominates in the weld fusion zone [4]. Tensile fracture was found to occur at the retreating side of the TMAZ, which coincided with minimum hardness and maximum tensile residual stress [44]. The thermal history, superplastic deformation and microstructural changes, such as grain refinement and dissolution/coarsening of precipitates, govern the microhardness and strength across the weld in friction stir welds [46]. Friction stir welding generates a region of low hardness in the weld centre which extends to the transition zone of the TMAZ and the HAZ. The decrease in HAZ hardness was ascribed to the coarsening and dissolution of the strengthening precipitates [47]. Hardness can be used as a 11

28 reliable method of estimating the yield and tensile strength of the HAZ [48]. Both tensile strength and hardness reduce to a minimum at a distance from the fusion line in the HAZ due to overaging consequent to the transformation of the strengthening metastable precipitates [49]. 2.5 Fatigue Crack Growth Characteristics Some welded components are subjected to cyclic loading in service. Therefore a thorough understanding of the underlying mechanisms of fatigue failure is critical for the safe design of structures and components. In engineering components, fatigue crack nucleation occurs predominantly from inclusions, dents or voids [50]. Constituent particles in Al alloys provide sites for fatigue crack nucleation. In high strength Al alloys such as AA 7075, there are mainly two forms of coarse constituent particles i.e. the Fe- or Si- containing particles. The Fecontaining particles are brittle and can be fractured during rolling or extrusion processes, forming initiation sites for fatigue cracks [5, 12-14, 31, 43, 44, 46, 50]. Research has shown that fatigue cracks can also be initiated from non-fractured particles in high strength Al alloys [50]. Researchers investigating an AA 7475 thermo-mechanically processed aluminium alloy found that crack initiation occurred in surface grains which are characteristic of high purity metals and alloys where there is a low concentration of impurities for crack nucleation [51]. The fatigue crack growth rate of a component subjected to constant amplitude stress reversal is expressed in terms of the crack length increment per cycle, [52]. When the applied stress is held constant, the rate of growth of a fatigue crack increases with increasing number of cycles. Linear elastic fracture mechanics is used for small applied cyclic stresses and a small plastic zone formed ahead of the crack is in an otherwise elastic field [53]. 12

29 Welding affects the microstructure of the material which in turn influences the fatigue properties of the material [43]. The crack initiation stage is absent in welded structures due to microimperfections such as voids, cracks, impurities and inclusions [38, 53]. For cyclic variation of the imposed stress field, the LEFM characterization of the rate of fatigue should be based on the stress intensity factor range, [53]: where K max and K min are the maximum and minimum values of the stress intensity factor. The most accepted approach to fatigue crack propagation by use of linear elastic fracture mechanics is the semi-empirical Paris power law [54]: ( ) where C and m are the material parameters. The values of m are between 2 and 10 for most metals. These constants are influenced by material microstructure, the environment, temperature and the load ratio. The stress ratio is defined as the minimum stress intensity factor ( ) divided by the maximum stress intensity factor ( ) [53]: Equation 2.3 is applicable for single mode of far field loading and fixed loading and fixed values of R. The fatigue failure normally involves two major stages. The first stage is the crack initiation, which involves microstructure growth based on persistent slip band mechanism whereas the second stage is crack propagation where LEFM principles can be applied [53]. The sum of the 13

30 cycles spent in these two stages constitutes the total fatigue life. A typical fatigue crack propagation curve is shown in Figure 2.1 [55] where three regions are distinct. In Region I there is a threshold stress intensity factor below which there is no crack growth or the crack growth is too small to be measured. The fatigue threshold also depends on the frequency of loading and environmental conditions and decreases with an increase in stress ratio R [56]. It has been postulated that the fatigue cracks cease to propagate when the extent of the plastic zone becomes comparable to some characteristic microstructural dimension [57]. Region II is linear in the log-log scale and the LEFM condition deals with crack propagation in this region by the Paris law. The fracture surface in this region is normally characterised by ripples or striations whose spacing may be correlated to average crack growth rates [53]. Fatigue crack growth becomes faster than predicted by the Paris law in Region III as the stress intensity factor range increases, approaching the value of fracture toughness (K c ). Forman et al. [58] proposed the following relationship to describe fatigue in regions I and III: ( ) ( ) where R is the stress ratio. The above relationship accounts for stress ratio effects, while the Paris law assumes that is dependent on only. 14

31 Figure 2.1: Typical fatigue crack propagation Curve [53] 2.6 Crack Closure Critical links have been sought between fatigue, fracture mechanics and crack closure mechanics [59, 60]. Elber proposed that premature contact between the crack faces can occur even during the tensile portion of the fatigue cycle because of a zone of residual plastic deformation left in the wake of the crack tip arising from prior plastic zones [60]. Crack closure and opening has been attributed to crack tip blunting and strain hardening. Elber postulated that crack closure decreased the fatigue crack growth rate by reducing the effective stress intensity range. He introduced a new stress intensity range, the effective intensity range to be used in the Paris equation:... [2.5] where is the maximum stress intensity, is the stress intensity at which the crack opens, U is the empirical relationship between and 15

32 The corresponding characterisation of fatigue crack growth rates can then be obtained by adjusting the Paris-Erdogan equation to [54]: ( ) ( )... [2.6] Under plane strain conditions obtained at near threshold stress intensities, plasticity induced crack closure (PICC) plays a less significant role [61]. Other than PICC, roughness induced crack closure, viscosity induced crack closure, oxide induced crack closure and transformation induced crack closure have been researched and documented [62] 2.7 Effect of Welding on Fatigue Crack Growth (FCG) Characteristics Welded components and machine parts find applications in areas where the stress is constantly varying in magnitude and direction. Research has shown that welding has a substantial effect on the microstructure of the base metal which in turn affects the fatigue properties [40]. Fatigue damage is generally described as the growth and nucleation of the small cracks to the final failure [53]. It is widely accepted and proven in research that the initiation stage is absent in welded structures generally due to micro-imperfections in welds [63, 64]. Research has shown that constant amplitude FCG properties in high strength aluminium alloys such as the AA 7075 is determined by the microstructure, composition, stress intensity ratio, the environment, grain size and yield stress [64]. Friction stir welding was found to impart inferior tensile properties which lowered the elongation of the welded joint and in turn imparts lower resistance to FCG [52]. The microstructure of the weld region plays a critical role in determining the mechanical properties of the joint. Mechanical properties of the FSW joints depend on the structural characteristics of the weld region which in turn depends on the specific thermal /mechanical cycles imposed during friction stir welding [40]. The crack like defects in weldment 16

33 begin to grow immediately when subjected to external cyclic fatigue loads, so that, for welded specimens the total fatigue life is dominated by crack propagation phase [65]. Slight diversity in microstructure around the notch tip is associated with the difference in fatigue crack initiation lives [66]. In friction stir welding, fatigue failures have been observed to initiate from surface roughness, tool marks, flash on the surface of the joint or typical defects such as lack of root penetration [67]. Studies have established that fatigue failure in friction stir welds occur in the HAZ or at the weld nugget in discontinuities in material flow pattern [68]. Low cyclic strength and an elongated grain structure coupled with coarse intermetallic particle distribution, favour crack propagation in the TMAZ [69]. Researchers investigating effects of weld geometry on fatigue properties of T-weld and cruciform joints showed that the ratio of the thickness of the main plate and the radius of weld toe are the critical parameters in fatigue behaviour of such welded joints [70]. Investigations done on the fatigue crack initiation and propagation behaviours of welded steel tubular structures reported the crack growth life of 75% to 89% of the total life for all the welded joints tested [71]. Residual stresses are inevitable in the welding process. They have two effects; firstly they produce distortions and secondly they can be the cause of premature fatigue failure [72]. Weld residual stresses, (WRS) are quite large and tensile in nature in the vicinity of the weld, where their magnitude is approximately equal to the yield strength of the weld metal [73]. Investigations have shown that effective stress intensity factor and the effective stress ratio can be applied to predict fatigue crack growth rate in both tensile and compressive stress field by using base material s crack growth rate data with different stress ratio [74]. During crack growth 17

34 opening, the crack opening load is required to offset compression at the crack tip caused by the superposition of clamping force attributed to the residual stress and crack closure [73]. In a review of residual stress and its effect on the fatigue behaviour of Al-5083 and Al-7020 welded components, it was reported that the magnitude of WRS reached up to 0.2 times the yield limit in the HAZ of as welded joints [75]. Hole drilling method was applied with strain gauge rosettes to measure WRS in a position approximately 1.0 mm from the weld toe. The stresses were found to be between 120 MPa and 140MPa [75]. Researchers investigating Al-5456-H117 butt welded plate reported that the difference in the FCG properties of the base metal, the HAZ and the weld metal was as a result of the difference in the WRS [76]. 2.8 High Cycle Fatigue The methods for characterising high cycle fatigue in terms of nominal stress amplitude emerged from the works of Woehler (1860) while working on fatigue of alloys in railroad bogey axles, as reported by Suresh [53]. Fatigue data is usually conveniently presented in the classical form of stress-life(s-n) or strain life curves as first adopted by Woehler. The graphical S-N curve is a plot of the number of loading cycles that a specimen can withstand before failure. The test specimens are machined to provide a waisted (hour glass) cylindrical gauge length which is tested in rotating bending mode. Some materials such as mild steels exhibit a knee in the S-N plot below which the specimen is assumed to have infinite life(i.e the curve decreases asymptotically to a horizontal constant amplitude stress value ) [77]. The endurance limit is ~35% to 50% of the tensile strength for most steels and copper alloys [53]. Many high strength steels and aluminium alloys do not exhibit a fatigue endurance limit. The endurance limit is therefore defined as the stress amplitude which the specimen can support for at least 10 7 fatigue cycles [53]. In most engineering designs 18

35 the assumption was made that any cyclic stress below the fatigue limit does not affect the component. Recent studies [78] have, however, indicated that fatigue failure in aluminium alloys can occur at 10 9 fatigue cycles, which contradicts the work of Woehler. High strength aluminium alloys such as AA7075 used in the aircraft or automotive industry are subjected to high cyclic loading in service, involving high frequency vibrations. Whereas it is time consuming and impractical to carry out fatigue tests to the gigacycle regime, most components and structures in service are often loaded beyond 10 8 cycles [79]. The fatigue strength of AA 7075, which was found to be superior among the high strength aluminium alloys, was attributed to the higher tensile strength [78]. An analysis of the fractured surface of AA7075 fatigued at 225 MPa to 7.6 x 10 6 cycles revealed micro-voids which coalesced to form macroscopic cracks [57]. For high cycle fatigue, Wang et al. [80] found the crack initiation sites in ferrous alloys were located at a porosity or inclusion in the interior of the specimen. There are two design philosophies applied by aerospace engineers in order to determine the operating lives for aluminium based aerospace alloys: damage tolerance and safe life approach. In damage tolerance approach, the S-N fatigue performance dictates initial inspection intervals. For repeat inspections, however, fatigue crack growth and fracture toughness properties are the primary factors taken into consideration [53, 81]. 2.9 Effect of Welding on the Fatigue Strength Although tensile tests are the most popular type of material test, fatigue strength is very important from the viewpoint of strength design because most of the fracture is due to fatigue failure. Welding introduces the presence of porosity and reduced mechanical properties in the weld metal [49]. The weld imperfections reduce the fatigue crack propagation life and fatigue 19

36 strength in welded components. Research has shown that undercuts and solidification cracks are more damaging than embedded porosity [65]. It was further demonstrated that decreasing the stress concentration and increase in fatigue strength could be achieved by improving the weld geometry. The residual stress, bending stress and dissimilar mechanical properties of weld metal and the HAZ decreases the fatigue strength of the welded structures. The effect of plate thickness on stress intensity factor and fatigue strength is less significant compared to that of weld flank angle and weld toe radius [82]. The fatigue strength at 10 7 cycles of FSW butt joints were about 40% of the ultimate tensile strength (UTS). Improved fatigue strength was obtained at a welding rotation speed of 1000 r/min and travel speed of 80 mm/min indicating that the energy input has a significant effect on fatigue behaviour of welds [83]. Some researchers have reported that high temperature cycle in welding process and the discontinuous microstructure between the base metal and weld area weakens the fatigue strength of the welded structure significantly [84]. The fatigue cracks were found to propagate directly from the welding defects for welded specimen compared to the parent metal [83]. In AA 7075 researchers have found that dissolution of strengthening precipitates ( ) in the HAZ lead to reduction in hardness [47]. Researchers investigating the effect of distributed porosity on the fatigue resistance of Al-5083 double V-groove butt weld joints found that welds with their reinforcement removed were more sensitive to porosity compared with the weld with weld reinforcement intact. When relatively higher level of weld inner defects such as inclusions, porosity, lack of penetration and fusion exist, welds with bead reinforcement removed exhibit shorter fatigue life than those with beads on. Biaxial loading and the combined effect of a solidification crack and a circular porosity was reported to reduce the fatigue crack propagation life [65]. Some studies concluded that the angle 20

37 at the toe of reinforcement in welded aluminium alloy joints was crucial in determination of the fatigue life [85]. The dissolution of strengthening precipitates during welding has been found to be the cause of softening and consequently the reduction of fatigue strength in welds [83]. Improvement of fatigue life can be done by toe grinding, TIG dressing, shot peening and hammer peening [84]. The dependence of fatigue life on grain size varies depending on the deformation mode. Grain size has its greatest effect on fatigue life in the low stress, high cycle regime in which stage I cracking predominates. In high stacking fault energy materials (such as Aluminium) cell structures develop readily and these control the stage I crack propagation [87]. The repair of welded joints by second weld does not necessarily extend the life of the joint above the original welded joints. The microstructure of weld repaired plate was found to exhibit abnormally coarse grains and weld porosity due to large heat input during the weld repair process [65]. In addition, more voids and porosity were found located along the grain boundaries of the weld repaired plates. This led to higher fatigue crack growth rates and also lowered fatigue lives of weld repaired specimens Effect of Welding on Impact Properties Welding process leads to deterioration of the mechanical properties of the base metal [88]. These properties are dependent on the microstructure change. Researchers have studied the role of grain size on toughness and hardness [89]. The excess grain growth can lead to reduced strength and increased crack growth and initiation. The weld metal (WM) fracture mode was found to change from ductile fracture (8.1kJ/cm) to brittle fracture mode (6.5kJ/cm) with increase in heat input [89]. This was attributed to grain coarsening as a result of high heat input which reduced the impact energy and toughness. Ductile fracture is characterized by appreciable plastic 21

38 deformation prior to and during the propagation of the crack whereas brittle fracture is characterized by rapid crack propagation with no gross deformation and very little microdeformation. The fracture stress of a material that undergoes plastic deformation before fracture is greater than that of a brittle material [87]. The impact energy decreases with increase in heat input in both the weld metal (WM) and the HAZ [90, 91]. However, the impact energy of the WM is lower compared to HAZ. This is caused by the defects in the weld metal that cause stress concentrations inducing cracks near defects at the time of impact tests. Charpy impact tests on the welded AA 6082 aluminium alloy showed that the weld metal had low impact energy of 6.9 J while the HAZ had 13 J [49]. Defects such as porosity in the weld metal caused lower and scattered impact properties than those of the HAZ [49]. Researchers studying the impact toughness of butt welded joints in a titanium alloy, found the WM to have the highest impact energy followed by the HAZ and the base metal [92]. The fractographic observation was found to concur with the findings. The weld metal had a tortuous appearance. The high energy absorbed by the weld metal was due to the long distance covered by the crack before fracture. The small dimples were observed in the WM compared to the HAZ and the BM. The WM exhibited a ductile fracture surface which was attributed to large grains. For a given metallurgical condition, the impact toughness is inversely related to the grain size [91]. The post weld heat treatment improves the impact properties. The toughness of the HAZ was found to increase with PWHT duration in P-91 steel welds [93]. 22

39 2.11 Summary of Literature Review In the utilization of aluminium alloys such as AA7075 in structural and components application, one difficulty to be overcome is the deterioration of the mechanical properties in the HAZ and the weldment compared to the base metal. The mechanical properties of the HAZ deteriorate due to welding thermal cycles. The resulting microstructure in the HAZ is heterogeneous and the properties differ too. The weakest region in a weldment has been found to be in the HAZ in most of the research that has been carried out. The thermal conditions and the peak temperature experienced govern and play a role in the resulting metallurgical structure, mechanical properties, residual stresses and distortion. Regulating heat input during welding is critical for the optimization of the HAZ strength. The hardness in the HAZ decreases as the fusion zone is approached. This is due to the coarsening of the microstructure, segregation of the strengthening precipitates into the grain boundaries and the dissolution of strengthening precipitates. The region with the lowest hardness is also found to be consistent with the lowest FCG resistance and also the fatigue strength. Welded components and structures experience cyclic loads in service and fail due to fatigue. A substantial part of the fatigue life may be taken up during the propagation of a crack under cyclic loading. The crack initiation stage is absent in welded structures due to micro-imperfections. The LEFM characterisation of the rate of fatigue is based on the stress intensity factor,. Paris Law is the most widely used and accepted approach to fatigue crack propagation analysis by use of LEFM. The applied stress ratio can have a significant effect on the crack growth rate. The Paris law can be modified to include the dependence on the crack growth rate of the stress ratio by using Elber s crack closure technique. 23

40 It has been demonstrated that the fatigue crack growth, fatigue strength, hardness and impact properties are governed by the microstructure. Whereas coarse microstructure has low hardness and yield strength, the ductility is increased which in turn improves the impact properties. The review reveals that welding affects the microstructure, mechanical properties and fatigue properties of the HAZ. Moreover, it is clear that the heat experienced by the HAZ is responsible for the observed changes. Since different parts of the HAZ attain different peak temperatures, it should be expected that the effects (on microstructure and properties) will be different in different regions of the HAZ. These differences have not been adequately addressed in the literature. It was the purpose of this study to address this shortcoming. 24

41 CHAPTER THREE EXPERIMENTAL TECHNIQUES 3.1 Introduction The experimental work that was undertaken involved: thermal heat treatment simulating heat flow condition during welding, microstructural characterization (Optical), and mechanical tests (fatigue crack growth, fatigue strength, Impact tests and hardness). The main aim of this work was to establish the fatigue and fracture properties across the HAZ of AA7075 in T651 temper condition. 3.2 The Materials and Preparation of Specimens The material used for this experiment was AA7075-T651.The material was purchased from the Good Fellow Cambridge Ltd in the UK and was supplied in form of plates and rods. The plates were supplied in 1/16-inch thick Alcad sheet of 250 mm by 100 mm. The rods had 12.7 mm diameter and a length of 500 mm. The material was supplied in T651 temper condition from the manufacturer. T651 temper condition is obtained by solution heat treatment at C for 2hrs followed by C for 14 to 20hrs and then water quenching. Aging treatment is done at room temperature for 12 to 24 hrs or at C for 20 hrs. The chemical composition of the materials in weight percent as provided by the supplier is given in Table 3.1. Table 3.1: Chemical composition of the tested material (wt %) Mtl Al Si Fe Cu Mg Mn Cr Ti Zn AA7075 Bal 0.09% 0.28% 1.4% 2.7% 0.03% 0.19% 0.03% 5.9% Since the whole specimen was inserted in the muffle furnace, the residual stresses were assumed to be negligible. The specimens heating cycles were simulated in a muffle furnace with a peak 25

42 temperature of 1300 o C. The simulations were done after machining of the specimens. The furnace was very accurate and could be programmed to attain different temperatures in a short time. The muffle furnace was heated up to the peak temperature and the specimens inserted for 60 seconds. It was then removed and left to cool at room temperature. The simulation done was one that resulted in as close as possible to the thermal cycle histories in Figure 3.1. The thermal histories had been obtained in an earlier study by Sakwa [21] by actual welding obtained using alternating current gas tungsten arc welding (GTAW) with the torch mounted on a carriage. Welding was done using a 2.6mm 2 percent thoriated tungsten electrode with a 50-degree tip with the arc gap preset to 2mm. Welding was done parallel to the rolling direction. Thermocouples were mounted at different distances from the weld centerline as shown in Figure 3.2. Figure 3.1: Thermal simulations cycles obtained by Sakwa [21] 26

43 Figure 3.2: Welding arrangement showing position of the thermocouples from the weld centerline. The regions of the HAZ were denoted with alphabetical letters A to F ; A was the base metal, B represented a distance of 15mm from the weld centerline with the peak temperatures of 250 C during the welding process (TC#6), C represented a distance of 11mm from the weld centerline with a peak temperature of 325 C during welding process (TC#5), D represented a distance of 6 mm from the weld centerline with a peak temperature of 498 C during the welding process (TC#3), E represented a distance of 5mm from the weld centerline with a peak temperature of 590 C during the welding process (TC#2) and F represented a distance of 4mm from the weld centerline with peak temperature of 650 C during the welding process (TC#1). Figure 3.3 shows the variation of the peak temperature with distance from the weld centerline. 27

44 Figure 3.3: Peak temperature vs distance from the weld centerline 3.3 Fatigue Crack Growth Testing Fatigue Crack Growth Testing Method Center cracked specimens were used for the BM and the five regions of the HAZ. The specimen dimensions are as shown in Figure 3.4. The specimens were prepared according to the ASTM E647 procedure [94]. The edges of the pre-crack were then polished to have a smooth scratch free surface so as to enable the viewing and measurement of the crack as it progressed during the fatigue test. Three specimens were prepared for each region of the HAZ and the BM. The fatigue crack growth test for this experiment was set to give a small stress ratio R, equal to 0.2. Two experiments were carried out during the fatigue crack growth test; constant load test and the 28

45 decreasing load test. The constant load test was carried out to determine values of C and m while the load shedding test was done to determine the threshold value of the stress intensity factor, K th. The load shedding test was carried out by decreasing P by not more than 10% after every three successive readings according to the ASTM E-647 standard [94]. The crack growth rate da/dn, in mm/cycle was obtained by the secant method as outlined in ASTM E-647 standard [92]. The corresponding stress intensity factor, K, was calculated from the average crack length during the interval a i -a i-1 (where a i is the final crack length and a i-1 is the initial crack length) according to the equation below [94]. (3.1) where is the difference between the maximum and minimum force in the cycle, B is the specimen thickness (1.6mm), a is the half crack length and W is the width of the specimen which is 100 mm. 29

46 Figure 3.4: Center Cracked Tension specimen for fatigue crack growth tests (dimensions in mm.) Fatigue Crack Growth Test Rig Fatigue crack growth test rig designed and constructed at the University of Nairobi mechanical engineering workshop by Rading [99] was used. The details of the crack growth test rig and assembly drawing are documented in the reference [99]. The crack growth test rig as shown in Figure 3.5 was used for FCG testing. The test rig is a cam operated eccentric arm machine. It comprises of load arms like those commonly used in tensile testing machines. One of the load arms is connected to a 5,000 N capacity load cell while the second load arm is connected to the cam mechanism for holding the specimen. The cam is driven by a 4 kw, 1,498 rpm motor. The eccentric arm connected to the cam converts the rotary motion of the motor into oscillatory motion of the driving force. The load was applied by tightening the nut located on one of the load arms. The load applied was displayed on a control panel fitted to the machine. The crack growth was measured using a travelling microscope accurate to 0.1 mm. 30

47 Figure 3.5: Fatigue crack growth test rig 31

48 3.4 Fatigue Strength Testing Fatigue Strength Testing Method High cycle fatigue was carried out to determine the approximate fatigue strength in the HAZ of AA7075. Step loading method was used to approximate the fatigue strength [95]. For purposes of this study, the fatigue strength was defined as the alternating stress that causes the material to fail in 10 7 cycles. At least 3 specimens were used for each case. The first was tested at a stress level estimated to cause failure at a value of stress close to 10 7 cycles. The second specimen was tested at a stress approximately 7% lower than the first. If the second specimen had a run-out (i.e did not fail after more than 10 7 cycles), the fatigue strength was determined by interpolation between the two results. The third specimen was then used to confirm the result. If the second specimen failed, the stress was reduced by a further 7%, and the process repeated until a run-out occurred. The 7% step was chosen due to the stepwise nature of the testing machine. The specimen used for fatigue strength testing is as shown in Figure 3.6. Figure 3.6: Cantilever specimen for fatigue testing (Dimensions in mm) 32

49 The specimen was machined at the Numerical Machining Complex Centre Ltd, Nairobi. CNC machines were used in turning the rod from 12.7 mm to 8.95 mm and consequently in reducing the neck to 4.05 mm Fatigue Strength Testing Machine Rotating bending fatigue testing machine (SM1090) supplied by TecQuipment Ltd in the UK and shown in Figure 3.7 was used to test the fatigue limit. It has two main parts: a main unit and a separate control and instrumentation unit. The main unit has a motor that rotates the test specimen under constant load (stress). The motor turns a coupling and a short driveshaft. The driveshaft turns the collet type chuck that grips the driven end of the test specimen with uniform pressure around its circumference. This is important because it gives an even stress distribution on the specimen and grips it with minimum eccentricity. At the loading end of the specimen, an adjustable dead weight applies a vertical (downwards) load on the specimen. It does this through a self-aligning bearing inside a gimbal. A sensor counts the rotations (cycles) of the specimen and a load cell measures the force that is applied to the specimen. A display panel shows the load, speed of rotations (cycle rate) and the cycle count. 33

50 Figure 3.7: Rotating bending fatigue testing machine (SM1090). 3.5 Microstructure Characterization Optical Microscopy The metallographic specimens were prepared from cut out pieces of the material from the BM and the five parts of the simulated HAZ. The preparation of the specimens involved sectioning, mounting, grinding, polishing and etching. The specimens for hardness were similarly prepared. The specimens were embedded in resin for ease of handling. The specimens were ground starting with coarse grinding and ending up with the fine grinding using a series of silicon-carbide (SiC) papers of increasing fineness (grades 240, 320, 400, 600) under a flow of running water. The specimen was cleaned in between each paper finish to remove any grit and ground in a different direction in the next. The samples were then successively polished on a rotating wheel covered with a suitable cloth impregnated with a polishing diamond paste of 6µm, 1µm and alumina 34

51 (Al 2 O 3 slurry) of 1/4µm. The specimens were then etched in proprietary etchant patented at the University of Nairobi and optical micrographs taken with the optical microscope Average ASTM Grain Size Number The optical micrographs of the BM and the five regions were obtained. The grain size was obtained using commercial software, Microstructure characterization software 3.0 (MIC 3.0). The software uses the ASTM E-112 linear intercept method [96] which is inbuilt in it. Universal Microscope (OPTIKA B-353 MET) with a maximum magnification of X500 (Figure 3.8) was used to obtain the optical micrographs. A digital camera of resolution 2592 x 1944 pixels connected with USB cable to the computer was used to take the micrographs. Figure 3.8: The Optical microscope connected to a computer [100]. Microstructural analysis was done using a microstructural characterizer 3.0 (MIC 3.0) software developed by TCR Advanced Engineering (P) LTD, India. The software has a micrographic 35

52 image solarizer for identification of phases and grain boundaries. The images can be taken directly from an attached camera or the images stored in the computer can be used. The MIC 3.0 then creates reports in specified layouts such as Joint photographic Experts Group (jpeg) format. It also has a file format saving system for all custom reports with Microsoft (MS-Word) document interface. 3.6 Hardness Testing The hardness profile was carried out using the Vickers macro hardness tester LV 800 with a test load of 10 kgf and a dwell time of 15 seconds. The test specimens were placed on the stage and focusing done to bring images to clear focus using the X20 objective. The hardness values were read directly from the LCD panel attached to the machine. The indentations were done at intervals of 1 mm apart and ten indentations were carried out on every specimen. Ten readings were taken for each sample to get the representative reading and an average was done to get the hardness value. The hardness test was conducted on a LV800 digital Macro Vickers Hardness Tester (Figure 3.9) with a pyramid indenter supplied by LECO Corporation in St. Joseph, Michigan, USA. The tester has a load capacity of 0.3 kgf to 10 kgf. It is fitted with 10x and 20x objectives and built with a video adapter. The hardness tester has an inbuilt conversion scale and hence the Vickers hardness number is directly read from its display panel. The test method is ASTM E92 [101]. 36

53 Figure 3.9: Vickers Macro hardness tester LV Tensile Testing The tensile specimens for the BM and the five regions of the HAZ were machined according to ASTM E8M standard [97]. The specimens were cut in the longitudinal direction. Figure 3.10 shows the tensile specimen which was used. The Hounsfield tensometer was used for tensile testing. The equipment was manufactured by Hounsfield Test equipment Ltd in England, UK. The tensometer has a capacity of 20 kn. The sample is loaded between two grips which are operated manually. 37

54 Figure 3.10: Tensile test specimen (Dimensions in mm) 3.8 Impact Testing The impact toughness test samples were machined and prepared from the AA7075 T651 rods according to the ASTM E23 [98]. The surfaces of the specimen were polished to eliminate any surface irregularities that could have affected the results. The impact energy was measured using Izod impact test machine. The readings were taken directly from the scale attached on the machine. Three readings were taken for each specimen and the average value calculated. The Izod impact tester was used to carry out the impact test. The equipment was supplied by Avery Birmingham Ltd, in UK. The total pendulum weight was 120 ft-lb. The distance between the impact centre to pendulum revolving center is 0.65 metres and a maximum lift angle of the machine is degrees. The impact energy data was taken by recording the lift and swing angles of the striking hammer during testing. 38

55 CHAPTER FOUR RESULTS 4.1 Introduction Thermal simulation of the different regions of the HAZ based on peak temperature made it possible to characterize the HAZ. The mechanical properties deteriorated closer to the weld centreline as a result of microstructural coarsening. The regions which had the lowest properties were 4 mm and 5 mm from the weld centreline. However slight improvement in mechanical properties was experienced closer to weld centreline at 4 mm from weld centreline. 4.2 Optical Microscopy Optical microscopy was carried out on the BM and the five regions of the HAZ of AA7075 which had been simulated and the results are as shown in Figures 4.1 to 4.6. The orientation of the specimen was in the longitudinal transverse direction (L-T). The grains were found to have an orientation parallel to the rolling direction. The process of rolling stretching to relieve stresses gives the material pan shaped grains as evident in the micrographs. The optical microstructure of the base metal (Figure 4.1) consists of fine insoluble second phase precipitates dispersed in various locations of the elongated grain. The insoluble precipitates have been identified in literature as the and the strengthening precipitate [40]. There is evidence of microstructural evolution for the regions B to F (Figures 4.2 to 4.6) which have more distinct grain boundaries compared to the base metal. Regions B to F show that the grain boundaries are enlarged as a consequence of the precipitation along the grain boundary. 39

56 50 µ Figure 4.1: Microstructure of the base metal 50 µ Figure 4.2: Microstructure for region B. 40

57 50 µ Figure 4.3: Microstructure for region C. Regions B and C of the HAZ have the finest grains among the five regions in the HAZ of AA7075-T651 which have been simulated as demonstrated in the ASTM grain size number in Figure 4.7. The grain size was found to increase as the distance from the weld centerline decreased. The maximum grain size was found to occur in region E (Figure 4.5) of the HAZ. This implies that grain growth has occurred with the maximum being evidenced at region E. 50 µ Figure 4.4: Microstructure for region D. 41

58 50 µ Figure 4.5: Microstructure for region E. 50 µ Figure 4.6: Microstructure of region F. 42

59 Figure 4.7: Average ASTM grain size number of the base metal and the five regions of the HAZ. The ASTM grain size numbers were found to be highest in regions B (11.430) and C (11.369) of the HAZ as shown in Figure 4.7. This implied that they had the smallest grain size in the HAZ. Regions D, E and F had ASTM grain size numbers , and respectively. 4.3 Fatigue Crack Growth (FCG). The FCG test was carried out at a stress ratio, R= 0.2 and the results are as shown in Figures 4.8 to The results represent the FCG behaviour of the BM and the five regions of the HAZ of AA 7075-T651. There are some variations in the FCG behaviour for the different regions across the HAZ as a consequence of the different thermal peak temperatures experienced. This is evident in the computed values of C and m shown in Table 4.1. There was no significant difference in the values of obtained in this study. 43

60 Figure 4.8: Fatigue crack growth curve for the base metal 44

61 Figure 4.9: Fatigue crack growth in region B. 45

62 Figure 4.10: Fatigue crack growth curve for region C. 46

63 Figure 4.11: Fatigue crack growth curve for region D 47

64 Figure 4.12: Fatigue crack growth curve for region E. 48

65 Figure 4.13: Fatigue crack growth for region F. 49

66 Figure 4.14: Comparison of the FCG curve for the base metal and region F of the HAZ 50

67 Figure 4.15: FCG curves for the base metal and various regions of the HAZ. 51

68 Table 4.1: FCG Parameters K TH, C and m for the BM and the different regions of the HAZ. Specimen K TH (MPa ) C (with da/dn in mm/cycle m and K in MPa ) BM E B E C E D E E E F E The FCG graphs of the base metal and the HAZ in Figure 4.15 clearly demonstrates that the HAZ has varying resistance to crack growth propagation and region F has the least resistance to FCG. There is no significant difference in the fatigue threshold, in the HAZ as shown in Table

69 4.4 Hardness Profile Hardness profile results across the selected regions of the HAZ and the BM of AA7075-T651 are as shown in Figure The general trend in the hardness profile shows that the base metal and region B have the highest hardness. Regions E and F which are close to the weld centerline exhibit low hardness values. These regions have coarse grained microstructure as is evident from the grain size number shown in Figure 4.7. The base metal had a hardness value of 192 HV while region B, C, D, E and F had hardness values of 191 HV, 186 HV, 131 HV, 112 HV and 123 HV respectively. The hardness slightly rises at region F. Figure 4.16: Graphical representation of Vickers Hardness Numbers across the HAZ of AA7075-T

70 4.5 Tensile Properties The results of the longitudinal tensile tests carried out on the base metal and the five regions of the HAZ are as shown in Table 4.2. Three specimens were tested for each region and the average value presented. Region B of the HAZ showed a significantly higher proof stress compared to region E and F which are closer to the fusion line. Table 4.2: Longitudinal tensile properties of the base metal and five regions in the HAZ of AA7075-T651. Region 0.2% Proof stress (MPa) Elongation in 25 mm gauge length (%) BM B C D E F Figure 4.17 shows that there is steep decline in the proof stress from region B of the HAZ to regions D, E and F. The high temperatures experienced in regions close to the weld centerline negatively affect the tensile properties. 54

71 Figure 4.17: Proof stress of the base metal and five regions of the HAZ in the longitudinal direction 4.6 Impact Test Izod Impact Energy The base metal had impact energy of 16 J which was close to that of region B. The results in Figure 4.18 show that the toughness across the HAZ varies with region B of the HAZ depicting higher toughness among the five regions of the HAZ selected for this study. The impact energy of this region was 15 J. Region D of the HAZ had the lowest toughness as indicated by the impact energy graph in Figure The toughness in region E and F were 12 J and 13 J respectively. 55