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Design and Development of Self-Passivating Biodegradable Magnesium Alloys Using Selective Element Oxidation

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Title:
Design and Development of Self-Passivating Biodegradable Magnesium Alloys Using Selective Element Oxidation
Creator:
Brar, Harpreet S
Place of Publication:
[Gainesville, Fla.]
Florida
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University of Florida
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english
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Thesis/Dissertation Information

Degree:
Doctorate ( Ph.D.)
Degree Grantor:
University of Florida
Degree Disciplines:
Materials Science and Engineering
Committee Chair:
Myers, Michele Viola
Committee Members:
Keselowsky, Ben
Dempere, Luisa A
Batich, Christopher D
Sarntinoranont, Malisa
Graduation Date:
5/5/2012

Subjects

Subjects / Keywords:
Alloying ( jstor )
Alloys ( jstor )
Binary alloys ( jstor )
Biomaterials ( jstor )
Corrosion ( jstor )
Grain size ( jstor )
Magnesium ( jstor )
Oxidation ( jstor )
Oxides ( jstor )
Thermodynamics ( jstor )
Materials Science and Engineering -- Dissertations, Academic -- UF
alloy-design -- biodegradable -- corrosion -- magnesium -- orthopaedic -- oxidation -- rare-earth
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bibliography ( marcgt )
theses ( marcgt )
government publication (state, provincial, terriorial, dependent) ( marcgt )
born-digital ( sobekcm )
Electronic Thesis or Dissertation
Materials Science and Engineering thesis, Ph.D.

Notes

Abstract:
Metallic biomaterials such as stainless steels, titanium alloys, and cobalt-chromium alloys have been used as structural implant materials for many years. However, due to their limitations in temporary implant applications, there has been increased interest in the development of a biodegradable structural implant device. Magnesium (Mg) alloys have shown great potential as a material for biodegradable structural implant applications. However, low strength and high degradation rate of Mg under physiological conditions are major limitations, causing the implant to lose its structural integrity before the healing process is complete. The main aim of this work was to investigate the possibility of designing Mg-based alloys with ability to form selective protective oxides, thereby aiding in the reduction of the initial degradation rate. A thermodynamics-driven design was utilized to select three elements, namely Gadolinium (Gd), Scandium (Sc) and Yttrium (Y), due to the low enthalpy of formation associated with their oxide species. First, binary alloys were cast under inert atmosphere, solution treated and investigated for degradation rate in Hanks' solution. The Mg-Gd binary alloy showed the fastest degradation rate whereas the Mg-Sc binary alloy showed the slowest degradation rate. The degradation of Mg-Gd and Mg-Y was 18 and 5 times faster than Mg-Sc alloy, respectively. The microstructural analysis of the alloys was performed using X-ray Diffraction (XRD), Optical Microscopy (OM) and Scanning Electron Microscopy (SEM). It was observed that the grain size of Mg-Sc alloys is significantly smaller than Mg-Gd and Mg-Y alloys and can be a contributing factor to the reduction in degradation rate. The hardness behavior of the alloys was also investigated using Vickers microhardness Testing. To understand the oxidation behavior and kinetics, samples were oxidized in pure oxygen environment and investigated using microstructural and thermogravimetric analysis (TGA). Auger electron spectroscopy (AES) and X-ray photoelectron spectroscopy (XPS) were also used to characterize the chemistry of the surface oxides. Selective oxidation of the alloying species was observed on the surface of all the alloys, indicating a strong driving force for their formation out of solid solution Mg. The degradation rate of the oxidized samples was also investigated in Hanks' solution and compared with the rate of freshly polished samples. The oxide formation on Mg-Y alloys was found to be most protective as it reduced the degradation rate by more than 50%. On the other hand, oxidized Mg-Sc samples did not show any appreciable decline in degradation rate as compared to polished samples. Based on the information gathered from the binary alloys, ternary alloy system was selected. The thermodynamic and empirical models were applied to predict the properties of the alloy. The models were validated using the techniques mentioned above. It was observed that the predictions matched with the experimental results. ( en )
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In the series University of Florida Digital Collections.
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Includes vita.
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Includes bibliographical references.
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Description based on online resource; title from PDF title page.
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This bibliographic record is available under the Creative Commons CC0 public domain dedication. The University of Florida Libraries, as creator of this bibliographic record, has waived all rights to it worldwide under copyright law, including all related and neighboring rights, to the extent allowed by law.
Thesis:
Thesis (Ph.D.)--University of Florida, 2012.
Local:
Adviser: Myers, Michele Viola.
Statement of Responsibility:
by Harpreet S Brar.

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Copyright Brar, Harpreet S. Permission granted to the University of Florida to digitize, archive and distribute this item for non-profit research and educational purposes. Any reuse of this item in excess of fair use or other copyright exemptions requires permission of the copyright holder.
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1 DESIGN AND DEVELOPMENT OF SELF PASSIVATING BIODEGRADABLE MAGNESIUM ALLOYS USING SELECTIVE ELEMENT OXIDATION By HARPREET SINGH BRAR A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 2012

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2 2012 Harpreet Singh Brar

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3 To my parents

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4 ACKNOWLEDGMENTS I would like to thank my advisor, Prof. Michele Viola M yers for providing me with an opportunity to work in her group and offering constant guidance, support and encouragement during my PhD. I would also like to thank her for the innumera ble scientific discussions that we had during the course of my study. I would also like to express my gratitude toward s my other committee members Prof. Benjamin Keselowsky Prof. Malisa Sarntinoranont Prof. Luisa Dempere and Prof. Christopher Batich for their contributions and their trust in my ability to perform this research I was also lucky to have wonderful group members who provided a very friendly and cordial atmosphere in the lab. I would like to thank Derek Hsu, Zachary Bryan, Hunter Hender son, Charles Fisher, Fatmata Barrie, Glenn Bean, Billy Valder ama, Ida Berglund, Ryan Hooper and Jack Tilka for the numerous thought provoking discussions and helpful suggestions throughout my PhD I wholeheartedly thank my friends Prateek, Reno, Tejas, Sam bhav, Sid, Aniket, Tara, Isis, Alexa, Amber and others who I may have missed, for keeping me entertained throughout my stay in Gainesville as well as for providing support and encouragement at times of need Their presence in my life greatly helped me on t he long and hard road to PhD. Last but not the least I would like to thank my family my father Mr. Nachhattar Singh, mother Harjit Kaur, sister Manpreet Kaur and brother Gurpreet Singh. W ithout their unwavering guidance and support, this work would never have been possible. My parents instilled in me the vir tues of honesty, humility, hard work and patience and stood by me through the ups and downs of my life. I can never be able to pay back their debt of unconditional love and care, and this work is d edicated to them.

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5 TABLE OF CONTENTS page ACKNOWLEDGMENTS ................................ ................................ ................................ .. 4 LIST OF TABLES ................................ ................................ ................................ ............ 7 LIST OF FIGURE S ................................ ................................ ................................ .......... 8 LIST OF ABBREVIATIONS ................................ ................................ ........................... 11 ABSTRACT ................................ ................................ ................................ ................... 12 CHAPTER 1 INTRODUCTION ................................ ................................ ................................ .... 14 1.1 Mo tivation ................................ ................................ ................................ ......... 14 1.2 Magnesium as a Potential Biomaterial ................................ .............................. 16 1.3 Historical Use of Magnesium as a Biomaterial ................................ .................. 18 2 BACKGROUND ................................ ................................ ................................ ...... 23 2.1 Corrosion of Magnesium and its Alloys ................................ ............................. 23 2.1.1 Thermodynamics of Magnesium Corrosion ................................ ............. 23 2.1.2 Oxidation Behavior and Surface Film Formation ................................ ..... 26 2.1.2.1 Oxidation of metals ................................ ................................ ........ 26 2.1.2.2 Surface film formation on magnesium ................................ ............ 27 2.1.3 Metallurgical Aspects of Corrosion of Magnesium and its Alloys ............. 30 2.1.3.1 Effect of impurities ................................ ................................ ......... 31 2.1.3.2 Effect of solid solution and secondary phases ............................... 32 2.1.3.3 Effect of grain size ................................ ................................ .......... 33 2.2 Strengthening of Magnesium Alloys ................................ ................................ .. 34 2.3 Summary ................................ ................................ ................................ .......... 36 3 DESIGN APPROACH ................................ ................................ ............................. 41 3.1 Curren t Trends in Biodegradable Magnesium Alloy Development .................... 41 3.2 Systems Design Approach ................................ ................................ ................ 41 4 DESIGN AND MICROSTRUCTURAL ANALYSIS OF BINARY MG RE ALLOYS .. 45 4.1 Design of Binary Alloys ................................ ................................ ..................... 45 4.1.1 Selection of Alloying Elements ................................ ................................ 46 4.1.1.1 Design considerations for surface oxide formation ........................ 46 4.1.1.2 Design considerations for increasing strength by solution strengthening ................................ ................................ .......................... 47

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6 4.1.1.3 Design considerations for grain size reduction ............................... 49 4.1.2 Selection of Alloying Compositions ................................ .......................... 49 4.2 Materials and Methods ................................ ................................ ...................... 52 4.2.1 Alloy Preparation ................................ ................................ ..................... 52 4.2.2 Microstructural Characterization ................................ .............................. 53 4.2.3 Hardness Testing ................................ ................................ .................... 53 4.3 Re sults and Discussion ................................ ................................ ..................... 53 4.4 Summary ................................ ................................ ................................ .......... 56 5 CHARACTERIZATION OF THE OXIDATION BEHAVIOR OF BINARY MG RE ALLOYS ................................ ................................ ................................ .................. 67 5.1 Thermodynamic Calculations and Prediction of Possible Oxidation Behavior .. 67 5.2 Materials and Methods ................................ ................................ ...................... 68 5.3 Results and Discussion ................................ ................................ ..................... 70 5.3.1 Oxidation Kinetics of Binary Mg RE Alloys ................................ .............. 70 5.3.2 Characterization of Surface Oxides on Mg RE Alloys ............................. 72 5.4 Summary ................................ ................................ ................................ .......... 74 6 IN VITRO DEGRADATION OF ALLOYS UNDER DIFFERENT SURFACE CONDITIONS ................................ ................................ ................................ ......... 87 6.1 Materials and Methods ................................ ................................ ...................... 87 6.2 Results and Discussion ................................ ................................ ..................... 88 6.3 Summary ................................ ................................ ................................ .......... 94 7 INVESTIGATION OF THE OXIDATION AND DISSOLUTION BEHAVIOR OF TERNARY MG SC Y ALLOY ................................ ................................ ................ 103 7.1 Design of Ternary Alloy ................................ ................................ ................... 103 7.2 Materials and Methods ................................ ................................ .................... 105 7.3 Results and Discussion ................................ ................................ ................... 107 7.4 Summary ................................ ................................ ................................ ........ 110 8 CONCLUSIONS ................................ ................................ ................................ ... 126 9 FUTURE WORK ................................ ................................ ................................ ... 128 LIST OF REFERENCES ................................ ................................ ............................. 130 BIOGRAPHICAL SKETCH ................................ ................................ .......................... 138

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7 LIST OF TABLES Table page 1 1 Summary of mechanical properties of natural and implant materials ................. 21 2 1 Chemical potential of Mg and its compounds in different states at 25C ............ 37 4 1 Maximum solid solubility of different alloying elements in Mg ............................. 57 4 2 Slope of liquidus line (m), Equilibrium distribution coefficient (k) and growth restriction parameter m(k 1) for different alloying elements in Mg ...................... 57 4 3 Nominal and actual compositions of binary Mg RE alloys ................................ .. 57 5 1 Parabolic rate constants for different alloys ................................ ........................ 76

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8 LIST OF FIGURES Figure page 1 1 Pourbaix diagram for the Mg H 2 O system at 25 o C ................................ ............. 22 1 2 Photograph of the tubular slot balloon expandable WE43 alloy stent ................. 22 2 1 Electrochemical force series ................................ ................................ ............... 38 2 2 Schematic illustrations of the main aspects of metal oxygen reactions and surface oxide growth. ................................ ................................ ......................... 39 2 3 Schematic diagram of a dislocation cutting through a particle ............................ 40 2 4 Schematic diagram of a dislocation passing between widely spaced ................................ .......... 40 3 1 Systems design chart for developing Mg based biodegradable alloys ............... 44 4 1 Schematic representation of the desired structure and the relative degradation rates ................................ ................................ ................................ 58 4 2 The free energy of formation of various oxides versus their volume per mole of metal at 25C ................................ ................................ ................................ .. 58 4 3 Solid solution strengt hening model ................................ ................................ ..... 59 4 4 Phase diagrams of the Mg rich region of the binary alloys ................................ 60 4 5 Optical micrographs of Mg 8Gd alloy ................................ ................................ 61 4 6 Polarization resistance versus immersion time plot for Mg Y alloys, HP Mg and LP Mg during 24 hour immersion in 0.1 M NaCl ................................ .......... 62 4 7 Glove box used for melting alloys under inert atmosphere ................................ 63 4 8 Optical micrographs of etched binary alloys ................................ ...................... 64 4 9 SEM micrographs of homogenized binary alloys ................................ ................ 64 4 10 The XRD plots of homogenized Mg 3RE alloys ................................ .................. 65 4 11 Experimental and calc ulated micro hardness of the binary Mg RE alloys .......... 66 5 1 The isothermal TGA plots of binary alloys at 500C ................................ ........... 77 5 2 Kinetic parameters for Mg 3Gd for the 3 different oxidation regimes .................. 77

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9 5 3 1/2 for the binary alloys ................................ ........................ 78 5 4 The XRD plots of the binary alloys after oxidation at 500C for 24 hours ........... 79 5 5 AES depth profiles of binary alloys oxidized a t 500C for 5 hours ...................... 80 5 6 SEM micrographs of the cross section of oxidized Mg 3Gd alloy ....................... 81 5 7 SEM micrographs of the cross section of oxidized Mg 3Y alloy ......................... 82 5 8 SEM micrographs of the cross section of oxidized Mg 3Sc alloy ....................... 83 5 9 AES depth profile of polished Mg 3Sc alloy ................................ ........................ 83 5 10 XPS survey scan of entire binding energy for binary alloys ................................ 84 5 11 XPS spectra of Mg 3Gd oxidized for 24 h ours ................................ .................... 84 5 12 XPS spectra of Mg 3Y oxidized for 24 hours ................................ ...................... 85 5 13 XPS spectra of Mg 3Sc oxidized for 24 hours ................................ .................... 86 6 1 Experimental setup for measuring hydrogen evolution of samples ..................... 96 6 2 Hydrogen evolution rate of different oxidized and non oxidized alloys ............... 96 6 3 Optical micrographs of corrosion propagation in Mg 3Sc alloys ......................... 97 6 4 Optical micrographs of corrosion propagation in Mg 3Y alloys ........................... 97 6 5 Optical micrographs of corrosion propagation in Mg 3Gd alloys ........................ 98 6 6 SEM micrographs of Mg 3Sc alloy surface after 96 hours of degradation .......... 98 6 7 SEM microg raphs of Mg 3Y alloy surface after 96 hours of degradation ............ 99 6 8 The SEM micrographs of Mg 3Gd surface after 96 hours of dissolution .......... 100 6 9 XRD plots of Mg RE binary alloys after 96 hours of immersion ........................ 101 6 10 Optical micrograph showing the selective corrosion of grains in Mg 3Y alloy .. 102 7 1 Effect of Sc additions on the grain size and the rate of hydrogen evolution in binary Mg Sc alloys ................................ ................................ .......................... 112 7 2 Yield strength vs hardness data fitted to a power law ................................ ....... 113 7 3 Optical micrograph of solution t reated Mg 3Sc 3Y alloy ................................ .... 113 7 4 XRD plot of Mg 3Sc 3Y alloy homogenized at 500C for 8 hours .................... 114

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10 7 5 E lemental mapping and ED S of solution treated Mg 3Sc 3Y alloy ................... 115 7 6 Plot showing the weight gain per cm 2 vs time ................................ ................... 116 7 7 Plot of mass gain per cm 2 vs square root of time ................................ ............. 116 7 8 SEM micrographs of the cross section of oxidized Mg 3Sc 3Y ........................ 117 7 9 XRD plot of Mg 3Sc 3Y alloy oxidized for 5 hours at 500C ............................. 118 7 10 AES depth profile of Mg 3Sc 3Y alloy oxidized for 5 hours at 500C ............... 118 7 11 XPS survey of Mg 3Sc 3Y alloy oxidized for 5 hours at 500 C ........................ 119 7 12 XPS multiplex peaks of the oxidized ternary alloy ................................ ............ 120 7 13 Hydrogen evolution behavior of Mg 3Sc 3Y alloy under different surface conditions ................................ ................................ ................................ ......... 121 7 14 XRD plot of Mg 3Sc ........ 122 7 15 SEM micrographs of the corroded surface of non oxidized Mg 3Sc 3Y alloy ... 123 7 16 SEM micrographs of the corroded surface of oxidized Mg 3Sc 3Y alloy .......... 124 7 17 The comparison between the predicted and experimental values of hardness and yield strength of Mg 3Sc 3Y alloy ................................ .............................. 125

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11 LIST OF ABBREVIATION S Al Aluminum Gd Gadolinium Gibbs Free Energy H 2 Hydrogen Gas HA Hydroxyapatite ICP Inductively Coupled Plasma Mg Magnesium MgCl 2 Magnesium Chloride Mg(OH) 2 Magnesium Hydroxide MgO Magnesium Oxide Ni Nickel OM Optical Microscopy P B Pilling Bedworth PDS Poly Dioxanone PGA Poly Glycolic Acid PLA Poly Lactic Acid RDA Recommended Daily Allowance RE Rare Earth Sc Scandium SEM Scanning Electron Microscopy UHMWPE Ultra Hig h Molecular Weight Poly Ethylene XRD X ray Diffraction Y Yttrium

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12 Abstract of Dissertation Presented to the Graduate School of the University of Florida in Florida in Partial Fulfillment of the Requirements for the Degree of Doctor of Philosophy DESIGN AND DEVELOPMENT OF SELF PASSIVATING BIODEGRADABLE MAGNESIUM ALLOYS USING SELECTIVE ELEMENT OXIDATION By Harpreet Singh Brar May 2012 Chair: Michele Viola M y ers Major: Materials Science and Engineering Metallic biomaterials such as stainless steels, titanium alloys, and cobalt chromium alloys have been used as structural implant materials for many years. However, due to their limitations in temporary implant applications, there has been increased interest in the development of a bio degradable structural implant device. Magnesium (Mg) alloys have shown great potential as a material for biodegradable structural implant applications. However, low strength and high degradation rate of Mg under physiological conditions are major limitatio n s causing the implant to lose its structural integrity before the healing process is complete. The main aim of this work was to investigate the possibility of designing Mg based alloys with ability to form selective protective oxides, thereby aiding in the reduction of the initial degradation rate. A thermodynamic s driven design was utilized to select three elements namely G adolinium (Gd) Sc andium (Sc) and Y ttrium (Y ) due to the low enthalpy of formation associated with their oxide species First, bin ary alloys were cast solution. The Mg Gd binary alloy showed the fastest degradation rate whereas the Mg Sc binary alloy showed the slowest degradation rate. The degra dation of Mg Gd and

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13 Mg Y was 18 and 5 times faster than Mg Sc alloy respectively The microstructural analysis of the alloys was performed using X ray Diffraction (XRD), Optical Microscopy (OM) and S canning Electron Microscopy (SEM). It was observed that the grain size of Mg Sc alloys is significantly smaller than Mg Gd and Mg Y alloys and can be a contributing factor to the reduction in degradation rate. The hardness behavior of the alloys was a lso investigated using Vickers m icrohardness Testing. To und erstand the oxidation behavior and kinetics, samples were oxidized in pure oxygen environment and investigated using microstructural and thermo gravimetric analysis (TGA). Auger electron spectroscopy (AES) and X ray photoelectron s pectroscopy (XPS) were als o used to characterize the chemistry of the surface oxides. Selective oxidation of the alloying species was observed on the surface of all the alloys indicating a strong driving force for their formation out of solid solution Mg. The degradation rate of t he compared with the rate of freshly polished samples. The oxide formation on Mg Y alloys was found to be most protective as it reduced the degradation rate by more than 50%. On the other hand, oxidized Mg Sc samples did not show any appreciable decline in degradation rate as compared to polished samples. Based on the information gathered from the binary alloys, ternary alloy system was selected The thermodynamic and empirical models were appli ed to predict the properties of the alloy. The models were validated using the techniques mentioned above. It was observed that the predictions matched with the experimental results.

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14 CHAPTER 1 INTRODUCTION 1.1 Motivation Advances in the field of medical science over the last century have significantly improved the life expectancy of people around the world. According to the US Census Bureau, 1 out of every 5 people in US will be over 65 yea rs of age by the year 2030. Additionally 80% of the people above the age of 65 have at least one chronic condition [1] Hence, the increasing aging population and the desire for an active lifestyle cre ate a driving force for the high performance biomedical implant materials. Orthopedic and cardiovascular implants constitute the two major sectors of the medical implant industry. The global market for orthopedic devices was an estimated $31.6 billion in 2 007 and the cardiovascular stent market stood around $5 billion [2] Metallic materials play an important role as biomaterials for load bearing applications such as stents and orthopedic implants. Of the total number of orthopedic operations performed per year in the United States, open reduction of f racture and internal fixation are the most comm on one s [3] Traditional implant materials like stainless steel, titanium all oys and cobalt chromium alloys are designed to stay in the body permanently. They are intended to remain neutral in vivo and not interact with the body [4] However, long term presence of the implant materia ls has its own specific drawbacks. In coronary stents, for example, bare metal stents can lead to thrombogenicity, restenosis, physical irritation inflammatory local reactions and mismatch in the mechanical behavior of the stented and non stented vessel a reas [5] Development in the drug eluting stents have reduced the incidence of restenosis, the other problems however, still persist. Similarly, use of permanent implants like screws

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15 and plates for secu ring serious fractures also have their side effects. Secondary surgeries maybe required to remove the implants once the bone is healed, thereby increasing the cost of treatment and patient morbidity [6] Additionally, the difference in the modulus of the present implant materials and the bone can cause stress shielding effects lea ding to reduction in strength of the healed tissue [7, 8] Biodegradable materials have the possibility of providing more physiological repair, with temporary, limited support and better tissue growth. Polymers were the first m aterials to be commercially used for biodegradable and bioabsorbable applications with poly glycolic acid (PGA), poly lactic acid (PLA) and poly dioxanone (PDS) being the earliest adopted and most commonly used absorbable materials [9] The use of biodegradable polymers in load bearing applications has been mainly limited by their low strength, as a large amount of material is be needed to attain the strength required for load bearing applications. Metals on the other hand, have desirable mechanical properties like high strength and fracture toughnes s. However, most of the metals are either non biodegradable or show toxicity to human body. Corrosion products of conventional metallic implant materials like stainless steel, cobalt chromium and nickel based alloy have been shown to be harmful to human bo dy [10 13] This need for a biodegradable material with improved mechanical properties and the ability to degrade in the body without releasing any harmful byproducts has lead to increased research interest on magnesium (Mg) and its alloys. Magnesium is an exceptionally lightweight material with high specific strength [14] It is one of the most reactive metals and degrades under physiological conditions [15] thereby providing an opportunity to develop it into a biodegradable implant materia l for

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16 l oad bearing applications. Additionally, i ts elastic modulus and strength are closer to the natural bone than other commonly used metallic implants [16] Table 1 1 compares the mechanical properties of bone and various materials being used for implant applications. Despite all these advantages, Mg has certain limitations like low strength and high degradation rate that pose a c hallenge in its use as a structural implant material. These limitations can be alleviated by addition of alloying elements to pure Mg or through use of surface coatings [16 18] However, alloy chemistry strongly drives microstructural evolution and desired properties often need conflicting complex microstructural features, making it difficult to find a balance of the properties by using a traditional, edisonian type approach and This work was performed with a focus on designing biocomaptible Mg alloys that is capability of growing a selective for implan t applications using systems design approach. Emphasis was given on utilizing a systematic approach to find a balance between the properties required for implant applications and optimizing the overall performance of the alloys. 1.2 Magnesium as a Potential Biomaterial Magnesium is the second most abundant element involved in cellular structures and is the fourth most abundant cation in human body. The body of an average adult contains around 21 28 g Mg, with more than half of it stored in the bones [19] Mg is an essential element for human metabolism as it acts as a cofactor for many enzymes and stabilizes the structure of DNA and RNA. The US recommended daily dietary allowances (RDA) for Mg are 320 mg/day for women and 420mg/day for men [20] Deficiency of Mg has been linked to various diseases like osteoporosis, decreased

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17 membrane integrity and function, increased risk of cardiovascular problems like cardiac arrhythmias, vasoconstriction of coronary arteries, increased blood pressure and accelerated aging [21] S tudies have also indicate d the possibility of stimulatory effects of Mg on the growth of new bone tissue due to its functional role and presence in bone [22, 23] It binds strongly to phosphates, ther eby influencing the mineralization of bony tissue by its control of hydroxyapatite (HA) formation As a result, Mg is often seen in dietary supplements and drugs. As compared to other metal ions, human body can accommodate higher amounts of Mg concentratio n levels in serum, with hypotension and respiratory distress occurring at levels exceeding 1.05 mmol/L [19] In ad dition, the concentration of Mg in human body is efficiently controlled by homeostatic mechanism and excess Mg excreted in urine, thereby making the incident of Mg toxicity extremely rare [19] In addition to its biocompatibility, Mg also has desirable mecha nical properties It is an exceptionally lightweight metal an d has a specific density of 1.74 g/cm 3 making it 1.6 and 4.5 times lighter than aluminum and steel respective ly [16] Furthermore, its density is similar to that of bone (1.8 2.1 g/ cm 3 ) [16] The specific strength of Mg is higher than conventional metallic implants, thereby reducing the amount of material requir ed for a given load application. As shown in Table 1 1, the elastic modulus of Mg (45 GPa) is closest to that of human bone when compared to other metallic implant materials This helps in reducing the modulus mismatch between the bone and the implant, the reby reducing the stress shielding effect and improving the bone strength. Mg is also one of the most reactive metals of the periodic table In aqueous environments, Mg dissolution generally proceeds by an electrochemical reaction with

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18 water to produce ma gnesium hydroxide (Mg(OH) 2 ) and hydrogen gas (H 2 ) [24] The overall reaction of Mg dissolu tion in water can be written as: (1 1) The overall reaction stated above can also be expressed as a sum of following partial reactions [17] : (anodic reaction) (1 2) (cathodic reaction) (1 3) (product formation) (1 4) The rate of degradation reaction depends on various factors like pH, solution compositi on, al loying elements, microstructure and presence of proteins. Figure 1 1 shows the Pourbaix diagram for Mg H 2 O system at 25C [25] It can be seen that under the physiological conditions of pH (7.2 7.4) Mg undergoes corrosion to produce Mg(OH) 2 Furthermore, pr esence of chloride ions in physiological media reduce the efficiency of Mg (OH) 2 as they convert it to magnesium chloride ( MgCl 2 ) which has a higher solubility. This leads to increased corrosion attack on Mg [17] 1.3 Historical Use of Magnes ium as a B iomaterial The first reported use of Mg as a biomaterial was by physician Edward C. Huse who used Mg wires as ligatures to stop bleeding vessels of 3 patients in 1878. He observed and documented the degradable properties of the material and the d ependence of degradation time on the size of Mg wire used [26] Austrian physic ian Erwin Payr started his first experiments on Mg resorption in 1892 [27] In 1900, he perf or med in vivo experiments on pigs and the femoral arteries of the dogs using vessel connectors made of Mg [27, 28] Albin Lambotte and his assistant Verbrugge used Mg

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19 for osteosynthetic application for the first time in 1906 when t he y used M g plate with 6 steel s crews to stabilize a fracture on the tibia [29] However, the implant degraded rapidly and resulted in extensive subcutaneous gas cavities, local swelling and pain to the patient. The small fragments of the Mg plate were left after 8 days and were retrieved. He concluded that the extensive dissolution of Mg took place due to the electrochemical reaction between Mg plate and the steel screws [29] After further investigation and animal studies [30] he continued using Mg screws and plates for healing fractures in children [30, 31] Based on his new encouraging results, he recommended using Mg without combining it with other elements to prevent galvanic co rrosion [30] critical reports on the use of Mg as implant material in bone applications were also published The major concerns in these studies were the production of abscess gas cavities and rapid degradation of Mg [32, 33] Troitskii and Tsitrin published one of the most extensive studies on the use of Mg alloys for osteosynthetic application in 1948. They reported the use of plates and screws mad e of Mg Cd alloy to treat of 34 cases of pseudoarthrosis in humans [34] There were 9 failures out of the 34 cases, which were attributed to infection or other factors. The corrosion process released hydrogen, which was drawn off using subcutaneous needles. No inflammatory responses or increase in serum levels of Mg was observed, and most of the implants maintained their mechanical integrity f or about 6 8 weeks [34] With improvements in corrosion protection, ther e has been a renewed interest in use of Mg for biodegradable implant applications. The most significant achievement of Mg as biomaterial has come from the use of Mg alloy as biodegradable coronary stents.

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20 Heublein and colleagues conducted the pioneering in vestigations on the suitability of AE21 alloy con taining 2% aluminum (Al) and 1% rare earth metals as a biodegradable stent in porcine model [5] The implantations demonstrated very low thrombogenic and inflammatory responses. Subsequently, a modified Mg alloy based Absorbable Magnesium Stent ( ABS) was developed by Biotronik. The stent was constructed using laser from a single tube of WE43 alloy containing Mg (>90%), Zirconium (Zr) (<5%), Yttrium (Y) (<5% ) and other rare earths (RE) (<5%) [35] Figure 1 2 shows the image of the stent [35] I nitial animal studies in porcine cor onary arteries showed promising results and le d the way for clinical trials of the stent. Initially, 20 patients with symptomatic critical limb isc hemia were treated with AB S [36] The preliminary analysis after stent implantation did not show any evidence of toxicity. A 3 month follow up showed 89.5% primary clinical patency and no major amputation was required [36] The positive results of this study resulted in initiation of first coronary clinical study. In this study, WE43 stents were implanted in 63 patients with anginal symptoms and single de novo lesions. Though the stents showed good initial results, final cross section area of the vessel was smaller than the non biodegradable stents. This was attributed to the early recoil as the stent completely degraded after 4 months [37]

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21 Table 1 1. Summary of m echanical properties of natural and implant m aterials Material Tensile Strength (MPa) Elastic Modulus (GPa) Bio degradable Natural Materials Collagen 60 b 1 b Yes Cortical bone 100 200 b 10 20 b Yes Inorganic Materials Magnesium 185 232c 41 45d Yes Stainless Steels 480 834 b 193 b No Cobalt Alloys 655 1400 b 195 210 b No Titanium Alloys 550 985 b 100 105 b No Platinum Alloys 152 485 b 147 b No Synthetic Hydroxy apatite 600 d,* 73 117d No/Yes Organic Materials L PLA 28 50 a 1.2 3 a Yes D,L PLA 29 35 a 1.9 2.4 a Yes UHMWPE 39 40 b 0.94 1.05 b No *Indicates compressive strength (MPa) (a) [ Engelbe rg I, Kohn J, 1991. Biomaterials. Vol. 12 Pages 292 304 ] (b) [Black J, 1988. Orthopaedic biomaterials in research and practice. New York: Churchill Livingstone] (c) [ Smit hells CJ, Gale WF, Totemeier TC, 2004. Smithells metals reference book. Amsterdam (various pagings) ] (d) [ Staiger M P, Pietak AM, Huadmai J, Dias G, 2006 Biomaterials Vol. 27 Pages 1728 34 ]

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22 Figure 1 1 Pourbaix diagram for the Mg H 2 O system at 25 o C [Adapted from Marcel Pourbaix, 1966. Atlas of Electrochemical Equilibria in Aqueous Solutions] Figure 1 2 Photograph of the tubular slot balloon expandable WE43 alloy stent [Reprinted with permission from John Wiley and Sons 2004. Journal of Interventional Cardiology Vol. 17 ( Pages 391 395 Figure 1) ]

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23 CHAPTER 2 BACKGROUND 2.1 Corrosion of Magnesium and its Alloys Currently, there is a significant effort to investigate various different alloys for structural biodegradable implant materials The high chemical activity of Mg allows it to corrode in aqueous environment, thereby making it a potential candidate for biodegradable applications. However, it is this very ability of Mg to corrode in an aqueou s environment that is also one of the biggest hurdles in its use as an implant material U nder physiological conditions, the corrosion rate is too rapid and the implants tend to loose their structural integrity before the tissue is properly healed, thereby leading to implant failure. It is therefore very important to improve and control the corrosion properties of the Mg alloys so that they can be used for biodegradable applications. This section focuses on the thermodynamics of pure Mg strengthening mech anisms, oxidation behavior and surface films and the effect of microstructure on the corrosion and oxidation of Mg alloys. 2.1.1 Thermodynamics of Magnesium Corrosion In order to understand the corrosion behavior of Mg and its alloys, it is vital to under stand the thermodynamic stability of pure Mg and its compounds in various environments This information can provide a valuable insight into the corrosion behavior of Mg and enable predictability of its behavior in those environments. As compared to other structural material s Mg has the lowest standard electrochemical potential, as shown in Figure 2 1 Pure Mg (Mg 2+ /Mg) in contact with solution containing Mg 2+ ions at 25C has an electrode potential of 2.37 V nhe [24] However, for practi cal purposes, actual corrosion potential of Mg in contact with dilute

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24 chloride solutions is 1.7 V nhe [24] This difference between the theoretical standard potential and actual corrosion potential is due to the formation of surface film of Mg(OH) 2 or MgO [24] Due to its reactivity, Mg has a strong driving force for oxidation. Under natural environments, Mg spontaneously transforms into its oxidized states. The driving force for these oxidation reactions is the reduction in Gibbs as mentioned below [17] : (2 1) (2 2) (2 3) As a result of the thermodynamic driving forces, exposure of Mg to environments containing oxygen or water results in the formation of oxides or hydroxides on the surface [40, 41] The stability of the oxidized states of Mg over its metallic state is also supported by the thermodynamic data of the Mg compounds and species as shown in Table 2 1. The more negative the chemical potential, the more stable the species is compared to others in the media. From the data in Table 2 1 it can be seen that corrosion of Mg is a spontaneous process as Mg 2+ MgO and Mg(OH) 2 have more negative potentials. The final product formed depends on the media to which Mg is b eing exposed. It can be further noted that since chemical potential of Mg(OH) 2 is more negative than MgO, exposure of Mg to aqueous solutions will lead to the preferential formation of Mg(OH) 2 This transformation of Mg to its hydroxide proceeds according to Equation 2 1 releasing H 2 as a byproduct.

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25 Pourbaix diagram and other E pH diagrams can be used to predict the theoretical stability of Mg in water. The p ourbaix diagram of Mg and pure water is shown in Figure 1 1. It is known that Mg has the tendency t o be oxidized into ions, oxides or hydroxides in most of the E pH regions, resulting in a very large corrosion domain. [14, 17, 25, 42] Surface passivity is only possible in the narrow negative potential region and high alkalinity regi on (pH > 10.5). Under high pH conditions, even though Mg(OH) 2 forms as the corrosion product film on the surface it does not provide complete protection since it is semi protective in nature [43] On the other hand, t he negative potential required for immunity is also significantly more negative than its equ ilibrium potential. As a result, even though the self passivity of Mg is theoretically possible, most of the common solutions do not fulfill the requirements for passivity. Though the E pH diagrams help us understand the activity and passivity of Mg surfa ce, they have some limita tions These diagrams offer limited information on the kinetics of the corrosion process and only predict the thermodynamic stability or driving force for corrosion. M a g nesium corrosion has been shown to be highly dependent on the kinetics [17] Another limitation of the E pH diagrams is that they assume uniformity of substances or phases. In practical systems this is seldom true, as variations in the co ncentration of ions, pH and phases exist. Also, typical E pH diagrams do not take into account the effect of chemical composition of the solution or the presence of aggressive species in the solution. Hence, extreme care should be taken while using these d iagrams to attempt corrosion prediction in solutions other than water.

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26 2.1.2 Oxidation Behavior and Surface Film Formation 2.1.2.1 Oxidation of metals A surface film can be described as a layer of corrosion products formed on the surface in presence of ambient environment When the ambient environment is oxygen or water rich the surface film formed constitutes of the metal oxide or hydroxide [17] Typically, t he formation of a surface film is thermodynamically and kinetically spontaneous on the surface of reactive metals. However, the actual oxidation mechanism and reaction path can depend on various factors. Figure 2 2 shows some of the stages and aspects of metal oxidation. T he initial step in metal oxygen reaction on a clean metal surface involves the absorption of oxygen on the metal surface. As the reaction proceeds, the oxygen can dissolve into the metal, leading to the formation of the oxide as a surface film or a separ ate oxide nuclei. When a continuous film covers the surface, it separates the metal from the gas, and the reaction can proceed only through solid state diffusion of reactants through the film. Some metals may form scales that are porous in nature having m icrocracks and porosity, and even macrocracks. Such scales do not act as solid state barrier between the reactants. In such cases, the oxide layers are non protective and may not passivate the surface [44] The passiv ity and protective nature of the surface oxide can be experimentally determined by measuring the reaction rate and kinetics of oxidation. One of the methods to do so is thermogravimetery, where the change in the weight of the metal is recorded over time. T his data can be used to identify the rate equations, which can then be used to classify the oxidation behavior of the metal [44, 45]

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27 In multi component alloys, oxidation becomes a complex phenomenon with the inclusion of higher order oxides, different oxygen affinities, and mobili ties. However, the oxide with the most negative free energy of formation under oxidation conditions has the highest probability to form preferentially compared to other oxides. Another important aspect is whether the oxide forms as sub surface precipitates known as internal oxidation, or as a scale on the surface due to external oxidation. If the outward flux of the selectively oxidized component is substantially smaller than the inward flux of oxygen, internal oxidation takes place. 2.1.2.2 Surface film f ormation on magnesium As Mg tends to oxidize or dissolve in mo st environments its surface is also covered with a surface film under normal conditions. A dditionally, a surface film on Mg can vary in composition and structure based on the alloying additions environment and formation conditions. In general, the spontaneous surface film formed on Mg in air is relatively thin (approximately 2 nm), though it can grow with time [46] However, the nature of the surface film on Mg is not very well understood [24] Based on thermodynamic free energy of formation, the surface film on Mg should consist of MgO in dry environments and Mg(OH) 2 in aqueous environments. However, under normal atmospheric conditions where some amount of water vapor is generall y present, surface films on Mg contain both Mg O and Mg(OH) 2 [47 49] Even if the original film is formed under dry conditions and consists only of MgO, under aqueous situations it converts to more stable Mg(OH) 2 as there is a thermodynamic drivin g force for the reaction [50] Howev er, it is believed that the microstructure of the surface film is more complex than uniform structure. I t has also been reported that even though Mg(OH) 2 layer grows, an ultra thin MgO layer is still maintained at the metal interface

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28 [51] It has been suggested that the multi layer structure consists of an outer platelet like structure (Mg( OH ) 2 ) and compact layer (MgO) next to the substrate [47] Other environmental and processing factors can also effect the composition and structure of the surface film. Additional gases present in the environment can also interact with the surface film to change its characteri s tics. Carbon dioxide (CO 2 ) present in the atmosphere can react with the Mg(OH) 2 to produce a hydrated carbonate [43] Upon immersion in aqueous solutions, the p resence of additional ions l ike chlorides or fluorides can also combine with the film resulting in formation of hydra ted salts containing respective ions [17] The p resence of alloying elements in the substrate can also have an effect on the composition of surface film formed. Elements present in the substrate also contribute to the sur face film and their oxides and hydroxides can become components of the surface film. The effect of the trace or alloying elements on the surface film depends on the mobility of the ions and their affinity for oxygen. Analysis of the surface films on Mg all oys containing Zn, Al, Y, Mn or Zr have shown the presence o f secondary constituents [46, 52, 53] The ratio of the different ions in the film as compared to substrate depends on the unique affinity of the ions for oxygen, hydroxyl groups and their respective mobility within the substrate and the surface film [17, 54] In one of the studies, it has been reported that if A l content in the substrate exceeds 4 wt %, it can lead to Al concentration of up to 35 wt% in the inner layer of surface film Above a critical Al concentration in substrate, a continuous Al 2 O 3 amorphous structure can be formed in the surface film [17] Despite the ability to form surface films comprising of oxides and hydroxides, Mg does not possess the corrosion resistance comparable to that provided by surface films

PAGE 29

29 of other elements [17] The lack of protection of the surface films is due to various contributing factors. One of the factors that determine the protective nature of the oxide is the Pilling Bedworth (P B) ratio The P B ratio is the ratio of the specific volume of a metal oxide to the specific volume of the metal. When this ratio is less than 1, it leads to porous or cracked film, thereby making it inadequate for protection fo r the substrate. When the P B ratio is higher than 1, the film is relatively passivating as it forms a protective barrier between the air and the substrate [55] Since the P B ratio of MgO /Mg is less than 1 (0.81) it is often mentioned as a reason for the lack of pro tective surface film [55] Under normal dry conditions, the oxide is thin and sufficiently ductile to provide limited protection. However, at higher temperature, the thickness of the oxide increases and it cracks due to small P B ratio [41, 56] On the oth er hand the lack of corrosion protection in aqueous solutions can not be attributed to P B ratio as Mg(OH) 2 has a P B ratio greater than 1. Therefore, there must be another mechanism driving the poor protection of Mg under aqueous solutions. Based on the E pH diagrams discussed in S ection 2.1.1 this can be explained by the electrochemical instability of Mg(OH) 2 in acidic, neutral or mildly basic solutions. Another reason for the lack of corrosion protection has been reported to be the presence of aggressive ion species like Cl In presence of these ions, Mg(OH) 2 is converted to MgCl 2 which has a significantly higher solubility in water than Mg(OH) 2 leadin g to an accelerated corrosion attack on the Mg metal [17] In aqueous environments, the conversion of MgO to Mg(OH) 2 can take place by one of the two potential path ways: (1) hy dration of MgO and (2) dissolution of MgO and deposition of Mg(OH) 2 According to the first approach, the oxide is converted to

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30 hydroxide on contact with water. This results in volume change from 14.02 cm 3 /mol for MgO to 24.33 cm 3 /mol for Mg(OH) 2 [17] This volume expansion is believed to cause the disruption of the surface film, thereby, leading to the porous and non protective nature of Mg(OH) 2 layer in the surface film [15, 57, 58] According to the second approach, exposure to aqueous conditions lead to the dissolution of MgO and the Mg substrate and deposition of Mg(OH) 2 on the surface due to low solubility of Mg(OH) 2 in the solution. The deposited Mg(OH) 2 l acks the compactness and thereby, the ability to protect the substrate from corrosion. Therefore, even if there is a compact MgO la yer on the surface, it can get partially dissolved and be converted to Mg(OH) 2 thereby initiating corrosion. Once the surfac e film breaks down its hard for it to repair itself as Mg(OH) 2 does not necessarily deposit on the areas of broken surface film. Additionally, the evolution of hydrogen from the dissolution of substrate reduces the chances of a compact layer deposition. A ll these factors lead to reduced deposition rate of Mg(OH) 2 as compared to the dissolution rate, thereby making the self inhibition of corrosion difficult [17] 2.1.3 Metallur gical Aspects of Corrosion of Magnesium and its Alloys Corrosion resistance of Mg and its alloys depends greatly on their metallurgy. Metallurgical factors affecting the ir corrosion are impurities, secondary phases and other microstructural features like g rain size and phase distribution [17] Understanding these factors and their influence on the corrosion can provide a vital tool in reducing the overall corrosion rate of exist ing alloys and developing new alloys with superior properties. Galvanic corrosion is an electrochemical process where two components with different electric potentials come in contact with each other in presence of an

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31 electrolyte resulting in the corrosi on of the component that is more electrochemically active. It i s one of the most important corrosion phenomenon of Mg alloys as they have high negative corrosion potential and act as anodes when in contact with more passive components. These components can be two different metals in contact with each other or two different phases present in a microstructure. While the former is known as macro galvanic corrosion, the latter is called m icrogalvanic corrosion [59, 60] Micro galvanic activity occurs as some grains, impurities, intermetallic particles etc. act as anodes while others act as cathodes. It has been generally observed that Mg matrix with lower amount of alloying elements acts as micro anode, and is therefore preferentially corroded. On the other hand impurity particles, secondary phases and intermetallic particles act as micro cathodes [17, 46, 61] The following sub sections discuss the effects of various constituents l ike impurities, secondary phases and other microstructural features on corrosion. 2.1.3.1 Effect of impurities Presence of different elements can decrease, increase or have negligible effect on the corrosion resistance of Mg alloys. The elements that are d etrimental for the corrosion resistance of Mg alloys are generally termed as impurity elements [24] There have been various studies to determine this effect on the corrosion resistance [43, 59, 62] Based on experimental evaluations, iron (Fe), nickel (Ni), cobalt (Co) and copper (Cu) are the most deleter ious imp urity elements for Mg. T hese elements were found to have an adverse effect on the corrosion rate even at concentrations less than 0.2% [17, 62] Subsequent studies have associated the deleterious effect of the Fe, Ni and Cu impu rities on the low solid solubility of these elements in Mg, as they provide active cathodic sites for propagation of corrosion [24, 63, 64] Another important factor

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32 regarding these impurities is the existence of critical conce ntration called tolerance limit [24] At concentrations below this limit, these impurities have minimal effect on corrosion but above it they accelerate corrosion greatly. Song et. al. have summarized that the impurity effect can be due to two main reasons: 1) a ccelerated galvanic corrosion of Mg matrix due to general dissolution of Mg alloy and re precipitation of metallic F e, Ni and Cu on the alloy surface, thereby actively forming new galvanic coupling sites 2) a cceleration of galvanic corrosion when any of the impurity elements Fe, Ni or Cu exceeds its solubility limit in Mg and precipitate as intermetallic phase [65] Increasing purity of the alloy can effectively eliminate this effect and has been shown to decrease the corrosion rate between 10 100 times [63] 2.1.3.2 Effect of solid solution and secondary phases Under the conditions of solid solution, there can be a change in the electrical potential of the Mg matrix depending on the concentration of alloying elements. In concentrated binary and ternary alloys, the concentration of alloying elements can vary subst antially at the center of the grain to the grain boundary and vicinity of second phase particles [17, 66] This c an make the area with higher concentration of alloying element more passive, and hence increase the anodic/cathodic activity and preferential corrosion of the centre of the grain [67] However, solid solution has only a minor effect on the corrosion rate in comparison to the effect of secondary phases as most of the alloying elements have significant effect on corrosion only after formation of secondary phases [24] Almost all the secondary phases in Mg alloys are nobler than the Mg matrix [65, 67 69] This makes them less susceptible to corrosion as compared to Mg matrix. On the other hand, since they are more inert than the matrix, they can act as an effective

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33 cathode, leading to higher corrosion rate of the Mg matrix [17, 70] Hence, the secondary phase is shown to have dual influence on the corrosion of Mg alloys, where it can act as a barrier to corrosion and a cathode at the same time. Which of the two influences dominate, depends on the volume fraction and distribution of secondary phases in the matrix. Presence of fine and continuously distributed secondary phase is more effective at inhibiting the corrosion in the alloy as compared to the p resence of small and discontinuous secondary phase [17] 2.1.3.3 Effect of grain size Even though the reduction in grain size is known to improve strength, ductility and wear resistance, there is lack of fundamental understanding of the effect of the grain size on the corrosion behavior of an alloy [71 73] There is a significant body of literature pertaining to the possible effects of grain size on cor rosion behavior but there is little consensus on the exact relationship between corrosion behavior and grain size. One of the major problems associat ed with understanding the effect of grain size on corrosion resistance of alloys is the difficulty to isolate the effect of grain size from other microstructural changes associated with processes used to obtain different grain sizes [72] However, for Mg alloys, a consistent trend in the literature can be seen that associates reduction in grain size with increase in corrosion resistance in neutral and alkaline so lutions [67, 71, 74, 75] Formation of a better passivating film on the surface of the fine grained alloys is said to be o ne the reasons for superior corrosion protection Since Mg has a P B ratio less than 1, increased grain boundary density might decrease the compressive stresses associated and help to compensate for the oxide/base mismatch [72] This claim has been substantiated by scratch tests [76, 77] and the observation of reduced cathodic kinetics in 0.1 M NaCl [71] Table 2 2 summarizes the

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34 data found in literature regarding the effect of grain size reduction on the corrosion suscep tibility of the alloys. 2.2 Strengthening of Magnesium Alloys This section focuses on the most significant ways by which the strength of crystalline solids can be increased by restricting the dislocation motion. Different obstacles like solute atoms, grain boundaries, precipitates or dispersions can be employed to inhibit the movement of dislocations and help in strengthening material s Most of the high strength materials are hardened by employing one or more of these mechanisms. Various studies have been d one to analyze the effects of grain boundary strengthening [78, 79] solid solution strengthening [8 0, 81] and p article strengthening [82 85] Grain boundaries provide effective barrier to the dislocatio n movement, as the crystallographic factors do not permit the passage of dislocation from one grain to adja cent one through the grain boundary. Hall Petch provided a relationship between the yield point of a material and the grain size as: (2 4) where y is the yield stren 0 is the frictional stress required to move dislocations, k is the Hall Petch coefficient and D is the grain size [86, 87] Researchers have provided v arious theories to explain this relationship [88 90] One of the accepted theory states that dislocations keep piling up at the grain boundary till the resulting stress concentration is sufficient to activate slip systems in the adjacent grains. [88] It has been observed that the degree of grain size hardening is dependent on the

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35 Petch coefficient and the extent of grain size refinement possible in materials [91] Addition of solute atoms can be another way to increase the strength of a material. Solid solution strengthening increases the yield strength of crystalline materials due to the interactions between solute atoms and the dislocation movement. Introduction of solute atom into a crystal produces distortion in the lattice, resulting in a symmetric str ess field around the solute [92] The interaction of these stress fields with the dislocation leads to solute atom dislocation interaction energy. Depending on the size of the solute atom, this interac tion energy can be negative (solute atom size is smaller than solvent) or positive (solute atom is bigger than the solvent atom) [92] Another potent source of strengthening can be the dispersion of second phase particles in the matrix, and is known as particle hardening. Particle strengthening can increase the strength markedly even with small volume fraction of dispersed particles as particles (aggregate of solute atoms) can resis t dislocation penetration to a greater extent than individual solute atoms [93] The amount of strengthening provided by particle strengthening depends on various factors like particle size, particle shape, volume fraction and n ature of interface between particle and the matrix. These particles can affect the movement of dislocations in two different ways. The smaller coherent and/or softer precipitates tend to hinder the dislocation motion by particle shearing, where as larger i ncoherent particles cause dislocation bowing, also known as Orowan looping. The transformation from shearing to bowing takes place when the stress required for dislocation bowing becomes less than that required for shearing [92]

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36 Figure s 2 3 and 2 4 describe the mechanisms of particle shearing and Orowan looping respectively. Even though having more strengthening mechanisms in a system can lead to higher strength, this study will be focused only on the solid solution strengthening due to the concerns about m icrogalvanic effect of precipitates and dispersoids [17, 65] The grain boundary effect will inevitably be present due to different grain refining abilities of different alloying additions, and will be discussed later cha pters. 2.3 Summary Different aspects of corrosion, oxidation and strength of Mg alloys were discussed in this chapter. The thermodynamic driving forces for the corrosion and oxidation were mentioned. Various structure property relationships were also anal yzed and the ones important to this study were identified.

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37 Table 2 1 Chemical p otential of Mg and its compounds in different states at 25C Species Oxidation State State (kcal/mol) Mg 0 Solid 0 Mg + +1 Ion 61 Mg 2+ +2 Ion 109 Mg(OH) 2 +2 Solid 199 MgH 1 Gas +34 MgH 2 1 Solid 8 MgO +2 Solid 136 [Adapted fr om Song, G. 2011. Corrosion of magnesium a lloys (Page 5, Table 1.1). Woodhead Publishing Limited, Cambridge] Table 2 2. Summary of literature related to the effect of changes in grain size on the corrosio n susceptibility of Mg alloys Material Environment Grain Size Range Effect of grain size reduction on Corrosion Susceptibility Processing Route AZ31B 3.5%NaCl Decreases HT, SPD (FSW) Mg Y RE Zr 3.5%NaCl + Mg(OH)2 Decreases HT Mg (99.9%) 0.1 M NaCl Decreases HT, SPD (ECAP, SMAT) Mg (99.9%) 0 .1 M NaCl Decreases HT, SPD (ECAP, SMAT) AZ91D 1 N NaCl Decreases C AZ91D 3.5% NaCl Decreases C AZ31 3.5% NaCl + Mg(OH)2 Decreases/Increases depending on processing SPD (ECAP) WE43 1% NaCl Increases SPD (E, ECAP) Mg (>99.9%) 0.1 M NaCl Decreases SPD (ECAP) (HT heat treatment, SPD severe plastic deformation, FSW friction stir welding, ECAP equal channel angular pressing, SMAT surface mechanical attrition treatment, C casting) [Adapted from Ralston K.D and Birbilis N. 2010 Effect of grain size on corro sion: A review. Corrosion Vol. 66 Issue 7 (Page 075005, Tab le 1)]

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38 Figure 2 1 Electrochemical force series [ R eprinted with permission from John Wiley and Sons 1999. Advanced Engineering Materials Vol. 1 (Pages 11 33, Figure 3] )

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39 Figure 2 2 Schematic illustrations of the main aspects of metal oxygen reactio ns and surface oxide growth The figures show the initial adsorption of oxygen on the surface, followed by nucleation of oxide and its growth and the scale growth by solid state diffusion. It is also shown that cracks and porosity in the oxide scale can lead to direct contact of the air with the substrate, thereby undermining the protective nature of oxide.

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40 Figure 2 3 Schematic diagram of a dislocation cutting through a particle Figure 2 4 Schematic diagram of a dislocation passing between widely spaced

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41 CHAPTER 3 DESIGN APPROACH 3.1 Current Trends in Biodegradable Magnesium Alloy Development A b rief overview of the current trends in the resear ch on biodegradable Mg alloys shows that there has been a significant increase in research and development in this field. However, a more detailed analysis of these trends shows that the majority of the initial research in the area was focused on the behav ior of commercial alloys that were initially designed for non biological applications like automobile and aerospace. A significant amount of work can be found on A l Z n [94 97] and rare e a r th alloying additions [35, 98 100] The alloying elements were p rimarily selected for their ability to increase their mechanical properties as well as their corrosion resistance. Thus at the time of their development, the biocompatibility of these alloying elements were not considered. Al though the preliminary in vitro and in vivo studies have not shown adverse toxicity, previous studies on alloying additions like Al and rare earth elements have shown them to be toxic. It is therefore important to design alloys targeted towards specific applications, keeping in mind all the properties required for the optimal performance in their working environment. This chapter details the design approach utilized to develop new alloys that will be discussed later in this body of work. 3.2 Systems Design Approach A system can be defin ed as a group of components that work coherently to achieve a certain goal and are often described by their complexity, hierarchy and multiple levels of inter related subsystems [101, 102] Based on this classification, materials inherently fall in the category of systems, as they have different components

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42 like matrix, secondary phases and surface films etc. that interact at the subsystem level in a hierarchical fashion t o output a required performance. This systems based approach [103] can b e applied to d esign and optimization of Mg based biodegradable materials. The combination of a systems based design strategy grounded in materials science where processin g structure property relationships dictate and drive an overall performance provides a methodolog y to strategically map the complex mechanistic relationships and understand their evolution from processing to performance [104] Figure 3 1 shows a generalized systems design chart for development of Mg based alloy for biodegradable applications. This chart identifies the performance requirements and then outlines the requisite property objectives required for achieving that performance g oal. The characteristic microstructures that determine the property objectives are outlined in the structure column, which are then connected to the processing steps necessary for obtaining those structures. The p roperty column of the chart outlines the objectives required for the optimal performance of the material. selection [105] can be a useful tool in clarifying these objectives as it helps us to identify the essential characteristics required for the desired application and comparing them with the available material set. This helps us in identify ing the r equirements that need to be fulfilled, and the consequent constraints on the material selection. The primary property requirements for biomaterials would be biocompatibility, corrosion resistance and strength. The systems design chart illustrates these pro perties in a hierarchal order, with the top tier having the highest weighting factor. After outlining the property objectives, the underlying microstructural subsystems that control these properties are

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43 identified. In a system, different subsystems can hav e different and often contradictory effect on the required properties An exhaustive and detailed listing of microstructural subsystems is therefore necessary for proper estimation of the final properties Finally, the processes required for driving the mi crostructural development is listed in a sequential manner. Interlinks connecting the different subsystems identify the potential processing stru cture or structure property relationships. The latter of which a re essentially physics based models that can be supported by simulation or empirical modeling. Lastly, processing structure relationships can be predicted using materials thermodynamics, enabling monitoring of microstructural evolution as a function of a particular processing condition.

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44 Figure 3 1. Systems design chart for developing Mg based biodegradable alloys

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45 CHAPTER 4 D ESIGN AND MICROSTRUCTURAL ANALYSIS OF BINARY MG RE ALLOYS 4 1 Design of Binary Alloys An optimal biodegradable implant materi al should be able to fulfill following essential requirements : ( 1) i t should have sufficient strength to provide required structural support ( 2) i t should corrode slow enough to maintain the structural support and its mechanical integrity while the tissue is healing and ( 3 ) i t should corrode at a great er rate upon completion of tissue healing to prevent complication s arising from long term implant presence in the body. These conditions have contrary microstructural requirements, and cannot be fulfilled using a homogenous alloy. Hence, to simultaneously satisfy all the requirements, the design methodology was focused towards developing a composite structure that has low initial degradation rate to maintain mechanical integrity until the tissue has healed and then have an increased dissolution rate to satisfy the third condition Various me thods of surface passivation have been investigated in literature, including polymer and ceramic coatings, microarc oxidation (MAO) plasma surface modificatio n etc [18, 95, 106, 107] However, these techniques often rely on either sophisticated machinery or additional che mical treatments of the alloys. Furthermore the fidelity of the se surface coatings is precarious especially when coating devices with complex geometries. For this study, we are proposing to use a systems design approach to design an alloy microstructure that can form a passivating oxide on its surface when heated duri ng the alloy processing, thereby eliminating the need for surface coating This overall objective scheme is outlined in Figure 4 1 .The proposed microstructural design has the benefits of eliminating the need of any complex

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46 instrumentation or chemical treat ments, thereby reducing the cost of material processing by designing a material that inherently self passivates with a pre determined scale that can demonstrate in vitro stability 4.1.1 Selection of Alloying Elements 4.1.1.1 Design considerations for surface oxide formation Under normal atmospheric conditions, most of the metals react with their environment to form a thin layer of oxide or hydroxide on the surface. Altering or replacing the porous or non protective surface layer with a protective one c an significantly alter the reactivity of the surface. As discussed in Section 2.1.2 the native oxide on the surface of Mg under dry conditions is MgO, which can convert into a mixture of MgO and Mg(OH) 2 upon coming in contact with water. Since MgO is semi protective, it does not provide sufficient protection against corrosion attack. To overcome the issue of non protective surface oxide, Mg can be alloyed with other elements to achieve selective oxidation of the alloying element and replace MgO with more t hermodynamically stable oxides. The process of selective oxidation of the elements depends on the thermodynamics and kinetic s of the system. Thermodynamically, the oxide with more f has higher thermodynamic stability and will have a strong driv ing force to form preferentially. Additionally, the oxide with its P B ratio closest to bulk material will have enhanced mechanical stability due to reduced mismatch between the oxide volume and the substrate volume. Figure 4 2 plots the f (negative Gib b s free energy of formation) values for the oxides of Mg alloying elements versus their molar volumes per mole metal [39] The horizontal dotted line on the plot shows the energy of formatio n of MgO and the vertical line represents the molar volume of elemental Mg. Oxides to the

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47 left of the vertical line have molar volume less than that of elemental Mg and will lack protection due to the porous nature of the oxide. The oxides below the horizo ntal dotted line are thermodynamically less stable than MgO and hence, will not form in presence of Mg. The arrow in the plot indicates the desired direction for the protective oxide, as the oxides too far to the right of vertical line can be prone to lack of mechanical stability due to compressive stress buildup. Based on these considerations, rare earth elements show promise of forming protective oxides, with Sc 2 O 3 being the most promising among them. On the other hand, Zn, Li Ca and Al can be eliminated as potential oxide forming elements due to the lower free energy of formation of their oxides and/or P B ratio being less than 1 A l though the thermodynamic data provides us information about the stabi lity of the oxides, the mobility of the cation and ox ygen anion through the matrix and oxide dictates rate of oxide formation Unfortunately, there is lack of kinetic data on the mobility and diffusivity of rare earth elements in Mg. To overcome this challenge a systematic detailed experimental plan was created and executed to gain a deeper understanding of the role of solute and oxygen mobility on the growth kinetics of the surface oxide. 4.1.1.2 Design considerations for increasing strength by solution strengthening Since low strength is one of the lim itations of pure Mg, the design process was targeted towards obtaining higher strength of the alloy by introducing potent solute atoms to the Mg matrix. The maximum solid solubility of different alloying additions i n Mg is mentioned in Table 4 1. It can be seen that Y, Sc and Gd show appreciable solid s olubility range. Even though Ce and Mn fit the requirements for oxide formation as mentioned in previous section, they have very limited solubility, and hence have a low

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48 potential for solid solution strengthe ning. Hence, they are not optimal choices for solid solution strengthening. Since different alloying elements can have different amount of strengthening effect, an empirical model was used to predict the strengthening efficiency of different alloying elem ents. For dilute solid solutions, F leischer [1 08] developed a model based on the elastic interaction of screw and edge dislocations with the s olute atoms, and calculated the relationship between solute concentration (c) and increase in resolved shear stress as (4 1) The statistical theory of F leischer was modified by Labusch [109] by using a statistical treatment of solute interactions with dislocations, and calculated the relationship between concentration and shear strength as (4 2) The experimental verification of these theoretical models showed that either model can be used depending on the investigated system [110] However, it was observed that the extrapolation of these models into concentrated alloys gave values lower than the actual stren g th. I n concentrated Mg Al and Mg Zn alloys (>0.1 at%), it was observed empirically that the solid solution effect had a linear relationship with increase in strength [80, 81] Similar model was prepared for different alloying additions and is shown in Figure 4 3. Furthermore, it was proposed by Nembach [111] that the strengthening mechanisms can be superimposed: (4 3)

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49 where ends on the strength of pure Mg ( ), the solid solution strengthening ( ) and the precipitation strengthening ( ). The value of k is derived empirically and typically lies between 1.0 and 2.0. Since we are not aiming for strength contribution from precipit ation, that term of the equation can be eliminated from the model. 4.1.1.3 Design c onsiderations for grain size reduction As described in Section 2.1.3.3 reduction in grain size has been associated with reduction in degradation rate and increase in streng th of Mg alloys [72] One of the most common ways for grain refinement is the addition of foreign nucleants and/or alloying elements [112] However, since this work is focused on solid solution based alloys, the effect of solute on grain size reduction is given more attention than the effect of different nucleating agents. The effect of solute on grain refinement can be explained in terms of growth restriction factor (GRF), which can be calcula ted using binary phase diagrams using the following equation: GR F = (4 4 ) where is the slope of the liquidus line, is the distribution coefficient and is the initial concentration of the element i Lee et. al. calculated the GRF for some of the alloying elements and the values are shown in Table 4 2 A mong the elements selected bas ed on their oxide forming abilit ies Sc had the highest GRF followed by Y [112 ] 4.1.2 Selection of Alloying Compositions Based on their highly stable oxides, large s olid solubility in Mg and grain refining capabilities, rare earths Gd, Y and Sc best fit the design parameters for alloying additions and were chosen for further inv estigation.

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50 After selecting the alloying elements, the next step in the design process was to calculate the nominal compositions of these alloys. There were three main considerations while deciding the alloy compositions: 1) single phase microstructure, 2 ) acceptable corrosion rate of the alloys and 3) minimum amount of alloying addition needed to form a protective selective oxide. In order to simplify the comparison and reduce the variables associated with the experiments, it was decided to have same amou nt of alloying additions in all three systems. First binary phase diagrams were used to investigate the solid solubility of the selected Mg RE systems. Phase diagrams were calculated using PANDAT [113] and its proprietary Mg database PANMAG [114] The Mg rich sections of Mg Gd, Mg Y and Mg S c alloys are shown in Figure 4 4 Mg Gd system has the lowest solid solution solubility at 500C (approximately 8.5 wt%). The temperature was chosen as 500C as it is sufficiently below the eutectic temperature of Mg Gd binary system to prevent any accidental liquification of eutectic due to any possible heating fluctuations in the furnace, and still sufficiently high to allow fast oxidation kinetics [115, 116] This comp osition was used for the first iteration of composition selection Mg 8wt%Gd alloy was prepared to check the possibility of achieving solid solution at this composition. Figure 4 5 (a) shows the as cast images of Mg 8Gd alloy. It can be seen that eutectic i s present in the microstructure of the as cast alloy. The al loy was then heat treated at 500 C for 48 hours and then quenched in water in an attempt to make it a single phase alloy. Figure 4 5 (b) shows the microstructure of Mg 8Gd alloy after this heat tre atment. It can be seen that the second phase precipitates are still present in the microstructure. In another attempt to get rid of these precipitates, an even longer heat treatment was

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51 carried out in an effort to provide sufficient time for dissolu tion of precipitates. Figure 4 5 (c) shows the microstructure even after 15 0 hours of homogenization at 500 C. However, the precipitates are still present in the microstructure. Based on these heat treatments, i t was concluded that the kinetics of homogenization a re too slow to attain thermodynamic equilibrium during the heat treatment processes. It was therefore decided to lower the concentrations of RE additions than the maximum solid solubility shown by the phase diagrams. Second consideration for composition selection was to keep a low degradation rate of the alloys. There has been limited amount of work on degradation properties of Mg Gd and Mg Y that could be used for guidance. Furthermore Mg Sc system lacks any corrosio n data in literature. Liu et. al. investigated the degradation behavior of a series of binary Mg Y alloys ranging from 2 7 wt%Y in 0.1M NaCl solutions and concluded that the degradation rate increases with increase in Y content [117] During electrochemical i mpeda n ce s pectroscopy (EIS) analysis, initial degradation behavior of Mg 2Y and Mg 3Y was similar to high purity Mg and they maintained low corrosion rate up to 24 hours after immersio n. On the other hand, Mg 4w t%Y showed severe degradation after only 5 hours i mmersion, as shown in Figure 4 6 [117] As a result, it was concluded that significant change in corrosion behavior and rate occurs when the Y concentration increases from 3wt% to 4wt% in an alloy. Thirdly, it is important for the alloying additions to be sufficient in concentration for formation of a selective surface oxide. Due to the lack of data on diffusion of rare earth materials in Mg, it was not possible to the oretically calculate the amount of alloying additions needed for formation of protective rare earth oxides on the Mg alloy surface.

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52 However, Wang et. al. investigated the oxidation behavior of four different Mg Y alloy s (Y = 0.82,1.09, 4.31 and 25 wt%). Th ey showed that there is a critical value of Y content above which, a dense and protective oxide can form. This value should be between 1.09 wt% and 4.31wt% as Mg 1.09wt%Y alloy shows linear oxidation behavior associated with non protective oxide and Mg 4.3 1wt%Y shows parabolic oxidation behavior associated with protective oxide formation. Based on the above considerations, it was decided to have 3 wt% alloying additions in the binary alloys as they were predicted to be well within the solid solubility rang e, were below the threshold level at which aggressive corrosion starts [117] and above the threshold level at which protective oxide formation occurs [116] 4.2 Materials and Methods 4.2 .1 Alloy Preparation Mg 3wt%Gd, Mg 3wt%Y and Mg 3wt%Sc alloys were made using pure Mg chips (99.98%) and respec tive Mg RE master alloys made from pure Mg (99.98%), Y (99.9%), Sc (99.9%) and Gd (99.9%) (Sigma Aldrich, St. Louis, MO). To minimize the oxidation during alloy preparation, t he entire process of mixing, melting, and casting was performed under argon atmos phere in a sealed glove box fitted with a resistance heating furnace as shown in Figure 4 7 To prepare the master alloys, t he raw materials were weighed and mixed in graphite crucibles and heated at 825C for 1 hour. The melts were st irred once using graphite rod poured into graphite molds and allowed to cool at room temperature. These master alloys were then used to make dilute binary alloys. The melting and casting process for desired binary alloys was similar to the one used for master alloys, exce pt that the furnace temperature was maintained at 750C. This two step process was used to ensure homogenous mixing and accuracy of

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53 composition. The composition of the alloys was analyzed using inductively coupled plasma (ICP AES, PerkinElmer Optima 3200RL ; Perkin Elmer, Waltham, MA). The nominal and actual compositions of binary Mg R E alloys are listed in Table 4 3 The alloys were then encapsulated under vacuum in quartz tubes, homogenized at 500C for 8 hours and quenched in water. 4.2 .2 Microstructural Characterization The samples for microscopic analyses were cut from the ingot using low speed saw (Allied Techcut 4; Allied Hig h Tech Products, CA ) using colloidal silica. The polished samples were etched using acetic picra l. The microstructure of the samples was investigated using light optic microscopy (LOM, Olympus PME3 ; Olympus Corporation, Tokyo, Japan ) and scanning electron microscopy (SEM, JEOL JSM 6400 ; JEOL Ltd., Tokyo, Japan ). Crystallographic phase identification was performed using X ray diffraction (XRD Philips APD 3720 ; Koninklijke Philips Electronics N.V., Amsterdam, Holland). Optical micrographs were analyzed using the Image J software program and grain size measurements were made according to the linea l inte rcept method outlined in ASTM Standard E112 [118] 4.2 .3 Hardness Testing The samples for hardness testing to eliminate the surface defects The micro hardness testing was performed using 300 gf load on the al loys for 15 s. Seven different measurements were taken for each sample and their average was use d. 4.3 Results and Discussion Figure 4 8 shows the optical micrographs of binary samples etched with acetic picral to reveal grain structure. It can be seen that Mg 3 Y and Mg 3 Gd grains show

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54 directional solidification and have columnar grains as compared to Mg 3 Sc which has finer equiaxial grains. Average grain size of Mg 3 Y Mg 3 Gd and Mg 3 Sc was 1032 322 230 119 6 The large standard deviation in the Mg 3Y and Mg 3Gd alloys is due to their direction solidification and difference in the longitudinal and transverse length s of the grains. Mg 3Sc had more equiaxed grain, and hence small standard deviation. Even though no studies on grain refinement effects of binary Mg Sc alloys has been found in literature, similar grain refining is seen with the addition of Zr to Mg alloys. Effects of Zr additions on the grain refinement of Mg and its alloys are well documented in literature [91, 119] It is widely accepted that grain refinement by Zr takes place by peritectic mechanism where Zr particles first precipitate as Zr rich Mg and promote nucleation of primary Mg grains through peritectic reaction. This mechanism can be identified by the presence of at least one Zr rich core in each grain of the alloy [91, 120] However, recently it has also been shown that addition of Zr lower than the peritectic composition can also lead to grain refinement of Mg alloys [119, 121] Zr has same crystal structure as Mg and similar lattice parameters, thereby making the un dissolved particles also effective nucleation sites for Mg alloy. It is also proposed that for low Zr additions (below peritectic composition), grain refinement is main ly caused by its high growth restriction effect during solidification [112] Peritectic nature o f Mg Sc system, along with its hexagonal crystal structure and its lattice parameter being very similar to Mg, present a compelling reason to believe that the grain refinements in Mg Sc alloys is of similar nature as Mg Zr alloys. However, since the alloy composition was lower than the peritectic composition, and no cores were observed at the centre of grains, its possible that the main contributor to

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55 grain refinement was the growth restriction effect of Sc and not the peritectic mechanism. The SEM images of solution treated Mg 3 Sc Mg 3 Y and Mg 3 Gd are shown in F igure 4 9 The microstructure consists primarily of Mg phase No second phase particles were observed in Mg 3 Y and Mg 3 Sc alloys. A few Gd rich particles were observed in the microstructu re and c an be seen in Figure 4 9 (a). The XRD analysis was performed to get crystallographic information of the alloys and the plots are shown in Figure 4 10 These plots did no t show any peaks associated with second phase in any of the alloys and only showed peak s associated with Mg Vickers micro hardness testing of solution treated alloys was also evaluated and the results are shown in Figure 4 11 Based on the strength model discussed in Section 4.1.1.2 and shown in Figure 4 3, the predicted relationships bet ween the amount of solute and hardness value (HV) are: For Mg Gd system: HV = 3.1 wt% + 29 (4 5 ) For Mg Y system: HV = 4.2 wt% +29 (4 6 ) For Mg Sc system: HV = 1.7 wt% + 29 (4 7 ) The value 29 represents the hardness of pure Mg as can be seen in Figure 4 11. Substituting the amount of solutes present in the alloys under investigation the pre dicted hardness values are 39 41.2 and 34 HV for Mg 3 Gd, Mg 3 Y and Mg 3 Sc alloy respectively. The ex perimental values for these alloys are 44.15.2, 41.73.2 and 42.42.0 respectively. It can be seen that for Mg 3 Gd and Mg 3 Y, the experimental and calculated values are within one standard deviation of the each other The experimentally obtained value for Mg 3 Sc is significantly higher than the predicted

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56 value. This could possibly be due to the higher amount of grain refinement in this alloy as compared to other alloys, thereby having greater grain boundary strengthening contribution in this alloy. Since t he strength model does not include the grain boundary strengtheni ng, it can lead to discrepancy in the predicted and actual strength values. 4.4 Summary A systematic design methodology was proposed to develop an alloy system with the ability to self passi vate. Step by step analysis of property requirements was done and the microstructural features required for achieving those requirements were identified. Three different alloying elements were selected for their ability to show selective oxidation behavior solid solution strengthening and grain refinement in Mg. Strength model to predict the strength of the designed a lloys was developed and then experimentally verified.

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57 Table 4 1. Maximum solid solubility of different alloying elements in Mg Alloying Element Maximum Solubility (wt%) Temperature (C) Aluminum 12.5 437 Scandium 24.6 710 Yttrium 11.4 587.4 Gadolinium 23.49 548 Cerium 0.52 592 Calcium 1.34 516.5 Zinc 6.2 340 Manganese 2.2 650 Table 4 2. Slope of liquidus line (m), Equilibrium distribution coefficient (k) and growth restriction parameter m(k 1) for different alloying elements in Mg Element m k m(k 1) System Ca 12.67 0.06 11.94 Eutectic Zn 6.04 0.12 5.31 Eutectic Al 6.87 0.37 4.32 Eutectic Sc 4.02 1.99 3.96 Peritectic Sr 3.53 0.006 3.51 Eutectic Ce 2.86 0.04 2.74 Eutectic Y 3.40 0.50 1.70 Eutectic [Adapted from Lee Y, Dahle A, StJohn D, 2000. Metallurgical and Materials Transactions A Vol. 31 Pages 2895 906] Table 4 3. Nominal and actual compositions of binary Mg RE alloys Nominal Compositions Actual Compositions Mg 3 wt% Gd Mg 3.2 wt% Gd Mg 3 wt% Y Mg 2.9 wt% Y Mg 3 wt% Sc Mg 2.9 wt% Sc

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58 Figure 4 1. Schematic representation of the desired structure and the relative degradation rates Figure 4 2 The free energy of formation of various oxides versus their volume per mole of metal at 2 5C

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59 Figure 4 3. Solid solution strengthening model showing the relationship between hardness and concentration of solute present in the alloys

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60 Figure 4 4 Phase diagrams of the Mg rich region of Mg Gd, Mg Y and Mg Sc binary alloys

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61 Figure 4 5 Optical micrographs of Mg 8Gd alloy (a ) As cast (b) Homogenized at 500 C fo r 48 hours (c) Homogenized at 50 0C for 150 hours

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62 Figure 4 6 Polarization resistance versus immersion time plot for Mg Y alloys, HP Mg and LP Mg during 24 hour immersion in 0.1 M NaCl [Reprinted with permission from Elsevier, 2010. Corrosion Science Vol. 52 Issue 11 (Pages 3687 3701, Figure 10)]

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63 Fi gure 4 7 Glove box used for melting alloys under inert atmosphere

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64 Figure 4 8 Optical micrographs of etched (a) Mg 3 Gd (b) Mg 3 Y and (c) Mg 3 Sc alloys Figure 4 9 SEM micrographs of homogenized (a) Mg 3 Gd (b) Mg 3 Y and (c) Mg Sc alloys

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65 Figure 4 10 The XRD plots of homogenized Mg 3RE alloys

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66 Figure 4 11 Experimental and calculated m icro hardness values of the binary Mg RE alloys. Hardness value of pure Mg is also shown for comparison

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67 CHAPTER 5 CHARACTERIZATION OF THE OXIDATION BEHAVIOR OF BINARY M G RE ALLOYS 5.1 Thermodynamic Calculations and Prediction of Possible Oxidation Behavior Based on the alloying elements present in the binary alloys, following reactions can take place between the alloying elements and oxygen : (5 1) (5 2) (5 3) (5 4) The Gibbs free energies ( ) of these reactions at oxidation temperature (500C) are 522.0 kJ, 1,675.4 kJ, 1,682.8 kJ and 1,583.9 kJ respectively. This shows that all the oxides have thermodynamic driving force for formation at the experimental conditions. In addit i on to these reactions, rare earth oxides can be formed by reduction of MgO. These reactions can be described as: (5 5) (5 6) (5 7) Under the experimental conditions, Gibbs free energy ( ) can be approximately calculated using available thermodynamic parameters as [39] : (5 8)

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68 where is the standard Gibbs free energy of the reaction at temperature T, is the activity of Mg and is the activity of respective rare earth metal. For simplification, the activities of the Mg, Gd, Y and Sc were replaced by their molar atomic concentrations and the activity of oxygen was rep laced by its volume fraction, which is equal to 1 as pure oxygen was used for oxidation Hence, the activity values for different constituents of the selected systems are: For Mg 3 Gd and ; For Mg 3 Y and ; For Mg 3 Sc and ; Substituting the values in Equation 5 8, we get (5 9) (5 10) (5 11) These calculations reveal that even though Y and Sc will reduce the MgO and form their respective oxides, whereas Gd will not be able to reduce MgO. Based on these calculations, it can be predic ted that binary Mg Gd alloy will have an oxide layer co nstituting both MgO and Gd 2 O 3 whereas Mg Sc and Mg Y alloy will have oxide layer s constituting predominantly of Sc 2 O 3 and Y 2 O 3 respectively. 5.2 Materials and Methods T he samples for oxidation analysi s were cut from the as cast ingots of Mg 3 Gd Mg 3 Y and Mg 3 Sc They were then encapsulated in P yrex tubes under vacuum and solution treated at 500C for 8 hours. Rectangular samples with nominal dimensions of 9.5 mm 9.25 mm 1.3 mm were cut from the so lution treated ingots. A hole was

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69 drilled through the samples to facilitate the use of quartz hooks for hanging them in the thermo gravimetric analyzer (T GA, Setsys Evolution TGA DTA/DSC SETARAM Inc., Hillsborough, NJ) This method was selected over using alumina pans for holding the sample because it reduces the buoyancy and/or drag from flow of O 2 through the specimen chamber. The samples were polished successively on emery papers of grit size 320, 600, 1200 and 4000, then cleaned and degreased ultrasoni cally in ethanol. To minimize the amount of oxide formed on the surface during room temperature exposure, the polished samples were placed in TGA immediately after degreasing in ethanol. The samples were inserted into the chamber at room temperature and he ated to 500C at 20C/min under vacuum. The ramping up of temperature was carried out under vacuum to prevent any oxidation during the ramp up. Once the temperature reached 500C, pure oxygen was introduced into the chamber at the rate of 196 ml/min. The s amples were held isothermally at 500C for up to 24 hours. The weight gain and temperature information was collected approximately every 17 seconds. Analyses on the weight gain data obtained were performed to determine the kinetic parameters. The dimensio ns of the samples were assumed to be constant and change in the dimensions of the samples before and after the oxidation were ignored. Since the initial O 2 flow rate changes cause fluctuations in buoyancy, only the data after the stabilization of the O 2 fl ow rate was considered and any weight gain during the stabilization period was neglected. The oxides on the samples were then characterized using XRD SEM auger electron spectroscopy (AES, Perkin Elmer PHI 660; Perkin Elmer, Waltham, MA) and x ray photoel ectron spectroscopy (XPS, Perkin Elmer PHI 5100 ESCA; Perkin Elmer,

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70 Waltham, MA). SEM analysis was performed to analyze the surface characteristics of the oxides. XRD analysis was done rate of 0.8/min for phase identification AES was then used to obtain the depth profile s of the surface oxides. The operating voltage and current of the instrument was kept at 10 kV and 100 nA A 3 kV argon beam was applied for sputtering and the sputter rate was estimated to be approximately 3.75 nm/min. Further, chemical analyses of the surface oxides were conducted using XPS. The instrument was operated at 15kV and ray source and a 4 kV argon beam was us ed for sputtering. 5.3 Results and Discussion 5.3.1 Oxidation Kinetics of Binary Mg RE Alloys Long term oxidation studies were performed on Mg 3 Gd Mg 3 Y and Mg 3 Sc alloys at 500C in pure O 2 atmosphere. The weight gain versus time plot of the alloys is sh own in Figure 5 1. The o xidation data show s that Mg 3 Gd has the fastest oxidation rate and Mg 3 Sc has the lowest oxidation rate. All the alloys show a high initial oxidation rate that reduces with time. A s mall amount of spalling was observed in Mg 3 Y and Mg 3 Sc that shows up as jumps/kinks in the weight gain versus time plots. As a result, only the data before spalling was used for calculation of oxidation kinetics. Based on the TGA data, all the alloys show parabolic oxidation behavior in the long run. Ho wever, after looking at different regions of the TGA curve, it is observed that the oxidation behavior of Mg 3 Gd and Mg 3 Y comprises of 3 different regions: (A) fast initial oxidation (linear oxidation), (B) transition period, and (C) second oxidation regi me following a parabolic oxidation mechanism. This behavior is illustrated in Figure 5 2,

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71 which shows the t ( time ) for Mg 3 Gd T he general rate law for oxidation kinetics can be given by (5 12) where k is the rate co nstant, t is the time, n is the time exponent of the rate law and c is a constant. For linear kinetics, the value of n is unity and for parabolic kinetics, the value of n is 0.5. However, it has been shown by various researchers that metals and alloys show oxidation behaviors that deviate from the traditional linear, parabolic or logarithmic rate laws [123 126] Often times, a combination of these laws can be seen in the oxidation behavior of a single material over a long period of oxidation. It has been observed that a steady state parabolic oxidation is often preceded by a faster oxidat ion due to poor protective nature of initial oxide formed on the surface [45, 124] oxidation of Mg 3 Gd and Mg 3 Y is 1. 5 and 1.8 respectively. Pierag gi [45] suggests that when t he initial oxide growth does not contribute to the steady state parabolic oxidation, 1/2 plot for evaluation of the parabolic rate law as 2 versus t plots as it helps in determination of true parabolic constant k p 1/2 for the alloys are shown in Figure 5 3 and the values for the rate constants of the alloys are mentioned in Table 5 1. It can be seen that the fit of linear regression for the curves lies between 0.946 and 0.998. I t is also observed that the oxidation rate of Mg 3 Sc is 16 times lower than Mg 3 Gd and 4 times lower than Mg 3 Y showing that Mg 3 Sc forms a compact and protective oxide However, this also implies that Mg 3 Sc forms a thinnest oxide scale whereas Mg 3 Gd fo rms the thickest scale.

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72 5.3.2 Characterization of Surface Oxides on Mg RE Alloys Figure 5 4 shows the XRD spectra of Mg 3 Sc Mg 3 Y and Mg 3 Gd alloys isothermally oxidized at 500C. It shows presence of Sc 2 O 3 Y 2 O 3 and Gd 2 O 3 in Mg 3 Sc Mg 3 Y a nd Mg 3 Gd alloys respectively. Additionally, p Mg are observed in all the alloys. However, o nly Mg 3 Gd shows the peaks for MgO and no such peaks were observed in Mg 3 Y and Mg 3 Sc To further characterize the oxide layers, AES was used to get depth profiles of the oxide layers. The AES depth profiles of alloys after 5 hours of oxidation are shown in Figure 5 5 and the thickness of the oxide layers is 3 Gd, Mg 3 Y and Mg 3 Sc alloys respectively. This fact implies hig he st oxidation rate of Mg 3 Gd and lowest oxidation rate of Mg 3 Sc which is in agreement with the TGA results presented in Figure 5 1. It is also observed that oxide profile of Mg 3 Gd consists mainly of Mg and O with minor gadolinium contribution. This result further indicates the presence of MgO and Gd 2 O 3 in the oxide layer of Mg 3 Gd, which is also in agreement with the XRD analysis shown in Figure 5 4. To visualize the structure of the oxide, the cross section of the oxidized samples was investigated u sing electron microscopy. Figure 5 6 shows the cross sectional micrograph of Mg 3 Gd. It can be seen that the oxide layer is non uniform and consists of two different oxides brighter Gd rich oxide and darker Mg rich oxide. On the other hand Figure 5 7 an d Figure 5 8 show that the oxide layer on the Mg 3Y and Mg 3Sc was uniform along the surface. Additionally, the AES depth profile of these alloys is different from Mg 3Gd profile, as it consists of two different regions, with a thin outer region comprising of a mixture of Mg and RE oxide Figure 5 5) and a thick RE rich oxide region 5) Though the

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73 total thickness of the oxide layer varies on all the alloys, the thickness of outer layer of Mg rich oxide is approx imately 20 30 nm in these alloys. Similar results have also been observed in other Mg RE containing alloys [115, 116] To verify if the region A of the oxide layer was the initially formed native oxide, AES analysis was perfo rmed on polished alloys that were kept in air at room temperature Figure 5 9 shows the AES profile of polished Mg 3 Sc alloy. It can be seen approximately 30 nm thick Mg rich native oxide is present on the surface. To account for the lack of MgO peak in t he XRD spectra and better understand the outermost layer of the oxides, XPS analysis was performed. All the energy positions were first adjusted by comparing the binding energy of C 1s peak to that of the standard binding energy of 2 84.8 eV, to compensate for the charging effect in the XPS measurements. F igure 5 10 shows the XPS spectra of the oxidation layer in the binding range of 0 1250 eV for all three alloys oxidized at 500C for 5 hours. It can be seen that the peaks of Mg, O, C and Ar exist in the sp ectra of all the alloys Mg KLL Auger peaks are also visible in the region of 300 380 eV and correspond to the Auger signal of kinetic energies in the region of 1106 1186 eV [127] Similarly, oxygen KVV Auger peaks around 980eV correspond to the kinetic energy of 506.7eV [128] Figure 5 10 (a), (b) and (c) contain XPS peaks of Gd, Y and Sc in their respective spectra High resolution XPS spectra for Mg 1s, O 1s and Gd 3 d are shown in F igure 5 11 The Mg 1S peak shown can be attributed to the MgO and fits well with the standard [129] Gd 4d peak for shown in Figure 5 10 is in agreement with the standard energy peak for Gd 2 O 3 showing that the Gd present in the outer layer is in the form of Gd 2 O 3 [130] Oxygen 1s is composed of two components with peaks at 529.6 and 531.7 eV.

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74 The binding energy of 529.6 eV co rresponds with the peak for MgO [131] while the other one can be attributed to Gd 2 O 3 [132] The XPS spectrum for Mg 3 Y is shown in Figure 5 12 The Mg 1s peak is similar to that in Mg 3 Gd, and corresponds to MgO [129] Y 3d peak can be fitted with two sets of doublets, corresponding to Y 3d 3/2 and Y 3d 5/2 spin orbit splitting. The peak at 156.8 eV corresponds well with the standard value of Y 3d 5/2 peak of Y 2 O 3 [133] whil e the peak at 158.8 eV is in agreement with the Y 3d 3/2 peak of Y 2 O 3 [134] This confirms the presence of yttrium in the form of oxide in the outer layer. T he O 1s peaks at 5 29.8 and 531.7 further confirm the presence of MgO [131] and Y 2 O 3 in th e outer layer [135] H igh resolution images of Mg 1s a nd O 1s are shown in Figur e 5 13 The Mg 1s peak is similar to ones seen in other alloys and corresponds to the MgO [129] The O 1s peak can be deconvoluted into two peaks. The peak at 531.5 fits with the standard MgO peak [136] while the other one at 529.6 corresponds to Sc 2 O 3 [137] The presence of Sc 2 O 3 in the outer layer is also supported by the presence of Sc 2p peak at 400.84 eV, which corresponds with the standard Sc 2 O 3 peak [137] 5.4 Summary The oxidation behavior of the Mg 3 Gd, Mg 3 Y and Mg 3 Sc alloy was calculated using thermodynamics a nd validated by experimentation. It was shown that the thermodynamic calculations were able to correctly predict the composition and nature of the surface oxide scales on these alloys. Oxidation kinetics were investigated using TGA and indicated that the alloys followed parabolic oxidation behavior, thereby f orming protective oxide scales The rate constants for oxidation were also determined and showed that Mg 3 Sc formed the thinnest and most protective oxide whereas Mg 3 Gd formed the thickest and least protective oxide. The oxides were characterized by

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75 diffe rent characterization methods namely SEM, XRD, AES and XPS. It was observed that Mg 3 Gd formed an oxide scale consisting of a mixture of MgO and Gd 2 O 3 whereas Mg 3 Y and Mg 3 Sc had their oxide scales consisting exclusively of Y 2 O 3 and Sc 2 O 3 respectively

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76 Table 5 1. Parabolic rate constants for different alloys Alloy Temperature Rate constant (mg/cm 4 .s) Mg 3Gd 500C 1.32 10 7 Mg 3Y 500C 3.57 10 8 Mg 3Sc 500C 7.71 10 9

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77 Figure 5 1. The isothermal TGA plots of Mg 3 Gd Mg 3 Y and Mg 3 Sc binary alloys at 500C Figure 5 2. Kinetic parameters for Mg 3 Gd for the 3 different oxidation regimes (A) linear oxidation (b) transition region and (c) parabolic oxidation

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78 Figure 5 3. 1/2 for the alloys (a) Mg 3 Gd (b) Mg 3 Y and (c) Mg 3 Sc showing long term parabolic oxidation behavior at 500C

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79 Figure 5 4 The XRD plots of the alloys after isothermal oxidation at 500C for 24 hours (a) Mg 3Sc, (b) Mg 3Y and (c) Mg 3Gd

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80 Figure 5 5 AES depth profile s of Mg 3Gd Mg 3Y and Mg 3Sc oxidized isothermally at 500C for 5 hours

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81 Figure 5 6. SEM micrographs of the cross section of oxidized Mg 3Gd alloy

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82 Figure 5 7. SEM micrographs of the cross section of oxidized Mg 3Y alloy

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83 Figure 5 8. SEM micrographs of the cross se ction of oxidized Mg 3Sc alloy Figure 5 9. AES depth profile of polished Mg 3Sc alloy

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84 Figure 5 10 XPS survey scan of entire binding energy for (a) Mg 3Gd (b) Mg 3Y and (c) Mg 3Sc Figure 5 1 1 XPS spectra of Mg 3Gd oxidized for 24 hours

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85 Figure 5 12 XPS spectra of Mg 3Y oxidized for 24 hours

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86 Figure 5 13 XPS spectra of Mg 3Sc oxidized for 24 hours

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87 CHAPTER 6 IN VITRO DEGRADATION OF ALLOYS UNDER DIFFERE NT SURFACE CONDITIONS 6 .1 Materials and Methods Magnesium degrades in physiologica l solutions according to the following reaction: (6 1) The measurement of hydrogen evolved can be conveniently used to calculate the degradation rate of Mg alloys and has been extensively used in literature [68, 138, 139] The samples for degradation studies were prepared from solution treated ingots of Mg 3Y Mg 3Gd and Mg 3Sc alloys by cutting r ectangular pieces of nominal dimensions 9mm 9mm 2mm using a low speed diamond saw. They were then polished up to 4000 grit surface finish using SiC emery papers. The polished samples were cleaned and degre ased ultrasonically in ethanol. To analyze the effect of oxidation on the degradation behavior of samples, some of them were oxidized in pure O 2 at 500C for 24 hours. The apparatus used for immersion testing of both oxidized and non oxidized al loy samples is shown in Figure 6 1. The immersion tests were carried out at 37C in 2 .2H 2 O, 0.40 g/l KCl, 0.06 g/l KH 2 PO 4 0.10 g/l MgCl 2 .6H 2 O, 0.10 g/l MgSO 4 .7H2O, 8.00 g/l NaCl, 0.35 g/l NaHCO 3 0.48 g/l Na 2 HPO 4 1.0 g/l D Glucose (Thermo Scientific Inc., Waltham, MA) The ratio of surface area to the volume of the solution was held approximately 150 ml/cm 2 The amount of hydrogen evolved and the change in pH was measured at 24 h our intervals. Due to the difference in the degradation rate and behavior of the alloys, degradation experiments were stopped at different times for each alloy The two parameters used for ending the degradation experiments were:

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88 1) if the amount of hydrogen released in 24 hours was more than the volume of the graduated cylinder used for measuring hydrogen evolved; 2) if there is visible disintegration of the pieces of sample due to corrosion, as it would affect the surface area. After taking out the samples from n, they were washed with ethanol and then dried in air. Extreme care was taken to not remove any of the corrosion products formed on the surface of the samples. The corroded samples were characterized using XRD and SEM (JOEL 6335F FEG JEOL Ltd., Tokyo, Ja pan ) 20 70 at the rate of 0.8 /min. T he samples for SEM analysis were carbon coated prior to analysis to reduce the charging of deposited corrosion products. 6.2 Results and Discussion Figure 6 2 shows the hydrogen evolution curves for the alloys in their oxidized and non oxidized states obtained from immersion tests Among non oxidized samples, Mg 3Gd has the highest average degradation rate ml/cm 2 /day) and Mg 3Sc h ad the lowest degradation rate ml/cm 2 /day) after 48 hours Hydrogen evolution measurement of Mg 3Gd was stopped after 48 hours, as the amount of hydrogen being evolved by the samples was more than the volume of the graduated cylinder used for measur ement. For Mg 3Y and Mg 3Sc it was seen that the rate of degradation increased with time. At the conclusion of the degradation analysis after 11 days, Mg 3Y 2 /day while Mg 3Sc had an average 2 /day. The corrosion propagation was directly observed as a function of time for all the alloys. Figures 6 3, 6 4 and 6 5 show the optical micrographs of corrosion propagation on polished surfaces of the alloys. Upon imm

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89 at various points all over the surface for these alloys. All of the alloys showed some areas of uniform corrosion, pitting corrosion and filiform corrosion attack. Though the corrosion mechanisms involved looked similar, the aggressiveness and the propagation rate of each of the corrosion processes varied Analysis of corrosion propagation with respect to time also supported the degradation analysis done using the hydrogen evolution studie s. Mg 3Sc had the slowes t corrosion propagation and showed only a small amount of pitting corrosion after 6 hours, with most of its surface being corrosion free. Mg 3Y had a little more corrosion attack after 6 hours, with some filiform and pitting corrosion occurring on the surf ace On the other hand Mg 3Gd showed the highest rate of corrosion and significant amount of pitting and filiform corrosion was seen on the surface. A l arge amount of degradation products were formed on its surface due to rapid degradation, which can be s een as cloudy grey areas in the optical images. It is also observed that the filiform corrosion predominantly starts from the edges of the samples and moves inwards, and can be seen in Figure 6 5. Similar observation was also made in other two alloys. SEM analysis of the corroded surfaces further illustrated the corrosion mechanisms in the alloys. The micrographs of the alloys after 96 hours of immersion are shown in Figure 6 6, Figure 6 7 and Figure 6 8. It is observed that different types of corrosion are occurring simultaneously on all the alloy surfaces. Figure 6 6 shows the surface of Mg 3Sc It can be seen that there are cracks in the surface layer (Figure 6 6b), presumably due to the volume change associated with conversion of MgO to Mg(OH) 2 [140] Figures 6 6 (c,d) show the progression and morphologies of filiform corrosion. Figure 6 7 (b) shows the pitting corrosion in Mg 3Y It was interesting to note

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90 that a large area had surface crack s oriented in a single direction, showing the possible weakness of oxide /hydroxide along certain crystallographic planes as shown in Figure 6 7(c). Figure 6 7(d) shows the tail of filiform corrosion with arrows marking the direction of propagation. It can be observed that the filiform corrosion in Mg 3Y is much more severe than in Mg 3Sc Mg 3Gd had the most severe degradation and the surface of the samples was s everely corroded after 96 hours Figure 6 8 shows that its surface is f ull of deep cracks and valleys and a large amount of deposition of corrosion products can be observed. The corrosion products were also identified using XRD. The XRD plots of the alloys after 96 hours of immersion are shown in Figure 6 9 It can be seen that the XRD spectra show the presence of Mg(OH) 2 on all the samples. To understand the corrosion mechanisms, it is essential to note that a spontaneous oxide/hydroxide film is formed on the surface of Mg and its alloys on exposure to air (native oxide) [54, 141] which has also been illustrated by XPS and AES analysis in previous chapter. The surface film is not necessarily compact and part of the substrate can b e easily exposed to the solution. The a ddition of different alloying elements affects the properties and composition of this film. This is one of the reason s why different alloying elements have different effect on the reduction in degradation rate, as the y can get incorporated into the surface film and change its protective properties. For Mg Y alloys, it is known that Y can incorporate into the surface film in the form of Y 2 O 3 and also as Y(OH) 3 in aqueous condition [98, 142] Similarly, Gd 2 O 3 has been found to be incorporated in surface film of Gd containing Mg alloys [115] and was also seen in the outer layer in Mg 3Gd described in previous chapter. Du e to the presence of these semi protective surface films, t he degradation reactions on Mg mainly

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91 occur at the bare parts of the substrate according to the following reactions [17, 143, 144] : Cathodic reaction: (6 2) Anodic reaction: (6 3) Chemical reaction NDE: (6 4) Overall reaction: (6 5) Corrosion product formation: (6 6) For Mg 3Gd Mg 3Y and Mg 3Sc following extra reactions can also occur during the degradation: Anodic reactions: (6 7) Corrosion product formation: (6 8) Addition of Cl ions f urther aggravates the problem of localized degradation as they are smaller than hydroxide ions and can penetrate the surface film [17, 140, 145] This can lead to the breakdown of it s protective nature, leading to a more active corrosion on local sites, and can lead to pitting corrosion [117] The formation of thicker oxide layers on the surface of the alloys can prevent this penetration of Cl ions by acting as a protective barrier. This is a possible explanation for the reduction in the degradation rate of oxidized samples as compared to the non oxidized samples in the immersion tests As shown in the section 5.2.2 of the previous chapter, Mg 3Y

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92 developed a compact and thick oxide layer on its surface, and showed the highest reduction in the degradation rate. On the other hand, Mg 3Sc had the thinnest oxide scale on the surface, and as a result showed the least amount of reduction in the degrad ation during the immersion testing. It is worthy to note that even though Mg 3Gd had the thickest oxide scale on the surface, it was non uniform and bi phase in nature. As a result, the protective nature of the surface oxide in Mg 3Gd was not as efficient as that of Mg 3Y For most of the metals, filiform corrosion is usually associated with coated surfaces and is driven by the oxygen reduction reaction [146] However, filiform corrosion has been observed on uncoated Mg and various other Mg alloys in NaCl containing soluti ons [147] It s presence indicates that there is a protective film on the surface of Mg which is in agreement with our previous results. The classical method of filiform corrosion involving oxygen reduction reaction i s not applicable in case of Mg and its alloys, as the cathodic reaction in these alloys is the generation of hydrogen [67, 77] The propagation of filiform corrosion in Mg alloys is not yet completely understood and different mechanisms have been proposed to explain its behavior in Mg and its alloys One of the proposed mechanisms is based on the acidification of filament tip due to the formation of hydrogen during degradation. According to this mechanism, Cl ca n penetrate the native oxide film on the surface due to its small size and attack the Mg substrate to form localized corrosion. This corrosion process leads to the release of H + ions due to hydrolysis of Mg and other alloying elements present. This leads t o the reduction of pH and acidification of the corrosion pits. These active sites act as the

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93 filament tips and progress along the material, followed by a tail covered with degradation products [1 40] An alternative mechanism attributes the filiform corrosion to the alkalization processes occurring at the bottom of the filiform attack. It assumes that during the initial localized corrosion, Mg(OH) 2 corrosion products are formed due to local pH i ncrease. Additionally, the transformation of Mg to Mg(OH) 2 almost doubles the volume, resulting in filling of space in the localized cell. However, due to non uniformity of the corrosion process, an active corrosion location is always present. As a result, a small active site is balanced with a large area of Mg(OH) 2 which acts as a cathode due to its high pH and passivates the surface. The active corrosion site always maintains a high concentration of Mg 2+ thereby attracting more Cl ions and keeps the lo cal site active [117] In the present study, t he initiation of the filiform corrosion from the sample edges indicates that the surface layer was less protective at the edges and it was easier for the Cl ions to penet rate it. This reduction in protectiveness can be a result of sudden change in the crystal orientation of the grains at the edges, which can lead to non uniformity or porosity in the oxide thereby acting as preferential sites for filiform corrosion initiat ion Another important observation during the corrosion propagation investigation was that some of the grains were less susceptible to corrosion as compared to their neighbor ing grains, as shown in Figure 6 10 One of the possible reasons for this preferen tial corrosion can be the crystallographic orientation of the grains. Since Mg has an anisotropic crystal structure, it has different atomic density along different crystallographic planes. For Mg, basal plane (0001) has the highest atomic density, followe d by plane and then the plane [148, 149] Since a close packed

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94 plane has a higher atomic coordination, it has higher binding energy and lower su rface energy. The surface energies (E s ) of (0001), and surfaces have been calculated to be 1.808, 2.156 and 1.868 eV/nm 2 respectively [150] It is known that the electrochemica l dissolution rate of a metal can be related to the activation energy (Q) required for a metallic ion to escape from the metal lattice and dissolve into the solution, and can be expressed as: (6 9) where n is the number of electrons involved in the electrochemical reaction, k is a s the gas constant, T is the absolute temperature and E is the electrode potential, respectively. Q is associated with the surface energy (E s ) according to the relation : (6 10) where Q o rates of different planes is dependent on their surface energy [149] This can explain the preferential corrosion of some grains as compared to their neighbors. 6.3 Summary The degradation behavior of the binary alloys was investigated in polished and oxidized conditions. Under polished conditions, Mg 3Sc has the lowest degradation rate, whereas Mg 3Gd has the highest degradation rate. Oxide scales led to varied amounts of re duction in the degradation rate. Mg 3Y was found to have the most protective oxide scale It was observed that the corrosion propagation takes place by filiform and pitting corrosion. Th e initi ation sites for filiform corrosion were predominantly

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95 located at the sample edges. Additionally, the preferential degradation due to grain orientation was observed.

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96 Figure 6 1. Experimental setup for measuring hydrogen evolution of samples Figure 6 2. Hydrogen evolution rate of different oxidized and non oxidized alloys

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97 Figure 6 3 Optical micrographs of corrosion propagation in Mg 3Sc allo ys (A) p itting corrosion ( B) f i liform corrosion Figure 6 4 Optical micrographs of corrosion propagation in Mg 3Y allo ys (A) pitting corrosion (B) f i liform c orrosion

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98 Figure 6 5 Optical micrographs of corrosion propaga tion in Mg 3Gd alloys (A) pitting c orrosion (B) f il iform c orrosion Figure 6 6 SEM micrographs of Mg 3Sc alloy after 96 hours of degradation

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99 Figure 6 7. SEM micrographs of Mg 3Y alloy s urface after 96 hours of degradation

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100 Figure 6 8 The SEM micrographs of Mg 3Gd surface after 96 hours of dissolution

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101 Figure 6 9. XRD plots of Mg RE binary alloys after 96 hours of immersion

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102 Figure 6 10 Optical micrograph showing the selective corro sion of grains in Mg 3Y alloy

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103 CHAPTER 7 INVESTIGATION OF THE OXIDATION AND DISSOL UTION BEHAVIOR OF TERNARY MG SC Y ALLOY 7.1 Design of Ternary Alloy Evaluatio n of the binary systems provided useful information that can be used in deciding the next step in the design process It was observed that no single alloying addition fulfilled all the desired design parameters. For example, although Gd was able to provide t he thickest surface oxide scale, it was not homogenous. Sc was able to provide a fine grain size and low bulk degradation rate, but only produce d a thin oxide layer that did not provide sufficient additional protection. On the other hand, Y additions were able to provide a compact and thick passivating oxide scale, but were not able to refine grains appreciably. Hence, as a next step to the design process, this chapter evaluates the design of a ternary alloy that incorporates the beneficial properties of tw o different alloying additions to achieve the design performance requirements. Based on the previous chapters, Sc and Y were chosen as alloying additions. Gd was not selected because it was not able to provide homogenous surface oxide or grain refinement. Both Y and Sc were able to produce homogenous oxide scales, and Sc showed significant grain refinement. To further narrow down the range of Sc compositions experimental investigation of its effect on grain size and degradation rate was conducted. Four Mg Sc binary alloys with compositions of Mg X Sc ( X =0.5,1,2,3 ) were investigated. Figure 7 1 shows the reductions in grain size and degradation rate with increasing Sc additions The a ddition of Sc to Mg reduces both the grain size and degradation rate. The i ncrease in Sc concentration from 0.5 to 1 wt% shows the largest reduction in corrosion rate. Further additions show minimal reductions in degradation rate. On the other hand, increasing Sc concentration from 0.5 to 1 wt% shows no

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104 change in grain size, but further increase leads to reduction in grain size. Based on these observations, 3 wt% Sc additions were chosen again as it showed the best properties out of all of the investigated alloys. The concentration of Y was kept at 3 wt% according to the design s t rategy described in Chapter 4. T hermodynamic calculations to evaluate the oxidation behavior of the ternary system were performed as described in Section 5 1 Since no ternary interactions have been observed in the literature [115, 151] all the calculations were based on binary systems The alloying elements present in the alloy can lead to the following oxidation reactions: (7 1) (7 2 ) (7 3 ) The Gibbs free energies ( ) of these reactions at the oxidation t emperature (500C) are 522.0, 1,675.4, 1,682.8 and 1,583.9 kJ respectively. Additionally, it is known that the oxides with higher thermodynamic stability are able to reduce the oxides with l ow er stability. This leads to the possibility of following reactions: (7 4 ) (7 5 ) Based on the relationship between the Gibbs free energy of these reactions, temperature and the activities of the constituents as described by equation below, (7 6 ) the values are:

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105 (7 7 ) (7 8 ) Hence, it can be seen that both Sc 2 O 3 and Y 2 O 3 have the thermodynamic driving force for simultaneous formation in the oxide scale. Also, the experimental data from the binary alloys was added to the strength model shown in Section 4.1.1.2, to recalculate the relationship between amount of alloying additions and the resulting hardness increase. For the solution treated ternary alloy, the total h ardness would be: (7 9 ) where HV is the total hardness of the alloy, is the hardness of pure Mg, and are the solid solution strengthening contributions of Y and Sc res pectively. Based on this model, the hardness value of the designed alloy was pre dicted to be 47 HV. Additionally, an empirical relationship was derived between hardness and yield strength by using the data from Gao et al [152, 153] A power law function was fit through the data as shown in Figure 7 2 and is represented as (7 10 ) y is the yield strength and HV is the Vickers hardness. Based on the predicted hardness value, the yield strength of the alloys was predicted to be 91 MPa. 7.2 Materials and M ethods Mg 3Sc 3Y alloy was prepared from pure Mg chips (99.98%) and the Mg RE master alloys used for making binary alloys in previous chapters. The method and apparatus for melting and casting was also similar to the binary alloy s. The composition of the alloy was analyzed using ICP AES. The alloy ingot were encapsulated under

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106 vacuum in quartz tubes, homogenized at 500C for 8 hours and quenched in water. The colloidal s ilica. The samples for grain size measurement were etched using acetic picral. L inea l intercept method outlined in ASTM Standard E112 [118] was used for grain size measurements. The microstructure characterization of the samples was investigated using optic al microscopy, SEM and XRD. For TGA analysis, r ectangular samples were cut from the ingots ground to 4000 grit and cleaned in ethanol. The clean samples were then oxidized in high purity oxygen in a tube furnace and weighed after 0.5, 1, 2, 3, 4, 5, 10, 15, 20 and 25 hours. Analyses on the weight gain data obtained were performed to determine the kinetic parameters. The dimensions of the samples were assumed to be constant, thus the dimensional changes before and after the oxidation were ignored. The oxides on the samples were then characterized using XRD, SEM AES and XPS SEM analysis was performed to analyze the surface characteristics of the oxides. 0.8/min for phase identifi cation. AES was then used to obtain the depth profiles of the surface oxides. The operating voltage and current of the instrument was kept at 10 kV and 100 nA A 3 kV argon beam was applied for sputtering and the sputter rate was estimated to be approximat ely 3.75 nm/min. Further, chemical analyses of the surface oxides were conducted using XPS. The instrument was operated at 15 kV and 200 mA. ray source and a 4 kV argon beam was used for sputte ring.

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107 Vickers microhardness testing was performed using 300 gf force f or 15 seconds. Compression testing was performed at a strain rate of 6% min 1 using Instron 5582 universal testing machine. 7.3 Results and Discussion ICP AES analysis showed the act ual composition of the alloy to be Mg 2.8Sc 3Y. Figure 7 3 shows the optical images of solution treated Mg 3Sc 3Y alloy. It can be seen that the microstructure has fine, equiaxed grains with an average grain size of 109 The XRD analysis of t his alloy is sh own in F igure 7 4 It can be seen that XRD spectra Mg only, and no peaks related to secondary phases were found. However, the SEM analysis of the microstructure did reveal some second phase p articles To identify the constituents of this phase, EDS was done on the samples. Figure 7 5 shows the SEM micrograph of the microstructure along with the EDS plot and elemental maps. It can be seen that the secon d phase particles present are Y rich binary phase. Based on the elemental analysis, binary phase diagram and the morphology of the particles, they can be identified as Mg 24 Y 5 phase particles. It can also be seen that Sc is evenly distribute d through out the matrix and no Sc rich secondary phase is pres ent. This also means that there is a higher amount of Sc present in the solid solution as compared to Y, as some of the Y is stored in the second phase precipitates. Figure 7 6 shows the thermo gravimetric analysis of the alloy at 500C The oxidation behavior of the alloy is parabolic in nature, with a higher initial rate of oxidation associated with the adsorp tion of oxygen and initial s cale formation, followed by slow oxide growth [44] The parabolic behavior of the oxi de growth as a function of time growth can be fitted to the parabolic equation as:

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108 (7 11 ) p is the parabolic rate constant, t is the oxidation time and c is the constant [44] time 1/2 was used to calculate the rate constant and i s shown in Figure 7 7 The parabolic rate constant for Mg 3Sc 3Y at 500C is 5.5710 8 mg/cm 2 /s and is close to that of binary Mg 3Y (3.5710 8 ) and about an order of magnitude higher than Mg 3Sc (7.7110 9 ). Figure 7 8 shows the SEM micro graphs o f the oxide cross section and shows a thick and compact surface oxide. XRD analysis showed presence of both Sc 2 O 3 and Y 2 O 3 and is shown in Figure 7 9 To further understand the relative positions of binary oxides in the scale, AES depth profiling was done on the a lloys and is shown in Figure 7 10 It can be seen that the oxide structure of the oxide is similar to the two region profile described in secti on 5.2.3, with a thin outer Mg rich oxide layer and a thick inner RE oxide rich layer It can be seen that both Sc and Y are present throughout the depth of the oxide scale. This is in agreement with the results of thermodynamic calculations described in S ection 7. 2. It is also observed that t he concentration of Sc in the oxide is higher than that of Y This shows that the oxide layer is has higher amount of Sc 2 O 3 than Y 2 O 3 This can be possibly due to higher Sc content in the solid solution as some of the Y was locked in second phase Mg 24 Y 5 particles (Figure 7 5 ) XPS analysis was also performed on the samples to characterize the phases present in the outer layer of the oxide. All the energy positions were first adjusted by comparing the binding energy pea k in the experimental spectra to that of the standard binding energy of 248.8 eV to account for the charging effect in the XPS measurements. The XPS spectrum of the oxide over the entire range of 0 1250 eV is shown in Figure 7

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109 11 It confirms the presence of Mg, O, Sc, and Y in the outermost layer. In addition to the XPS peaks, Mg KLL and O KLL Auger peaks are present in the spectrum in the regions of 300 380 eV and 1106 1186 eV [127] High resolution XPS spectral peaks for constituent e lements are shown in Figure 7 12 The positions of the peaks are similar to those found in the binary alloys. The position of Mg 1s peak is 1304.4 eV, which corresponds to the standard Mg 1s pe ak for MgO. Y 3d peak can be split into a set of doublet peaks, corresponding to 3d 3/2 and 3d 5/2 spin orbital splitting and the peak at 157.7 aligns with the Y3d 5/2 peak [130] This confirms the presence of Y in the outer layer of the oxide as Y 2 O 3 The Sc 2p peak was observed at 401.8 eV, which corresponds with Sc 2 O 3 peak reported in the literature [154] Similarly, the deconvoluted peaks of O 1s correspond well with the peaks for MgO and Sc 2 O 3 [136, 137] Degradation studies were performed on the oxidized and non oxidized to evaluate the surface passivation properties of the oxide. The hydrogen evolution curves of the alloy s urface with different treatments are shown in F igure 7 13 It can be seen that polished sample has a corrosion rate of approximately 1.1 ml/cm 2 /day, which is comparable to that of the binary Mg 3Sc and lower than the rate of both binary Mg 3Gd and Mg 3Y al loys It was observed that oxide layer formed by oxidation at 500C for 5 hours led to the initial reduction in the degradation rate to 0.16 ml/cm 2 /day. This rate was maintained up to 7 days, after which the degradation rate increased to that comparable of polished sample. This shows that the oxide layer provides protection from the degradation and once the oxide layer is corroded, the degradation rate jumps back to that of the bulk material. To see the effect of oxide thickness samples oxidized

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110 able to reduce the degradation rate to 0.01 ml/cm 2 /day for up to 23 days. This rate is lower than most of t he commercially available Mg alloys and is comparable to that of high purity Mg [17, 68] As sho wn in Fig ure 7 14 the XRD analysis of the corroded surfaces revealed the presence of Mg(OH) 2 on the surface. For non oxidized surface, uniform corrosion with localized regions of pitting and filiform cor rosion was observed. Figure 7 15 shows two different regions of the same sample. Figure 7 15 (a) shows the area around a pit, which has higher amount of corrosion cracks and degradation pro ducts as compared to Figure 7 15 (b). On the other hand, the SEM analysis of the surface of oxidized Mg 3Sc 3Y alloy samp les after immersion shows no corrosion cracks or any pitting corrosion (Figure 7 16) The darker regions in the micrographs show that the formation of hydroxide on the surface has st arted, but is not yet complete. The microhardness evaluation of the sampl es was performed to validate the strength model prediction. The measured value of Vickers hardness of the alloy was 48.53.9 HV. This calculated value was within half a standard deviation of the experimental value, showing a good predictability of the mode l. The experimental yield strength of the alloys was found to be 100 3.5 MPa. The experimental value of yield strength is higher than the predicted value, because it was based on predicted hardness values, which are lower than the experimental values. The comparison of the predicted and experimental hardness and strength values is shown in Figure 7 17. 7.4 Summary Ternary Mg 3Sc 3Y alloy was designed based on thermodynamic calculations and the data obtained from the analysis of binary alloy systems. The ox idation kinetics were evaluated using TGA and the oxidation rate constant was determined. The

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111 characterization of the oxide showed that the oxide layer consisted of Sc 2 O 3 and Y 2 O 3 which was in agreement with the thermodynamic calculations. It was also seen that the surface oxide scale was passivating in nature, and that increasing the oxidation time increased the passivating behavior of the oxide. It was shown that 24 hours of oxidation at 500C was enough to passivate the surface for 23 days, thereby r educing the degradation rate to less than 0.01 ml/cm 2 /day of hydrogen evolved during this period Finally, the microhardness and compression testing showed that t he modified strength model was able to predict the hardness and yield strength values with app reciable accuracy

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112 Figure 7 1 Effect of Sc additions on the grain size and the rate of hydrogen evolution in binary Mg Sc alloys

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113 Figure 7 2. Yield strength vs hardness data f itted to a power law to derive E quation 7 10 Figur e 7 3 Optical m icrograph of solution treated Mg 3Sc 3Y

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114 Figure 7 4 XRD plot of Mg 3Sc 3Y alloy homogenized at 500 C for 8 hours

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115 Figure 7 5 SEM micrographs showing elemental mapping and EDX plot of solution treated Mg 3Sc 3Y alloy

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116 Figure 7 6 Plot showing the weight gain per cm 2 vs time Figure 7 7 Plot showing the calculation of parabolic rate constant by plotting mass gain per cm 2 vs square root of time

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117 Figure 7 8 SEM micrographs of the cross section of Mg 3Sc 3Y oxidized for 5 hours

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118 Figure 7 9 XRD plot of Mg 3Sc 3Y alloy oxidized at 500C for 5 hours Figure 7 10 AES depth profile of Mg 3Sc 3Y alloy oxidized for 5 hours at 500C

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119 Figure 7 11 XPS survey of Mg 3Sc 3Y alloy oxidized for 5 hours

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120 Figure 7 12 XPS multiplex peaks for Mg, O, Sc and Y, showing the presence of MgO, Sc 2 O 3 and Y 2 O 3 in the outermost layer of oxide scale of Mg 3Sc 3Y alloy

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121 Figure 7 13 Hydrogen evolution behavior of Mg 3Sc 3Y alloy under differ ent surface conditions

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122 Figure 7 14 XRD plot of Mg 3Sc 3y (polished)

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123 Figure 7 15 SEM micrographs showing the corroded surface of no n oxidized Mg 3Sc 3Y alloy (a) area of severe pitting corrosion (b) area of relatively low corrosion r ate on the same sample

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124 Figure 7 16 SEM micrographs of the corroded surface of 24 hour oxidized Mg 3Sc 3Y after 96 hour s

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125 Figure 7 17. The comparison between the predicted and experimental values of hardness and yield strength of Mg 3Sc 3Y alloy

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126 CHAPTER 8 CONCLUSIONS Mg alloys have the potential to be a viable alternative to the traditional metallic implants for temporary structural implant applications. However, some of the limitations like high degradation rate and low strength need to be overcome before they can play a competitive role as a biodegradable implant material. The focus of this research was to use a systematic thermodynamics based approach to design and develop new alloys with a potential for controlled degradation by self passivation. A detailed analysis of property requirements for the design goal was described and the required microstructural features were identified. Based on the design parameters, the oxide stability of the alloying elements and soli d solution solubility were chosen to be the most important requirements for their selection. From the available pool of alloying elements, Gd, Sc and Y were selected due to the high thermodynamic stability of their oxides >1.0 P B ratio high solid soluti on solubility and grain refining abilities Additions of 3.0 wt% were chosen for binary alloys based on their respective phase diagrams and the effect of concentration on the oxidation and corrosion behavior. Additionally, hardness was used as a measure fo r strength and a model was developed to predict the strength of the alloys. Microstructural analysis of the alloys showed that the processing steps were able to obtain the desired microstructure, with some minor deviations in Mg 3Gd alloy. It was also sho wn that the predicted values of hardness were within one standard deviation of experimental values for Mg 3Gd and Mg 3Y alloys. The values obtained were used to refine the model. Then based on the thermodynamic calculations, the oxidation behavior of the a lloys was predicted. Additionally, insight into the kinetic behavior of the

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127 oxide species was obtained using TGA to get a complete understanding of the oxidation process in these alloys. The oxidation behavior was analyzed and the oxides were characterized. It was found that Mg 3Gd alloy has the fastest oxidation kinetics and the thickest oxide scale, whereas Mg 3Sc had the slowest kinetics and the thinnest oxide scale The SEM, AES and XP S analysis revealed that Mg 3Gd developed a non homogenous oxide scale that comprised of MgO and Gd 2 O 3 On the other hand, Mg 3Sc and Mg 3Y developed mostly homogenous oxide scales consisting of their respective RE 2 O 3 with a thin outer layer rich in MgO. The oxidation behavior of the alloys was in agreement with the thermodynamic calculations. In the next step, the degradation behavior of these alloys was investigated and the effects of different alloying additions as well as surface oxides were investigat ed. It was seen that Gd addition led to the highest degradation rate whereas Sc addition led to the slowest degra dation rate. However, due to its low thickness, Sc 2 O 3 did not provide any passivation. On the other hand, Y had a thick and uniform oxide layer which showed the best passivation behavior among the three alloys. It was also observed that the corrosion propagation took place by similar mechanisms in all the alloys and differed only in the rate at which these reactions progressed. Finally, all the information gathered from the binary alloys was used to design a ternary alloy with a mixed surface oxide consisting of Sc 2 O 3 and Y 2 O 3 Degradation analysis of oxidized and non oxidized ternary alloy showed that the oxide scale was protective and 24 hour o xidation was able to passivate the surface for up to 25 days, after which the corrosion rate changed to match that of non oxidized sample.

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128 CHAPTER 9 FUTURE WORK The current study used a systems design approach to select an alloy system predict the struc ture property relationships and validate the m with experimental analysis. It was observed that selective ox idation of alloying additions to Mg can be a potent way to control the initial degradation rate. Even though this study was able to fulfill its desig n goals, further work is needed to optimize the composition and the processing parameters to achieve better properties To improve the yield and ultimate strength of the alloy, additional mechanisms like precipitation strengthening should be added to the design approach and their effect on the strength should be investigated. A n empirical model similar to the solid solution strengthening model presented in this work should be made and used to predict the relationship between alloying additions and precipit ation strengthening contribution. However, presence of precipitates would also affect the degradation behavior of the alloy due to microgalvanic corrosion. Therefore, further studies are needed to find a balance between higher strength and low degradation rate. Additional ly, we saw a relationship between the oxidation time /oxide thickness and the surface passivity T he kinetic and degradation data obtained in this study can be combined with results from additional oxidation studies at different temperature s to understand the relationships between time, temperature, oxide thickness and the duration of surface passivity. These empirical relationships can help us control the passivity of surface oxide and predict the length of initial passivity of the material Finally in order to for this work to reach its ultimate design goal of producing a completely biodegradable and biocompatible alloy with controlled and predictable

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129 degradation, in vivo analysis will be essential. They would aid in correlating the in vit ro and in vivo results, thereby providing us with better insight into the material behavior in its actual working environment as it i s not possible to completely replicate the in vivo conditions in vitro. The in vivo data for hydrogen evolution, toxicity of the ions and the mechanical stability of the alloys can then be used in next iterations of the design to improve the performance.

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138 BIOGRAPHICAL SKETCH Harpreet Singh Brar was born in 1984 in Bathinda, India. After completing his secondary school education in Bathinda, he moved to Chandigarh where he completed his senior secondary education from D.A.V. College. He earned his B.E in metallurgical and materials engineering from Punjab Engineering College (PEC) C handigarh. In addition to academics, he also pursue d several of his other interests like soccer and basketball. He then joined the Materials Science and Engineering Department at the University of Florida to pursue graduate studies. Here, he was introduced to design and development of magnesium alloys and w ent on to pursue a Ph.D in this field.