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Processing, mechanical behavior and biocompatibility of ultrafine grained zirconium fabricated by accumulative roll bonding

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Processing, mechanical behavior and biocompatibility of ultrafine grained zirconium fabricated by accumulative roll bonding

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Content PROCESSING, MECHANICAL BEHAVIOR AND BIOCOMPATIBILITY OF ULTRAFINE GRAINED ZIRCONIUM FABRICATED BY ACCUMULATIVE ROLL BONDING by Ling Jiang ___ ________ A Dissertation Presented to the FACULTY OF THE GRADUATE SCHOOL UNIVERSITY OF SOUTHERN CALIFORNIA In Partial Fulfillment of the Requirements for the Degree DOCTOR OF PHILOSOPHY (MECHANICAL ENGINEERING) December 2008 Copyright 2008 Ling Jiang ii ACKNOWLEDGEMENTS It is a pleasure to give my thanks to my academic advisor Professor M E Kassner for his continued guidance, critical advice and financial support during the entire period of my doctorate study. I am grateful to him for sending me through three different countries (Spain, Italy and the United States) to complete my dissertation. While studying for my doctorate degree I had wonderful opportunities to experience and be involved in different cultures. I am also grateful to him for offering vacation time when I was able to visit my family in China and show my love to them. Especially, I sincerely appreciate a long vacation in the past summer of 2008 granted by him, which allowed me to get married, an important step forward in my life. I would like to take this opportunity to show my sincere thanks to Professor M T Perez-Prado, my co-advisor from CENIM, Madrid, Spain. This thesis would not have been possible without her help, support and friendship. Prof. Perez-Prado not only invited me to study in CENIM by using the facilities in her laboratory but also guided me in many scientific discussions. Her optimistic spirit gave me much encouragement when I was stuck on my research during these years. Her assistance for my life made my stay in Spain much easier. I want to thank Professor J L Gonzalez-Carrasco at CENIM, Spain for the collaboration work on the biocompatibility study. I also want to thank Professor O A Ruano and Dr. J A Del Valle for our discussions. I want to show my sincere gratitude towards Professor M Cabibbo at Marche Polytechnic University, Ancona, Italy for his help in conducting experiments of measuring grain boundary misorientations in TEM. iii I am very grateful to all of my friends who made my daily life easier through my doctorate study. I am especially thankful to my colleagues in CENIM, Spain: Javier Corrochano, Paloma Hidalgo Manrique, Hassan Abdollah Pour and Carmen Estevan Pastor at CENIM, Spain. Their kindness and help made me have a good time during my stay in Spain. I also want to thank Jianhua Chen, Kevin Zhu, Peter Geantil and Erik Tolmachoff at USC, without whom California sunshine would not have been enjoyable. Finally, I would like to thank my parents and my parents-in-law and my brother and sister for their continuous encouragement, support and inspiration. Especially thank my mother for her dedication to our family with patience and love, which let me remain focused on my academic assignment. Also, this thesis would have taken longer if my newly married wife had not supported me morally and materially during the last several years. Thank you to all the people who have helped me. iv TABLE OF CONTENTS ACKNOWLEDGEMENTS ii LIST OF TABLES vii LIST OF FIGURES viii ABSTRACT xviii CHAPTER 1. INTRODUCTION 1 1.1 Nanostructured Materials 1 1.2 Processing of Nanostructured Materials 1 1.3 Severe Plastic Deformation 2 1.3.1 Equal Channel Angular Pressing (ECAP) 2 1.3.2 High Pressure Torsion (HPT) 4 1.3.3 Accumulative Roll Bonding (ARB) 5 1.4 Mechanical Properties of Ultrafine Grained Materials 16 1.4.1 Yield Stress 16 1.4.2 Ductility 20 1.4.3 Strain Hardening 22 1.4.4 Strain-rate Sensitivity 23 1.5 Texture 26 1.5.1 Introduction 26 1.5.2 Representation of Texture 26 1.5.3 Texture Measurements 31 1.6 Recrystallization 41 1.6.1 Discontinuous Dynamic Recrystallization (DRX) 41 1.6.2 Geometric Dynamic Recrystallization (GDX) 43 1.6.3 Continuous Dynamic Recrystallization (CDX) 45 1.6.4 Geometric Necessary Boundaries (GNBs) 45 1.7 Zirconium 52 1.7.1 Crystallographic Properties 52 1.7.2 Deformation Mechanisms 52 1.7.3 Recrystallization Mechanisms 55 1.7.4 Texture in Zr and Zr Alloys 56 CHAPTER 2. EXPERIMENTAL PROCEDURE 58 2.1 Accumulative Roll Bonding of Pure Zr 58 2.1.1 Selection of the Severe Plastic Deformation to Be Used 58 2.1.2 Material Selection 58 2.1.3 ARB Process 59 2.1.4 Microstructure Characterization 62 2.1.5 Mechanical Testing 68 v 2.2 Biocompatibility of Zr Processed by Large Strain Rolling 73 2.2.1 Fabrication and Characterization of Materials 74 2.2.2 Cell Culture 77 2.2.3 Attachment Assays 77 2.2.4 Cell Spreading 78 2.2.5 Immunofluorescence Assays 78 2.2.6 Fibronectin Levels 79 2.2.7 Cell Viability 79 2.2.8 Alkaline Phosphatase Activity and Mineralized Nodule Formation 80 2.2.9 Statistical Analysis 80 CHAPTER 3. RESULTS 81 3.1 Accumulative Roll Bonding of Pure Zr 81 3.1.1 Initial Material 81 3.1.2 Annealed Material 83 3.1.3 Microstructure Evolution at the Sheet Mid-thickness during ARB 90 3.1.4 Through-thickness Homogeneity in the ARBed Samples 100 3.1.5 Mechanical Behavior of the ARBed Samples 104 3.1.6 Effect of Rolling Temperature and Thickness-reduction per Pass 107 3.1.7 Assessment of the Bonding Quality 108 3.2 Biocompatibility of Zr Processed by Large Strain Rolling 112 3.2.1 Material Characterization 112 3.2.2 Cell Attachment 117 3.2.3 Cell Spreading and Arrangement of Tubulin and Actin Cytoskeletons 118 3.2.4 Fibronectin Analysis 121 3.2.5 Cell Viability 124 3.2.6 ALP Activity and Degree of Mineralization 125 CHAPTER 4. DISCUSSION 127 4.1 Effectiveness of ARB as a Grain Refinement Technique for Zr 127 4.2 Texture Homogeneity 130 4.3 Comparison with Other HCP Materials 132 4.4 Bonding Strength of ARBed Samples 133 4.5 Biocompatibility of Zr Processed by Large Strain Rolling 136 CHAPTER 5. CONCLUSIONS 141 REFERENCES 143 ALPHABETIZED BIBLIOGRAPHY 153 vi APPENDIXES 163 Appendix I 163 Appendix II 171 Appendix III 173 vii LIST OF TABLES Table 1: Mechanical properties of various materials processed by several cycles of accumulative roll bonding 15 Table 2: The crystallography of Zirconium 52 Table 3: Dimensions of the tensile specimens 69 Table 4: Dimensions of the specimens for bonding strength testing 71 Table 5: Composition of the Zr material under study (ppm) 81 Table 6: Mechanical property data of annealed and ARB Zr 106 Table 7: Bond strength for three specimens in group #1 111 Table 8: Grain size and hardness values for the different materials 114 Table 9: R q and R a roughness values at length scale of 30 μm for the different materials 115 Table A1: The explanation on the functions of each window in the block diagram of LJInstron program. 164 Table A2: Technical data of the rolling machine used in the present study. 171 Table A3: Dimensions of the specimens and corresponding strain and strain rate for each rolling pass during ARB process. 172 viii LIST OF FIGURES Fig 1: Schematic illustration of Equal Channel Angular Pressing (ECAP) [1]. 3 Fig 2: The four fundamental processing routes in ECAP [20]. 4 Fig 3: Fig 3: Schematic illustration of a High Pressure Torsion (HPT) die [1]. 5 Fig 4: Schematic representation of the Accumulative Roll- Bonding (ARB) process [17]. 6 Fig 5: TEM microstructure of the 1100 Al alloy processed by 6 cycles of ARB (ε = 4.8). The microstructure was observed along the transverse direction (TD). The crystallographic orientation was determined by Kikuchi-line analysis in TEM. The boundary misorientation is superimposed in the map in degrees [22]. 8 Fig 6: Mean grain thickness and fraction of high angle boundaries through the thickness of the UFG IF steel processed by 7 cycles of ARB (ε = 5.6) at 500 o C [22]. 9 Fig 7: TEM micrographs showing the microstructure of commercial purity Ti processed by ARB using 4 cycles (ε = 3.2) at room temperature (a). High magnification micrographs of the lamella (b) and equiaxed structure (c), respectively [34]. 11 Fig 8: Fig 8: TEM micrographs of commercial purity Ti ARB-processed by (a) 6 cycles (ε = 4.8) and (b) 8 cycles (ε = 6.4) [34]. 12 Fig 9: True tensile strength of several ultrafine grained materials processed by ARB and comparison with the strength of the starting materials [16, 17, 33, 34]. 13 Fig 10: Strength and elongation of (a) IF steel ARB-processed by various cycles at 500 o C [17] and (b) commercial purity Ti ARB-processed by various cycles at room temperature [34]. In (b), broken lines indicated typical tensile strength and total elongation for a conventional Ti-6Al-4V alloy. 14 ix Fig 11: Inverse Hall-Petch trend for Cu and Pd as shown by Chokshi et al. (from hardness measurements) [41]. 18 Fig 12: Yield stress versus grain size plot for Cu. Data complied from various sources [1]. 18 Fig 13: Positive Hall-Petch slope for Cu with higher values for compressive (from hardness measurements) than for tensile strengths [45]. 19 Fig 14: Variation of the yield stress with grain size for Ti. Comparison to the linear Hall-Petch response [1]. 20 Fig 15: Compilation of yield strength versus % elongation of various ultrafine grained metals [13]. 22 Fig 16: Strain rate sensitivity plot for Cu as a function of grain size [53, 54]. 24 Fig 17: Tensile stress-strain curves for ball milled (3 h) Zn tested at 20, 40 and 200 o C at a constant strain rate of 10 -4 s -1 [14]. 25 Fig 18: Strain-rate jump tests (compression) performed on ball milled (3 h) Zn at 20 and 60 o C [14]. 25 Fig 19: Reference sphere, where the origin, ‘north pole’ and ‘south pole’ labelled O, N, S respectively. A crystal is imagined to lie at O. The normal of a plane in the crystal intersects the reference sphere at P. P is projected from the south pole onto the equatorial plane at p; this is the stereographic projection of P [58]. 28 Fig 20: Construction of pole figures. (a) projection of <001> directions on the reference sphere; (b) projection from south pole of <001> directions on the equatorial plane; (c) repeat the process with other crystals with similar orientations; (d) and (e) representing the density of poles through contours of intensity. 29 x Fig 21: Stereographic projection on a wolfnet, where the two points located near the center are separated by 20 o and the two points located at the perimeter of the projection are also separated by 20 o . It should be noted that while the angular separation distance is the same in both cases, the physical distances on the projection are not [59]. 30 Fig 22: Schematic showing the effect of sample rotation on the arrangement of the lattice planes. (a) Untilted position (α=β=0 o ); (b) sample tilted such that the lattice planes satisfy the Bragg condition (α > 0 o , β > 0 o ) [58]. 32 Fig 23: Schematic representation of all electron scattering localized at a single point in (a). In (b) some of the diffusely scattered electrons undergo Bragg diffraction from certain atom planes. In (c) Kossel cones are formed and projected onto the Ewald sphere, creating parabolas which approximate to straight Kikuchi lines [62]. 34 Fig 24: A schematic drawing showing the experimental devices for EBSD. 36 Fig 25: EBSD system used for determining microtexture in this work (OIM TM , TexSEM Laboratories, Inc., Provo, Utah). 37 Fig 26: Determination of orientations. (a) Kikuchi patterns obtained in SEM (b) detecting the Kikuchi lines with greater intensity and (c) allocation of indexes. 38 Fig 27: Methodology used for determining individual orientation and studying the nature of grain boundaries present in materials. 40 Fig 28: Schematic representation of the strain hardening versus stress for dynamically recrystallized (ABC) and dynamically recovered (ABD) material [68]. 43 Fig 29: Transmission electron micrograph and schematic representation of geometric necessary boundaries (GNBs) and incidental dislocation boundaries (IDBs) in pure Ni, rolled at a reduction of 20% [96]. RD represents the rolling direction. 46 xi Fig 30: Schematic three dimensional drawing of the GNBs (solid planar boundaries) and IDBs (short grey boundaries) bridging between them. The misorientation (θ) and the boundary spacing (D) are illustrated for each type of boundary [97, 98]. Bold lines represent high angle boundaries (HABs) and thin lines represent low angle boundaries (LABs). 47 Fig 31: Power law relationship between the average misorientation angle and applied von Mises strain for high purity cold rolled aluminum and nickel. The average angle increases more rapidly with strain for GNBs than for IDBs [96]. 49 Fig 32: Average misorientation angles of both IDBs and GNBs versus strain in high-purity aluminium (99.996%) [89]. 49 Fig 33: Misorientation angles across IDBs and GNBs as a function of strain in aluminium alloy AA5005 [100]. 50 Fig 34: The c/a axial ratio for various hcp metals. 53 Fig 35: Slip systems in α-zirconium. 54 Fig 36: Most common rolling and annealing textures in Zr and Zr alloys [121,123-130]. (a) Rolling texture: c-axis are tilted 25−40 o from the ND toward the TD, along the ND-TD plane, and > < − 0 1 10 directions are aligned with the RD. This texture is retained when annealing at T < 600 o C. (b) Annealing texture of a rolled sheet (approx. 95% cold rolling). Annealing treatments at T ≈ 600 o C for approximately 1 h lead to a rotation of the crystallites around the c-axis so that a > < − 0 2 11 direction becomes parallel to the RD. More severe annealing (higher T, longer time) leads to the gradual rotation of the c-axis toward the ND. 57 Fig 37: Schematic of the processing route designed to fabricate bulk ultrafine grained Zr. 61 Fig 38: The two-high rolling mill used for accumulative roll bonding. 62 xii Fig 39: The schematic drawing for measuring the misorietation angle of grain boundaries in TEM. 65 Fig 40: X-ray Diffractometer available at the National Center for Metals Research (Madrid, Spain). 67 Fig 41: Instron tension machine used in this study (Centro Nacional de Investigaciones Metalúrgicas, Madrid, Spain). 68 Fig 42: Schematic illustration of the dogbone flat tensile specimens. 69 Fig 43: Schematic diagram of the test sample to measure the shear strength of the bond. 70 Fig 44: Specimens used to measure the strength of the bond. (a) Front view and (b) Side view. 72 Fig 45: Gripping system (a) and set-up (b) utilized to measure the strength of the bond. 72 Fig 46: Plate of pure zirconium. 81 Fig 47: Optical micrographs showing the microstructure of the as-received Zr. 82 Fig 48: Grain size distribution in the initial material. Twin boundaries are included. 83 Fig 49: Macrotexture of the initial material. 83 Fig 50: EBSD map viewed along the ND (a) and along the RD (b) showing the microstructure after the as-received material was annealed at 600 o C for 10 minutes. Thick lines correspond to boundaries with misorienations higher than 15°, while the thinner lines correspond to boundaries with misorientations between 4° and 15° (boundaries less than 4 o can not be reliably detected here). 85 xiii Fig 51: EBSD map viewed along the ND (a) and along the RD (b) showing the microstructure of the as-received Zr after annealing at 600 o C for 40 minutes. Thick lines correspond to boundaries with misorienations higher than 15°. 86 Fig 52: Microstructure of the initial material, annealed at 600 o C for 40 minutes, examined by optical microscopy. 87 Fig 53: Grain size distribution of the as-received material without counting twins (a) and after annealing at 600 o C for 10 min. (b) and for 40 min. (c). 88 Fig 54: Direct pole figures corresponding to rolled Zr after annealing at 600 o C for 10 minutes (a) and 40 minutes (b). 89 Fig 55: Macrotexture of the initial material annealed for 40 min. examined by X-ray diffraction. 90 Fig 56: EBSD maps showing the microstructure of pure Zr at different stages of the process and the corresponding ND and RD inverse pole figures. Mapping was performed in areas close to the mid-thickness of a layer at the RD-ND plane. (a) After one pass at room temperature with 25% reduction (ε = 0.33); (b) ARB1 (ε = 1.97, T = 327 o C); (c) ARB2 (ε = 3.57, T = 327 o C); (d) ARB3 (ε = 5.17, T = 327 o C). 94 Fig 57: X-ray pole figures corresponding to Zr following (a) 1, (b) 2, and (c) 3 cycles of ARB. 95 Fig 58: (a) Evolution of grain and (sub)grain size with strain. The lines (dotted and dashed) indicate the expected reduction in grain length and thickness, respectively, with strain, due to geometrical thinning alone. (b) Evolution of the grain aspect ratio along the RD-ND plane. 96 Fig 59: TEM micrographs of Zr after (a) ARB1 (ε = 1.97), (b) ARB2 (ε = 3.57) and (c) ARB3 (ε = 5.17). The plane perpendicular to the transverse direction and all areas are close to the mid-thickness of the rolled sheet. 97 xiv Fig 60: TEM micrographs of Zr after (a) ARB1 (ε = 1.97), (b) ARB2 (ε = 3.57) and (c) ARB3 (ε = 5.17). The plane is parallel to the rolling plane and all areas are close to the mid-thickness. 98 Fig 61: (a) Fraction of high angle boundaries (HABs) with increasing strain. Note that only boundaries with θ > 2 o are observed by EBSD. (b) Frequency distribution of boundaries at lower misorientation (2 o < θ < 15 o ) for different processing steps. 99 Fig 62: Microstructure and texture at the surface of the Zr sheet rolled at room temperature using a reduction of 25%. (a) EBSD map along the rolling plane. (b) ND, RD, and TD inverse pole figures. 101 Fig 63: Through-thickness texture gradient in the ARB2 sample (ε = 3.57). Inverse pole figures illustrating the texture at several locations: (a) surface; (b) mid- thickness of the first layer; (c) interface formed during ARB1 (ε = 1.97); (d) mid-thickness of an inner layer; (e) interface formed during ARB2. 103 Fig 64: Microstructure and texture of pure Zr in the vicinity of an interface (close to mid-thickness) after the ARB3 step. Measurements performed by EBSD. 104 Fig 65: Stress-strain curves for annealed and rolled (ARB1 and ARB3) Zr. The tensile tests were performed at room temperature and 10 -3 s -1 . Dashed curves represent engineering stress-strain data and solid curves represent true stress-true strain data, respectively. Beyond the UTS, the true stress-true strain values were extrapolated until the true strain to failure. The latter was calculated taking into account the final cross-section of the specimen. 105 Fig 66: TEM micrographs showing the microstructure of pure Zr rolled using (a) one rolling pass with 75% thickness reduction at 300 o C; (b) one rolling pass with 75% thickness reduction at 600 o C; (c) multiple rolling passes with 5% thickness reduction, summing up to a total reduction of 75% at 300 o C. 108 xv Fig 67: Optical micrographs showing the layered structure of pure Zr after 1, 2, and 3 cycles of accumulative roll bonding. The observation plane is perpendicular to the transverse direction (TD). 110 Fig 68: Fracture of two specimens (a) as a result of shear test (specimen #1) and (b) as a result of tensile test (specimen #4). 111 Fig 69: BEI (A,B) and TEM (C) micrographs revealing the microstructure of Ti64 (A), Zr (B) and UFG Zr (C). Grain size distribution corresponding to UFG Zr is illustrated in D 113 Fig 70: AFM topographical profile of zirconium surfaces. Representative profile (A), and profiles containing the highest (B) and lowest (C) wavelength components, on the Zr sample. Representative profile (D), profile containing the highest (E) wavelength component, (F) complementary profile to (E), and a subsequent decomposition of (F) into a high (G) and low (H) wavelength components, on the UFG Zr sample. 116 Fig 71: Cell attachment on zirconium surfaces. (A) Saos-2 cells were cultured on Ti64 ( ), Zr ( ) and UFG Zr ( ) for 0.5, 3 and 24 h. The results are expressed as the percentage of the absorbance measured on Ti64 at 0.5 h, which was given the arbitrary value of 100. Each data represents the mean ± SD of four independent experiments. * p<0.05 compared to 0.5 h, # p<0.05 compared to 3 h. (B) Actin staining of Saos-2 cells cultured for 0.5 h on Ti64, Zr and UFG Zr. Bars= 30 μm. 118 Fig 72: Cell spreading on zirconium surfaces. Saos-2 cells were cultured on Ti64 ( ), Zr ( ) and UFG Zr ( ) for 24 h. (A) Actin staining of cells cultured on Ti64, Zr and UFG Zr. Representative images of three independent experiments with similar results. (B) Spreading of cells cultured for 24 h on Ti64, Zr and UFG Zr. Bars = 50 μm. Results represent the mean ± SD of cell area measured in three different experiments. 120 xvi Fig 73: β-tubulin and actin cytoskeletons on zirconium surfaces. Saos-2 cells were cultured on Ti64, Zr or UFG Zr for 24 h. Bars = 30 μm. 121 Fig 74: Fibronectin analysis on zirconium surfaces. (A) Levels of fibronectin were detected in cell layers and media from Saos-2 cells cultured on Ti64 ( ), Zr ( ) or UFG Zr ( ) for 24 h. The data are expressed as the percentage of the fluorescence measured in cell layers of cells cultured on Ti64, which was given the arbitrary value of 100 and is equivalent to 70 pg/μg total proteins. Each value represents the mean ± SD of four independent experiments. (B) Fibronectin staining of Saos-2 cells cultured for 24 h on Ti64, Zr or UFG Zr. Bars = 50 μm. 123 Fig 75: Double staining of actin (red) and fibronectin (green) on zirconium surfaces. Saos-2 cells were cultured on Ti64, Zr or UFG Zr for 24 h. Bars = 20 μm. 124 Fig 76: Cell viability on zirconium surfaces. Saos-2 cells were cultured on Ti64 ( ), Zr ( ) and UFG Zr ( ) for 1, 4 and 7 days. The data are expressed as the percentage of the fluorescence measured on Ti64 at 1 day, which was given the arbitrary value of 100. Each value represents the mean ± SD of four independent experiments. * p<0.05 compared to 4 and 7 days. 125 Fig 77: ALP activity and degree of mineralisation on zirconium surfaces. hMSC cells were cultured on Ti64, Zr or UFG Zr in growth ( ) or osteogenic media ( ). The data are expressed as percentage of the absorbance measured on cells cultured on Ti64 in growth medium, which was given the arbitrary value of 100. Each value represents the mean ± SD of three independent experiments. 126 Fig 78: Schematic of geometric dynamic recrystallization. 130 Fig 79: The schematic diagram of the test sample to measure the shear strength of the bond in Krallics’s work [164]. 134 xvii Fig A1: The front panel of “LJInstron” program. The path to save data files, the strain rate of testing and gage length of specimens can be assigned in this panel. 165 Fig A2: The block diagram of the LJInstron program. 166 xviii ABSTRACT The aim of this study is to produce large quantities of bulk zirconium with an ultrafine grained microstructure and with enhanced properties. Accumulative roll bonding (ARB), a severe plastic deformation technique based on rolling, is chosen due to its availability in industrial environment. The texture, microstructure and mechanical behavior of bulk ultrafine grained (ufg) Zr fabricated by accumulative roll bonding is investigated by electron backscatter diffraction, transmission electron microscopy and mechanical testing. A reasonably homogeneous and equiaxed ufg structure, with a large fraction of high angle boundaries (HABs, ~70%), can be obtained in Zr after only two ARB cycles. The average grain size, counting only HABs (θ>15 o ), is 400 nm. (Sub)grain size is equal to 320 nm. The yield stress and ultimate tensile stress (UTS) values are nearly double those from conventionally processed Zr with only a slight loss of ductility. Optimum processing conditions include large thickness reductions per pass (~75%), which enhance grain refinement, and a rolling temperature (T ~ 0.3T m ) at which a sufficient number of slip modes are activated, with an absence of significant grain growth. Grain refinement takes place by geometrical thinning and grain subdivision by the formation of geometrically necessary boundaries. The formation of equiaxed grains by geometric dynamic recrystallization is facilitated by enhanced diffusion due to adabatic heating. Optical microscopy examination and shear testing suggest accepted bonding quality compared to that achieved in materials processed by diffusion bonding and that obtained in other ARB studies. Biocompatibility of ultrafine grained Zr processed by large strain rolling is studied by evaluating the behavior of human osteoblast cells. It is xix suggested that ultrafine grained Zr has a similar good biocompatibility as Ti6Al4V alloy and conventional Zr with a large grain size have. The improved mechanical properties together with an excellent biocompatibility make ultrafine grained Zr a promising biomaterial for surgical implants. 1 CHAPTER 1. INTRODUCTION 1.1 Nanostructured Materials Nanostructured materials may be classified into two categories: ultrafine grained (UFG) materials, with grain sizes between 100 nm and 1 μm, and nanomaterials, where the grain sizes lie below 100 nm. Due to their exceptional properties, nanostructured materials have attracted the widespread attention of researchers over the past decades, and especially in the last few years. Materials are, in general, formed by two types of atoms [1]: crystal atoms with neighbor configuration corresponding to the lattice, and boundary atoms with a variety of interatomic spacings. With decreasing grain size, the fraction of atoms ascribed to the grain boundaries increases, i.e., the volume fraction of interfaces increases [2]. This may significantly alter the physical, mechanical and chemical properties of nanostructured materials in comparison with conventional coarse-grained polycrystalline materials. With this incentive, widespread research has been carried out in this area and it has been found that nanostructured materials may exhibit higher strength/hardness, increased toughness, reduced elastic modulus and ductility, enhanced diffusivity, increased specific heat, enhanced thermal expansion coefficient (CTE), and superior soft magnetic properties in contrast to conventional polycrystalline materials. Nanostructured materials present an attractive potential for technological applications with these novel properties. 1.2 Processing of Nanostructured Materials The techniques used for the fabrication of nanostrucutred materials may be classified into “top-down” and “bottom-up”. The top-down approach consists on building 2 up a nanostructure by breaking down the microstructure of a polycrystalline bulk material of conventional grain size (d>1μm). The main synthesis methods that belong to this approach are severe plastic deformation (SPD) techniques, mechanical alloying and cryomilling. The bottom-up approach consists on arranging a nanostructure atom-by- atom, layer-by layer. The most important synthesis methods in this category are inert gas condensation, electrodeposition, sputtering, chemical vapor deposition and crystallization from amorphous materials. Bottom-up techniques give rise to smaller grain sizes, down to less than 100 nm, whereas the main advantage of top-down methodologies is the possibility of fabricating bulk structures. In the present study a technique based on applying severe plastic deformation is employed to fabricate bulk ultrafine grained materials. 1.3 Severe Plastic Deformation (SPD) Severe plastic deformation (SPD)-based techniques are nowadays widely used for the processing of bulk ultrafine-grained materials which will be investigated in this study [3-19]. Very large plastic strains are introduced in metals in order to develop finer and finer grains. Although there are many means of introducing large plastic strains, three principal methods have been invented to subject a material to severe plastic deformation: equal channel angular pressing (ECAP) [3-9], high pressure torsion (HPT) [10-12] and accumulative roll bonding (ARB) [15-19]. The three techniques are described in detail in the following sections. 1.3.1 Equal Channel Angular Pressing (ECAP) ECAP was first proposed in the Soviet Union in the 80s. Figure 1 shows the schematic illustration of an ECAP die [1, 5]. Two channels, equal in cross-section, 3 intersect at an angle Ф that is generally close to 90 o . The test sample, machined to fit within these channels, is pushed down from the upper die by a piston and is forced around the sharp corner. The strain induced on the sample after each pass depends on both the angle Ф between the two channels, and the angle Ψ defining the outer arc of curvature at the point where the two channels intersect. The ECAP process can be conducted following four different routes, termed A, B A , B C and C, that differ in the rotation of the sample between two subsequent passes, as shown in Figure 2 [20]. No rotation is applied in route A. In route B A , a 90 o rotation in alternating sense (counterclockwise-clockwise) is applied after each pass. Route B C comprises a 90 o rotation, always in the same sense. Route C comprises an interpass rotation of 180 degrees. Fig 1: Schematic illustration of Equal Channel Angular Pressing (ECAP) [1]. 4 Fig 2: The four fundamental processing routes in ECAP [20]. 1.3.2 High Pressure Torsion (HPT) The schematic sketch of a high pressure torsion die is shown in Fig 3. A small sample with the form of a disk is held under an applied pressure of several GPa and then subjected to torsional straining. HPT can provide very large values of strain (true strains ≥ 10 [3]) and has the advantage of producing exceptionally small grain sizes, often in the nanometer range (<100 nm), and even down to the 10 nm range. Furthermore, this method has the ability to process brittle materials such as intermetallics and semiconductors. A disadvantage of HPT is that the samples are very small, with maximum disk diameters of ~20 mm and thickness of ~1 mm. 5 Fig 3: Schematic illustration of a High Pressure Torsion (HPT) die [1]. 1.3.3 Accumulative Roll Bonding (ARB) Accumulative roll bonding was invented by N. Tsuji et al. in 1998 [21]. Compared with conventional rolling, ARB is a severe plastic deformation technique that allows a large amount of strain to the materials. This process consists on cutting a rolled sheet into several pieces, which are then stacked together and rolled again (Fig 4) [17]. The stacked layers are bonded together during rolling and thus a bulk material is obtained. Then the whole process is repeated until the desired strain is achieved. For achieving good bonding, surface treatments, such as degreasing by acetone and wire-brushing, are carried out before each stacking. This process allows to introduce unlimited plastic strain without any significant change in the sample geometry. The rolling in the ARB process is not only a deformation process but also a bonding process [22]. The bonding quality is a 6 decisive factor to evaluate whether accumulative roll bonding is a suitable technique to fabricate an ultrafine grained structure for a specific material. In order to achieve good bonding, stacked sheets are heated at elevated temperature below recrystallization temperature. Furthermore, large reduction in each pass is employed in order to achieve sufficient bonding. Tsuji et al. [23] reported that there is a critical rolling reduction per pass, below which it is difficult to achieve sufficient bonding. This critical reduction is approximately 35%, and it depends on the materials and the processing temperature [23]. An important advantage of ARB is that only a rolling mill is required, an equipment that is readily available in industry. Fig 4: Schematic representation of the Accumulative Roll-Bonding (ARB) process [17]. 7 ARB has widely been used to fabricate ultrafine grained cubic materials such as Al alloys [16, 21, 24-28], steels [17, 25, 28-30] and OFHC-Cu [28]. Also, some limited ARB research on HCP materials such as Mg alloys [31-33] and Ti alloys [34] was reported. The ARB process has been utilized successfully to fabricate ultrafine grained microstructures in most metallic metals. The poor bonding quality in Mg alloys did not allow obtaining bulk samples with mechanical stability [31]. Some characteristics of the microstructures and the mechanical properties of materials processed by ARB will be discussed in the following sections. 1.3.3.1 Microstructures of Materials Processed by ARB Figure 5 shows a TEM micrograph of the microstructure of the 1100 Al alloy processed at 200 o C by 6 cycles of ARB (ε = 4.8). The (sub)grain misoreintation map of the same region of the TEM image is also shown [22]. A large fraction of elongated ultrafine (sub)grains with low dislocation density is observed. This ultrafine lamellar (or pancake) boundary structure, formed uniformly in the ARB processed materials after several cycles, was also reported in numerous studies of FCC and BCC materials [24, 25, 35]. Thus, it appears that an elongated ultrafine grained structure is a typical microstructure of cubic materials fabricated by ARB. Those pancake shaped grains normally have boundary spacing or grain thickness much smaller than 1 μm. The misorientation measurements shown in Figure 5 clearly confirm that most of the elongated grains were surrounded by high-angle grain boundaries (HABs), with misorientation larger than 15 o . A large fraction of high angle boundaries, more than 70%, was obtained in this ARB processed 1100 Al alloy [24]. A large fraction of HABs is a characteristic feature of ARB processed materials. 8 Fig 5: TEM microstructure of the 1100 Al alloy processed by 6 cycles of ARB (ε = 4.8). The microstructure was observed along the transverse direction (TD). The crystallographic orientation was determined by Kikuchi-line analysis in TEM. The boundary misorientation is superimposed in the map in degrees [22]. Li et al. [28] investigated the microstructure homogeneity through the sheet thickness by electron back scattering pattern (EBSP) analysis for various kinds of metallic materials (Al, Cu, 36% Ni steel and IF steel) deformed to an equivalent strain of 9 4.8 by accumulative roll bonding. Figure 6 presents the mean grain thickness and the fraction of high angle boundaries at various thickness locations in the IF steel processed by 7 cycles of ARB at 500 o C (ε = 5.6) [22]. It can be seen that the grain size and the fraction of high angle boundaries are homogenous throughout the thickness. A homogeneous elongated ultrafine grained structure through the ARB processed sheets was observed in all the studied materials. This through thickness uniform microstructure was also reported by Kamikawa et al. [35]. However, several studies [21, 35-37] showed that a complicated texture gradient developed through the thickness of ARB processed materials. Fig 6: Mean grain thickness and fraction of high angle boundaries through the thickness of the UFG IF steel processed by 7 cycles of ARB (ε = 5.6) at 500 o C [22]. 10 There is very limited research on ARB processed HCP materials. It is well known that the deformation behavior of HCP metals is significantly different from that of cubic metals due to their limited active slip systems. Thus it is likely that the microstructure evolution of hexagonal metals during ARB is different from that in cubic metals. Pérez- Prado et al. [32] studied the effectiveness of accumulative roll bonding for the fabrication of ultrafine grained structure of Mg-Al-Zn alloys with different Al content (from 3% to 9%). It was found that an increase in the Al content allows obtaining ultra-fine grain sizes (d < 1 μm) via accumulative rolling bonding. The homogeneity of the microstructure could be improved by increasing the number of rolling passes [32]. Del Valle et al. [31] investigated the effectiveness of accumulative roll bonding to refine grain sizes for a Mg- based AZ61 alloy. In their study, AZ61 alloy was initially rolled using a first pass of 30% and a second pass of 50% reduction, and it was subsequently rolled using one final pass of various reductions (25%, 50%, 66% and 80%). It was reported that the grain size could be reduced to the submicron range only when thickness reductions larger than 66% were used [31]. ARB processing has also been applied to other hcp materials. Terada et al. [34] deformed commercial purity titanium (ASTM grade 2) by accumulative roll bonding using up to 8 cycles (ε = 6.4) at room temperature. Figure 7 shows the microstructures of the commercial purity Ti ARB-processed by 4 cycles (ε = 3.2) at room temperature by means of TEM micrographs. It can be clearly seen that two kinds of ultrafine microstructures were formed: a lamellar boundary structure along the rolling direction and an equiaxed grained structure. The area fraction of equiaxed grains increased as the number of cycles increased, finally reaching a value of 90% after 8 cycles (Figure 8). The 11 formation of these equiaxed grains was attributed to recovery at inhomogeneously deformed regions (e.g. shear bands), where very large strain localizes. The recovery might be enhanced by local adiabatic heating enhanced by the low thermal conductivity of Ti [34]. It should be emphasized that an equiaxed grained structure has not been observed in cubic metals processed by ARB even if the metals were subjected to very large strains. Thus, equiaxed grains may be a characteristic feature of microstructures of ARB-processed HCP materials. Fig 7: TEM micrographs showing the microstructure of commercial purity Ti processed by ARB using 4 cycles (ε = 3.2) at room temperature (a). High magnification micrographs of the lamella (b) and equiaxed structure (c), respectively [34]. 12 Fig 8: TEM micrographs of commercial purity Ti ARB-processed by (a) 6 cycles (ε = 4.8) and (b) 8 cycles (ε = 6.4) [34]. In summary, the development of an elongated ultrafine grained microstructure is the characteristic feature of cubic materials processed by ARB. However, equiaxed grains form in commercial purity Ti, a HCP metal, processed by ARB. More research is needed to clarify whether an equiaxed grained structure is a characteristic microstructure feature of other ARB-processed HCP metals. The microstructures fabricated by ARB process normally have a large fraction of high angle boundaries. Furthermore, a homogenous grain size and boundary distribution through the thickness is frequently observed in ARB processed materials. However, a complex through-thickness texture gradient may appear. 1.3.3.2 Mechanical Properties of ARB Processed Materials In general ARB-processed materials have high mechanical strength due to their small grain size. Figure 9 shows a summary on the tensile strength of various metals 13 (1100 Al alloy, IF steel, CP-Ti and AZ91 Mg alloy) fabricated by ARB. For comparative purposes, the tensile strength of the starting materials is also shown. The concrete data is tabulated in Table 1. It is to note that ARB-processed materials possess a high tensile strength, which is normally 2-4 times higher than that of the starting annealed materials. Simultaneously, the ductility at room temperature decreases significantly. Figure 10 shows the decrease in elongation of both IF steel and CP-Ti processed by ARB. Clearly, the elongation drops dramatically after the first ARB pass. This can be explained in terms of early plastic instability in the UFG metals due to high flow stress and lack of strain- hardening capability [25, 38]. It is envisaged that ultrafine grained materials would potentially be good candidates for superplastic forming, since grain boundary sliding is stimulated due to the large fraction of high angle boundaries. Tsuji et al. [39, 40] reported that an ARB-processed 5083-Al alloy with UFG microstructure undergoes low-temperature superplasticity at 200 o C. 84 274 400 159 304 870 860 405 0 200 400 600 800 1000 1100 Al Alloy IF Steel CP-Ti AZ91 Mg Alloy True Tensile Strength (MPa) Starting Materials ARB Processed Fig 9: True tensile strength of several ultrafine grained materials processed by ARB and comparison with the strength of the starting materials [16, 17, 33, 34]. 14 (a) (b) Fig 10: Strength and elongation of (a) IF steel ARB-processed by various cycles at 500 o C [17] and (b) commercial purity Ti ARB-processed by various cycles at room temperature [34]. In (b), broken lines indicated typical tensile strength and total elongation for a conventional Ti-6Al-4V alloy. 15 Table 1: Mechanical properties of various materials processed by several cycles of accumulative roll bonding. Materials ARB Process Thickness reduction per pass Strain True tensile strength of as-ARB material (MPa) True tensile strength of the starting material (MPa) Elongation of as-ARB material Elongation of the starting material Reference 1100-Al 8 cycles at 200 o C 50% 6.4 304 84 8% 42% [16] IF steel 7 cycles at 500 o C 50% 5.6 870 274 1.5% 57% [17] CP-Ti 8 cycles at room temperature 50% 6.4 860 400 7% 45% [34] Mg alloy AZ91 4 cycles at 375 o C 80% 7.4 405 159 2.5% 2% [33] 16 1.4 Mechanical Properties of Ultrafine Grained Materials The promising mechanical properties of nanostructured materials are one of the incentives for the widespread research in this field. In this section, the current knowledge on the mechanical properties of nanostructured materials is reviewed. In particular, properties such as yield stress, ductility, strain hardening and strain-rate sensitivity will be discussed in separate sections. 1.4.1 Yield Stress It is well known that the grain size has a significant effect on the mechanical behavior of materials. In the conventional polycrystalline range (d>1μm) the relationship between yield stress, σ y , and grain size, d, is found to follow the Hall-Petch equation: 2 / 1 0 − + = kd y σ σ (1) Where σ 0 is the friction stress and k is a constant. The Hall-Petch relationship predicts that the yield stress increases when the grain size decreases. This dependency has been proved right in materials with conventional grain sizes (d>1μm). However, when the grain size is smaller than 1 μm, or especially for grain sizes close to 100 nm, some deviations from the Hall-Petch law have been observed. Chokshi et al. [41] first reported an inverse Hall-Petch effect by investigating nanocrystalline Cu and Pd samples made by inert gas condensation. Both metals exhibited a negative slope, shown in Fig 11. Fig 12 shows the Hall-Petch plots for Cu taken from different sources [1]. These plots show ambiguity in the dependency of the yield strength on grain size as d falls down to a value below ~25 nm (d -1/2 =0.2). While some results predict a plateau, others show a decrease. The breakdown in the Hall-Petch 17 law has been attributed to different deformation mechanisms that become dominant once the grain size is reduced below a critical value [13]. Chokshi et al. [41] attributed the negative Hall-Petch slope to enhanced diffusional creep in nanocrystalline samples at room temperature. That is, diffusional creep (Coble creep) is likely to play an important role in the deformation of nanocrystalline copper for two reasons: (1) the specimens have a small grain size (~10 nm), which enhances the diffusion creep rate and (2) grain boundary diffusion in nanocrystalline copper is substantially more rapid than polycrystalline coarse grained copper [42, 43]. Therefore, enhanced diffusional creep results in a decrease in the strength of nanocrystalline copper when the grain size is decreased. Weertman et al. [44] obtained a positive H-P slope when performing hardness measurements in samples processed by inert gas condensation followed by ambient temperature densification through uniaxial pressing. However, tensile tests in those same samples resulted in a negative H-P slope (Fig 13). The decrease in σ y with decreasing grain size was attributed by Weertman et al. [44] to the presence of flaws. Weertman [44] argued that the negative slope obtained by Chokshi et al. [41] is an artifact due to the use of a single sample subjected to repeat anneals to change the grain size. Grain boundaries in the processed samples in Chokshi’s study are in a non-equilibrium state and may appear “transparent” to dislocations. However, after annealing, these non-equilibrium boundaries are converted to an equilibrium condition, without grain growth. Thus, the processed material could be strengthened by annealing, resulting in the occurrence of the negative slope in the Hall-Petch relation. 18 Fig 11: Inverse Hall-Petch trend for Cu and Pd as shown by Chokshi et al. (from hardness measurements) [41]. Fig 12: Yield stress versus grain size plot for Cu. Data complied from various sources [1]. 19 Fig 13: Positive Hall-Petch slope for Cu with higher values for compressive (from hardness measurements) than for tensile strengths [45]. Figure 14 shows the variation of the yield strength in Ti with decreasing grain size [1]. For grain sizes smaller than 400 nm, the slope of the Hall-Petch curve becomes increasingly smaller than that from the conventional H-P prediction. The causes for this deviation are not clear. Additionally, there is insufficient evidence to validate the existence of this effect, since information on the H-P relation for small grain sizes in other hcp materials is lacking. It is thus very timely to investigate the Hall-Petch effect in nanostructured hcp metals such as Zr. 20 Fig 14: Variation of the yield stress with grain size for Ti. Comparison to the linear Hall- Petch response [1]. 1.4.2 Ductility In general, the ductility decreases as the grain size is reduced down to the ultrafine and nanograin regime (Fig 15). However, several strategies have been devised over the years in order to enhance ductility in very fine grained materials without losing significant strength. First, ductility can be enhanced by developing a bimodal grain size distribution, i.e., a microstructure with a number of coarse micron-sized grains embedded in a matrix of ultrafine or nanograins. The small grains contribute to enhance the strength and the large grains allow significant dislocation activity and, thus, strain hardening. Zhang et al. [14, 46-48] fabricated bimodal microstructures by milling. During the early milling stages (≤1 h), defect/dislocation generation dominates and therefore, grain size is continuously refined. However, for longer milling time (1 to 3 h) dynamic recovery 21 dominates and heterogeneous grain growth is facilitated. Thus, they found out that after 1 h the instability of the fine grains formed at short milling times and their subsequent coarsening result in the occurrence of the bimodal microstructure. Plastic deformation accommodated by large grains enhances the ductility of the material milled at room temperature about 3 h. Wang et al. [38] produced microstructures consisting of a combination of nanoscale/ultrafine grains (80-200 nm) along with 25% volume fraction of coarser grains (1-3 μm) by rolling pure Cu to a 93% thickness reduction at liquid nitrogen temperature and then annealing it at temperatures below 200 o C. Growth twins in the larger grains and preferential accommodation of strain in the larger grains result in the excellent combination of strength and ductility. An alternative method for enhancing the ductility in very fine grained materials is by introducing non-equilibrium grain boundaries. These boundaries store an excess number of dislocations that may slip readily. Also, due to the disordered nature of the boundaries, diffusion is favored and the occurrence of grain boundary sliding is significantly enhanced, even at low temperature. Non-equilibrium grain boundaries are formed during processing by severe plastic deformation techniques, especially when processing takes place at rather low temperatures, due to the very large strains applied. Strain rate sensitivity, m, is another factor contributing to and determining ductility. Equation (2) shows that the total ductility achieved in a material depends, at least in part, on the value of m [49]: m m m A B P dt dA / ) 1 ( / 1 ) / ( − = − (2) 22 Where, P is the tensile force, A is the cross-sectional area of the specimen, t is the time and B is a constant which incorporates experimental conditions such as temperature and grain size. It follows from equation (2) that the value of dA/dt tends to a common level which is independent of A as m→1. The possibility of failure by necking therefore becomes increasingly less important at higher values of m, so that ductility tends to increase with increasing m. However, the strain rate sensitivity of nanostructured metals has not been well studied; this issue will be discussed later. Fig 15: Compilation of yield strength versus % elongation of various ultrafine grained metals [13]. 1.4.3 Strain Hardening Ultrafine grained materials exhibit, in general, no strain hardening and even softening at room temperature after an initial stage of rapid strain hardening over a small plastic strain regime (~1-3%) [7, 8, 50]. As mentioned earlier, strain hardening is deeply 23 related to ductility and its absence causes localized deformation leading to low ductility. In ultrafine grained materials dislocation density saturates after small strains due to enhanced dynamic recovery caused by the annihilation of dislocations into the grain boundaries. This premature saturation leads to a low strain hardening rate or even the absence of strain hardening. In nanomaterials with grain sizes smaller than 50 nm deformation is suggested to be controlled by other mechanisms such as grain-boundary sliding [51, 52]. Here, rapid strain hardening at the first stages of deformation leads to early fracture. 1.4.4 Strain-rate Sensitivity There have been reports of both increased and decreased strain rate sensitivity with decreasing grain size in metals. The study by Wei et al. [53] shows that the strain rate sensitivity is increased at grain sizes below a critical value (Fig 16). Results by Gray et al. [54] show relatively high rate sensitivity on ultrafine grained FCC metals produced by ECAP (Cu, Ni and Al-4Cu-0.5Zr) which is attributed to a high pre-existing dislocation density. A similar effect is observed for aluminum produced by ECAP. Ultrafine grain sized Al has a slope m over three times than that of conventional polycrystalline Al [55]. The increased strain rate sensitivity can be explained by a change in the rate controlling mechanism for plastic deformation. Cutting of forest dislocations is a mechanism that operates at coarse grain sizes. However, at the scale of ultrafine grain sizes grain- boundary diffusion starts to predominate. Hoppel et al. [19] observed that in ECAP copper there was a significant increase in strain-rate sensitivity with the number of passes, which was attributed to the switch to thermally-activated dislocation annihilation at grain boundaries. 24 Fig 16: Strain rate sensitivity plot for Cu as a function of grain size [53, 54]. A decrease in the strain rate sensitivity with decreasing grain size has been reported as well, especially in BCC metals. Jia et al. [56] reported that in BCC iron the calculated strain rate sensitivity values m were 0.009 for a grain size of 80 nm, 0.012 for a grain size of 138 nm, 0.023 for a grain size of 268 nm, 0.045 for a grain size of 980 nm and 0.08 for a grain size of 20 μm. Malow et al. [57] prepared nanocrystalline Fe using ball milling and consolidation, and found a low m~0.006 at d~20 nm. In HCP metals, Zhang et al. [47] reported the effect of temperature on the flow behavior of UFG and nanocrystalline structured Zn by ball-milling. Tensile test results for 3 h ball-milled Zn samples tested at different temperatures, but the same strain rate (10 -4 s -1 ), are shown in Fig 17. Strain hardening was low at all temperatures; at 200 o C, it ceased to exist. The results of jump tests performed at 20, 40 and 60 o C on the ball milled Zn samples are illustrated in Fig 18. The calculated m values were 0.15 for tests at 20 and 40 o C and about 0.17 for test at 60 o C. 25 Fig 17: Tensile stress-strain curves for ball milled (3 h) Zn tested at 20, 40 and 200 o C at a constant strain rate of 10 -4 s -1 [14]. Fig 18: Strain-rate jump tests (compression) performed on ball milled (3 h) Zn at 20 and 60 o C [14]. 26 1.5 Texture 1.5.1 Introduction Most engineering materials are polycrystalline, and their component units are referred to as crystals or grains. Each individual grain in a material possesses a discrete crystallographic orientation. The presence or absence of a preferred orientation among numerous individual grains in the material defines the texture of the material. The importance and significance of texture to materials lies in the fact that many properties of materials are texture-dependent. The most established method for measuring textures is by X-ray diffraction. The texture that is obtained by this technique is an average value for the whole sampled volume, typically comprising thousands of grains. This texture is termed “macrotexture”. The X-ray texture tells us what volume fraction of the specimen has a particular orientation. However, it does not tell us how these grains are spatially distributed throughout the material. Another approach to texture which deals with the orientation of individual grains, and also analyzes the spatial location of these grains, has been termed “microtexture”. Microtexture measurements are usually performed using electron backscatter diffraction (EBSD) methods and orientation imaging microscopy (OIM). 1.5.2 Representation of Texture The main orientation representations are direct and inverse pole figures, which will be discussed in the following paragraphs. The stereographic projection is a means of representing three-dimensional angular relationships in two dimensions. The stereographic projection utilizes a sphere that is defined in a reference axis system (Figure 19). The axis system normally is chosen to 27 coincide with a set of directions fixed on the sample rather than specific crystallographic directions. In this study the rolling, transverse and normal directions (RD, TD and ND) have been chosen to be the characteristic directions. Direct pole figures are constructed by, first, projecting a given family of crystal directions from an origin O, located at the center of the reference sphere, shown in Figure 19, on the surface of the sphere. Points such as P on the sphere surface are then projected onto the equatorial plane to obtain the pole figure corresponding to the specific family of crystallographic directions. The projection plane is usually normal to one of the fixed axes of the reference frame, i.e. normal direction (ND), so that the plane contains the other two axes. In this work the families of low index crystallographic directions chosen were <0001>, > < − 0 1 10 and > < − 0 2 11 . Figure 20 represents the process of building a pole figure. When examining pole figures it is important to bear in mind that the stereographic projection method causes the distance (or angle) between two points near the center of the plane to appear smaller than that of the same points located further from the center of the plane. This can be seen in Figure 21, wherein the two sets of points are separated by 20 o . 28 Fig 19: Reference sphere, where the origin, ‘north pole’ and ‘south pole’ labelled O, N, S respectively. A crystal is imagined to lie at O. The normal of a plane in the crystal intersects the reference sphere at P. P is projected from the south pole onto the equatorial plane at p; this is the stereographic projection of P [58]. 29 Fig 20: Construction of pole figures. (a) projection of <001> directions on the reference sphere; (b) projection from south pole of <001> directions on the equatorial plane; (c) repeat the process with other crystals with similar orientations; (d) and (e) representing the density of poles through contours of intensity. 30 Fig 21: Stereographic projection on a wolfnet, where the two points located near the center are separated by 20 o and the two points located at the perimeter of the projection are also separated by 20 o . It should be noted that while the angular separation distance is the same in both cases, the physical distances on the projection are not [59]. Inverse pole figures are constructed in a similar manner except that the reference axes are now chosen to be parallel to specific crystallographic directions. In this study, which focuses on the study of hcp materials, the inverse pole figures will be represented using a unit triangle with the apices being the <0001>, > < − 0 1 10 and > < − 0 2 11 crystal directions. Orientations may be represented in a number of ways, using either discrete points for each measurement or contour plots with each contour level corresponding to a density of points above the normalized background level. Randle and Engler [58] have provided a through introduction to these topics as well as a background on texture representations that have been used in this study. 31 1.5.3 Texture Measurements 1.5.3.1 Macrotexture Measurements X-ray diffraction was first employed to investigate the texture of materials by von Laue in 1912 [60, 61]. The principle of pole figure measurements by means of X-ray is based on Bragg’s law which can be described by the following equation: θ λ sin 2d n = (3) where n is the order of reflection, λ is the X-ray wavelength, d is the interplanar spacing of the diffracting planes and θ is the corresponding Bragg angle. To derive the crystallographic orientation of a given crystallite, the arrangement of a set of lattice planes (hkl) has to be determined with respect to an external axis system. The axis system normally is chosen to coincide with a set of directions characteristic of the sample, i.e. RD, ND and TD. Figure 22 illustrates the principle of texture measurements by X-ray diffraction. The crystal is irradiated with monochromatic radiation at the proper Bragg angle for reflection at the lattice planes (hkl), and the detector is set at the angle 2θ with respect to the incident beam. A reflected intensity is only measured if the corresponding lattice planes are arranged such that they lie parallel to the sample surface in order to satisfy the Bragg’ law, as shown in Figure 22(b). In most cases no reflection is obtained (Figure 22(a)). In order to ensure reflection from other lattice planes, the sample has to be rotated and/or tilted until the planes are in reflection condition, i.e. parallel to the sample surface. The rotation and tilt angles are a measure for the arrangement of the lattice planes with the crystal and illustrate the orientation of the crystal with respect to the external axis system. In practical applications, the sample is systematically rotated in a texture goniometer about well-defined angles in such a way 32 that all possible lattice planes are successively brought into the reflection condition and reflected intensities are recorded as a function of these rotation angles so that they can be readily be represented in a pole figure. Fig 22: Schematic showing the effect of sample rotation on the arrangement of the lattice planes. (a) Untilted position (α=β=0 o ); (b) sample tilted such that the lattice planes satisfy the Bragg condition (α > 0 o , β > 0 o ) [58]. As mentioned previously, X-ray diffraction measurements are conducted over large surface areas, typically encompassing thousands of grains. This technique is the most common technique to determine the macrotexture of a material. However, the computation of grain-to-grain misorientations is practically impossible since the X-rays illuminate many grains simultaneously, making the distinction of individual grain orientations unfeasible. An alternative technique to measure macrotextures is neutron diffraction, which is out of the scope of this study. 33 1.5.3.2 Microtexture Measurements The use of electron diffraction in the transmission electron microscopy (TEM) and scanning electron microscopy (SEM) to determine microtexture has evolved more recently. For the vast majority of microtexture work in both the TEM and SEM a Kikuchi diffraction pattern has been used to obtain the complete crystallographic information corresponding to each individual crystallite. The formation of a Kikuchi pattern is depicted schematically in Figure 23. An electron beam is diffusely scattered in all directions when it enters a crystalline solid, shown in Figure 23(a). There must always be some electrons arriving at the Bragg angle θ B at every set of lattice planes, and these electrons can then undergo elastic scattering to give a strong, reinforced beam (Figure 23(b)). Since diffraction of the electrons through the Bragg angle is occurring in all directions, the locus of the diffracted radiation is the surface of a cone (Kossel cone). These Kossel cones will intersect the Ewald sphere, and be observed as lines, shown in Figure 23(c). Since the radius of the Ewald sphere is large, the projection of the cones will be in the form of straight lines rather than parabolas. These are Kikuchi lines. Each pair of parallel lines represents a lattice plane and the width between lines is an angular distance of 2θ B which in turn is proportional to the interplanar spacing by following the Bragg’s law shown in equation (3). 34 Fig 23: Schematic representation of all electron scattering localized at a single point in (a). In (b) some of the diffusely scattered electrons undergo Bragg diffraction from certain atom planes. In (c) Kossel cones are formed and projected onto the Ewald sphere, creating parabolas which approximate to straight Kikuchi lines [62]. 35 In TEM, the diffracted beams pass through the specimen, which means that the material must be in the form of a thin foil to avoid absorption of the diffusely scattered electron. In the SEM diffraction occurs from the interaction of primary backscattered electrons with lattice planes close to the specimen surface. Thus, Kikuchi patterns will be observed only for regions near the free surface. Diffraction patterns obtained from the backscattered electrons in an SEM following Kikuchi diffraction have been termed electron backscatter pattern (EBSP) or electron backscatter diffraction (EBSD). SEM EBSD methods have been used extensively in this research. Figure 24 shows an outline of the experimental device used for EBSD. It consists basically of a scanning electron microscope, a TV camera and a computer with the software to analyze the required data. After polishing, samples are placed into the SEM in a special holder that allows placing the sample with a high inclination (usually 70 o ) to the horizontal. This reduces the penetration of the beam and therefore, absorption, and the diffracted signal is more intense. The inclination of the holder has been optimized, moreover, to guide the diffracted signal towards the lens of the camera and to get Kikuchi patterns with as much information as possible and with maximum contrast. The diffracted electrons impinge upon a phosphor screen, thus illuminating it. The image of Kikuchi patterns on the phosphor screen is captured by a camera and an analog signal is sent to the control unit of the chamber, where the light intensity can be measured and adjusted. Simultaneously, it is possible to see the image of Kikuchi patterns in real time on a television screen. After being integrated and filtered, the signal is digitized and sent to a work station provided with a program that helps to determine the orientation and save the data analyzed in a file. 36 Fig 24: A schematic drawing showing the experimental devices for EBSD. An example of an EBSD system used to measure the microtexture of samples is shown in Figure 25. It includes a scanning electron microscope, a CCD camera fitted with a lens coated with phosphorus, a control unit of the CCD camera that allows us viewing images, correcting, integrating and filtering the analog signal input, and producing a digital signal output, and a workstation in which a computer program has been installed to perform the information analysis. 37 Fig 25: EBSD system used for determining microtexture in this work (OIM TM , TexSEM Laboratories, Inc., Provo, Utah). Figure 26 presents three steps in the process of interpreting Kikuchi patterns. Once the image of Kikuchi patterns is captured (Fig. 26(a)), the software detects the lines with greater intensity (Fig. 26(b)) and associates them with the corresponding crystallographic planes (Fig. 26(c)), assigning the various indexed poles. A very detailed description of the process of the calibration and geometric considerations necessary for the allocation of indexes can be found in [58]. The software also allows analysis and representation of data, as will be explained later. 38 a) b) c) Fig 26: Determination of orientations. (a) Kikuchi patterns obtained in SEM (b) detecting the Kikuchi lines with greater intensity and (c) allocation of indexes. The foundations of the process of acquiring data on an EBSD system are explained by Randle et al. [58]. In practice, the electron beam is displaced from a point to another point over a preselected area on the sample surface following a route. Initially, an incident beam impinges on a crystalline grain (“a”), shown in Figure 27. The arrow indicates the Kikuchi lines obtained in this position, which will be indexed automatically later. This Kikuchi pattern is displayed in real time on a TV screen. Then the beam is moved to an adjacent position. When Kikuchi pattern obtained in this point changes, it is indicated that the beam now impinges on an adjacent crystalline grain (“b”), whose Kikuchi pattern is re-analyzed. The number of times that this operation is repeated and the route of the electron beam on the sample must be chosen by a user taking into account the nature of the material and microstructural information that is sought. The point-point step size is another crucial parameter to control and is based on the expected microstructure and size of the region being examined. 39 The route, followed by the electron beam in Figure 27, has been drawn by a white line. As the beam sweeps the sample along this path, data on orientation will be acquired and recorded from adjacent positions. When comparing two consecutive orientations it is possible to obtain information about orientation distinction between those two positions or the nature of the grain boundary that separates them. Then the software is employed to produce orientation image mapping, pole figures representing the texture, disorientation distribution histograms, etc. 40 Fig 27: Methodology used for determining individual orientation and studying the nature of grain boundaries present in materials. 41 1.6 Recrystallization On annealing a cold worked metal at an elevated temperature or during hot-working, the microstructure and the properties of materials may be partially restored to their original values by recovery. This phenomenon includes all processes decreasing the stored energy that do not require movement of high angle boundaries. A more pronounced restoration process is called recrystallization. By the latter, new dislocation- free grains are formed within the deformed or recovered structure. Recrystallization can be generally divided into static recrystallization (SRX), that occurs when a heat treatment is applied to a deformed material [63], and dynamic recrystallization (DX), that takes place during deformation. The latter is more relevant to the present study, where processing at temperatures above room temperature (~0.3T m ) will be used to fabricate nanostructured materials. Doherty et al. [63] reviewed dynamic recrystllization mechanisms. Recently, Kassner et al. [64] rationalized the types of dynamic recrystallization as follows: (1) discontinuous dynamic recrystallization (DRX); (2) geometric dynamic recrystallization (GDX); (3) continuous reactions or continuous dynamic recrystallization (CDX); (4) formation of geometrically necessary boundaries (GNBs) (not generally considered as a recrystallization phenomenon), all of which will be discussed subsequently. 1.6.1 Discontinuous Dynamic Recrystallization (DRX) Discontinuous dynamic recrystallization describes the formation and migration of high-angle boundaries (HABs, with a misorientation, θ > 15 o ) driven by the stored energy of deformation. DRX is frequently observed in materials with low or medium stacking fault energy where dynamic recovery may be retarded due to inhibition of dislocation 42 cross-slip. DRX can be simply described as follows. Nearly dislocation-free new grains can nucleate at the old grain boundaries and start growing. As the material continues to deform, the dislocation density of the new grains increases, thus reducing the driving force for further growth. A critical strain (ε c ) is necessary in order to initiate DRX and this strain steadily decreases with decreasing stress [65]. Another general characteristic of DRX is that the stress-strain curves (σ vs. ε) of materials which undergo DRX usually present a broad peak before a plateau appears with increasing strain. Under conditions of low Zener-Hollomon parameter, Z, defined as: ) exp( RT Q Z • =ε (4) (here, • ε is low strain-rate; T is temperature; Q is activation energy and R is a constant), multiple peaks or undulations can be observed. DRX may be also evident in hardening curves (θ=dσ/dε versus σ). Figure 28 illustrates schematically the onset of DRX in strain hardening curves. In the absence of DRX and only under dynamic recovery, the hardening rate θ decreases dramatically on the Stage II which is followed by Stage III with a “linear-like” decrease in θ as the stress increases (path ABD on Fig 28). However, sometimes, at a stress σ c , corresponding to a critical strain ε c, the hardening rate changes abruptly indicating the onset of DRX (path ABC on Fig 28). Kassner et al. [66, 67] observed this saturation (θ=0) in association with DRX in high purity (99.999%) aluminum subjected to low strain rate (10 -5 s -1 ) compression. Stage IV may be also observed, in which a constant hardening rate is evident. This can preclude the steady state as a balance between dislocation hardening and dynamic recovery. 43 Fig 28: Schematic representation of the strain hardening versus stress for dynamically recrystallized (ABC) and dynamically recovered (ABD) material [68]. 1.6.2 Geometric Dynamic Recrystallization (GDX) Geometric dynamic recrystallization has been usually observed in materials deformed to large strains at elevated temperatures. GDX was first described by McQueen and coworkers in pure aluminum and in aluminum magnesium alloys [69-71]. McQueen and coworkers [69, 72, 73] and Kassner and coworkers [70, 74] observed strain hardening to a broad maximum at equivalent strains ε ≈ 2 in torsion of aluminum under constant strain-rate conditions. Then, the aluminum softens by about 20% to reach a steady-state where stress is independent of strain. McQueen and coworkers observed a decrease in the average Taylor factor which is consistent with the observed change in flow stress. They attributed this softening to texture change rather than substructural changes suggested by other investigators [75-77]. Regarding the microstructural 44 evolution, Kassner et al. [70] reported the development of serrations during the deformation. This is characteristic of GDX. The serrations originate at the original grain boundaries and generally consist of a triple junction between an original high angle boundary and a low angle boundary. As the deformation continues, the original grains elongate and subgrains develop constantly. The subgrains remain equiaxed and of constant size through the ultra-large deformation, which suggests that the subgrain boundaries must migrate and/or annihilate to maintain the equiaxed morphology [78]. At large strains, the original grains thin down to the dimensions of the subgrain diameter and grain boundary serrations begin to pinch off by coming in contact with each other. Another microstructural characteristic observed in many investigations [69, 70, 72, 73, 79] is that nearly one third of the subgrain facets become high angle boundaries at these large strains. Montheillet and coworkers [75] suggested that the continued accumulation of dislocations in initially low misorientation subgrain boundaries leads to the formation of high angle boundaries, akin to “continuous dynamic recrystallization” (CDX) which will be discussed in a subsequent section, and thus result in an increasing number of high angle boundaries observed in [69, 70, 72, 73, 79]. Thus, GDX has been confused with CDX. However, Kassner et al. [70] investigated the evolution of the misorientation distribution with various steady-state strains up to 16.33 in pure aluminum deformed in torsion at 644 K. A bimodal distribution in the misorientation of 220 boundaries was observed and over the strain of 3.11 it was composed of a fraction (up to one-third) of high angle boundaries (θ > 10 o ) and of the remaining fraction (two-thirds) of very low angle boundaries (θ < 4 o ). The absence of boundaries with a misorientation between 4 and 10 degrees suggested that the misorientation angle of low-angle boundaries does not 45 continuously increase to form the high-angle boundaries, as in other gradual grain refinement mechanisms, such as CDX and rotational dynamic recrystallization (RDX). 1.6.3 Continuous Dynamic Recrystallization (CDX) Continuous dynamic recrystallization has been suggested to describe the refinement of grain sizes for various fcc metals, such as commercial purity Al [80-84], stainless steel [85, 86], and pure copper [87]. CDX is usually associated with the recovery process where the formation of subgrain boundaries appears. As deformation proceeds, subgrain boundaries, normally having misorientations of just a few degrees, migrate and annihilate with other subgrain boundaries, forming equiaxed subgrains. Their misorientation remains constant although some have suggested a continous increase. However, it has been suggested that sometimes due to the pinning by precipitates or other particles, subgrain boundary migration and annihilation is retarded. Dislocations thus accumulate in low-angle subgrain boundaries, causing the increase of the misorientation across the boundaries. Eventually, these subgrain boundaries may transform into high angle boundaries (θ >10 o −15 o ). 1.6.4 Geometric Necessary Boundaries (GNBs) The formation of geometric necessary boundaries is also related to the appearance of a new HAB structure during deformation, but they are, in general, not considered the result of recrystallization phenomenon. GNBs have been observed during both cold [88- 90] and warm [91, 92] deformation. GNBs are usually referred to as dense dislocation walls (DDWs), cell block boundaries (CBB), extended planar boundaries (EPB), micro bands (MB) or carpet structures [93, 94]. GNBs separate regions that deform by different slip system combinations, with different strain amplitudes, and various strains [64]. 46 GNBs are different from incidental dislocation boundaries (IDBs), which form by trapping of glide dislocations [95]. Figures 29 and 30 present GNB and IDB arrangements in pure Ni along with a schematic drawing. Fig 29: Transmission electron micrograph and schematic representation of geometric necessary boundaries (GNBs) and incidental dislocation boundaries (IDBs) in pure Ni, rolled at a reduction of 20% [96]. RD represents the rolling direction. 47 Fig 30: Schematic three dimensional drawing of the GNBs (solid planar boundaries) and IDBs (short grey boundaries) bridging between them. The misorientation (θ) and the boundary spacing (D) are illustrated for each type of boundary [97, 98]. Bold lines represent high angle boundaries (HABs) and thin lines represent low angle boundaries (LABs). One of the main distinctions between GNBs and IDBs is the difference in misorientation and the corresponding misorientation evolution with strain. GNBs are formed in order to accommodate lattice rotations and separate cell blocks of different operating slip systems [96, 99]. IDBs result from the entanglement and mutual trapping of glide dislocations formed during homogeneous deformation. These dislocation boundaries do not contribute to the misorientation between adjacent cell blocks. Thus, the 48 misorientation across GNBs is larger than that across IDBs [96, 99]. Figure 31 presents the average misorientation angle of GNBs and IDBs for high purity cold rolled aluminum and nickel [96]. It shows that in both materials, the average misorientation of GNBs is larger than that of IDBs. It also shows that the average angle increases more rapidly with increasing strain for GNBs than for IDBs. This is another marked difference between GNBs and IDBs. Liu and Hansen [89] analyzed 1500 boundaries in high purity (99.996%) aluminum, cold-rolled to 30% (ε = 0.41), in order to study the change in the angle of misorientation across dislocation boundaries which were classified as GNBs and IDBs with increasing strain. It was found that a much more rapid increase in misorientation angle across GNBs is evident with increasing strain compared to IDBs [89], which is shown in Figure 32. Cizek et al. [100] also reported large misorientation increases for GNBs with strain, but average misorietation across IDBs remaining constant during straining, in a commercial Al alloy alloy deformed in tension at 250 o C and a strain rate of 2 × 10 -4 s -1 up to strains of 0.5. The results reported in Cizek et al’s investigation are presented in Figure 33. 49 Fig. 31: Power law relationship between the average misorientation angle and applied von Mises strain for high purity cold rolled aluminum and nickel. The average angle increases more rapidly with strain for GNBs than for IDBs [96]. Fig 32: Average misorientation angles of both IDBs and GNBs versus strain in high- purity aluminium (99.996%) [89]. 50 0.1 0.2 0.3 0.4 0.5 0 1 2 3 4 5 Misorietation Angle (deg) Strain IDB GNB Al alloy 5005 Fig 33: Misorientation angles across IDBs and GNBs as a function of strain in aluminium alloy AA5005 [100]. Another characteristic parameter for GNBs and IDBs is the spacing between the boundaries. Liu and Hansen [89] reported that the spacing between dislocation boundaries decreases with strain and that this change is more rapid for GNBs than for IDBs in a cold-rolled high purity Al mentioned above. A simple relationship between average misorientation angles (θ av ) and average spacing between boundaries (D av ) was derived [101, 102] and it was found that the product of the average spacing and the average misorientation angles, compensated by the Burgers vector, b, is equal to a constant C, C b D av av = ⋅θ (5) 51 However, it appears that Eq (5) is only valid for boundaries of the same type [102], confirming the difference of character between GNBs and IDBs. Hughes et al. [103] determined the relationship between the scaling parameters (D av and θ av ) and strain. For GNBs this relation is 3 / 2 3 / 2 ~ , ~ − ε ε θ GNB GNB D (6) and for IDBs 2 / 1 2 / 1 ~ , ~ − ε ε θ IDB IDB D (7) According to equations (6) and (7) the reduction in spacing between GNBs with increasing strain is significantly faster than the reduction in spacing between IDBs, which is consistent with the observations reported by Liu and Hansen [89]. In summary, it has been reported that geometric necessary boundaries and incidental dislocation boundaries form by different processes during plastic deformation. The boundary misorietation angles and boundary spacings of GNBs and IDBs evolve differently as a function of strain. For each type of boundary, the coefficient b D av av θ ⋅ remains constant and can be used to differentiate each type of boundary. 52 1.7 Zirconium 1.7.1 Crystallographic Properties Zirconium has two allotropic phases at atmospheric pressure. The alpha phase is close packed hexagonal and is stable up to 862 o C. The beta phase, which is stable from this temperature to the melting point (approximately 1860 o C), is body centered cubic. Their crystallographic parameters are summarized in Table 2. Table 2: The crystallography of Zirconium Alpha Beta Crystal structure Close packed hexagonal Body centered cubic Schoenflies space group 4 6h D 9 h O Coordination number 6,6’ 8’ Lattice constants a 3.230±0.002 Å c 5.133±0.003 Å 3.62 Å – Axial ratio c/a 1.593 – 1.7.2 Deformation Mechanisms The deformation mechanisms of hexagonal-close packed (hcp) metals are more complex than those of face-centered cubic (fcc) and body-centered cubic (bcc) metals. It is well known that in fcc metals, 12 possible slip systems, belonging to the same primary slip mode, {111} < 0 1 1 − >, are distributed symmetrically. Slip in these planes is the main deformation mechanism. Although hcp metals normally form what is considered a single class, the c/a axial ratio varies from one metal to the other and can attain a value smaller or larger than the ideal spherical packing, as shown in Fig 34. Because of the various c/a 53 ratios, the angular relationships between the corresponding crystallographic planes are different in various hcp metals, which thus have different primary slip systems. On the other hand, owing to the small number of possible slip systems within one slip mode and to their asymmetrical distribution over the reference sphere, various deformation mechanisms (including different slip modes and twinning) usually become operative in order to achieve compatibility deformation between neighboring grains. Fig 34: The c/a axial ratio for various hcp metals. 1.7.2.1 Slip Modes In the α structure of zirconium, in which the c/a ratio (1.593) is smaller than the ideal value (1.633), the prism plane ( 0 1 10 − ) is, on average, more densely packed than the basal plane. < 0 2 11 − > directions are the favorable slip directions. That is, within the class of perfect dislocations, those dislocations with Burgers vectors of the 1/3< 0 2 11 − > type ( a dislocations) are favored over dislocations of the <0001> type ( c dislocations) and these, in turn, over dislocations of the 1/3< 3 2 11 − > type ( a c+ dislocations). Therefore, in 54 α zirconium slip takes place usually on the { 0 1 10 − } first-order prism planes along the < 0 2 11 − > a directions [104-106], as shown in Fig. 35a. Slip has, however, also been observed in the same direction on the (0001) basal plane [106-108] (Fig. 35b). Under constrained deformation and at high deformation temperatures, slip occurs on the first and second-order pyramidal planes { 1 1 10 − } and { 1 2 11 − } along in < 3 2 11 − > directions [109-111] (Fig. 35c). Fig 35: Slip systems in α-zirconium. 55 1.7.2.2 Twinning Modes Twinning in zirconium can be divided into two categories: tensile twinning and compression twinning. Tensile twins include { 2 1 10 − }< 011 1 − > twins and, less commonly, { 1 2 11 − }< 26 1 1 − − >, activated under a tensile stress along the direction of the c axis [107, 112-114, 116-120]. { 2 2 11 − }< 23 1 1 − − > twinning [104, 107, 113-116, 117-120] and, at elevated temperatures, {1011}<1012> twinning [110, 119], are observed under compression in a <0001> − c direction. Additionally, an as yet incompletely defined twinning mode on the { 3 2 11 − } pyramidal plane is mentioned in the literature [105, 107, 112-114]. Although twinning can, in general, not provide large strain, it may play an important role in deformation. Lattice regions initially unfavorably oriented for slip may be reoriented more favorably by twinning, thus allowing the possibility of larger strain. Moreover, twinning could activate new slip systems in regions of local stress concentrations such as twin boundaries. 1.7.3 Recrystallization Mechanisms Zhu et al. [121] carried out a study on SRX of a Zr-Hf alloy sheet, deformed by plane strain compression and annealed, and revealed that small recrystallized grains appeared in highly deformed zones by the combination of continuous or discontinuous recrystallization. An important indication distinguishing contiounus recrystallization and dynamic recrystallization is the misorientation of small recrystallized grains from the parent structure. In the case of continous recystallization, the deformation microstructure 56 evolves through extended recovery reactions. Thus, subgrain boundaries form within the starting polycrystals with misorientations of just a few degrees of orientations. A gradual misorientation increase takes place with increasing deformation. However, in the case of dynamic recrystallization, the nucleation and growth of grains with new orientations, i.e., the formation and long-range migration of high angle boundaries, takes place. In general, the recrystallized grains are rotated 30 o around the (0001) basal pole with respect to the matrix. For example, a typical rolling texture is formed by basal planes slightly rotated away from the rolling plane and < 0 1 10 − > directions parallel to the RD. If this rolled material is heated at high temperatures, so that dynamic recystallization occurs, in the final texture < 0 2 11 − > directions will be parallel to the RD. Kassner et al. [64] suggested that GDX was the predominant recrystallization mechanism on pure Zr creep deformed up to large strains at high temperature (600 o C~800 o C, 0.5T m ). This statement is based on their observation of the continual thinning of the original grains, of the serration of the initial high-angle boundaries, of bimodal misorientation distributions and of the presence of boundaries with very low misorientations at high strains [122]. 1.7.4 Texture in Zr and Zr Alloys The texture of Zr and Zr alloys has been studied previously, and the main components are now well established [121,123-130]. Fig 36 illustrates the main orientations present in cold rolled and rolled and annealed Zr. The texture of cold rolled Zr and Zr alloys consists on the alignment of the > < − 0 1 10 directions along the rolling direction (RD) and that the c-axes are tilted 25 o to 40 o from the normal direction (ND) 57 toward the transverse direction (TD) along the ND-TD plane. This deformation texture is represented by the texture component, > < − 0 1 10 } 0001 { ({rolling plane}<rolling direction>, see Fig 36a). This texture is retained after heat treatments at T < 600 o C. Annealing at T ≈ 600 o C for about 1 h leads to a 30 o rotation of the lattice around the c- axis, so that > < − 0 2 11 directions become parallel to the RD (Fig 36b). This annealing texture is represented by the texture component, > < − 0 1 10 } 0001 { ({rolling plane}<rolling direction>). More severe annealing causes a rotation of the c-axis toward the ND. Fig 36: Most common rolling and annealing textures in Zr and Zr alloys [121,123-130]. (a) Rolling texture: c-axis are tilted 25−40 o from the ND toward the TD, along the ND- TD plane, and > < − 0 1 10 directions are aligned with the RD. This texture is retained when annealing at T < 600 o C. (b) Annealing texture of a rolled sheet (approx. 95% cold rolling). Annealing treatments at T ≈ 600 o C for approximately 1 h lead to a rotation of the crystallites around the c-axis so that a > < − 0 2 11 direction becomes parallel to the RD. More severe annealing (higher T, longer time) leads to the gradual rotation of the c-axis toward the ND. 58 CHAPTER 2. EXPERIMENTAL PROCEDURE 2.1 Accumulative Roll Bonding of Pure Zr 2.1.1 Selection of the Severe Plastic Deformation Method to Be Used Top-down techniques are the most appropriate to fabricate large quantities of nanostructured materials. The various SPD methods described in Section 1.3 have different advantages and disadvantages. Equal Channel Angular Pressing (ECAP) is a complex processing route that usually requires at least 6 passes in order to fabricate ultrafine structures. For example, Zhu et al. [131] have succeeded to produce ufg Ti by applying 8 to 12 ECAE passes at high temperature followed by several cold rolling steps. Moreover, ECAP does not allow fabricating large parts. High Pressure Torsion (HPT) has the ability to provide very large strains. However, it is also not possible to produce large parts using this technique. More recently, Dinda and coworkers [132] synthesized nanocrystalline Ni, Ti and Zr foils at room temperature by a rolling and folding technique (R&F) involving 60 to 80 low strain cold rolling passes. Although very small grain sizes were obtained, this R&F process was very complicated to carry out and bonding between the layers was not achieved. In this work, the method selected for the fabrication of ultrafine grained materials is accumulative roll bonding. There are several reasons that justify this choice: (a) Large quantities of bulk ultrafine grained materials with sheet shape can be fabricated. (b) ARB could be easily scaled up to meet industry needs. 2.1.2 Material Selection The material chosen for this investigation is the hexagonal closed packed (hcp) Zr. This material has important applications in today’s world. For example, zirconium and 59 zircaloy are widely used in nuclear reactors as fuel rods. They have excellent biocompatibility and are potentially candidate materials for implants. Accumulative roll bonding has been used almost exclusively to produce ultrafine grained cubic materials, such as Al and Fe-based alloys [15-17, 24]. The efficacy of this process to fabricate bulk ufg hcp materials is, however, still not well known. Good candidate materials for ARB should be ductile and easy to bond at moderate temperatures (T < 0.4T m ), since high processing temperatures would cause recrystallization and grain growth, thus cancelling the effect of the accumulative strains. Zr has good potential to be processed by ARB due to its good low temperature ductility. Other hcp materials, such as Mg, have been discarded due to their brittleness at low temperatures and their rapid oxidation, which prevents good bonding [31]. 2.1.3 ARB Process The complete ARB processing route designed to fabricate bulk ufg Zr is illustrated in Fig 37. The initial material was first annealed at 600ºC for 40 min in order to homogenize the microstructure (for further details see section 3.1.2). Two sheets of the annealed material, with dimensions of 200 ×21 × 5.2 mm 3 , were first separately rolled at room temperature with a reduction in thickness of 25% (equivalent true strain of 0.33). The final thickness of the rolled sheet was 3.9 mm. Subsequently, two pre-strained sheets were stacked and roll bonded using one pass with approximately 75% reduction (ε=1.6~1.64, ε=28 s -1 ) at 327ºC (~0.28 Tm). The material was able to accommodate this very large strain without any evidence of cracking. The final sheet thickness was 1.9 mm. Delamination was not observed. The sheet produced after this first ARB pass, termed ARB1, was then cut into four pieces, which were stacked and subjected to a second 60 rolling pass with approximately 75% thickness reduction at 327ºC (ARB2). The full cycle was repeated (ARB3). The true strains accumulated after the ARB1, ARB2 and ARB3 cycles were 1.97, 3.57 and 5.17, respectively. Rolling was performed in a Carl Wezer rolling machine (Fig 38), furnished with two 13 cm-diameter rolls that rotate at 49 rpm (Centro Nacional de Investigaciones Metalúrgicas, Spain). Before each pass, the samples were heated at the rolling temperature for 10 min. The surface of the sheets was ground using a 400 grit SiC paper and degreased by acetone before stacking for ARB. The rolls were neither lubricated nor heated. After each roll-bonding step, the sheet was cooled in water. The main novelty of the present ARB route with respect to that devised by Tsuji and coworkers is the use of large thickness reduction (up to 75%) per pass, in order to speed up grain refinement and reduce the total number of ARB cycles. 61 Fig 37: Schematic of the processing route designed to fabricate bulk ultrafine grained Zr. 1.9 mm 327ºC ARB 2 327ºC ARB 3 Stacking 327ºC ARB 1 Stacking Annealed material Pre-strained material Room temperature 5.2 mm 3.9 mm 62 Fig 38: The two-high rolling mill used for accumulative roll bonding. 2.1.4 Microstructure Characterization An exhaustive characterization of the microstrucutre of the samples processed by ARB was carried out in order to evaluate their quality, i.e., to investigate the final grain size, the nature of the boundaries present, as well as the through-thickness grain size and texture homogeneity. A good knowledge of the main microstructural features is also crucial in order to understand the mechanical behavior of the bulk ultrafine grained materials processed. A combination of techniques was used for microstructural characterization. Optical microscopy (OM) was employed to examine the microstructure of the initial material, with coarse grain size, as well as to assess qualitatively the bonding quality after various 63 ARB passes. Transmission electron microscopy (TEM) allowed the measurement of the (sub)grain size, which was estimated by the linear intercept method, counting all boundaries, as well as the boundary misorientations. The disadvantage of TEM is that it only allows us to obtain information from a very small area. Therefore, electron backscatter diffraction (EBSD) was performed in order to get more statistically significant microstructural information. In particular, EBSD provides a good description of the “real” grain size, i.e., that is calculated by counting only boundaries with misorientations larger than 15 o , of the boundary misorientation distributions, and the microtexture. The drawback of EBSD is that it does not detect subgrain boundaries with misorietation smaller than about 2 o . These very low misorietations must be analyzed by TEM. OM was performed on an Olympus microscope (Centro Nacional de Investigaciones Metalúrgicas, Spain). Sample preparation consisted of grinding on increasingly finer SiC papers of 600, 1200 and 1600 grit sizes, followed by polishing with 9 μm and 6 μm diamond paste. Final surface finishing was performed with a suspension of 0.05 μm silica particles. The samples were additionally etched with a solution of nitric acid (45%), distilled water (45%) and hydrofluoric acid (10%) for optical microscopy. Electron backscattered diffraction (EBSD) were carried out both in a W-filament JEOL SEM operated at 20 kV using the Oxford Instruments INCA (Crystal) software (Universidad Politécnica de Valencia, Spain) and in a LEO FEG SEM equipped with the HKL software (Max Planck Institute for Metals Research, Stuttgart, Germany). Sample preparation for EBSD examination was identical to that for OM. Much care was 64 taken in order to get a very good surface quality, with very low roughness, which was crucial to obtain good EBSD data. TEM was performed in a 200 KV 2010 JEOL microscope (Centro Nacional de Investigaciones Metalúrgicas, Spain). TEM samples were usually made on two planes, the ND-RD plane and RD-TD plane, in order to verify the homogeneity of the microstructure. The sample preparation procedure for TEM was complex and included many steps. First, a piece of sample was both-side ground down to 0.5 mm using 600 grit paper. Second, 1200 and 1600 grit papers were used to thin the sample down to a thickness of 0.3 mm. Then, disks with 3 mm in diameter were cut and cleaned using ethanol. A final thinning step to electron transparency was carried out at room temperature in a double jet Struers Tenupol electropolisher at 25 V and using as an electrolyte of 10% perchloric acid and 90% acetic acid. In all cases, special care had to be taken during the cutting and grinding steps to avoid introducing mechanical damage (dislocations and microtwins). Grain boundary misorietation measurements were performed by Kikuchi line analysis in TEM. The misorietation angle between two crystallites can be qualitatively determined using the following equation [133]: L R Δ = Δθ (8) Where, Δθ is the misorietation angle between two crystallites; ΔR is the distance between two main poles in Kikuchi patterns before and after tilting the electron beam which is located on the grain boundary; L is the camera length of transmission electron microscope. 65 The measuring procedure consists of the following steps. First, the electron beam is moved on a grain boundary, the misorietation angle of which is to be measured. Then, a Kikuchi pattern with a main pole located at A can be obtained, schematically shown in Figure 39(a). Here, the main pole means the pole with the most number of Kikuchi bands intersecting. Subsequently, the electron beam is tilted slightly (Figure 39(c)) in order to obtain another Kikuchi pattern in which the main pole is located on B (Figure 39(b)). The position of A and B are recorded and the distance (ΔR) between these two positions can be calculated. Compare with the value of ΔR corresponding to the case of Δθ =15 o , the grain boundary can be classified as a high angle grain boundary (misorietation angle greater than 15 o ) or a low angle grain boundary (misorietation angle smaller than 15 o ). The method to precisely determine the misoritation angles of grain boundaries is described in Appendix III. Fig 39: The schematic drawing for measuring the misorietation angle of grain boundaries in TEM. 66 2.1.4.1 Texture Characterization Special emphasis was placed in analyzing the texture of the initial and processed samples, since it provided very useful information about the grain refinement mechanisms that took place during processing. Moreover, since Zr is an hcp material, the texture may have a significant influence on the mechanical properties. Texture analysis was performed by two complementary techniques. X-ray diffraction allowed to measure the “macrotexture”, i.e., the orientation that predominates in a large volume of material (approximately 3mm 2 × 50μm in depth), as well as the approximate volume fraction associated with each component. In this work, macrotexture measurements were performed in an X-ray diffractometer (SIEMENS D5000) with an open Euler ring. Figure 40 shows a diagram of the device and a photograph of diffractometer available at Centro Nacional de Investigaciones Metalúrgicas (Mardid, Spain). The sample with a flat surface is placed in a special holder in the goniometer so that the surface has always radiated an X-ray line tangent to the circle which goes through the X-ray source and the detector. The sample can be rotated about three perpendicular directions φ, χ and ω (shown in Figure 40) to ensure that the angle of incident X-ray beam on the sample surface satisfies the Bragg reflection condition for (h,k,l) planes. The detector receives the reflected diffraction intensity in each of positions of the sample. The information is processed by software DIFFRAC PLUS that calculates the various pole figures from the diffracted intensity data in each exact Bragg position. Sample preparation for X-ray diffraction was identical to that for EBSD, explained in the previous section, except that the etching step was not necessary. 67 EBSD, in turn, provides information about the spatial distribution of orientations in the material (“microtexture”). Sample preparation for EBSD was described in the previous section. Fig 40: X-ray Diffractometer available at the National Center for Metals Research (Madrid, Spain). 68 2.1.5 Mechanical Testing 2.1.5.1 Tensile Testing Room temperature tensile tests were performed at 10 -3 s -1 in an Instron testing machine (Centro Nacional de Investigaciones Metalúrgicas, Spain), shown in Figure 41. Tensile specimens consisted of dogbone flat tensile coupons of 15 mm gage length machined out of the annealed and rolled samples. The specimen dimensions are listed in Table 3. A schematic illustration of tensile specimens is represented in Fig 42. Strain was recorded by a linear variable displacement transducer (LVDT) attached to the load frame. The elastic part of the strain was subtracted to the total strain during the data analysis. Fig 41: Instron tension machine used in this study (Centro Nacional de Investigaciones Metalúrgicas, Madrid, Spain). 69 Table 3: Dimensions of the tensile specimens Dimensions (mm) Material Annealed Zr specimen from the initial material Specimen after one ARB cycle Specimen after three ARB cycle Gage length 10 10 10 Thickness 1.82 1.85 1.8 width 4.03 4.02 4.07 Fig 42: Schematic illustration of the dogbone flat tensile specimens. 2.1.5.2 Shear Testing The strength of the bond between different layers was measured by shear testing. A schematic of testing method is shown in Fig 43. Tests were performed in dogbone flat specimens, subjected to the whole processing route up to the ARB3 step (ε = 5.25), which have 32 layers through the thickness. Two narrow slots were milled at about 14 mm from 10 16 34 12 4 r = 3 Thickness 9 70 each end of the coupons to carefully controlled depths. The layers between those two layers marked by arrows, shown in Fig 43, were subjected to nearly pure shear deformation. Variation in the depths of two slots in specimens allows testing the bond strength at the mid-thickness (Group #1) and near the surface (Group #2). The specimen dimensions are listed in Table 4 and real specimens are shown in Fig 44. Tension tests, conducted at a strain rate of 10 -3 s -1 , were then performed at room temperature on an Instron 8521 (University of Southern California, USA) using a gripping system and a set- up shown in Fig 45 (a) and (b) respectively. The force on the load cell was recorded and stored in a personal computer until fracture. A control and data acquisition program for this system using Labview codes is attached in Appendix I. Fig 43: Schematic diagram of the test sample to measure the shear strength of the bond. 71 Table 4: Dimensions of the specimens for bonding strength testing Dimensions (mm) Group #1 Group #2 Specimen #1 #2 #3 #4 Depth of slot 1 (L1) 1.13 1.14 1.1 1.69 Depth of slot 2 (L2) 1.21 1.14 1.13 0.5 Width of slot 1 (W1) 0.65 0.49 0.48 0.56 Width of slot 2 (W2) 0.62 0.49 0.47 0.68 Thickness of the specimen (t) 2.04 1.98 2 2.01 Width of the specimen (w) 1.93 1.98 1.99 1.93 Spacing between two slots (L0) 1.9 2 2.01 1.99 Length of the gage section 6 9.6 9.1 6 72 (a) (b) Fig 44: Specimens used to measure the strength of the bond. (a) Front view and (b) Side view. (a) (b) Fig 45: Gripping system (a) and set-up (b) utilized to measure the strength of the bond. 73 The strength of the bond can be understood as the maximum engineering shear stress between two layers, which can be calculated using the following equation (9): 0 * max max L w F = τ (9) Where, τ max represents the maximum engineering shear stress that a layer can hold; F max is the maximum load during the testing; w is the width of the gage section and L0 is the spacing between two slots. 2.1.5.3 Hardness Testing Microhardness measurements were performed through the thickness of the ARBed samples at steps of 50 μm using a Misawa Seiki Vickers hardness tester (Centro Nacional de Investigaciones Metalúrgicas, Spain). A load of 10 g was applied for 15 s at each location. The sample surface was etched prior to performing the indentations to remove damage or artifact. The measurements were done immediately after etching in order to avoid the formation of an oxide layer. 2.2 Biocompatibility of Zr Processed by Large Strain Rolling In general SPD-processed microstructures fall within the UFG regime. The enhanced corrosion resistance [134] and mechanical [3, 5, 135, 136] properties of UFG metallic materials have been reported in previous investigations. However, little is known of the effect of grain size on the biological response to metals. Enhanced osteoblast adhesion has recently been reported in “nanophase” Ti, Ti6Al4V, and CoCrMo produced by traditional powder metallurgy procedures, compared to conventionally sized metals [137]. The resulting compacts exhibited high porosity (5- 74 10%), with interparticular voids of up to about 1 μm. These microstructural features alter the surface roughness and, therefore, may result in the change of biocompatibility of ultrafine grained materials. Conventional Zr metal shows acceptable mechanical strength and good biocompatibility; thereby it is a material of interest for surgical implants [138-142]. In vivo evidence has indicated that zirconium implants exhibit good osseointegration [143, 144] and studies comparing zirconium and titanium implants showed that the degree of bone-implant contact is higher in the case of zirconium [145, 146]. In vitro biocompatibility of zirconium surfaces has scarcely been explored [147]. In the present work, the in vitro biocompatibility of severely deformed UFG Zr and conventional Zr were investigated by evaluating the behavior of human osteoblastic cells. For comparative purposes, a commercial Ti6Al4V alloy, which is known to be a highly compatible metallic biomaterial, was also investigated. 2.2.1 Fabrication and Characterization of Materials The material used for this biocompatibility study is commercially pure Zr. Its chemical composition and microstructure will be described in Section 3.1.1. Two different microstructures (coarse and fine) were prepared out of the initial material. First, a coarse microstructure was developed by annealing the initial pure Zr at 700 o C for 2h. Next, an UFG structure was fabricated by severely rolling the initial material at 330 o C (~0.27 T m ) with a 75% reduction in one pass. This is equivalent to applying a true strain of approximately 1.6 at a strain rate of 28 s -1 . Hereafter, the conventional and UFG structure specimens will be described as Zr and UFG Zr, respectively. For comparative purposes, discs of Ti6Al4V alloy, hereafter Ti64 alloy, were used. 75 Discs of about 20 mm diameter and 1-2 mm thick were used for this investigation. Surface preparation consisted of consecutive grinding on 600, 1200 and 1600 grit SiC paper, followed by polishing using 6 and 1 μm diamond paste. Final surface polishing was performed with a suspension of silica particles, 0.05 μm in size. Optical (OM) and scanning electron microscopy (SEM) was used to reveal the grain size of the conventional coarse materials. The grain size was measured in the optical micrographs by the linear intercept method. Sample preparation for OM included mechanical grinding, as described above, and final etch-polishing with an etching solution (45% HNO3, 45% distilled water and 10%HF). In the case of the Ti64 alloy, the grain size was determined on specimens etched with a mixture of HF (30%) and HNO 3 (70%) at around 0 o C. SEM micrographs were obtained using a field emission gun (FEG) emitter using both the secondary (SEI) and the backscattered electron signals (BEI). A 200KV 2010 JEOL transmission electron microscope with a double tilt stage was used to characterize the microstructure of UFG Zr. Sample preparation was almost the same as what was described for TEM sample preparation in Section 3.1.3. However, a Struers Accutom 5 saw (Struers, Willich, Germany) using a 400-μm-thick diamond blade was used to cut 300-μm slices. The specimens were cut using the lowest settings for the blade rotation speed and the sample advancement, at 300 rpm and 0.005mm/s, respectively, in order to minimize mechanical damage. Three mm-diameter discs were cut from the slices with a South Bay brass core-drill abrasive cutter and mechanically ground with 1600 grid abrasive paper to thicknesses ranging between 120 and 200 μm. Final thinning was carried out by electropolishing at ambient temperature with a Struers Tenupol-3 double jet electro-polisher with a voltage of 25 V, a medium-high flow rate, 76 and using an electrolyte of 90% acetic acid and 10% perchloric acid. The grain size was determined on several TEM micrographs by applying the linear intercept method and counting all boundaries. An Autoprobe CP commercial atomic force microscopy (AFM) (Veeco Instruments, CA, USA) was used to analyze surface topography. Both the cantilever and the tip are made of silicon (Supersharp tips, Mikromasch, Estonia). The tip and both sides of the cantilever are coated by a continuous film of Cr (first layer) and Au (second layer), each of 20 nm width. The nominal (triangular) cantilever force constant is 0.35 N/m. The resonance frequency of the cantilever is 28 kHz. The tip apex has a radius of curvature < 50nm, having a full cone tip angle < 30 o . The same tip was used throughout the entire study. All measurements were performed in non-contact mode, that is, by choosing a working frequency of the oscillating cantilever slightly higher than the resonance frequency. Water contact angles on samples were measured with an automatic goniometer G2 (Krüss, Germany). Images of the drop were taken with a video camera, and the contact angles of symmetric drops were evaluated with the instrument’s drop shape analysis software. Water used for measurements was distilled and deionized using a water purification system (Milli-Q Plus Millipore, Billerica, MA, USA). Contact angles values are the average result of at least eight different symmetrical drops. The hardness of the investigated materials was measured using a Wilson Wolpert Vickers indenter (Wolpert, Heerlen, The Netherlands). Indentations were performed on polished surfaces with a 1-kg load and 15-s dwell time. 77 Discs of Zr, UFG Zr and Ti64 of 20 mm diameter were used for cell culture experiments. All specimens were washed in running water and rinsed extensively in ethanol. The discs were sterilized under ultraviolet light before use in cell culture experiments. 2.2.2 Cell Culture Cells were maintained at 37 o C under 5% CO 2 and 95% air in a humidified incubator. Human osteosarcoma Saos-2 cells (ECACC, Salisbury, Wiltshire, UK) were grown in DMEM medium supplemented with 10% (v/v) heat-inactivated foetal bovine serum (FBS), 500 UI/ml of penicillin and 0.1 mg/ml of streptomycin. Human mesenchymal stem cells (hMSC) from bone marrow were purchased from Cambrex Bio Science (Verviers, Belgium). These cells were maintained in growth medium and switched to the osteoblastic phenotype by incubation in osteogenic induction medium (both from Cambrex). 2.2.3 Attachment Assays Saos-2 cells were seeded on the three metallic surfaces in 12-well plates (1.5 × 10 5 cells/well) for 0.5, 3 and 24 h. Attached cells were incubated in culture medium containing 1.2 × 10 -3 M (=mol/liter) 3-(4,5-dimethylthiazol-2-yl)-2,5-diphenyltetrazolium bromide (MTT, Sigma, Madrid, Spain) for 4h and after carefully removing the medium, the purple formazan product was dissolved in dimethyl sulfoxide. The absorbance was read at 570 nm using a spectrofluorimeter (Victor2 Wallac 1420 Multilabel Counter, Perkin Elmer, Turkin, Finland). For actin staining, attached cells cultured for 3h were fixed with 4 % formaldehyde in phosphate buffered saline (PBS) and permeabilized with 78 0.1% Triton X-100 in PBS. To visualize actin filaments, cells were then stained with PBS containing 4 × 10 -7 M Phalloidine-TRITC (Sigma) and observed using a fluorescence microscope (Leica AF6000, Wetzlar, Germany). High resolution, fluorescence images were captured and digitally deconvolved using Huygens software (Scientific Volume Imaging, Hilversum, The Netherlands). 2.2.4 Cell Spreading Saos-2 cells were seeded on the three metallic surfaces in 12-well plates (1.5 × 10 5 cells/well) and cultured for 24 h. Cell spreading was determined after staining cells with phalloidine-TRITC. High-resolution fluorescence images from representative fields of each sample were captured and digitally deconvolved using Huygens software. A total of 50 cells randomly selected from five representative images per sample were manually outlined, and cell areas were measured using ImageJ v1.34 image analysis software. 2.2.5 Immunofluorescence Assays Saos-2 cells were seeded on the three metallic surfaces in 12-well plates (1.5 × 10 5 cells/well) and cultured for 24 h. Cells were fixed with 4% (w/v) formaldehyde in phosphate buffered saline (PBS) and permeabilized with 0.1% Triton X-100 in PBS. For immunostaining, cells were blocked in PBS containing 2% bovine serum albumin (BSA) and 0.05% Tween 20 and then stained with a mouse anti-human fibronectin monoclonal antibody (Chemicon, Harrow, UK) or a rabbit anti-human β-tubulin (Santa Cruz Heidelberg, Germany) diluted 1:50 in 1% BSA in PBS. After washing with 0.05% Tween in PBS, cells were incubated with goat anti-mouse Alexa-Fluor 488 or goat anti-rabbit Alexa-Fluor 594 secondary antibodies (both from Molecular Probes, Leiden, Holland) di- 79 luted 1:1000 (v/v) in PBS containing 1% BSA. For double-staining, cells were additionally incubated with phalloidine-TRITC as described above. After washing with 0.05% Tween in PBS, cells were examined using a fluorescence microscope. High resolution, fluorescence images from representative fields were captured and deconvolved using Huygens software. 2.2.6 Fibronectin Levels Saos-2 cells were seeded on the three metallic surfaces in 12-well plates (1 × 10 5 cells/well) and cultured for 24 h. Media were collected, centrifuged at 1200g for 10min, supplemented with a mixture of proteases inhibitors (17.5 μg/ml phenylmethylsulfonyl fluoride, 1 μg/ml pepstatin A, 2 μg/ml aprotinin, 50 μg/ml bacitracin, all from Sigma) and frozen at −80 o C. Cell layers were washed exhaustively with PBS and extracted with 5 × 10 -2 M Tris-HCl pH 8.0, 5 × 10 -1 M NaCl and 1% Triton X-100. A specific enzyme immunoassay kit (Takara, Gennevillears, France), with a sensitivity of 12.5 ng/ml, was used to measure human fibronectin in cell layers and media. Data were normalized to the total protein amount in cell layers, measured by the Bicinchoninic Acid (BCA) protein assay (Pierce, Rockford, IL, USA), using BSA as a standard. 2.2.7 Cell Viability Saos-2 cells were seeded on the three metallic surfaces in 12-well plates (1 × 10 5 cells/well) and cultured for 1, 4 and 7 days. Cell viability was assessed using the alamarBlue assay (Biosource, Nivelles, Belgium). This assay incorporates a redox indicator that changes colour in response to metabolic activity. After washing with PBS, cells were incubated in DMEM containing 10% alamarBlue dye for 4 h and after 80 excitation at 530 nm, the fluorescence emitted at 590 nm was quantified using a spectrofluorimeter. 2.2.8 Alkaline Phosphatase Activity and Mineralized Nodule Formation Human mesenchymal stem cells (hMSC) were expanded to passage 7 (7 th generation) in growth medium. Then, cells were seeded on the three metallic surfaces in 12-well plates (1.1 × 10 4 cells/well) and cultured in osteogenic induction medium for 18 days. The media were replaced every 3 or 4 days. hMSC cultured in growth medium were used as controls. Cell layers were washed exhaustively with PBS and extracted as described above. Alkaline phosphatase (ALP) activity was assessed in cell layers by determining the release of p-nitrophenol from p-nitrophenylphosphate at 37 o C at a pH of 10.5. Data were normalized to the total protein amount in cell layers. The degree of mineralization was determined using Alizarin Red staining. Briefly, cells were fixed with ethanol and stained with 4 × 10 -2 M Alizarin Red in deionised water (adjusted to pH 4.2). Following rinsing with PBS, the bound stain was eluted with 10% (w/v) cetylpyridinium chloride and the absorbance at 562 nm was measured using a spectrofluorimeter. 2.2.9 Statistical Analysis The data are presented as mean ± SD (standard deviation) of several independent experiments. Analyses of variance (ANOVAs) with Bonferroni’s correction were used for post-hoc comparisons. The p-values<0.05 were considered to be statistically significant. 81 CHARPTER 3. RESULTS 3.1 Accumulative Roll Bonding of Pure Zr 3.1.1 Initial Material The initial material is Zr of commercial purity (99.859%). Its chemical composition is shown in Table 5. The dimensions of the initial plate, bought from the company Haines & Maassen, Germany, are 200 cm in length, 140 mm in width and 5.2 mm in thickness, as shown in Fig 46. Table 5: Composition of the Zr material under study (ppm) Fe Mn Hf S Nd Zr 330 27 452 <550 <50 Rest Fig 46: Plate of pure zirconium. Fig 47 illustrates the microstructure of the initial material. A large number of twins is observed in the grains. The average grain size, measured by the linear intercept method, 30 mm 82 is 13 μm if twin boundaries are counted and 31 μm without taking into account twin boundaries. Fig 48 shows the grain size distribution in the initial material. As can be seen, a unique peak can be identified near 2.5 μm. The material possesses a weak rolling texture, with c-axes rotated approximately 30 o from ND towards TD (Fig 49). The main texture component is > < − 0 1 10 } 0001 { ({rolling plane}<rolling direction>). This is consistent with the initial material, having undergone previous rolling. Fig 47: Optical micrographs showing the microstructure of the as-received Zr. 83 0 10 20 30 40 50 60 1 3 5 7 9 11 13 15 17 19 21 23 25 27 Linear intercept (μm) Number of grains Fig 48: Grain size distribution in the initial material. Twin boundaries are included. Levels: 1.00 2.00 3.00 4.00 5.00 6.00 Fig 49: Macrotexture of the initial material. 3.1.2 Annealed Material Separate initial tablets were annealed at 600 o C during 10 minutes and 40 minutes in order to obtain a homogenous and appropriately coarse grain size. Figures 50 and 51 illustrate the microstructure resulting from both annealing treatments using EBSD 84 orientation maps. After annealing at 600 o C for 10 minutes, the microstructure shows evidence of the occurrence of both recovery and recrystallization. Very few twins remain present (the remnants of twins are marked by arrows A in Fig 50) after annealing 600 o C during 10 min. The thin lines present in some grains appear to reveal the rearrangement of dislocations and the formation of subgrains (misorientation angle <15 o ). Additionally, recrystallized grains, smaller in size than those present in the initial material, are clearly visible in Fig 50. These small recrystallized grains have two main orientations, i.e. > < − 0 1 10 } 0001 { and > < − 0 2 11 } 0001 { . After annealing at 600 o C for 40 minutes the material is almost fully recrystallized, as shown in Fig 51. No evidence of twinning is apparent after this heat treatment. Fig 52 shows that the microstructure, examined by optical microscopy, after annealing for 40 min. is homogenous, with equiaxed grains and without twins. 85 (a) View along ND: (b) View along RD: Fig 50: EBSD map viewed along the ND (a) and along the RD (b) showing the microstructure after the as-received material was annealed at 600 o C for 10 minutes. Thick lines correspond to boundaries with misorienations higher than 15°, while the thinner lines correspond to boundaries with misorientations between 4° and 15° (boundaries less than 4 o can not be reliably detected here). 86 (a) View along ND: (b) View along RD: Fig 51: EBSD map viewed along the ND (a) and along the RD (b) showing the microstructure of the as-received Zr after annealing at 600 o C for 40 minutes. Thick lines correspond to boundaries with misorienations higher than 15°. 87 Fig 52: Microstructure of the initial material, annealed at 600 o C for 40 minutes, examined by optical microscopy. The grain size of the annealed samples was measured by the linear intercept method using only high angle grain boundaries (misorientation higher than 15 o ). The grain size corresponding to Zr annealed for 10 min. and 40 min. are both 17 μm, which is less than 31 μm, the grain size for as-received Zr, without counting twin boundaries. The grain size distributions are depicted in Fig 53. It is clear that there are no small grains (≤4 μm) present in as-recieved Zr. However, small grains (≤ 4 μm) appear after annealing for 10 min. and 40 min. This suggests that some new grains are nucleated by static recrystallization. The fraction of high angle grain boundaries (misorientation greater than 15 o ) increases from 61% (10 min.) to 92% (40 min.), suggesting that, after the longer treatment, subgrain boundaries annihilate. 88 Fig 53: Grain size distribution of the as-received material without counting twins (a) and after annealing at 600 o C for 10 min. (b) and for 40 min. (c). 89 Fig 54 shows the texture of the Zr after annealing at 600 o C for 10 min. and 40 min. In both cases, the two texture components, > < − 0 1 10 } 0001 { and > < − 0 2 11 } 0001 { , are observed. The macrotexture of the initial material after annealing for 40 min. is shown in Fig 55. The texture of the annealed material is very similar to that of the initial material, i.e., a rolling texture, but with a lower intensity. As concluded from the results obtained above, a homogenous microstructure with an appropriate grain size of 17 μm was obtained after annealing the initial material for 40 min. This annealed material was chosen to be subjected to the following ARB process. Fig 54: Direct pole figures corresponding to rolled Zr after annealing at 600 o C for 10 minutes (a) and 40 minutes (b). 90 Levels: 1.00 2.00 3.00 4.00 5.00 6.00 Fig 55: Macrotexture of the initial material annealed for 40 min. examined by X-ray diffraction. 3.1.3 Microstructure Evolution at the Sheet Mid-thickness during ARB The annealed pure Zr was processed using the route described in Fig 37. The microstructural evolution of the regions located at the mid-thickness of the sheet during all the ARB steps (Figs 56-60) will be described in detail in the following paragraphs. The subsequent section will examine the through-thickness microstructure and the presence of texture gradients in the ARBed sheets. Fig 56 shows a series of EBSD maps of the ND-RD plane in the sample rolled at ambient temperature with a 25% reduction (ε = 0.33) and in the ARB1 (ε = 1.97, T = 327 o C), ARB2 (ε = 3.57, T = 327 o C) and ARB3 (ε = 5.17, T = 327 o C) samples. With increasing rolling strain and, in particular, after the ARB1 step, c-axis become tilted approximately 25 o from the ND and in all cases > < − 0 1 10 directions become parallel to the RD, as is commomly observed in rolled Zr [121,123-130]. The stabilization of a rolling texture in the areas examined is consistent with the fact that they were mostly subjected to plane strain (strain tensor: ε RD,RD > 0; ε TD,TD = 0; ε ND,ND = −ε RD,RD ; other ε ij = 91 0). The texture of the ARB processed samples, measured on a section parallel to the rolling plane and located approximately at the mid-thickness of the sheets, is shown in Fig 57 by means of the (0002), (10-10) and (1120) X-ray pole figures. A typical rolling texture, with {0002} planes slightly rotated away from the normal direction (ND) towards the transverse direction (TD) and > < − 0 1 10 directions parallel to the rolling direction (RD), was formed after ARB1 step (Fig 57a). With increasing rolling strain, the intensity of the rolling component increases further in the ARB2 and ARB3 samples (Figs 57b and 57c), which is quite consistent with observations in EBSD discussed previously. The intensification of a typical rolling texture during successive cycles of accumulative roll bonding suggests that grain refinement takes place mainly by geometrical thinning and grain subdivision [148] as is often the case in materials processed by severe plastic deformation methods [149]. After the first room temperature pass with 25% reduction (Fig 56a), the grains elongate along the RD. Very few twins are apparent and crystallographic slip seems to be the predominant deformation mechanism during the first stages of deformation. A large fraction of equiaxed grains has already formed after ARB1 step (Fig 56b). A reasonably homogeneous microstructure develops following the ARB2 and ARB3 steps (Fig 56c and d). The variation of the grain thickness (along the ND) and the length (along RD) with increasing strain, in the areas illustrated in Fig 56, is plotted in Fig 58(a) together with the values predicted when considering only geometrical thinning. It must be noted that “grain” refers to a region fully surrounded by HABs (θ> 15 o ). At all strain levels the grain length is significantly smaller than the predicted values (Fig. 58(a)). This suggests 92 that grain subdivision, possibly by the formation of geometrically necessary boundaries (GNBs), occurs. The grain thickness and length remain fairly constant after the ARB2 step (ε = 3.57). After the ARB3 step a bulk structure with an average grain size of about 400 nm is obtained. The grain aspect ratio increases up to 1.6 after the first pass with 25% reduction at room temperature and then decreases, reaching a value close to 1 after ARB2. Finally, it increases slightly up to 1.4 after ARB3. These values are significantly smaller than those reported for face centered cubic (fcc) materials processed by accumulative roll bonding to similar strains [150]. Figs 59 and 60 are TEM micrographs that show the (sub)grain structure after one, two and three cycles of ARB along the RD-ND plane (Fig 59) and along the RD-TD plane (Fig 60), respectively. During the early stages of deformation (sub)grains appear elongated when viewed along the RD-ND plane, and have an equiaxed cross-section at the RD-TD plane. However, after ARB3 the (sub)grain structure becomes fairly equiaxed. The average (sub)grain size stabilizes at 320 nm. The grain size obtained after the ARB3 step (~400 nm) is quite similar to the measured (sub)grain size. Fig 61 describes the grain boundary structure in the areas shown in Fig 56. All the boundary misorientation measurements shown in Fig 61 were performed by EBSD and thus only boundaries with misorientation larger than 2 o ~ 4 o are detected. The evolution of the fraction of HABs with strain is shown in Fig 61a and distribution of LABs (2 o < θ < 15 o ) is plotted in Fig 61b. The fraction of HABs, equal to 92% in the starting material (Fig. 61a), decreases down to about 30% after one pass with 25% reduction at room temperature due to the development of a substructure (Fig. 61b). With increasing strain the fraction of HABs increases and thus, the relative number of LABs decreases. After 93 the ARB3 step, the fraction of high-angle boundaries is 88%, almost the same as that present in the initial annealed material. Boundary misorientation measurements were also performed in the ARB3 sample by Kikuchi line analysis in the TEM. Here, boundaries with misorientations between 0 o and 2 o can be detected. The fraction of HABs is found to be 66%, that of boundaries with 2 o < θ < 15 o is 29% and, finally, the fraction of boundaries with θ < 2 o is just 5%. Thus, it may be concluded that (in the areas that were subjected mostly to plain strain deformation during the ARB process) a bulk ufg microstructure, with fairly equiaxed (sub)grains, about 320 nm in size, and a large fraction of HABs (88% from EBSD measurements and 66% from TEM), develops. If only HABs are considered, the grain size is about 400 nm. A typical rolling texture, with {0002} planes slightly away from the ND and > < − 0 1 10 directions parallel to the RD, becomes stable. The variations in the texture and microstructure through the thickness of the processed sheets are analyzed in the following sections. 94 Fig 56: EBSD maps showing the microstructure of pure Zr at different stages of the process and the corresponding ND and RD inverse pole figures. Mapping was performed in areas close to the mid-thickness of a layer at the RD-ND plane. (a) After one pass at room temperature with 25% reduction (ε = 0.33); (b) ARB1 (ε = 1.97, T = 327 o C); (c) ARB2 (ε = 3.57, T = 327 o C); (d) ARB3 (ε = 5.17, T = 327 o C). 95 Fig 57: X-ray pole figures corresponding to Zr following (a) 1, (b) 2, and (c) 3 cycles of ARB. 96 0.01 0.1 1 10 100 1000 10000 0 1 2 3 4 5 6 True strain Size (microns) Grain length (EBSD) Grain thickness (EBSD) (Sub)grain size (TEM) Predicted grain thickness Predicted grain length (a) 0 1 2 3 4 5 0 2 4 6 True strain Aspect ratio Grain (Sub)grain (b) Fig 58: (a) Evolution of grain and (sub)grain size with strain. The lines (dotted and dashed) indicate the expected reduction in grain length and thickness, respectively, with strain, due to geometrical thinning alone. (b) Evolution of the grain aspect ratio along the RD-ND plane. 97 Fig 59: TEM micrographs of Zr after (a) ARB1 (ε = 1.97), (b) ARB2 (ε = 3.57) and (c) ARB3 (ε = 5.17). The plane perpendicular to the transverse direction and all areas are close to the mid-thickness of the rolled sheet. 98 Fig 60: TEM micrographs of Zr after (a) ARB1 (ε = 1.97), (b) ARB2 (ε = 3.57) and (c) ARB3 (ε = 5.17). The plane is parallel to the rolling plane and all areas are close to the mid-thickness. 99 Fig 61: (a) Fraction of high angle boundaries (HABs) with increasing strain. Note that only boundaries with θ > 2 o are observed by EBSD. (b) Frequency distribution of boundaries at lower misorientation (2 o < θ < 15 o ) for different processing steps. 100 3.1.4 Through-thickness Homogeneity in the ARBed Samples 3.1.4.1 Microstructure and Texture at the Surface (RD-TD Plane) Fig 62 illustrates the microstructure and texture at the surface (RD-TD plane) of the sample rolled at room temperature to a reduction of 25%. A rather equiaxed structure develops in the surface regions at this small strain. Some twinning is observed. The average grain length and width are 4.9 and 4.1 μm, respectively, i.e. the aspect ratio is 1.2. The aspect ratio at the mid-thickness was 1.6 (Fig. 56a). The decrease of the aspect ratio and of the grain size when approaching the sheet surface is a common observation in materials processed by rolling and ARB [28]. Grain refinement is enhanced by additional shear at the surface due to friction with the rolls [36], which results in more pronounced grain subdivision. Additionally, a texture gradient develops during rolling. The added shear at the surface leads to the alignment of > < − 0 2 11 directions with the RD. At the mid-thickness regions, > < − 0 1 10 directions align with the RD (Fig 56), characteristic of Zr rolling texture (Fig 36). 101 Fig 62: Microstructure and texture at the surface of the Zr sheet rolled at room temperature using a reduction of 25%. (a) EBSD map along the rolling plane. (b) ND, RD, and TD inverse pole figures. 3.1.4.2 Through-thickness Texture Fig 63 illustrates the texture of the ARB2 samples at several depths. The average grain size at each location is also indicated. The grain size values are very similar at the surface and within the bulk, both close to and away from interfaces. At the surface (Fig. 63a), > < − 0 2 11 directions are clearly aligned with the RD. At the mid-thickness of one of the outer layers (Fig. 63b) the RD is still preferentially aligned with a > < − 0 2 11 prismatic direction, but it tends to spread along the > < − 0 2 11 − > < − 0 1 10 symmetry boundary of the stereographic triangle, probably as a result of the combined plane strain and shear deformation in this area. The interfaces between the different layers can be classified in two types: those (a total of four) formed during the previous ARB1 step (termed here ARB1 interfaces) and those (a total of three) formed during the ARB2 step (termed here 102 ARB2 interfaces). Fig 63c-e illustrates the texture in the vicinity of an ARB1 interface, in the mid-thickness of an inner layer and in the vicinity of an ARB2 interface, respectively. The regions close to ARB1 interfaces were sheared during the first pass at room temperature, with 25% reduction, and were then subsequently deformed by plane strain during the ARB1 and ARB2 cycles. As a consequence of the large plane-strain deformation imposed, a typical rolling texture predominates, with > < − 0 1 10 directions parallel to the RD (Fig 63c) and with {0002} planes tilted away from the ND. The areas at the mid-thickness of an inner layer (Fig 63d), as already described, were subjected to plain strain deformation during the entire processing route, and also possess a typical rolling texture. The areas in the vicinity of the ARB2 interfaces were subjected to shear strains during the first pass at room temperature as well as during the ARB1 step and then to plain strain deformation during ARB2. The latter is apparently not sufficient to “erase” the influence of the previous shear strain, and this is reflected in the texture as the maximum in the RD inverse pole figure is now rotated 17 o away from the > < − 0 1 10 pole along the > < − 0 2 11 − > < − 0 1 10 symmetry boundary. There are a total 31 interfaces in the ARB3 sample, three of which are newly created during this step and 28 of which were present in the previous cycles. Thus, a rolling texture will be present throughout the thickness of the ARB3 sheets at all depths except in the surface regions and in the vicinity of the new ARB3 interfaces, where a “rotated rolling texture” will predominant. Fig 64 illustrates the texture and microstructure at an area traversed by an ARB2 interface in the ARB3 sample. As expected, a clear rolling texture develops. In summary, a reasonably homogeneous microstructure and texture develop during ARB of pure Zr 103 already after two cycles under the conditions described in this study. The microstructure remains stable and homogeneous upon further ARB cycles. Fig 63: Through-thickness texture gradient in the ARB2 sample (ε = 3.57). Inverse pole figures illustrating the texture at several locations: (a) surface; (b) mid-thickness of the first layer; (c) interface formed during ARB1 (ε = 1.97); (d) mid-thickness of an inner layer; (e) interface formed during ARB2. 104 Fig 64: Microstructure and texture of pure Zr in the vicinity of an interface (close to mid- thickness) after the ARB3 step. Measurements performed by EBSD. 3.1.5 Mechanical Behaviors of the ARBed Samples Fig 65 illustrates the stress-strain curves for annealed and rolled (ARB1 and ARB3 samples) Zr. The tests were performed at room temperature and 10 -3 s -1 . The mechanical properties are summarized in Table 6. ARB results in a large increase in the yield strength and UTS without any significant decrease in ductility. In order to investigate the variation of the hardness through the thickness of the ARB1, ARB2 and ARB3 samples, Vickers indentations were made at the mid-thickness, at one-quarter of the thickness and near the surface. No significant variation in hardness was observed. This is attributed to the homogenization of the texture and the microstructure described above. The hardness of the ARB1, ARB2 and ARB3 samples were, respectively, 135 ± 8, 152 ± 7 and 159 ± 8 HV. The corresponding approximate 105 UTS values [151] are 405, 503 and 524 MPa, respectively. These values are in good agreement with those obtained from tensile tests (difference within 5−14%). Fig 65: Stress-strain curves for annealed and rolled (ARB1 and ARB3) Zr. The tensile tests were performed at room temperature and 10 -3 s -1 . Dashed curves represent engineering stress-strain data and solid curves represent true stress-true strain data, respectively. Beyond the UTS, the true stress-true strain values were extrapolated until the true strain to failure. The latter was calculated taking into account the final cross- section of the specimen. 106 Table 6: Mechanical property data of annealed and ARB Zr Yield Stress (MPa) Engineering Stress at the UTS (MPa) True Stress at the UTS (MPa) Engineering Elongation to the UTS True Strain to the UTS Annealed Zr 185 268 324 0.207 0.188 1ARB 350 458 481 0.051 0.05 3ARB 428 513 550 0.074 0.071 Engineering Stress at Failure (MPa) True Stress at Failure (MPa) Elongation at Failure True strain at Failure Annealed Zr 224 435 0.403 0.662 1ARB 305 643 0.296 0.744 3ARB 341 651 0.30 0.648 107 3.1.6 Effect of Rolling Temperature and Thickness-reduction per Pass The ARB processing route described was designed to obtain a bulk homogeneous ufg structure using just a few rolling passes. Processing variables such as rolling temperature and strain rate were optimized in order to achieve this goal. Higher rolling temperatures give rise to coarser microstructures, since the average (sub)grain size (minimum grain size achievable) increases with temperature. Fig 66 illustrates the microstructure of pure Zr rolled using one pass with a 75% reduction at 300 o C (Fig 66a) and at 600 o C (Fig 66b). Rolling at ambient temperature led to a microstructure formed mainly of dislocation cell walls; the development of new high- or low-angle boundaries is significantly retarded due to the absence of recovery. Thus, even though room temperature ARB could potentially give rise to a finer grain size than rolling at 327 o C (the temperature chosen for the process described in this work), a significantly larger number of rolling cycles appears to be required in order to obtain a bulk, homogeneous, equiaxed microstructure. The thickness reduction per pass is also a critical parameter that must be optimized for effective ARB processing. Fig 66c shows the microstructure of pure Zr rolled at 300 o C using 19 passes with 5% thickness reductions, summing to a final reduction of 75%. The micrograph was taken at a location very near the sheet surface. An elongated structure develops although the total reduction (75%) is the same as in Fig 66a. Thus, a very large thickness reduction per pass was utilized (75%) in this work. Shear strains are imposed on the surface regions due to friction with the rolls. The combination of shear and plane strain at the surface regions, together with recovery possibly facilitated by adiabatic heating, leads to the formation of ultrafine equiaxed grains in these areas after 108 just one rolling pass (this finding inspired an idea to study the effect of grain size on biocompatibility in the surface regions of ultrafine grained Zr processed by one pass large strain rolling, which will be discussed subsequently). The regions with equiaxed grains are then placed in the interior of the stack upon subsequent ARB cycles. Fig 66: TEM micrographs showing the microstructure of pure Zr rolled using (a) one rolling pass with 75% thickness reduction at 300 o C; (b) one rolling pass with 75% thickness reduction at 600 o C; (c) multiple rolling passes with 5% thickness reduction, summing up to a total reduction of 75% at 300 o C. 3.1.7 Assessment of the Bonding Quality 3.1.7.1 Optical Microscopy Figure 67 consists of a series of optical micrographs showing the RD-ND plane of the Zr samples following the ARB1, ARB2, and ARB3 steps. Final surface polishing using 0.05 μm colloidal silica particles revealed the interfaces between the stacked layers. After ARB1, the interface between the two layers is clearly visible. The added strain applied during the ARB2 steps enhances the bonding quality at the interfaces created during ARB1, termed here “ARB1 interfaces”, which become almost invisible. The new 109 interfaces created during ARB2 could still be revealed, albeit not as clearly as in the ARB1 case. After ARB3 all interfaces are almost indistinguishable. The gradual improvement of the bonding quality in the “new interfaces” during successive ARB steps may be attributed to enhanced mechanical bonding caused by the presence of finer grains with increasing deformation (see the following sections for a detailed description of the microstructure of the roll bonded samples). This gradual improvement of the bonding quality with increasing strain is consistent with what was reported in other ARB studies for cubic materials [16, 26]. 3.1.7.2 Shear Testing The specimens in “group #1” (see designation in section 2.1.5.2), which are designed to test the bond strength in the mid-thickness of the ARB3 samples (32-layered strip), broke into two parts by the separation at the weakest layer as shown in Figure 68(a). The bond strength for three specimens of group #1 is listed in Table 7. The average shear stress to separate the bond, calculated using equation (9), is about 70 MPa. The shear fracture stress is calculated by measuring the true fracture strength of the bonded assembly (in the loading direction parallel to the rolling plane and direction) divided by 3 . The measurement of true fracture strength of the bonded assembly is shown in Section 3.1.5. The test performed on the ARB3 specimen #4 (group #2) was designed to measure the strength of the bond near the surface of the strip. Specimen #4 broke due to a tensile stress concentration at the smallest section, rendering the test invalid (Fig 68(b)). 110 ARB1 ARB2 ARB3 Fig 67: Optical micrographs showing the layered structure of pure Zr after 1, 2, and 3 cycles of accumulative roll bonding. The observation plane is perpendicular to the transverse direction (TD). 200 μ μ μ μm ARB1 interface ARB2 interface ARB1 interface RD ARB2 interface 111 Table 7: Bond strength for three specimens in group #1 Group Specimen Shear stress to separate the bond (τ B ) [MPa] Shear fracture strength of the bonded material (τ A ) [MPa] The ratio of τ B / τ A Group #1 (interior) Specimen #1 64 376 17% Specimen #2 64 376 17% Specimen #3 83 376 22% Average 70 376 19% (a) (b) Fig 68: Fracture of two specimens (a) as a result of shear test (specimen #1) and (b) as a result of tensile test (specimen #4). 112 3.2 Biocompatibility of Zr Processed by Large Strain Rolling 3.2.1 Material Characterization Figure 69 illustrates the typical microstructures of the investigated materials and Table 8 indicates the measured values of grain size and hardness. Ti64 exhibits the expected biphasic microstructure, consisting of α (dark zones) and β (bright zones) phases arranged in a cellular way. An average grain size of about 5 μm was found. The microstructure of annealed Zr is characterized by the presence of equiaxed grains with a homogenous size of about 42 μm. TEM examination reveals that severe deformation of the as-received Zr results in a refinement of the grain size (Fig 69C). Fig 69D shows the grain size distribution for this material. A reduction in grain size of several orders of magnitude occurs. An average size of 240 nm was determined. 113 Fig 69: BEI (A,B) and TEM (C) micrographs revealing the microstructure of Ti64 (A), Zr (B) and UFG Zr (C). Grain size distribution corresponding to UFG Zr is illustrated in D. 114 Table 8: Grain size and hardness values for the different materials Grain size (μm) Hardness (HV1) Zr 42 98±5 UFG Zr 0.24 172±3 Ti64 5 293±9 The roughness values corresponding to a scale comparable to the size of the osteoblastic cells were investigated by AFM. Results are summarized in Table 9. Both Zr samples showed similar R a (average roughness) and R q (root-mean-square roughness) roughness parameters in the order of nanometers and higher than those of the Ti64 sample. These parameters are scale dependent, meaning that cells with different sizes sense surface patches having different roughness parameters [152-154]. Although R a and R q are the most frequently used parameters for describing the surface roughness in the biomaterials field, a general concern is arising about their usefulness, since they do not contain information about the spatial distribution of surface features. To extract from the AFM images the information about the lateral distance between the topographical features, the obtained profiles were expressed as a sum of harmonic terms (Fourier series), which can be associated with different components of the surface roughness. A representative profile of the Zr sample has been plotted in Fig 70A. In Fig 70B, the surface topography was plotted taking only the large spatial wavelength components of the above decomposition. The overall appearance of the original profile has been reproduced without its finest features. Fig 70C shows the surface profile without the previously used large spatial wavelength components, that is, showing only the finest surface features. Note the different y-scales of these figures. Thus, the original complex 115 topography can be expressed as the sum of two independent profiles, each of them showing a specific component of the topography. This procedure has also been applied to the UFG Zr sample. The result of applying an analogous treatment to the representative profile shown in Fig 70D is presented in Figs 70E and 70F. As can be readily appreciated, the general profile of Fig 70F contains small irregularities, which are indicative of the presence of an even smaller roughness component. To resolve both components, the profile in Fig 70F has been decomposed into two independent profiles (Figs 70G and 70H), following the same procedure as before. Once the decomposition was done, the peak-to-peak distances of the different roughness components were quantified. The obtained values of averaged peak-to-peak distances and the height of the peaks for the Zr sample (Fig 70C) were 1.0±0.1 μm and 3±1 nm, respectively. For the UFG Zr sample, the results were 1±0.3 μm and 3±1 nm (Fig 70G), and 0.4±0.1 μm and 1.0±0.4 nm (Fig 70H). It can be observed that both samples display a very similar sub-microscopic topography except for the minute contribution depicted in Fig 70H, taking into account both the amplitude roughness parameters and the study conducted of the lateral distribution of peaks. Table 9: R q and R a roughness values at length scale of 30 μm for the different materials R q (nm) R a (nm) Zr 7 ± 2 4.7 ± 0.9 UFG Zr 9 ± 2 7.0 ± 1.0 Ti64 3 ± 3 2.3 ± 0.1 116 Fig 70: AFM topographical profile of zirconium surfaces. Representative profile (A), and profiles containing the highest (B) and lowest (C) wavelength components, on the Zr sample. Representative profile (D), profile containing the highest (E) wavelength component, (F) complementary profile to (E), and a subsequent decomposition of (F) into a high (G) and low (H) wavelength components, on the UFG Zr sample. The values of the water contact angle measured on the Zr sample was θ w = 73 ± 2 o and on the UFG Zr sample was θ w = 68 ± 3 o , indicating that both solids are mainly hydrophobic. A lack of significant differences between values measured on both materials indicates that the procedure on Zr to get the UFG Zr does not introduce any clear change in its surface hydrophobicity. Surfaces of both Zr samples share identical chemistry and similar roughness and therefore, different values in water contact angle would not be expected. 117 3.2.2 Cell Attachment Cell attachment assays were conducted by incubating cells for 0.5, 3 and 24 h on Ti64, Zr and UFG Zr (Fig 71A). The number of attached cells increased with time on the three tested surfaces. Cell attachment at the assessed times was similar on both conventional metallic surfaces of Ti64 and Zr. A reduction in the grain size of zirconium did not affect the number of attached cells. We next examined cell morphology in a short culture period of 0.5 h (Fig 71B). Morphology of cells attached at 0.5 h was similar on Ti64 and Zr and was not affected by the reduction in the grain size of Zr. 118 Fig 71: Cell attachment on zirconium surfaces. (A) Saos-2 cells were cultured on Ti64 ( ), Zr ( ) and UFG Zr ( ) for 0.5, 3 and 24 h. The results are expressed as the percentage of the absorbance measured on Ti64 at 0.5 h, which was given the arbitrary value of 100. Each data represents the mean ± SD of four independent experiments. * p<0.05 compared to 0.5 h, # p<0.05 compared to 3 h. (B) Actin staining of Saos-2 cells cultured for 0.5 h on Ti64, Zr and UFG Zr. Bars= 30 μm. 3.2.3 Cell Spreading and Arrangement of Tubulin and Actin Cytoskeletons Microscopic examination of actin-stained cells suggested a similar level of spreading on the three tested surfaces (Fig 72A). The quantified area of cells cultured on Ti64 was similar to those measured on both Zr-based materials, independent of the grain size of zirconium (Fig 72B). We next examined the organization and density of tubulin 119 and actin cytoskeletons of cells cultured for 1 day on the various materials (Fig 73). A tubulin network appeared, radiating from the microtubule organizing centre, on all tested samples. There were no changes either in organization or in density of the tubulin cytoskeleton between cells cultured on conventional Ti64 and Zr. Tubulin pattern was unaffected by a reduction in the grain size of conventional zirconium. Examination of actin was focused on the ability of cells cultured on the various materials to form stress fibers. On the three tested surfaces, actin filaments appeared organized in well-defined stress fibers and mostly oriented in a parallel direction following the main cellular axis. The pattern of actin distribution was similar on both conventional metal surfaces, Ti64 and Zr, and was independent of the grain size of zirconium. 120 Fig 72: Cell spreading on zirconium surfaces. Saos-2 cells were cultured on Ti64 ( ), Zr ( ) and UFG Zr ( ) for 24 h. (A) Actin staining of cells cultured on Ti64, Zr and UFG Zr. Representative images of three independent experiments with similar results. (B) Spreading of cells cultured for 24 h on Ti64, Zr and UFG Zr. Bars = 50 μm. Results represent the mean ± SD of cell area measured in three different experiments. 121 Fig 73: β-tubulin and actin cytoskeletons on zirconium surfaces. Saos-2 cells were cultured on Ti64, Zr or UFG Zr for 24 h. Bars = 30 μm. 3.2.4 Fibronectin Analysis The secretion and cellular distribution of fibronectin after culturing cells for 24 h on the samples was studied. To this end, fibronectin levels were quantified in media and cell 122 layers. On all tested substrates, amounts of this extracellular matrix protein detected in media were significantly higher than in cell layers (Fig 74A). No differences in fibronectin levels were found on any substrate. Fibronectin organization was examined by immunofluorescence assays (Fig 74B). Cells were stained positively for fibronectin on both conventional Ti64 and Zr and reduction in grain size of zirconium samples did not alter the staining pattern. A closer examination of the images suggested that fibronectin was organized in fibrils on some cells cultured on any substrate. Fibronectin was clearly detected on the apical site of such cells. When the distribution of fibronectin fibrils and actin fibers by double immunofluorescence staining was investigated, merge images revealed that fibronectin fibrils on the cell surface paralleled the alignment of the actin filaments, independently of the substrate (Fig 75). 123 Fig 74: Fibronectin analysis on zirconium surfaces. (A) Levels of fibronectin were detected in cell layers and media from Saos-2 cells cultured on Ti64 ( ), Zr ( ) or UFG Zr ( ) for 24 h. The data are expressed as the percentage of the fluorescence measured in cell layers of cells cultured on Ti64, which was given the arbitrary value of 100 and is equivalent to 70 pg/μg total proteins. Each value represents the mean ± SD of four independent experiments. (B) Fibronectin staining of Saos-2 cells cultured for 24 h on Ti64, Zr or UFG Zr. Bars = 50 μm. 124 Fig 75: Double staining of actin (red) and fibronectin (green) on zirconium surfaces. Saos-2 cells were cultured on Ti64, Zr or UFG Zr for 24 h. Bars = 20 μm. 3.2.5 Cell Viability Cell viability assays were conducted by culturing cells for 1-7 days (Fig 76). The number of viable cells increased for 1-4 days and remained constant for up to 7 days on all tested substrates. There were no differences in cell viability among conventional Ti64 and Zr materials at any point in time. Compared to conventional Zr, UFG Zr did not affect this cellular parameter. 125 Fig 76: Cell viability on zirconium surfaces. Saos-2 cells were cultured on Ti64 ( ), Zr ( ) and UFG Zr ( ) for 1, 4 and 7 days. The data are expressed as the percentage of the fluorescence measured on Ti64 at 1 day, which was given the arbitrary value of 100. Each value represents the mean ± SD of four independent experiments. * p<0.05 compared to 4 and 7 days. 3.2.6 ALP Activity and Degree of Mineralization The ability of hMSC to differentiate into osteogenic cells when cultured on the tested materials was evaluated by measurement of ALP activity and mineralization degree in cell layers. ALP activity and mineralized nodule formation increased greatly when hMSC were cultured in osteogenic induction medium on all tested materials (Fig 77). On Ti64, cells cultured under osteogenic conditions exhibited similar ALP activity and degree of mineralization as was detected on conventional Zr, which were unaffected by the reduction in the grain size of UFG Zr under the same conditions. 126 Fig 77: ALP activity and degree of mineralisation on zirconium surfaces. hMSC cells were cultured on Ti64, Zr or UFG Zr in growth ( ) or osteogenic media ( ). The data are expressed as percentage of the absorbance measured on cells cultured on Ti64 in growth medium, which was given the arbitrary value of 100. Each value represents the mean ± SD of three independent experiments. 127 CHARPTER 4. DISCUSSION 4.1 Effectiveness of ARB as a Grain Refinement Technique for Zr Accumulative roll bonding is a suitable technique to fabricate bulk ufg pure Zr. A combination of microstructural characterization techniques was used to investigate the ARB microstructure and to evaluate the grain size and grain shape. First, EBSD was used to estimate “true” average grain size values, measured by the linear intercept method with only HABs (θ > 15 o ). An average grain size of 400 nm was determined. Second, the fraction of HABs, measured both by EBSD and by Kikuchi line analysis in the TEM, was 88% and 66%, respectively. The discrepancy between these two values is characteristic of severely deformed materials, in which there is a large fraction of boundaries with misorientations smaller than 2 o , which can not be detected by EBSD [155]. A minimum value for the fraction of HABs that is necessary for a material to be called “ultrafine-grained” has not been proposed. Instead, ufg materials are still loosely defined as structures in which a “majority” of grain boundaries have high angles of misorientation [5]. Previous studies have reported HAB fractions of 70−80% (measured by EBSD) in fcc materials processed by multiple ECAP passes and in Ni processed by HPT with five or more revolutions at relatively low temperatures (T = 0.3T m ) [4]. Also, HAB fractions between 50% and 90% (measured by EBSD) were obtained in several fcc materials (OFHC Cu, 36% Ni steel, IF steel and 1100 Al) processed by six ARB cycles to a final strain of 4.8 [28]. Finally, pure Al processed by ARB up to a strain of 4.8 (six cycles) had HAB fractions of 71% (EBSD) and 52% (TEM) [155]. Thus, the fraction of HABs in the ufg pure Zr developed in the present study is ge- 128 nerally larger than that obtained in previous investigations in fcc metals for equivalent strains. Furthermore, the strain required to achieve such a large fraction of HABs is smaller. The possibility of performing large strain passes in Zr, thereby reducing the number of ARB cycles needed to achieve a bulk ufg structure, may be due to the high ductility of Zr at low temperatures. It is well known that several slip systems are active in Zr, even at low temperatures. In particular, slip takes place preferentially in > < − − 0 2 11 } 0 1 10 { and > < − 0 2 11 } 0002 { slip modes [130]. Each one has only two independent slip systems. Additionally, (c+a) slip on pyramidal } 1 2 11 { − and } 1 1 10 { − planes has been observed in Zr at temperatures between 25 and 75 o C [130]. The activity of non-prismatic slip systems increases with increasing temperature. A noteworthy observation of the present study is the formation of equiaxed grains (in the bulk material) after only two ARB cycles (ε = 3.57). In fact, the grain aspect ratio is smaller than 1.6 throughout the entire ARB process. This is remarkable, since previous studies on fcc materials processed by ARB [28] have reported grain aspect ratios higher than 3 even after strains as high as 5.1 (six cycles). Moreover, elongated grains were observed in AA8011 Al even after two ARB cycles (ε = 9.6) [26]. It has not been possible to produce ultrafine equiaxed grains in Al or steels by ARB (in the absence of subsequent annealing). The presence of elongated grains may result in an anisotropic mechanical behavior. In contrast, an equiaxed structure forms readily in Zr even at small strains. The equiaxed structure was also observed in commercial purity Ti processed by accumulative roll bonding process up to 8 cycles (equivalent strain of 6.4) at ambient 129 temperature [34]. It may be suggested that equiaxed structure is the characteristic of microstructures for hcp materials processed by ARB compared with cubic materials processed by ARB. Several factors could explain this observation of equiaxed structure. First, it was previously reported that pure Zr “recrystallizes” by geometric dynamic recrystallization [122]. This phenomenon, described schematically in Fig 78, involves grain elongation and thinning, undulation of the original HABs to form distinct “crystallites” when the grain thickness approaches the subgrain size. Zr has a very low thermal conductivity (≈17 W m K -1 ) compared with other metals such as Al (210 W m K -1 ). Thus, adiabatic heating during high strain rate (28 s -1 ) processing might contribute particularly to enhanced diffusion, facilitating the recovery process and thus leading to equiaxed grains already after a small number of passes. A simple calculation of the temperature increase during rolling under the conditions of the present study using βσε = ρCpΔT (where ε = 1.6 = strain at the mid-thickness during one pass, β = 0.9, ρ = 6.53 g cm -3 , σ = 610 MPa [156] and Cp = 302 J K -1 kg -1 ) yields maximum temperature close to 0.5Tm (ΔT = 485 K). Of course this temperature may only be maintained for a short time. Second, the shear strain imposed in the ARB process described in this study is significantly larger than that used in previous studies on fcc materials [12, 19, 27, 28, 150, 155, 157-162]. Lee et al. [36] achieved shear stains as high as 8 during ARB of 1100 Al using 50% reduction per pass. In this study the reduction per pass during ARB is 75% and this results in larger shear strains. In summary, the microscopic mechanism by which an ufg structure develops in pure Zr during ARB may be described as follows. At the first stages of deformation, i.e. 130 before the ARB2 step (ε < 3.57), the (sub)grain and grain sizes are smaller than the grain thickness predicted by geometrical thinning (Fig 58). This suggests the presence of significant grain subdivision, possibly by the formation of GNBs. At larger strains (ε > 3.57) the (sub)grain and grain sizes remain constant at a value smaller than the grain thickness predicted by geometrical thinning. This is attributed to geometrical dynamic recrystallization (GDX), where the pinching-off effect is enhanced by adiabatic heating, and to local boundary migration. The presence of a through-thickness texture gradient might also influence to some extent the rate at which GDX occurs at different depths within the bulk. Finally, it is expected that if ARB is performed in Zr of higher purity than that of the present study, higher grain boundary mobility would result in a larger grain size after processing. Fig 78: Schematic of geometric dynamic recrystallization. 4.2 Texture Homogeneity Various studies have investigated the evolution of texture in fcc materials during successive ARB cycles, as well as the through-thickness texture homogeneity [27, 30, 157]. Reis et al. [30] described the development of a clear shear texture at the outer surface of a Ti-interstitial free (IF) steel with increasing ARB passes. However, they did 131 not observe any trace of this shear texture in the inner layers. Chowdury et al. [27] followed the evolution of the texture in a region close to an inner interface in an Al 8090 alloy. After the first cycle, the texture was rather asymmetrical, resembling neither typical shear nor rolling components. With an increasing number of cycles, a typical rolling texture developed close to the interface. Finally, Kim et al. [157] studied the evolution of texture in an Al 8011 alloy processed by ARB. They reported the development of a clear shear texture at the outer surfaces and the gradual disappearance of this shear texture once the outer layers were placed in the interior of the ARB sample. The results obtained in the present study are in agreement with previous studies. A pronounced basal texture, with > < − 0 2 11 directions parallel to the RD, develops at the outer surfaces. The regions below the surface are deformed by a combination of shear and plain strain, with the contribution of shear strain to the total strain decreasing as the depth increases. Thus, the RD is spread along the > < − 0 2 11 − > < − 0 1 10 boundary of the inverse stereographic triangle. Once the outer surfaces are placed in the interior of the ARB stack and are deformed by pure plain strain upon successive ARB cycles, rotation of the RD from > < − 0 2 11 toward > < − 0 1 10 is observed. The rotation is completed after two ARB cycles. No DRX texture was observed. The absence of any trace of a shear texture in the measurements performed by Reis et al. [30] in the interior of the ARB processed samples may be due to the fact that their plane of observation was parallel to the rolling plane, probably close to the mid-thickness of one layer, where a pure rolling texture develops. The observation of an asymmetric texture by Chowdury et al. [27] might indeed correspond to an intermediate orientation between the pure shear texture 132 and the pure rolling texture. The present work constitutes a more comprehensive investigation, in which the evolution of the texture at various locations through the thickness of the ARBed sheets is followed and related qualitatively to the processing history of that specific location. It is also the first study of this kind in an hcp metal. 4.3 Comparison with Other HCP Materials ARB was previously utilized for other hcp materials, such as several Mg-Al alloys, but with only limited success [31, 32]. In order to prevent significant grain growth, ARB must be carried out at temperatures that are lower than about 0.5T m . Due to the high value of the c/a ratio in Mg alloys, only basal slip systems are activated at these temperatures and, therefore, large thickness reductions per pass result in cracking. It was previously reported that reductions of 75% per pass may be applied to Mg-Al alloys (3- 6% Al) at temperatures equal or higher than 0.6T m but only when the materials have a random texture [163]. Any deviations from these conditions (lower rolling temperature, presence of texture components) significantly reduce the maximum thickness reduction per pass that can be applied. Therefore, ARB is not an effective method to produce ufg Mg alloys because of the two following reasons: first, when rolling is performed at low rolling temperatures, at which rapid grain growth is avoided, only small reductions may be applied. This results in poor bonding quality [31]. Second, if ARB is performed at higher temperatures, at which a sufficient number of slip systems become active, rapid grain growth takes place and thus it is not possible to obtain ultrafine grains. Additionally, Mg is a highly reactive material, which oxidizes readily. Thus, the quality of the interface bonding is rather poor even when relatively large thickness reductions are applied [32]. 133 In contrast, Zr is an excellent candidate for ARB. Due to its low c/a ratio, it possesses high ductility at low temperatures. Thus, large thickness reductions per pass can be applied, leading to a homogeneous ufg microstructure after a small number of passes and to accepted interlayer bonding quality. Other metals, such as Ti, a structural metal with properties similar to those of Zr, would be attractive for further study. 4.4 Bond Strength of ARBed Samples In Krallics’s sample design [164], shown in Fig 79, two narrow slots were milled by the designated layer in order to apply a shear stress on that specific layer. However, the location of the designated layer may vary slightly from the expected location (especially if the layer thickness is very small). Furthermore, the application of a tensile stress on the layer was nondeductible even if the depths are controlled precisely. Thus, sample design to measure the shear strength of the bond in Krallics’s work [164] may cause the interface, where bonding strength is to be tested, to be subjected to a tensile stress rather than a shear stress. In contrast, in the present study a new sample design was put forward and explained in the experimental section. The essential change was that two narrow slots were milled to overlap along the direction through the thickness (shown in Fig 43), which guarantees the tested area was nearly subjected to shear stress. 134 Fig 79: The schematic diagram of the test sample to measure the shear strength of the bond in Krallics’s work [164]. The parameters that influence the bond strength in ARB alloys were discussed in detail by Krallics [164]. According to this work, the main three parameters influencing the bond strength are the ductility, the thickness reduction per pass and the rolling temperature. Zr is an hcp material that has excellent low temperature ductility [123]. Due to its low c/a ratio, slip can be activated on the { 0 1 10 − } first-order prism planes and on the {0001} basal planes along < 0 2 11 − > directions, as well as on the first and second- order pyramidal planes { 1 1 10 − } and { 1 2 11 − } along > < − 113 2 directions [130]. This large low-temperature ductility may favor the formation of stronger solid-state deformation bonds than those in materials with lower ductility. ARB was previously utilized in other hcp materials, such as several Mg-Al alloys, but with only limited success [31, 32]. Due to the high value of the c/a ratio in Mg alloys, only basal slip systems are activated at low temperatures (< 0.5 Tm) and, therefore, large thickness reductions per pass result in 135 cracking. Higher reductions promote a larger degree of bonding. Del Valle et al. [31] reported that only the samples of a Mg-based AZ61 alloy with reductions higher than 50% were bonded successfully (based on optical inspection). Lower reductions were not sufficient to bond the slabs and delamination occurred while cutting the samples. Tsuji et al. [23] pointed out that there is a critical rolling reduction, below which it is difficult to achieve sufficient bonding (optical inspection) during one-pass roll-bonding. In the present study, the bond strength in the middle of a Zr strip, undergoing three ARB cycles, is 70 MPa (about 66% of the original yield strength in shear and 19% of the shear fracture strength). This ratio (66%) is larger than the value (50%) obtained by Krallics et al. [164] in an ultra-low carbon steel strip also subjected to three ARB cycles. Also, 19% is higher than the 14% of the shear fracture strength of Krallics [164]. More successful bonding in the present work may be due to a larger thickness reduction per pass than that utilized by Krallics (75% vs. 50%), as well as material differences. The rolling temperature is another important parameter influencing the bonding quality. Del Valle et al. [31] investigated the degree of bonding of a Mg-based AZ31 alloy by optical microscopy. They found that the degree of bonding increased with increasing temperature, being better at 400 o C than at 300 o C. This may be related to the increasing ductility at higher temperatures. As mentioned earlier, the high ductility of Zr allows favorable bonding quality even at the moderate temperatures (327 o C = 0.28 T m ). However, ARB is not an effective method for grain refinement in Mg alloys because rapid grain growth occurs if ARB is performed at higher temperatures. Due to the very limited quantitative studies on the bond quality of ARB processed materials, the bonding strength values obtained in the present study are compared with 136 those in materials processed by diffusion bonding. Somekawa et al. [165] examined the diffusion bonding of an AZ31 magnesium alloy with an average grain size of 8.5 μm. They reported that the shear strength was 68.5 MPa after a diffusion bonding at a pressure of 20 MPa after 1 h at 300 o C. Ravishankar et al. [166] reported that the shear strength of a Ti-6Al-4V alloy with a grain size of 9.3 microns, processed by diffusion bonding (at a pressure of 12 MPa for 15 min at 900 o C), was 80 MPa. The shear strength in the present study (64-83 MPa) is of the same order or better than that achieved in these diffusion bonding studies. 4.5 Biocompatibility of Zr Processed by Large Strain Rolling The enormous potential of nanostructured materials for a wide range of applications is the driving force for the increasing activity in the development of these materials. Nanocrystalline materials are structurally characterized by a large volume fraction of grain boundaries, which confers them different mechanical, physical or chemical properties in comparison with conventional coarse-grained polycrystalline materials, whose grain sizes usually are in the range 10-300 μm. In particular, their usefulness for biomedical applications has logically triggered considerable interest in a number of devices and systems [167, 168]. Thin films of pure Zr with grain size down to 100 nm have recently been developed using a rolling and folding technique, which requires close to 100 rolling steps [169]. The final product was in the form of thin layers of Zr, with an average grain size of 80 nm, which were only a few microns in thickness. Due to the complexity of their current method of production, it is not predicted that those films would be of commercial interest. In this work, we have investigated the biological response of Zr with two different 137 microstructures. The finest grain condition is obtained by SPD (severe rolling). As a consequence of the high strain per pass imposed, an UFG microstructure is readily obtained. This is possible because Zr is very ductile at rather low homologous temperatures (0.27 T m ), at which a large number of slip systems become active [130]. Additionally, the superposition of plain strain in the surface regions and shear deformation due to friction with the rolls [36], leads to the formation of equiaxed ultrafine grains at the surface regions after only one pass. Adiabatic heating during rolling contributes to enhance recovery and thus to the acceleration of boundary formation. “Nanophase” Ti, Ti64 and CoCrMo composed of particles smaller than 100 nm enhance osteoblast adhesion rate, as compared to conventional metals with micrometric grain sizes, an effect mainly related to the altered topography of the resulting “nanophase” metals [137, 170]. On these substrates, cells preferentially attach to the particle boundaries that cover the surface as result of the processing of the metal particles. Elongated cell protrusions were observed on those topographical features. We have not observed any difference in the attachment rate of osteoblastic cells incubated on UFG Zr, compared to conventional Zr. Cells cultured for a short time period of 0.5 h displayed an expected poor degree of actin cytoskeleton organization on both surfaces, without evidence of filopodial guidance. This result is consistent with the fact that both polished materials share an identical surface chemistry while only marginally differing in their nanotopographies. A few studies have demonstrated that osteoblasts could “sense” surface topography at the nanometric scale although the lowest dimension that these cells may distinguish in early interaction periods has not yet been accurately defined. A comparison study between nanosized niobium oxide coatings indicated that adhesion 138 strength was higher on topographies having an R a of 15 nm, compared to surfaces with 7 nm [171]. According to our data, nanometric changes in roughness below an R a of 7 nm or subtle differences in the lateral distribution of topographical features do not seem to have any significant impact on this initial attachment phase. The effect on cell behavior after longer interaction periods has been investigated in osteoprogenitor cells that react by enhanced spreading and expression of stress fibers, a larger tubulin network and filopodial guidance to hemi nanofeatures down to a defined height size of 10 nm and spaced by 180 nm [172]. Filopodial protrusions were not observed in osteoblastic Saos-2 cells cultured either on UFG Zr or on Zr, which displayed a well-organized actin and tubulin cytoskeleton. It could be then speculated that the smallest features that osteoblast- lineage cells may sense are about 10-15 nm, which might explain the similar response that we have observed on Zr-based substrates. Organization of these surface features, rather than their amplitude, has a strong influence in cell adhesion [173]. Thus, it is plausible to speculate that high peak-to-peak distances (0.5−1 μm) of the small features detected in our work also contribute to the lack of differential response on Zr and UFG Zr surfaces. Extracellular matrix assembly confers the mechanical properties of the new tissue produced at the implant interface. Fibronectin is secreted by cells as a soluble dimer and is subsequently assembled into an insoluble dense meshwork of interconnected fibrils at the cell surface which provide a dynamic environment for cells [174, 175]. Fibril formation is mediated by cells and depends on the interaction between soluble fibronectin and integrin receptors on the cell surface [176]. Integrins link fibronectin fibrils to the actin cytoskeleton and this connection influences the alignment of the fibronectin fibrils 139 and the cytoskeleton, leading to variations in fibril organization [176, 177]. Microtopography of the material, which greatly influences integrin reorganization during osteoblast adhesion to surfaces, also modulates the alignment of the fibronectin fibrils and the stress fibers. Fibril distribution was clearly parallel to actin bundles on polished titanium while micrometric alterations in roughness disrupted such fibril organization [178]. The typically aligned distribution of fibronectin and actin was also observed on cells cultured on both Zr-based surfaces. UFG Zr did not affect production and secretion of the extracellular matrix protein, providing additional evidence that changes in low scales of nanotopography may not affect the process of fibronectin matrix formation. Cells proliferate on the UFG Zr surface to the same extent as on conventional Zr as shown in our viability assays performed up to 7 days. In concordance with our data, it has been recently published that nanometric changes in roughness of titanium films, ranging from 2 to 21 nm, do not influence the proliferation or viability of human osteoblastic cells at identical observation periods as those employed by Cai et al [179]. In agreement with the data in the present work, subtle nanometric changes in roughness of those films did not affect the hydrophobic behavior of titanium surfaces [179]. Finally, an essential requirement for a candidate material to tightly bind to living bone is its ability to promote bone matrix maturation and mineralization. In this work, we have examined the behavior of human mesenchymal stem cells from bone marrow to exhibit the osteoblast-specific activities on Zr-based surfaces. After induction in osteogenic media, ALP activity and degree of mineralization were unaffected on UFG Zr as compared to conventional Zr or Ti64. These in vitro results suggest that this novel 140 biomaterial could favour bone formation and contribution to successful osseointegration to the same extent as previously observed in vivo for conventional Zr [143, 144]. The lower mechanical properties of Zr are the main reason for their limited development. The introduction of novel technologies for processing, however, opens a window to improvement in the mechanical properties associated with the grain size decrease. A reduction in the grain size is accompanied by an increase in hardness (98 HV1 for Zr and 172 HV1 for UFG Zr). We noticed that the hardness of UFG Zr is comparable to that of commercially pure Ti (cpTi) (160 HV1) with an average grain size of about 45 μm. Improved mechanical properties together with its excellent in vitro biocompatibility make UFG Zr a promising biomaterial for dental implants. A previous study indicates that osteoblastic cells respond to Zr in a similar fashion as they do to cpTi [147]. Although it has been speculated that oxide titanium film that develops on cpTi may be different from that formed on Ti64, human osteoblastic cell response to both polished metals has been found to be essentially identical [180,181]. In this work, the in vitro biocompatibility of Zr-based materials has been compared to Ti64 and the several cell parameters assessed indicate an equivalent osteoblast response. Based on its higher strength, Ti64 is especially suitable for load-bearing implants whereas the use of cpTi has been reduced to dental implants. Although several other issues remain to be investigated, our present results suggest that performance of UFG Zr-based dental implants might be comparable to that achieved by the in-use of cpTi implants. 141 CHARPTER 5. CONCLUSIONS 1. Accumulative roll bonding is a suitable technique to fabricate bulk ufg pure Zr. A homogenous, equiaxed microstructure, with an average grain size of 400 nm, a (sub)grain size of 320 nm and a large fraction of high angle boundaries (HABs) (66% using TEM and 88% using EBSD), may be obtained after only two ARB cycles. Optimum processing conditions include a rolling temperature of about 300 o C and large thickness reductions per pass, in excess of 75%. 2. Grain refinement during processing occurs by grain subdivision through the formation of geometrically necessary boundaries, and by geometric dynamic recrystallization (grain thinning). 3. A large volume fraction of equiaxed grains is readily obtained after two ARB cycles (ε = 3.57). This is attributed to the large shear strain imposed and to the occurrence of recovery, possibly assisted by adiabatic heating. 4. A through-thickness texture gradient develops during rolling. A typical rolling texture develops in the areas subjected to pure plain-strain deformation, with basal planes rotated slightly away from the ND toward the TD and with > < − 0 1 10 directions aligned with the RD. In the outer layers, due to the macroscopic shear strain imposed by friction with the rolls, a {0002} > < − 0 2 11 texture develops. The shear texture evolves toward a rolling texture after two ARB cycles when the sheared surfaces are placed in the interior of the ARB stack. 5. The yield strength and UTS are nearly double that of conventionally processed Zr with considerable ductility. 142 6. The shear stress necessary to break the bond in ARB3 samples is on the order of 70 MPa. This strength is considered favorable, as it is comparable to that achieved in materials processed by diffusion bonding and higher than that obtained in other ARB studies. This high bonding strength may be attributed to the excellent low temperature ductility of Zr as well as to the high reductions per pass utilized during processing. 7. Osteoblast response to a novel ultrafine grained Zr, displaying significantly higher hardness values than conventional coarse Zr, was evaluated. Cell attachment, spreading and viability, cytoskeleton reorganization and fibronectin production were not affected by the grain size reduction of Zr. After induction in osteogenic media, osteogenic differentiation and bone matrix mineralization of human mesenchymal stem cells proceed to the same extent on both Zr-based materials, independently of the grain size of zirconium. 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This program consists of three main parts as the following: (Part I) requesting control from the machine; (Part II) controlling the machine; and (Part III) data acquisition. This program can be applied to do tension tests and shear tests in the current research. Furthermore, this program can also be employed to do compression tests and rotation tests, etc. if it is slightly modified. Figure A1 represents the front control panel of “LJInstron” program. The strain rate of tests, the gage length of specimens and the path to save data files can be set in this panel. The block diagram of “LJInstron” program is shown in Fig A2. It can be seen that windows W1~W5 constitute the Part I of requesting control from the tension machine. Windows W6~W13 are used to set parameters for a tensile test including strain rate, gage length of specimens and some security parameters, which constitute the second part of this program (Part II). Window W14 is the core part for data acquisition and recording. Finally the control will be returned to the machine by commands in window W15. The detail information for each window is tabulated, shown in Table A1. 164 Table A1: The explanation on the functions of each window in the block diagram of LJInstron program. Block Name Function W1 To check machine and computer are connected by GPIB W2 A command to request computer control W3 To do loops until control is given by machine W4 To clear system setting parameters in machine W5 To set watchdog timer in order to keep control by computer W6 To set ramp waveform parameters in order to act a moving. Strain rate is set here. W7 To set maximum and minimum load limits W8 To set actions if load limits are exceeded W9 The minimum or maximum limit set by command #122 can be directed to OFF or to be ARM W10 A command to request data requisition on displacement, load and ramp time W11 To set the interval time to update data W12 To check the initial load and position W13 To start the action W14 Data acquisition and saving W15 A command to return control to the machine 165 Fig A1: The front panel of “LJInstron” program. The path to save data files, the strain rate of testing and gage length of specimens can be assigned in this panel. 166 Fig A2: The block diagram of the LJInstron program. 167 Fig A2, Continued 168 Fig A2, Continued 169 Fig A2, Continued 170 Fig A2, Continued 171 Appendix II: Calculation on Strain and Strain Rate In the present study, accumulative roll bonding was performed using a Carl Wezer rolling machine located in Madrid, Spain (CENIM), furnished with two rolls. Its technical data is tabulated in Table A2. Table A2: Technical data of the rolling machine used in the present study. Width of roll body 200 mm Diameter of roll body 130 mm (R=65 mm) Rolling speed 20m/min (Ώ=V/2πR=49 r.p.m) Capacity of motor 15 kW Motor speed 1500 r.p.m Strain and the strain rate during the ARB process in this study can be calculated by using the following equations: 1. Strain: 2 1 ln 3 2 h h • = ε (A1) Where, h 1 : Initial thickness, h 2 : Final thickness. 2. Strain rate: r h r R m − • ⋅ • Ω • = 1 1 ln 60 2 3 2 1 ' π ε (A2) Where, Ω: Rotating speed (how many rounds per minute) R: Radius of roll body 172 h 1 : Initial thickness h 2 : Final thickness r: Reduction ((h 1 -h 2 )/h 1 ) Dimensions of the specimens for each rolling pass and corresponding strain and strain rate calculated using equations (A1) and (A2) are listed in Table A3. Table A3: Dimensions of the specimens and corresponding strain and strain rate for each rolling pass during ARB process. Initial thickness (mm) Final thickness (mm) Thickness Reduction Strain imposed Strain rate per this pass (s -1 ) Rolling pass at room temperature 5.2 3.9 25% 0.33 12 ARB1 7.8 1.88 76% 1.65 28 ARB2 7.52 1.87 75% 1.6 28 ARB3 7.48 1.85 75% 1.6 28 173 Appendix III: Boundary Misorientation Measurements Carried Out Through Kikuchi Bands in TEM With the aid of a TEM double-tilt holder, a method for determining grain boundary misorietation angle is detailedly explained by Liu et al [Q Liu, X X Huang and M Yao, Ultramicroscopy 41 (1992) 317]. Here, this method will be briefly described. This method is based on the fact that the orientation of a crystallite can be determined by recording three zone-axes from a Kikuchi pattern, and this is true for a double-tilt holder only. In a TEM with a double-tilt holder, it is easy to adjust the thin foil normal-to-beam specimen position to obtain the exact zone axis patterns [u 1 v 1 w 1 ], [u 2 v 2 w 2 ] and [u 3 v 3 w 3 ], and meanwhile, determine the corresponding tilt angles (α 1 -β 1 , α 2 -β 2 , α 3 -β 3 ) of the two axes of the double-tilt holder, directly from the tilt holder meter reading. Then, the orientation [u v w] of the crystal specimen where the double-tilt holder is at a position (α- β) can be determined using the following method. The angle θ i (i=1,2,3) is calculated by equation (A3) as follows [Q Liu, Micron Microsc Acta 20 (1989) 261]: )] cos( cos cos sin arccos[sin i i i i β β α α α α θ − + = (A3) where i=1,2,3 Also, the angle between two orientations can be calculated using the following equation [A4]: i i i w v u uvw i i I I L = θ cos (A4) 174 Here i=1,2,3, and C vu uv ab B uw wu ac A wv vw bc ww c vv b uu a L i i i i i i i i i i cos ) ( cos ) ( cos ) ( 2 2 2 + + + + + + + + = 2 / 1 2 2 2 2 2 2 ) cos 2 cos 2 cos 2 ( C ab B cawu A bcvw w c v b u a I uvw + + + + + = a, b, c, A, B and C are lattice parameters of the crystal. With the known position (α-β) of the double tilt holder, equation (A3) and (A4) allow the determination of the orientation [u v w] of the crystal specimen. In contrast, the tilt angles (α-β) corresponding to an orientation [u v w] of the crystal specimen being parallel to the beam direction can be established by calculating the angles θ i by equation (A4) firstly and then using equation (A3) to obtain α and β. Based on the principles discussed above, in general, computer software can be developed to determine the crystal orientation [u v w] at any tilt position (α-β) in a double-tilt holder, or determine the tilt position (α-β) when the orientation [u v w] of the crystal specimen is coincident with the beam direction. In practice, the easiest and more accurate way for taking advantages of these principles, is to build the rotational matrix R between two adjacent crystallites. The axis and angle of misorientation between two adjacent crystallites can be calculated from the rotational matrix R which specifics the rotation of the crystal coordinate system from one crystallite to the other. Three zone-axis directions from each of two adjacent crystallites are aligned to the beam direction and the corresponding tilt angles are directly, on-line, read from the meter. Then, we call crystallite A the one chosen to be oriented as to have one of the crystallographic directions (cubic Al-alloy) [100], [010] and [001] (zone axis) aligned to the beam direction. [u i v i w i ] (i=1,2,3) directions in the adjacent crystallite B 175 are assumed to be parallel to directions [100], [010] and [001] in the crystallite A. Thus, the relationship between the two groups of directions from two adjacent crystallites A and B can be related as follow: A B R w v u           =           0 0 1 1 1 1 ’ A B R w v u           =           0 1 0 2 2 2 ’ A B R w v u           =           1 0 0 3 3 3 The related misorientation matrix R is obtained as:           = 3 2 1 3 2 1 3 2 1 w w w v v v u u u R from which it is possible to determine the angle of misorientation across the specific boundary [F F Lange, Acta Metall, 15 (1967) 311, C T Young, J H Steele, Jr. and J L Lytton, Metall Trans 4 (1973) 2081]. The last step is to build the R matrix. First, three Kikuchi pole directions, e.g. [111], [200] and [220], of crystallite A are aligned to the beam direction with reference to the Kikuchi pattern, and the corresponding tilt angles are recorded. Then, three other Kikuchi pole directions, e.g. [011], ] 121 [ − and ] 11 1 [ − , of the crystallite B are aligned to the beam direction and the corresponding tilt angles are recorded. The corresponding tilt angles can 176 be calculated to (α i -β i ) using equations (A3) and (A4) when the directions [100] A , [010] A and [001] A are aligned to the beam direction. Then the three directions of grain B which are parallel to the directions [100] A , [010] A and [001] A can be determined using above tilt angles (α i -β i ) being [u 1 v 1 w 1 ], [u 2 v 2 w 2 ] and [u 3 v 3 w 3 ]. Now the misorientation matrix R is known and can be used to calculate the misorentation angle between the two adjacent crystallites.

Abstract (if available)

Abstract The aim of this study is to produce large quantities of bulk zirconium with an ultrafine grained microstructure and with enhanced properties. Accumulative roll bonding (ARB), a severe plastic deformation technique based on rolling, is chosen due to its availability in industrial environment. The texture, microstructure and mechanical behavior of bulk ultrafine grained (ufg) Zr fabricated by accumulative roll bonding is investigated by electron backscatter diffraction, transmission electron microscopy and mechanical testing. A reasonably homogeneous and equiaxed ufg structure, with a large fraction of high angle boundaries (HABs, ~70%), can be obtained in Zr after only two ARB cycles. The average grain size, counting only HABs (θ>15°), is 400 nm. (Sub)grain size is equal to 320 nm. The yield stress and ultimate tensile stress (UTS) values are nearly double those from conventionally processed Zr with only a slight loss of ductility. Optimum processing conditions include large thickness reductions per pass (~75%), which enhance grain refinement, and a rolling temperature (T ~ 0.3T_m) at which a sufficient number of slip modes are activated, with an absence of significant grain growth. Grain refinement takes place by geometrical thinning and grain subdivision by the formation of geometrically necessary boundaries. The formation of equiaxed grains by geometric dynamic recrystallization is facilitated by enhanced diffusion due to adabatic heating. Optical microscopy examination and shear testing suggest accepted bonding quality compared to that achieved in materials processed by diffusion bonding and that obtained in other ARB studies. Biocompatibility of ultrafine grained Zr processed by large strain rolling is studied by evaluating the behavior of human osteoblast cells. It is suggested that ultrafine grained Zr has a similar good biocompatibility as Ti6Al4V alloy and conventional Zr with a large grain size have.

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Asset Metadata

Creator Jiang, Ling (author)

Core Title Processing, mechanical behavior and biocompatibility of ultrafine grained zirconium fabricated by accumulative roll bonding

School Viterbi School of Engineering

Degree Doctor of Philosophy

Degree Program Mechanical Engineering

Publication Date 11/09/2008

Defense Date 09/29/2008

Publisher University of Southern California (original), University of Southern California. Libraries (digital)

Tag accumulative roll bonding,biocompatibility,OAI-PMH Harvest,shear testing,texture,ultrafine grained materials,zirconium

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Kassner, Michael E. (committee chair), Perez-Prado, Maria-Teresa (committee chair), Goo, Edward K. (committee member), Nutt, Steven R. (committee member), Phares, Dennis (committee member)

Creator Email lingjian@usc.edu,ljiang79@gmail.com

Permanent Link (DOI) https://doi.org/10.25549/usctheses-m1759

Unique identifier UC1127948

Identifier etd-Jiang-2496 (filename),usctheses-m40 (legacy collection record id),usctheses-c127-131002 (legacy record id),usctheses-m1759 (legacy record id)

Legacy Identifier etd-Jiang-2496.pdf

Dmrecord 131002

Document Type Dissertation

Rights Jiang, Ling

Type texts

Source University of Southern California (contributing entity), University of Southern California Dissertations and Theses (collection)

Repository Name Libraries, University of Southern California

Repository Location Los Angeles, California

Repository Email cisadmin@lib.usc.edu