oxide, which was used to obtain an estimate of the hydrogen concentration in niobium. This estimate was found to be 40% atomic H in a non heat

Transcription

1 ABSTRACT MAHESHWARI, PRATEEK. Surface Characterization of Impurities in Superconducting Niobium for Radio Frequency (RF) Cavities used in Particle Accelerators. (Under Dr. Dieter Griffis and Dr. James Rigsbee). Niobium (Nb) is the material of choice for Superconducting Radio Frequency (SRF) Cavities used in particle accelerators owing to its high critical temperature (T c = 9.2 K) and critical magnetic field ( 200mT). However, niobium tends to harbor interstitial impurities such as H, C, O and N, which are detrimental to cavity performance. Since the magnetic field penetration depth (λ) of niobium is 40nm, it is important to characterize these impurities using surface characterization techniques. Also, it is known that certain heat treatments improve cavity efficiency via interstitial impurity removal from the surface of niobium. Thus, a systematic study on the effect of these heat treatments on the surface impurity levels is needed. In this work, surface analysis of both heat treated and non heat treated (120 o C-1400 o C) large grain (single crystal) bulk niobium samples was performed using secondary ion mass spectrometry (SIMS) and Transmission Electron Microscopy (TEM). Impurity levels were compared on the surface using SIMS after various types of heat treatments expected to improve cavity performance, and the effect of these heat treatments on the surface impurities were examined. SIMS characterization of ion implanted standards of C, N, O, D showed that quantification of C, N and O impurities in Nb is achievable and indicated that H is very mobile in Nb. It was hence determined that quantification of H in Nb is not possible using SIMS due to its high diffusivity in Nb. However, a comparative study of the high temperature heat treated (600 o C o C) and non heat treated (control) samples revealed that hydrogen levels decreased by upto a factor of 100. This is attributed to the dissociation of the niobium surface oxide layer, which acts as a passivating film on the surface, and subsequent desorption of hydrogen. Reformation of this oxide layer on cool down disallows any re-absorption of hydrogen, indicating that the oxide acts as a surface barrier for absorption/desorption of hydrogen and that hydrogen does not diffuse in the oxide. Subsequent ion implantation of hydrogen in an anodized niobium sample thus provided a quantification factor of hydrogen in niobium

2 oxide, which was used to obtain an estimate of the hydrogen concentration in niobium. This estimate was found to be 40% atomic H in a non heat treated niobium sample. Such high levels of hydrogen observed in Nb before heat treatment ensures that is the main contributor to cavity degradation. TEM analysis was performed to study the effect of heat treatment on the surface oxide thickness of niobium. Results showed a continuous oxide layer with a sharp metal-oxide interface. No significant changes in the oxide thickness were seen after heat treatment. Time of Flight (TOF)-SIMS imaging was used to characterize the grain boundaries of large grain niobium bicrystals, since it was believed that impurity segregation at the grain boundaries of Nb might deteriorate cavity performance. Images showed segregation of carbon at the grain boundaries after 800 o C heat treatment of the samples, while no segregation of hydrogen and oxygen were seen for both non heat treated and heat treated samples. An important aspect of this study was the record-performance improvement of the 1400 o C heat treated cavity, which showed a 200% increase in the cavity efficiency. SIMS analysis of the surface of this sample showed high levels of titanium, down to 1µm depth, and it is speculated that this Ti might be responsible for high performance of the cavity by affecting the distribution of impurities within the penetration depth.

3 Copyright 2012 by Prateek Maheshwari All Rights Reserved

4 Surface Characterization of Impurities in Superconducting Niobium for Radio Frequency (RF) Cavities used in Particle Accelerators by Prateek Maheshwari A dissertation submitted to the Graduate Faculty of North Carolina State University in partial fulfillment of the requirements for the degree of Doctor of Philosophy Materials Science and Engineering Raleigh, North Carolina 2012 APPROVED BY: Dr. Dieter Griffis Committee Co-Chair Dr. James Rigsbee Committee Co-Chair Dr. Alan Batchelor Dr. Steve Shannon Dr. Ganapati Myneni

5 DEDICATION To My Parents, Family and My Wife, For their love and support ii

6 BIOGRAPHY Prateek received his Batchelor s in Engineering Degree in Metallurgical Engineering from Punjab Engineering College (PEC), Chandigarh, India. After completing his studies, he applied to the University of Florida (UF), USA, for Master s in Sciences in Materials Engineering, and was admitted with a College of Engineering Achievement Scholarship. At UF, he worked with Assistant Professor Dr. Jacob Jones on study of Ferroelectric and Piezolelectric materials using X-Ray Diffraction. He was then recommended for a doctoral degree at the Dept. of Materials Science and Engineering at the North Carolina State University, for a project under Dr. Dieter Griffis and Dr. James Rigsbee, working on Surface Characterization of Niobium. iii

7 ACKNOWLEDGMENTS I would like to express my sincere gratitude to the following: Dr. Dieter Griffis and Dr. James Rigsbee, for providing me with this excellent opportunity, guidance, and support. Fred Stevie for his expert guidance and mentorship, with his encouragement and support throughout my research. Dr. Ganapati Myneni and Dr. Gigi Ciovati, for providing me the opportunity to work on this project and helping me through out my PHD. Dr. Dale Batchelor and Dr. Steve Shannon for being my committee members and evaluating my work. Dr. Elaine Zhou, Roberto Garcia, and Chuck Mooney for their encouragement, guidance, and technical assistance. The Analytical Instrumentation Facility graduate students for their support and assistance. My parents and my wife for their love and support. iv

8 TABLE OF CONTENTS LIST OF FIGURES......xi LIST OF TABLES...xxii 1. Introduction Particle Accelerators Forces used in particle acceleration Electrostatic Accelerators Oscillating Field Accelerators Radiofrequency (RF) Cavities RF Cavity Materials SRF Technology for Accelerator Cavities Basic Properties of Superconductors Niobium Niobium SRF Cavity Performance Parameters Surface Impurities in Niobium Cavity Fabrication Surface Treatment Need for Niobium Surface Characterization Single Crystal vs Polycrystalline Nb Objective of this Thesis References Secondary Ion Mass Spectrometry (SIMS) Introduction to SIMS Static and Dynamic SIMS The sputtering process v

9 2.1.3 Introduction to SIMS Instrumentation Mass Analyzers used in SIMS Magnetic Sector Mass Analyzer Quadrupole Mass Analyzer Time of Flight (TOF) Mass Analyzer CAMECA IMS-6f Magnetic Sector Dynamic SIMS Instrument Primary Column Sample Chamber The Secondary Column Charge Neutralization ION TOF V SIMS Instrument References Experimental Sample Preparation Types of Heat Treatments Nanopolishing SIMS: Experimental Parameters Primary Beam Primary Beam Energy and Angle of Incidence Sputtering Rate and Detected Area SIMS Quantification SIMS Analysis of Residual Gas Species Raster Reduction SIMS Characterization of impurities in Niobium High Energy SIMS analysis Low Energy SIMS analysis vi

10 3.6 TOF-SIMS Characterization of Nb Bicrystals Sample Preparation TOF-SIMS Imaging: Analysis Conditions References Hydrogen Source of Hydrogen Contamination in Nb Cavities Effect of Hydrogen on properties of SRF Nb Residual Losses from Hydrides Critical Temperature (T c ) Electrical Resistance Q disease Solubility, Diffusion and Site Occupancy of Hydrogen in Nb Solubility Diffusion Site Occupancy SIMS Characterization of H in Nb Heat Treatments Discussion Conclusion from Results Low Energy SIMS Characterization of H in Nb Results and Discussion Conclusion SIMS Characterization of Hydrogen in Niobium Oxide Experimental Results and Discussion Conclusion vii

11 4.7 SIMS Characterization of Non-Nanoplished Nb TOF-SIMS Imaging of Hydrogen in Large Grain Nb Bicrystals Results and Discussion Conclusion Effect of Hydrogen on SIMS Nb Secondary Ion Intensities References Oxygen Source of Oxygen Contamination in Nb Cavities Effect of Oxygen on the properties of SRF Nb Residual Losses due to Oxides Oxide as a Source for Magnetic Flux Pinning Sites Effect on Critical Temperature (T c ) and Resistivity (10K) Solubility and Diffusion of Oxygen in Nb Solubility Oxidation of Nb Pressure-Temperature relationships in the Nb-O system Diffusion of Oxygen in Nb Oxygen Pollution Model SIMS Characterization of Oxygen in Nb Heat Treatments Results and Discussion Conclusion Experimental Results and Discussion Conclusion Low Energy SIMS Characterization of Oxygen in Nb viii

12 5.7.1 Results and Discussion Conclusions TEM Characterization of the Surface Oxide of Niobium Experimental Results and Discussion Conclusion TOF-SIMS Imaging of Oxygen in Large Grain Nb Bicrystals Results and Discussion Bicrystals having different crystallographic orientation combinations Conclusion References Carbon and Nitrogen Effect of Carbon and Nitrogen on properties of SRF Nb Solubility, Diffusion of Carbon and Nitrogen in Nb Solubility of C in Nb Pressure-Temperature Relationships in the Nb-N system Diffusion of Carbon and Nitrogen in Nb SIMS Characterization of Carbon and Nitrogen in Nb Results Nitridation Experiments Conclusions Low Energy SIMS Characterization of Carbon and Nitrogen in Niobium Results and Discussion Conclusion TOF-SIMS Imaging of Carbon in Large Grain Nb Bicrystals Results and Discussion ix

13 6.5.2 Conclusion References Cavity Performance Results, Conclusions and Future Work Cavity Performance Technique for Cavity Performance Measurements Cavity Performance Results Hydrogen: Primary Factor in Cavity Performance Conclusions Hydrogen Oxygen Carbon and Nitrogen Recommendations for Future Work Nb Thin Films for SRF Applications In situ High Temperature TEM Studies of Nb Surface Oxide Effect of an Oxidizer after Buffer Chemical Polishing Bulk vs Surface Hydrogen in Nb References x

14 LIST OF FIGURES Figure 1.1: The methods to investigate the micro-world... 1 Figure 1.2: General principle of an electrostatic accelerator... 4 Figure 1.3: The linear accelerator... 6 Figure 1.4: RF muti-cell cavity made of Niobium at Jefferson Laboratories... 8 Figure 1.5: Working of an RF cavity : As an electron moves forward, the electric field polarity is switched to provide and accelerating push... 8 Figure 1.6: (a) and (b): Cubic (BCC) structure of Niobium, (c): Effectively 2 atoms per unit cell Figure 1.7: Tetrahedral (left) and Octahedral sites (right) of Nb Figure 1.8: Deep drawing of a Nb sheet to cavity shape Figure 2.1: SIMS sputtering process Figure 2.1: Block diagram of basic components of a SIMS instrument Figure 2.3: Double focusing mass spectrometer Figure 2.4: Schematic of a quadrupole mass spectrometer Figure 2.5: TOF mass spectrometer with a reflectron Figure 2.6: Schematic of a CAMECA IMS-6f SIMS instrument Figure 2.7: Parts of the primary column in a CAMECA IMS-6F instrument Figure 2.8: Schematic of a duoplasmatron source Figure 2.9: Accel/Decel system for the CAMECA IMS-6f xi

15 Figure 2.10: CAMECA IMS-6F Cs Microbeam source Figure 2.11: Secondary ion extraction and optics Figure 2.12: Normal trajectories of varying energies in the mass spectrometer Figure 2.13: Non-normal trajectories of the same energies in the mass spectrometer Figure 2.14: Working of an electron multiplier Figure 2.15: Schematic of a faraday cup detector Figure 2.16: Schematic of the working of the electron gun Figure 2.17: Schematic of an ION TOF V Time of Flight SIMS instrument Figure 2.18: ION TOF V Time of Flight SIMS instrument at AIF Figure 2.19: Schematic of a Liquid Metal Ion Gun Figure 2.20: Schematic of a Liquid Metal Ion Source Figure 3.1: Flowchart for Nb sample preparation for surface analysis Figure 3.2: Schematic diagram of the induction heating system used Figure 3.3: Optical images of niobium surface (a) unpolished and (b) after nano-polishing. 73 Figure 3.4: The effect of primary ion beam energy on the secondary ion yield Figure 3.5: Effect of primary beam energy on sputter yield for normal and 60 angle of incidence from the normal Figure 3.6: Effect of angle of incidence on secondary ion yield using an O 2 and Ar beam xii

16 Figure 3.7: The effect of a smaller raster/detected area ratio on the depth profile of As in Si : A denotes the larger raster area (220 x 220 µm) and B denotes a smaller raster area (80 x 80 µm), shaded area denotes the detected area (60µm diameter) Figure 3.9: Quantification procedure followed in SIMS Figure 3.10: Raster reduction Cts vs Time profile for NbN in Nb Figure 3.11: EBSD results for control (a) and heat treated (800C/3hrs, 120C/24hrs) (b) bicrystal samples. Color coding similar for both samples Figure 4.1: Nb-H phase diagram: Hydride phase formation at low temperature and H concentration Figure 4.2: T c vs H conc. observed by various groups using the DC resistivity method Figure 4.3: Increase in resistance of Nb with increasing H conc. at 285K (a); Similar for (b) but at 9.5 K Figure 4.4: Faster cooling rates can avoid Q o degradation by surpassing hydride precipitation temperatures Figure 4.5: Increasing solubility of H in Nb with decreasing temperatures at various pressures (mm Hg) Figure 4.6: P-composition isotherms for different concentrations of H in Nb Figure 4.7: Nb-H and Nb-D diffusion coeff. vs Temperature Figure 4.8: Diffusion processes of hydrogen at different temperatures Figure 4.9 : D in Si implant shows an implant peak Figure 4.10: D in Nb implant shows a constant D signal xiii

17 Figure 4.11: Lattice expansion in Nb due to H Figure 4.12: Intense niobium hydride peaks in a non heat treated sample Figure 4.13: Less intense hydride peaks in a 800C/3hrs, 400C/20min heat treated Nb sample as compared to control, higher hydrides are not seen after heat treatment Figure 4.14(a): Raster reduction for H and Nb for a control sample Figure 4.14(b): Raster reduction for H and Nb for a heat treated sample Figure 4.15: SIMS H/Nb vs Depth data for 600C-1400C heat treated Nb samples vs contro Figure 4.16: SIMS H/Nb vs Depth data for different high temperature heated-low temperature baked Nb samples vs control Figure 4.17: SIMS H/Nb vs Depth data for high temperature heating/low temperature heated Nb samples vs control Figure 4.18: Hydrogen average partial pressure in the furnace vs H/Nb ratio measured by SIMS after heat treatment Figure 4.19: SIMS analysis (cts vs time) of 800C/3hrs, 400C/20 min heat treated sample: (a) before and (b) after etching in HF for 15 min and rinsing with water, H levels seen to increase to control sample H levels in (b) Figure 4.20: H/Nb levels remain low even after high pressure water rinsing of a heat treated sample owing to the surface oxide layer Figure 4.21: H/Nb Ratio vs Depth for control vs 120C/48hrs baked large grain Nb samples xiv

18 Figure 4.22: H/Nb Ratio vs Depth for control vs 1200C/6hrs heat treated large grain Nb samples Figure 4.23: H/Nb Ratio vs Depth for control and 1200C-1400C, 120C/12hrs heat treated large grain Nb samples. No H in the 2-6nm depth region Figure 4.24: H/Nb Ratio vs Depth for control vs 800C/3hrs, 400C/15min in 10 mbar N 2, 120C/6hrs heat treated large grain Nb samples. No H on the Nb surface Figure 4.25: H and D implant peaks seen in niobium oxide; insignificant change in the Nb intensities from in the oxide to in the substrate; D implanted to show peak at the interface 130 Figure 4.26: H concentration estimate for H in Nb in a control niobium sample, the atomic concentration is found to be 2x10 atoms/cm, which when divided by the Nb atomic density (5.44x10 atoms/cm) provides the concentration estimate to be 37 atomic % H in Nb Figure 4.27: H/Nb vs time (sec) for nanopolished and non-nanopolished control samples. 133 Figure 4.28: Dynamic SIMS results for control (left) and heat treated (800C/3hrs, 120C/24hrs) (right) bicrystal samples. The sputtering rate differences may be attributed to the different crystal orientations. As seen in previous data, lower H levels observed in the heat treated sample Figure 4.29: TOF SIMS ion images for control (a) and heat treated (800C/3hrs, 120C/24hrs) (b) bicrystal sample grain boundaries. No segregation of H seen at the grain boundary for both samples Figure 4.30: TOF SIMS ion images for control (a) and heat treated (1200C/6hrs) (b) bicrystal sample grain boundaries with similar crystal orientation of the upper crystal for both samples and lower crystal for both samples Figure 4.31: Nb intensities a factor of 3 higher for a control sample vs a heat treated sample (HT= 800C/3hrs, 400/20min) xv

19 Figure 4.32: Both Nb and H follow similar trends in the control to heat treatment intensity ratios of secondary ions vs heat treatment,. H intensity ratios have been scaled for comparison Figure 4.33: HT= 1000C/2hrs,120C/12hrs; Comparison of the non heat treated (control) and HT oxygen implants; O RSF seen to decrease by a factor of 3 for the heat treated implant 140 Figure 4.34: Mass spectra of Control (left) and 800C/3hrs, 400C/20min heat treated Nb (right) samples Figure 4.35: Possible secondary ion reactions relating to increased Nb intensities in SIMS anlysis of a control Nb sample Figure 5.1: Dependence of T c on solute (O) concentration; O seen to decrease T c linearly with increasing concentration Figure 5.2: Nb-O phase diagram Figure 5.3: Oxygen solubility in Nb follows an Arrhenius plot Figure 5.4: Kinetics of oxidation of Nb, 3hr periods; q (mg/cm) is the oxidation characteristic (ratio of the gain in weight to the initial surface area) Figure 5.5: Pressure-Temperature diagram for Nb-O system Figure 5.6: Diffusion Coeff. vs Temp (K) for oxygen in Nb Figure 5.7: X-Ray diffuse scattering results: only 2nm diffusion of oxygen seen after 145C/48hrs baking Figure 5.8: SIMS Depth profile of O in Nb implant Figure 5.9(a): Raster reduction for O and Nb for a control sample, some background contribution seen indicating that oxygen is near detection limit of the instrument xvi

20 Figure 5.9(b): Raster reduction for O and Nb for a heat treated sample, showing oxygen above detection limit and the measured oxygen levels have no background contribution Figure 5.10: Oxygen concentration for heat treated samples (600C-1400C) and control Figure 5.11: Oxygen concentration for various high temperature heated/low temperature baked heat treated samples vs. non heat treated sample (control) Figure 5.12: Oxygen concentration for 800C/3hrs, 400C/20min heat treated sample as compared to a control Nb sample Figure 5.13: Oxygen concentration for 800C/3hrs, 400C/20min (N 2 at 10mbar), 120C/6hrs; and the 1000C/2hrs, 800C/10min (N 2 at 10mbar) heat treatments as compared to a control sample Figure 5.14: TOF-SIMS mass spectrum of 1400C /3hrs heat treated sample. High Ti levels detected on the surface Figure 5.15: Concentration profile of Ti in Nb implant standard; implant peak clearly seen Figure 5.16: Concentration profile of Ti in Nb for 1400/3hrs and 1200C/2hrs, 120C/12 hrs heat treated samples in comparison to a control sample Figure 5.17: Concentration profile of O in Nb for 1400/3hrs and 1200C/2hrs, 120C/12 hrs heat treated samples. O concentration seems to match the Ti conc. profiles seen in figure Figure 5.18: Concentration profile of Ti in Nb for 1400/3hrs, 120C/12 hrs heat treated sample, using low energy O 2 primary ion beam (1.25keV). Ti concentration estimated to be about 1% atomic fir 5-50nm Figure 5.20: Concentration profile of O in Nb for 1200C/6 hrs heat treated sample in comparison to a control sample (Low energy SIMS, 6keV Cs) xvii

21 Figure 5.21: Concentration profile of O in Nb for 800C/3hrs, 120C/12hrs and 1200C/2hrs, 120C/12hrs heat treated samples in comparison to a control sample (Low energy SIMS, 6keV Cs) Figure 5.22: Concentration profile of O in Nb for 1400C/3hrs, 120C/12hrs heat treated sample in comparison to a control sample (Low energy SIMS, 6keV Cs) Figure 5.23: Concentration profile of O in Nb for (800C/3hrs, 400C/15min at 10 mbar N 2, 120C/6hrs) heat treated sample in comparison to a control sample (Low energy SIMS, 6keV Cs) Figure 5.24: TEM image of a control Nb sample Figure 5.25: TEM image of a control (a) and heat treated (800C/3hrs, 140C/3hrs) Nb (b) sample Figure 5.26: TOF SIMS ion images of control (a) and heat treated (800C/3hrs, 120C/24hrs) (b) bicrystal samples Figure 5.27: TOF SIMS ion images for control (a) and heat treated (1200C/6hrs) (b) bicrystal sample grain boundaries with similar crystal orientation combination for both samples Figure 6.2: Nb-C phase diagram Figure 6.3: Phase diagram of the N-Nb system Figure 6.4: Equilibrium pressure-temperature-concentration diagram of the Nb-N system 192 Figure 6.5: Diffusion Coeff. vs Temp (K) for nitrogen in Nb Figure 6.6: Depth profile of C and N in Nb implant Figure 6.7(a): Raster reduction for C and NbN for a control sample xviii

22 Figure 6.7(b): Raster reduction for C and NbN for a control sample Figure 6.8: Carbon concentration for heat treated samples (600C/-1400C) versus non heat treated sample (control) Figure 6.9: NbN concentration for heat treated samples ( C) versus non heat treated sample (control) Figure 6.10: Carbon concentration for various high temperature heated/low temperature baked heat treated samples versus non heat treated sample (control) Figure 6.11: NbN concentration for various high temperature heated/low temperature baked heat treated samples versus non heat treated sample (control) Figure 6.12: NbN concentration for 1200C/2hrs, 120C/12hrs; and the 1400C/3hrs, 120C/12hrs heat treatments as compared to a control sample Figure 6.13: C concentration for 800C/3hrs, 400C/20min heat treated sample as compared to a control Nb sample Figure 6.14: C concentration for 800C/3hrs, 400C/15min (N 2 at 10mbar), 120C/6hrs; and the 1000C/2hrs, 800C/10min (N 2 at 10mbar) heat treatments as compared to a control sample205 Figure 6.15: N concentration for 800C/3hrs, 400C/20min heat treated sample as compared to a control Nb sample Figure 6.16: NbN concentration for 800C/3hrs, 400C/15min (N 2 at 10mbar) and 800C/3hrs, 400C/15min (N 2 at 10mbar), 120C/6hrs heat treatments, as compared to a control sample; (SIMS 6keV Cs) Figure 6.17: N concentration for 1000C/2hrs, 800C/10min (N 2 at 10mbar) as compared to a control sample xix

23 Figure 6.18: Concentration profile of C in Nb for 120C/48 hrs baked sample in comparison to a control sample. High C levels seen after baking (Low energy SIMS, 6keV Cs) Figure 6.19: Concentration profile of NbN for 120C/48 hrs baked sample in comparison to a control sample. No difference observed after baking (Low energy SIMS, 6keV Cs) Figure 6.20: Concentration profile of C in Nb for 1200C/6 hrs heat treated sample compared with a control sample. High C levels seen after baking (Low energy SIMS, 6keV Cs) Figure 6.21: Concentration profile of N for the 1200C/6 hrs heat treated sample compared with a control sample. (Low energy SIMS, 6keV Cs) Figure 6.22: Concentration profile of C in Nb for 1200C/2hrs with 120C/12hrs and 1400C/3hrs with 120C/12hrs heat treated samples in comparison to a control sample (Low energy SIMS, 6keV Cs) Figure 6.23: Concentration profile of NbN for 1200C/2hrs with 120C/12hrs and 1400C/3hrs with 120C/12hrs heat treated samples in comparison to a control sample (Low energy SIMS, 6keV Cs) Figure 6.24: C concentration for 800C/3hrs with 400C/15min (N 2 at 10mbar) and 800C/3hrs with 400C/15min (N 2 at 10mbar) and with 120C/6hrs heat treatments, as compared to a control sample; (SIMS 6keV Cs) Figure 6.25: TOF SIMS ion images of C for control (a) and heat treated (800C/3hrs with 120C/24hrs) (b) bicrystal sample grain boundaries Figure 6.26: TOF SIMS ion images of C for control (a) and heat treated (1200C/6hrs) (b) bicrystal sample grain boundaries with similar crystal orientation of the upper crystal for both samples and lower crystal for both samples Figure 7.1: Cavity with RF probes Figure 7.2: SIMS Analysis of Nb thin film on c-plane sapphire (14.5 kev Cs) xx

24 Figure 7.3: H/Nb ratio vs time for water rinsed and nitric acid rinsed Nb Figure 7.4: DSC data for a control large grain Nb sample. High H levels seen using SIMS should lead to a phase transformation in the material on heating up to 200C, which is not seen in this data Figure 7.5: XRD peak overlap for control and 1250C/6hrs heat treated polycrystalline Nb samples. No expansion seen in the control sample, indicating H not present at significant levels in the bulk xxi

25 LIST OF TABLES Table 1.1: Comparison of superconducting cavities vs normal conducting cavities Table 1.2: Properties of Niobium Table 2.1: Different mass analyzers used in SIMS Table 3.1: Table showing the various types of heat treatments performed Table 3.2: Table relating the heat treatments to the respective furnaces Table 4.1: Comparison of H and D diffusion coefficients in Nb and Si respectively at 300K Table 4.2: Heat treatments and H 2 pressure in the furnace Table 4.3: Heat treatments and H 2 pressure in the furnace Table 4.4: Heat treatments and H 2 pressure in the furnace Table 4.5: Average H/Nb ratios obtained from SIMS data after various heat treatments Table 5.1: Decreasing critical temperature with increasing O concentration in Nb. Resistivity increase in Nb with increase in the O concentration Table 7.1: Cavity performance results for cavities heat treated with the Nb samples; next to the Q o values (in parentheses) is the magnetic field for the values reported. All measurements taken at 2 K temperature xxii

26 1. Introduction The effort to understand the basic building blocks of matter is an area of much study in modern physics. High energy particle beams have proven to be excellent tools for these studies and the current understanding of elementary particle physics would not be possible without them. 7 It is known from the development of science that matter is made of atoms which are composites of a nucleus and electrons while nuclei are composites of protons and neutrons. The latter consist of even smaller particles called quarks. As a microscope for the microworld, the resolution of particle beams is limited by the debroglie wavelength λ 1 : λ = h/p (1.1) where p is the momentum and h is the planck s constant The above equation can be rephrased, in the form of energy E and relativistic speed of the particle beam v: λ = hc/ev (1.2) where c is the speed of light. It can be seen from equation 1.2 that the smaller the observed substance, the higher a beam energy is needed. 2 The methods to investigate the micro-world are mentioned in fig.1.1. Figure 1.1: The methods to investigate the micro-world 2 1

27 Another important aspect of elementary particle physics is the study of new, mostly very short lived particles, produced via collision of high energy particle beams on to target materials. These collisions help answer some of the fundamental open questions in physics, concerning the basic laws governing the interactions and forces among elementary objects. The energy needed to produce a particle via collision follows directly from the relation: 3 E= mc 2 (1.3) Most particles can only be produced along with their anti-particles in pairs. For example, electrons and positron can only be produced as a pair using high energy γ rays and their interaction with a heavy nucleus. As a result of conservation of momentum, a reaction of this kind can only take place near a heavy nucleus. The nucleus itself gains some momentum and energy, reducing the energy available for particle production. The γ ray energy needed for particle production is therefore always higher than that, given by equation 1.3 : 3 E gamma ray > 2m e c 2 = x J 1MeV...(1.4) The above discussion shows that high energies are needed to study atomic and sub-atomic building blocks of matter. Achievement of high energy particles beams for this purpose is why devices known as particle accelerators are extremely important. 1.1 Particle Accelerators A particle accelerator is a device that uses electromagnetic fields to propel charged particles to high speeds and to contain them in well-defined beams. There are two basic classes of accelerators, known as electrostatic and oscillating field accelerators. Electrostatic accelerators use static electric fields to accelerate particles. A small-scale example of this class is the cathode ray tube in an old television set. Other examples are the Cockcroft Walton 4 generator and the Van de Graaf generator 5. The achievable kinetic energy for particles in these devices is limited by electrical breakdown. Oscillating field accelerators, on the other hand, use radio frequency electromagnetic fields and circumvent the breakdown problem. 6 Facilities such as CERN, Jefferson Laboratories, Fermilab etc. utilize this technology for high energy physics. 2

28 1.1.1 Forces used in particle acceleration 7 In order to reach high kinetic energies, a directional force must be exerted on a particle to achieve acceleration. While electrostatic force is commonly used for the initial stages of particle acceleration, electromagnetic force, a combination of electric and magnetic forces, is employed in modern high energy accelerators to provide the acceleration needed to reach the very high particle velocities required. When a particle (example an electron) of velocity v passes though a volume containing a magnetic field B and an electric field E, the force acting upon it is given by the Lorentz equation: F = e (v x B + E).(1.5) As the particle moves from point r 1 to r 2, its energy changes by the amount: ΔU = F. dr = e (v x B + E). dr ; (from r 1 to r 2 ).(1.6) As the particle traverses the field, the path element dr is always parallel to the velocity vector v. The vector v x B is thus perpendicular to dr, thereby: (v x B). dr = 0 Hence the magnetic field B does not change the energy of the particle. Acceleration involving an increase in energy can thus only be achieved by the use of electric fields. The gain in energy thus follows directly from the equation 1.6 : ΔU = e E. dr ; (from r 1 to r 2 ) = ev.(1.7) where V is the voltage crossed by the particle. Although magnetic fields do not contribute to the energy of the particle, they play a very important role when forces are required which act perpendicular to the particle s direction of motion. Thus, magnetic forces are used to steer and focus particle beams Electrostatic Accelerators 8 The simplest particle accelerators use a constant electric field between two electrodes, produced by a high voltage generator. The source of the charged particles is electrically biased to act as the first accelerating electrode, which in the case of electron beams, is a thermionic cathode. Protons, as well as light and heavy ions, are extracted from the gas phase by using an additional DC or high frequency voltage to ionize a rarified gas, often, but not always, by producing plasma inside the particle source. Charged particles are then 3

29 continuously emitted from the plasma inside the particle source, and are accelerated by the electric field as in equation 1.7. Fig. 1.2 schematically illustrates the operation of such an accelerator. In electrostatic accelerators, the maximum achievable energy is directly proportional to the maximum voltage which can be developed across the accelerator, and this determines achievable particle acceleration and thus the achievable energy. Thus, the maximum achievable energy is limited by breakdown and voltages of a only few MV are technically possible. Hence, substantially higher energies useful to study the building blocks of matter cannot be achieved from this technology (equation 1.2). The utilization of these accelerators is prevalent in various analytical instrumentation, both ion beam and electron beams (SEM, TEM, SIMS, FIB etc), commercial technologies such as cathode ray tubes, and insertion sources for oscillation field accelerators discussed below. Figure 1.2: General principle of an electrostatic accelerator Oscillating Field Accelerators To overcome the insufficiently high acceleration limit imposed on electrostatic accelerators by electrical discharge, techniques involving a series of oscillating, high voltage sources are used for acceleration of particles to the higher energies desired. The electrodes in these 4

30 oscillating, high voltage sources can either be arranged to accelerate particles in a line or in an arc, depending on whether or not the particles are deflected by a magnetic field while they are accelerated. Linear oscillating field accelerators or linacs consist of a series of metal drift tubes arranged along the beam axis which connected, with alternating connection polarity, to a radiofrequency (RF) supply. The RF supply delivers a high frequency alternating voltage of the form U(t) = U max sin(ωt). The acceleration process goes as follows. The particles exit the ion source and enter the first RF energized drift tube with a velocity v 1. During a the first half period of the RF, the voltage applied between the first drift tube and the source acts to accelerate the particles leaving the ion source. The particles reach the first drift tube with a velocity v 1. Then they pass through this drift tube, which acts as a faraday cage which shields them from external fields. Meanwhile, the direction of the RF fields is reversed without the particles feeling any effect since they are within the drift tube. When the particles reach the gap between the first and second drift tubes, they again undergo acceleration. This process is repeated for the series of drift tubes. During the acceleration, the velocity increases monotonically, but the frequency of the alternating voltage must be kept constant to keep costs low. 16 This requires that the size of the gaps between the drift tubes must increase (Fig.1.3). Thus, for a linac, the energy imparted to the particles is proportional to the number of linear stages traversed by the particles. This poses a limitation on the structure of these accelerators, since higher beam energies require a longer linac which increases structural and operating costs. The largest contemporary electron linac is situated at the Stanford Linear Accelerator Center (SLAC) in California, which is about 3km long and reached final energies of 50GeV. 9 5

31 Figure 1.3: The linear accelerator 10 In the circular accelerator, particles move in a circle until they reach sufficient energy. The particles are typically deflected into a circle using electromagnetic fields. The advantage of circular accelerators over linear accelerators is that the ring topology allows continuous acceleration, and that a circular accelerator is smaller than a linear accelerator of comparable accelerating capability. Depending on the energy and the particle being accelerated, circular accelerators suffer a disadvantage in that the particles emit synchrotron radiation. Synchrotron radiation results from the property that when any charged particle is accelerated, it emits electromagnetic radiation and secondary emissions. 11 As a particle traveling in a circle is always accelerating towards the center of the circle, it continuously radiates towards the tangent of the circle. This radiation is called synchrotron light and depends highly on the mass of the accelerating particle. For this reason, many high energy electron accelerators are linacs while certain accelerators (synchrotrons) are built specially for synchrotron light (Xrays). 11 Almost all the oscillating field accelerators currently use radio frequencies (3kHz-300GHz) to accelerate particles to extremely high velocities. In order to switch the electric fields at these high frequencies, devices known as Radiofrequency Cavities are 16 used instead of the earlier discussed drift tubes. These RF cavities exploit the fact that, upon reaching energies of a few MeV, the particles have already reached velocities very close to the speed of light. As they are accelerated, the electrons increase in mass but their velocity remains relatively 6

32 constant, allowing cavity structures of the same size to be used along the whole length of a linear accelerator Radiofrequency (RF) Cavities A radio frequency (RF) or microwave cavity is a special type of RF resonator, consisting of a closed (or largely closed) metal structure that confines electromagnetic fields in the microwave region of the spectrum (3kHz-300GHz). The structure is either hollow or filled with dielectric material. A resonant microwave cavity acts 13 similarly to a resonant circuit with extremely low loss at its frequency of operation, resulting in extremely high quality factors (energy stored/ power dissipated) compared to circuits made with separate inductors and capacitors at the same frequency. In addition to particle accelerators, they are used in oscillators and transmitters to create microwave signals, and as filters to separate a signal at a given frequency from other signals, in equipment such as radar facilities, microwave relay stations, satellite communications, and microwave ovens. 13 The cavity's interior surfaces reflects waves of a specific frequency. When a wave that is resonant with the cavity enters, it bounces back and forth within the cavity, with low loss, like a standing wave. As more wave energy enters the cavity, it combines with and reinforces the standing wave, increasing its intensity. 13 In particle accelerators such as the Continuous Electron Beam Accelerator Facility (CEBAF) at Jefferson Laboratory, about 330 RF cavities are used to accelerate the beam to 6 GeV. 40 The beam passes through the apertures of the cavities, often tunable wave reflection grids, in succession. A collector electrode is provided to intercept the beam after passing through the cavities. The first cavity causes bunching of the particles passing through it. The bunched particles travel in a field-free region where further bunching occurs, then the bunched particles enter the second cavity giving up their energy to excite it into oscillations. This forms a particle accelerator that works in conjunction with a specifically tuned cavity by the configuration of the structures. On the beam-line of such an accelerator system, there are specific sections that are RF cavities, especially in the linac. 14 7

33 Fig.1.4 and 1.5 show an RF cavity structure and the operation of an RF cavity respectively. 15 Such cavities are multi-cell cavities and their specific shapes are designed to minimize power losses. Apart from their geometry, the material used to make these cavities also plays a critical role in minimizing power losses. Due to the critical role of electric field forces in particle acceleration, and the need for extremely high energies, materials for such devices have to provide minimum resistance to minimize losses to provide maximum efficiency. 15 Figure 1.4: RF muti-cell cavity made of Niobium at Jefferson Laboratories 15 Figure 1.5: Working of an RF cavity : As an electron moves forward, the electric field polarity is switched to provide and accelerating push 15 8

34 1.3 RF Cavity Materials As mentioned in section 1.2, the ideal cavity material should impart maximum energy to the particle, thus minimizing losses. The material should have very low resistance and high conductivity. Conventionally, copper was used to manufacture cavities due to its low resistivity and easy formability to the desired cavity shape. 16 Although, for all practical applications, copper (Cu) is an excellent conductor, it s relatively high resistivity limits the maximum electric field due to thermal and electrical breakdown. This limitation has been addressed via the use of specialized materials: superconductors, to make RF cavities. The use of these new materials gave this technology a new name: Superconducting Radiofrequency (SRF) technology SRF Technology for Accelerator Cavities 16 The strongest incentive to use superconducting materials for accelerator cavities is for devices that operate in a continuous wave mode. For such operation, the power dissipation or loss in the walls of a copper structure is substantial, and this is where superconducting materials are important. The microwave surface resistance of a superconductor is typically five orders of magnitude lower than that of copper, and therefore superconducting cavity efficiencies are five orders of magnitude higher. The dissipated power/ unit length is given by: P/L = E 2 acc/ [(r a /Q o )Q o ]...(1.8) Here r a /Q o is the geometric shunt impedance in Ω/m which depends primarily on the geometry of the structure. Q o is the cavity quality factor defined as the ratio of the energy stored in the cavity to the energy lost in one rf period. The typical Q o of a superconducting cavity is in the range of 10 9 to Table 1.1 shows the performance characteristics of a 500MHz cavity made of copper compared to a cavity made of a superconducting material (Niobium). 9

35 Table 1.1 : Comparison of superconducting cavities vs normal conducting cavities 16 Characteristics Superconducting Cu Conductor Q o 2 x x 10 4 r a /Q o (500MHz) 330 (ohm/m) 900 (ohm/m) P/L for E acc = 1MV/m 1.5 (Watt/m) (Watt/m) AC Power, E acc = 1MV/m 0.54 (kw/m) 112 (kw/m) AC Power, E acc = 5MV/m 13.5 (kw/m) 2800 (kw/m) It is evident from the above data that the dissipated power per meter for niobium cavities is reduced by a factor of 4 x The gain in performance for this type of cavity versus a Cu cavity does not quite reach this level of improvement since the dissipated power for a superconducting cavity also has to take into account the power required to cool down the cavity to liquid helium temperature, currently a requirement for nearly all superconducting materials for optimum performance. From thermodynamic calculations, the net power dissipated for a superconducting cavity is about a factor of 200 lower than a conventional copper cavity. The advantage of the power savings provided by superconducting cavities is clear and, since the dissipated power increases with the square of the operating field (equation 1.8), only these cavities can economically provide the needed voltages to achieve desired accelerations. Hence, almost all SRF cavities these days are made of the superconducting material, Niobium Basic Properties of Superconductors 17 The physical properties of superconductors vary from material to material. These properties include heat capacity and the critical temperature, critical field, and critical current density at which superconductivity ceases. On the other hand, there is a class of properties that are independent of the underlying material. For instance, all superconductors have exactly zero resistivity to low applied currents when there is no magnetic field present or if the applied field does not exceed a critical value. The existence of these "universal" properties implies 10

36 that superconductivity is a thermodynamic phase, and thus possesses certain distinguishing properties which are largely independent of microscopic details. a) Zero Resistvity 17 In a normal conductor, an electric current may be visualized as a fluid of electrons moving across a heavy ionic lattice. The electrons are constantly colliding with the ions in the lattice, and during each collision some of the energy carried by the current is absorbed by the lattice and converted into heat, which is essentially the vibrational kinetic energy of the lattice ions. As a result, the energy carried by the current is constantly being dissipated giving rise to the phenomenon of electrical resistance. However, in a conventional superconductor, the electronic fluid cannot be resolved into individual electrons. Instead, it consists of bound pairs of electrons known as Cooper pairs. This pairing is caused by an attractive force between electrons from the exchange of phonons. Due to quantum mechanics, the energy spectrum of this Cooper pair fluid possesses an energy gap, meaning there is a minimum amount of energy ΔE that must be supplied in order to excite the fluid. Therefore, if ΔE is larger than the thermal energy of the lattice, given by kt, where k is Boltzmann's constant and T is the temperature, the fluid will not be scattered by the lattice and thus would lead to zero resistance. b) Critical Temperature and Critical Magnetic Field 18 The characteristics of superconductivity appear when the temperature is lowered below a critical temperature T c. The value of this critical temperature varies from material to material. Conventional superconductors usually have critical temperatures ranging from around 20 K to less than 1 K. Solid mercury, for example, has a critical temperature of 4.2 K. Similarly, at a fixed temperature below the critical temperature, superconductivity ceases to exist when an external magnetic field is applied which is greater than the critical magnetic field. This is because the Gibbs free energy of the superconducting phase increases quadratically with the magnetic field while the free energy of the normal phase is roughly independent of the magnetic field. If the material super-conducts in the absence of a field, then the superconducting phase free energy is lower than that of the normal phase and so for 11

37 some finite value of the magnetic field the two free energies will be equal and a phase transition to the normal phase will occur. Niobium has a critical magnetic field of 200mT. c) Meissner Effect and the London Penetration Depth 19 When a superconductor is placed in a weak external magnetic field, and cooled below its transition temperature, the magnetic field is ejected. The Meissner effect does not cause the field to be completely ejected but instead the field penetrates the superconductor only to a very small distance, characterized by a parameter λ, called the London penetration depth, decaying exponentially to zero within the bulk of the material. The Meissner effect is a defining characteristic of superconductivity. For most superconductors, the london penetration depth is about 100 nm. A superconductor with little or no magnetic field within it is said to be in the Meissner state. The Meissner state breaks down when the applied magnetic field is too large. Superconductors can be divided into two classes according to how this breakdown occurs. In Type I superconductors, superconductivity is abruptly terminated when the strength of the applied magnetic field rises above a critical value. In Type II superconductors, raising the applied magnetic field past a critical value leads to a mixed state (also known as the vortex state) in which an increasing amount of magnetic flux penetrates the material, but there remains no resistance to the flow of electric current as long as the current is not too large. At a second critical magnetic field strength, superconductivity is terminated. Most pure elemental superconductors, except niobium, technetium, vanadium and carbon nanotubes, are Type I, while almost all impure and compound superconductors are Type II Niobium Niobium (Nb) has been the material of choice for SRF cavities due to its excellent superconducting properties and easy formability. Niobium is a lustrous, grey, ductile, paramagnetic metal in group V of the periodic table, although it has an atypical configuration in its outermost electron shells compared to the rest of the members of the same group. Niobium becomes a superconductor at cryogenic temperatures and, at atmospheric pressure, Nb has the highest critical temperature of the elemental superconductors: 9.2 K. It has a high critical magnetic field (200mT) with a London penetration depth of 40nm. In addition, it is 12

38 one of the three elemental Type II superconductors, along with vanadium and technetium. The superconductive properties are strongly dependent on the purity of the niobium metal. When very pure, it is comparatively soft and ductile, but impurities increase hardness. Fig. 1.6 shows the lattice structure of Niobium, which is Body Centered Cubic (BCC) with effectively 2 atoms per unit cell. Figure 1.6: (a) and (b): Cubic (BCC) structure of Niobium, (c): Effectively 2 atoms per unit cell The BCC structure of Niobium allows the possibility of 12 tetrahedral interstitial sites and 6 octahedral interstitial sites depicted in Fig Nb atoms Figure 1.7: Tetrahedral (left) and Octahedral sites (right) of Nb 21 13

39 The size of the octahedral site is given by the relation, R oct = 0.155r Nb and that of the tetrahedral site is given by: R tet = 0.291r Nb. Thus, unlike FCC crystals, Nb has a larger tetrahedral site size than the octahedral site. 21 Table 1.2 shows the various atomic, physical and thermodynamic properties of Niobium: Table 1.2: Properties of Niobium 20 Properties Value Structure Body Centered Cubic (BCC) Atomic No. and Mass 41, g Electronic Configuration [Kr] 4d 4 5s 1 Density; Atomic Density 8.57 g/cm 3 ; 5.44 x atoms/cm 3 Melting; Boiling Point 2750 K; 5710 K Heat of Fusion; Vaporization 30 KJ/mol ; KJ/mol Molar Heat Capacity J/mol/K Oxidation States 5, 4, 3, 2, -1 Atomic Radius 1.46 Å Electrical Resistivity (273 K) 152 nω-m Ionization Energies (1 st, 2 nd, 3 rd ) KJ/mol, 1380 KJ/mol, 2416 KJ/mol Brinell; Vickers Hardness 1320 MPa, 736MPa Thermal Expansion 7.3 µm/(m-k) Niobium SRF Cavity Performance Parameters Niobium cavity performance is characterized by some basic properties which impact the quality factor (Q o ) of the cavity resonator. The main performance parameters are: a) Surface Resistance 22 The condensation of charge carriers in Niobium to cooper pairs below T c (9.2 K) provides Nb its superconducting capabilities. At T = 0 K, all charge carriers are condensed, with the fraction of unpaired carriers increasing with temperature until none of the carriers are paired at T > T c. In the case of RF currents, dissipation occurs for 0 < T < T c, albeit very small 14

40 compared to the normal conducting rate. This dissipation is termed as the surface resistance and is given by: R s = A (1/T) f 2 exp (- Δ(T)/kT) + R o.(1.9) Here A is a constant dependent on materials parameters such as the niobium London penetration depth, 2Δ is the energy gap of the superconductor, f is the RF frequency and R o is the residual resistance of the material, discussed later in this section. The first term in the above equation is also known as the BCS resistance of the material and is temperature related. Needless to say, the surface resistance must be kept minimal for better cavity performance. b) Field Limits 23 The accelerating field, E acc, in an RF accelerator cavity is proportional to the peak electric field (E pk ) as well as the peak magnetic field (B pk ) on the surface of the cavity. Thus, besides the phenomenally low RF surface resistance, another important fundamental cavity performance parameter is the maximum surface fields that can be tolerated by niobium without increasing the microwave surface resistance substantially. The ultimate limit to the accelerating field is the theoretical RF critical magnetic field of niobium which is about 200mT. These surface fields translate to a maximum accelerating field of 55 MV/m for v/c = 1 Nb structure and about 30MV/m for v/c < 1 Nb structure. However the typical cavity performance is significantly below the theoretically expected surface fields. The limitation is the result of thermal breakdown of superconductivity, originating at sub-mm size regions of high RF loss, called defects. When temperature outside the defects exceeds T c, the losses increase, as large regions become normally conducting. Thus, purification of the material is important to obtain defect free structures. c) Residual Resistance and Residual Resistance Ratio (RRR) 24 The second term of equation 1.8 is a temperature independent resistance which is typically material purity related, known as the residual resistance. The operating temperature of a superconducting cavity is usually chosen so that the first term in equation 1.8 (BCS resistance) is reduced to an economically tolerable value. Residual losses can arise from several sources e.g. residues from chemical etching, foreign material inclusions, or 15

41 condensed gases. Thus, it is important to maintain cleanliness during forming, welding and surface preparation of niobium. 20 Since superconductivity occurs at cryogenic temperatures for niobium, the thermal conductivity of the cavity is equally important for its performance. The electrical conductivity of niobium increases with decreasing temperature because of less scattering of electrons from lattice vibrations. The improvement in conductivity finally saturates to a value determined by impurity scattering, and since electrons are also the dominant heat carriers, these impurities also limit the heat carrying capability of the cavity, thus influencing the performance. The purity of niobium is hence characterized in terms of its residual resistance ratio (RRR) which is given by: RRR = Resistance at 300 K / Resistance at low temperature (non superconducting state). It can be seen that the electrical and thermal conductivities are closely related. High purity niobium has higher RRR values, with the theoretical limit being 35,000. d) Surface Roughness 25 Another important factor in the cavity performance is the roughness of the surface of niobium. Edgy and uneven surfaces can lead to magnetic field enhancement in the region, thus lowering the breakdown field of the cavity. Multi-cell cavities like those made at Jefferson Laboratories for CEBAF fabricated by welding, and it is this equator weld that is the roughest part of the cavity, which poses a problem in cleaning the surface for impurity free niobium material. It is for this reason that chemical etching and polishing techniques like Buffered Chemical Polishing (BCP) and Electropolishing are used for large grain niobium materials (single crystal) and polycrystalline niobium respectively. 26 e) Quality Factor Q o The cavity performance or efficiency of an SRF cavity is generally expressed in terms of its quality factor (Q o ). By definition, Q o is the ratio of the energy stored in the cavity to the power dissipated by it. Mathematically: Q o = ω o U/P c...(1.10) 16

42 Since the time averaged energy in the electric field equals that in the magnetic field, U is given by: U = 1/2μ o V B 2 dv = 1/2ϵ o V E 2 dv..(1.11) Where the integral is taken over the volume of the cavity, and μ o and ϵ o are magnetic permeability and electric permittivity of free space respectively. The dissipated power is a factor of surface resistance and is given by : P c = 1/2R s s B 2 ds (1.12) Combining 1.10, 1.11 and 1.12, we get a specific expression for Q o : Q o = ω o [1/2μ o V B 2 dv]/[1/2r s s B 2 ds]...(1.13) or Q o = G/R s.(1.14) where G = ω o μ o [ V B 2 dv]/[ s B 2 ds] is a geometric factor Thus, equation 1.14 relates the Q o inversely proportional to R s, which is in turn affected by the residual resistance and surface roughness. The purity of niobium has been established to be an important and limiting factor to the cavity performance, since this affects the surface resistance. The purity can be hampered significantly by the affinity of niobium for surface interstitial impurities and residual gas elements like H, C, O, and N , Surface Impurities in Niobium As discussed earlier, the purity of Nb is important for cavity performance, especially in terms of surface impurity content, since the RF penetration depth of Nb (λ) is 40nm. Inclusions on the Nb surface act as normal conducting sites for thermal breakdown of the cavity, since dissolved impurities serve as scattering sites for electrons not condensed into Cooper pairs. These impurities also lower the thermal conductivity and thereby limit the maximum tolerable surface magnetic field before the onset of thermal breakdown. The main interstitial impurities in Nb are H, C, O and N and these impurities are particularly detrimental to the mobility of electrons. Impurities such as refractory elements like W, Zr, Hf, Ti and Ta are also prevalent in commercial Nb, but these impurities are not substantially detrimental to the electronic properties of Nb, since they are substitutional impurities. It is thus important to understand the sources of the interstitial impurities (H, C, O, N). 17

43 1.4.1 Cavity Fabrication After extraction from the ore, the most common method for consolidation and refinement for Nb is electron beam melting in a furnace. This procedure drives out all the volatile impurities. After a satisfactory ingot is cast, it is forged into a thick rectangular slab, annealed and rolled in stages to the final sheet thickness. The number of steps involved contributes to the inclusion of impurities in Nb. Manufacturing steps such as rolling increase the bulk impurity content of the material. After rolling, the fabrication of the superconducting cavity takes place either by deep drawing or spinning, following which the cavity cells are trimmed and electron beam welded together to form a multi-cell cavity. Deep drawing is a forming process whereby a Nb disk cut from a sheet is pressed into shape using a set of dies machined out of 7075-A6 Aluminum alloy, because of the alloy s high yield strength, ease of machining and low cost. Like all sheet forming methods, deep drawing is sensitive to niobium s mechanical properties. Fig 1.8 shows the deep drawing of a Nb disk into a half cell shape of a cavity. Cavity half cells can also be made by a procedure known as spinning, where the Nb sheet is slowly pushed in to the desired shape on a mandrel. This method eliminates the need for a high tonnage hydraulic press. Nb sheet Figure 1.8: Deep drawing of a Nb sheet to cavity shape 27 After deep drawing or spinning, the half cells are trimmed to the final size for electron beam welding. Trimming is performed either on a lathe machine or a CNC (Computer Numerical Method) milling machine. Electron beam welding of the trimmed half cavity cells to obtain 18

44 the desired cavity is the final step of the fabrication process. The weld parameters are chosen to achieve full penetration butt welds with a smooth underbead by using a defocused electron beam, with the pressure in the vacuum chamber being less than 10-5 torr. As mentioned before, all these manufacturing steps make the material prone to contamination. For example, care has to be taken while trimming the formed Nb parts, due to its high reactivity to oxygen Surface Treatment 29 After their fabrication, the inner surface of the RF cavities must be etched in order to remove damaged layers and any surface contamination resulting from the fabrication process. The most common technique used for this purpose is Buffered Chemical Polishing (BCP), which is typically achieved using a solution of HF : HNO 3 : H 3 PO 4, 1:1:1-1:1:4, by volume. HF provides the F - ions which are able to complex with Nb 5+ ions and allow them to go into solution, while NO - 3 acts as an oxidant toward the Nb metal, to transform it into Nb + ions (oxidized). The role of H 3 PO 4 is to modify the viscosity and/or etching rate of the mixture. Another method of surface treatment gaining popularity is Electropolishing (EP). Niobium is used as the positive electrode (anode) and the cathode is made from aluminum. In this case, the etching takes place in HF-H 2 SO 4 solution, 15%:85% by volume, and the oxidation of niobium originates from an anodic polarization of the cavity (8-10 Volts). H 2 SO 4 plays the role of a buffer while the HF plays the same complexing role. 29 These surface treatments, though important, incorporate a substantial amount of impurities on the surface of niobium, especially hydrogen. Hydrogen is known to be responsible for the degradation of Q o in the cavities, termed as Q- disease, which is discussed in section 3.1. Niobium s tendency for impurities such as H, C, O, N, especially on the surface, affects the cavity performance Contamination from the various manufacturing steps (rolling, annealing, trimming and etching) is the major source of these impurities. 1.5 Need for Niobium Surface Characterization It has been established that surface interstitial impurities, especially H, C, O and N result in reduced cavity performance by influencing the superconducting properties of Nb, e.g. surface resistance, critical magnetic field and critical temperature. 30 The tendency of Nb to harbor 19

45 these impurities makes their surface characterization important. It is also expected that certain heat treatments involving low temperature baking for long periods of time, and high temperature heating for relatively shorter times, can improve the purity of Nb, especially within its London penetration depth (λ = 40nm). Another issue related to the basic understanding of the superconductivity of high-purity bulk Nb in strong RF fields is the occurrence of a sharp increase of the RF losses when the peak magnetic field, B pk, reaches about 90 mt, which is much lower than the critical magnetic field of Nb (200mT) and consequently limits the operational accelerating gradient of the cavities to about 25 MV/m. This phenomenon was discovered in , 32. Experiments showed that the onset of the newly discovered anomalous losses, which are commonly referred to as high-field Q o slope or Q o drop, occurs between mt, depending on how the materialhad been processed. An empirical cure for the Q o drop had been discovered in 1998 and consisted of a low-temperature ( o C, 48 hrs) baking of the cavities in ultrahigh vacuum. 33,34 Later experiments showed that the benefit of baking and the baking parameters (time and temperature) depend significantly on the cavity material processing. These findings are summarized by Ciovati. 35 The reasons for this Q o -degradation are still not well understood, and explanations range from magnetic field enhancements at grain boundaries in polycrystalline Nb to aspects of the metal-oxide interface. 36,37 Although some progress has been made to understand the performance losses, the effect of surface interstitial impurities, especially hydrogen, on the cavity efficiency has not been studied in detail. Moreover, there is a need for a comprehensive evaluation of the effect of both low temperature baking and high temperature heat treatments on impurity concentrations (H, C, O and N), with subsequent correlation of these impurity levels with SRF cavity performance. It is hoped that the detailed understanding of these issues will provide information which will allow development of new, more economical methods for improving the overall performance of Nb cavities. 1.6 Single Crystal vs Polycrystalline Nb As explained in section 1.4.1, traditionally the cavities have been made from polycrystalline niobium sheet. These cavities were fabricated by electron beam melting, hot forging to slab, 20

46 hot rolling to plate, and cold rolling to sheet followed by a final vacuum recrystallization anneal. This process provides sheet Nb with a reasonably equiaxial grain structure. However, it is proposed that the grain boundaries of the polycrystalline sheet Nb contribute to a reduction in the performance of the cavities, due to impurities that nucleate on the grain boundaries. Grain boundaries are "weak" areas in a niobium surface that can easily be contaminated via segregation of impurities. The contaminated grain boundaries form "weak links" which are prime candidates for causing the degradation in the RF cavity performance. Other effects include reduce thermal conductivity, and enhanced penetration of external magnetic fluxes. To overcome these grain boundary issues, intense research is underway to these performance detriments by making sheet material from a single crystal of high RRR niobium..thus, by fabrication and testing of cavities with fewer (large grain) or no grain boundaries (single crystal), the influence of grain boundaries on cavity performance and their impact on the Q- drop can be determined. A study on the benefits of large grain Nb was performed by Myneni et. al.. 39 It was found that large grain or single crystal niobium provided a high performance alternative to polycrystalline niobium for RF cavities. The following conclusions were reached: 1. Large grain Nb may provide a less expensive alternative as a result of the elimination of the sheet fabrication process and the associated reduction in the rate of defect creation could eliminate laborious procedures such as eddy current scanning. 2. Large grain Nb provided very smooth and shiny surfaces after BCP. The measured surface roughness over an area of 0.2 mm x 0.2 mm was 27 nm, a factor of 10 smoother than electropolished polycrystalline material and 50 times smoother than BCPed polycrystalline Nb. 3. Electropolishing as a final surface treatment step for high performance cavity might not be necessary for such material, eliminating a very costly preparation step. 4. There were some indications that the spread in cavity performance might be narrower than that previously experienced with electropolished cavities fabricated using polycrystalline niobium. In several cases, very low residual resistances (high Q o va- 21

47 lues) were measured even at high fields which are desirable for continuous wave mode applications such as in energy recovery linacs. 1.7 Objective of this Thesis Large grain Nb has the tendency to harbor surface interstitial impurities such as H, C, O and N following the various manufacturing steps used in SRF cavity fabrication and surface treatments. A detailed analysis of these impurities is thus required to understand the extent of this contamination and to suggest ways to minimize impurity effect on cavity performance. One such procedure is post-purification heat treatments, which largely consists of low temperature baking ( o C) or high temperature heating ( o C) of the cavity material in vacuum. Determination of the concentration and location of these interstitial surface impurities, both before and after heat treatments, may provide information that will increase understanding of their role in cavity performance and that can be used to guide further development of this material and its use in the fabricatin of SRF cavities. Thus, surface analysis of these interstitials may provide information critical to reducing operational costs and increasing the overall efficiency of SRF cavities. The following enumerates the objectives of this work: 1. To characterize the Nb surface for the presence of impurity elements (H, C, O, N). 2. To investigate the effect of various heat treatments on the surface impurity levels in Nb. 3. To understand the factors resulting in the improvement of cavity efficiency resulting from certain heat treatments. 4. To establish primary causes for relatively poor cavity performance with respect to surface impurities and to gain understanding of these causes. The major analytical techniques used to characterize the Nb surfaces in this work are Secondary Ion Mass Spectrometry (SIMS) and Transmission Electron Microscopy (TEM). The former was used to obtain mass spectra and concentration-depth profiles of the impurities in Nb before and after various types of heat treatments. The major advantages of SIMS are its extremely high sensitivity (ppm) and its ability to measure all elements, including H. A detailed description of SIMS is provided in chapter 2. TEM was used to study 22

48 the inherent surface oxide layer of Nb, which is about 5-7nm thick. With angstrom level resolution, TEM is an excellent technique for observation of any changes in the Nb surface oxide after heat treatment. Sample preparation and all heat treatments were performed in an induction furnace, at the Thomas Jefferson National Accelerator Facility (Jefferson Labs), VA, which provided support for this project. 23

49 References 1. K. Wille: The Physics of Particle Accelerators, Oxford University Press, NY (2000) 2. A. Wu et. al: Accelerator Physics, Technology and Applications, World Scientific Publications, Singapore (2004) 3. U. Amaldi: Proc. EPAC Vienna, p3 (2000) 4. J. Takacs et. al.: Energy stabilization of electrostatic accelerators, John Wiley and Sons (1996) 5. J. Cockroft et. al.: Proc. Royal Soc. London, A137, p229 (1932) 6. M Livingston et. al.: Particle Accelerators, Mcgraw Hill, NY (1962) 7. H. Weidemann et. al..: Particle Accelerator Physics, Springer, Berlin Heidelberg, NY (1993) 8. H Greinacher: Z. Physik, 4, p195 (1921) 9. R. Helm et. al.: The Stanford two mile Accelerator, W A Benjamin Inc. (1968) 10. R. Wideroe: Arch. Electrotech, 21, p387 (1933) 11. E.Courunt et. al.: Phys. Rev. 88 (5), p1190 (1952) 12. W. Kennedy et. al.: Rev. Sci. Instrum., 19, p89 (1948) 13. D. Pozar : Microwave Engineering, 2 nd edition, Wiley, NY (1998) 14. Jefferson Labs SRF Technology brochure (2012) 15. E source: - J Lab website (2012) 24

50 16. H. Padamsee et. al.: RF Superconductivity for Accelerators, 2 nd edition, Wiley, NY (2008) 17. J. Gallop: SQUIDS ; The Josephson Effects and Superconducting Electronics, CRC press, NY (1990) 18. H. Johnston: Physics World, Institute of Physics (2009) 19. L Landau et. al.: Course of Theoretical Physics, Oxford Butterworth, NY (1984) 20. H Piel et. al.: Nuc. Sc. 32 (5), p3610 (1985) 21. T. Wong: Nb Properties and Production and Applications, Nova Science Publishers, NY, (2011) 22. H. Padamsee et. al.: Annu. Rev. Nuc. Part. Sci., 43B, p635 (1993) 23. H. Padamsee et. al.: AIP Conf. Proc. 249, p1402 (1992) 24. K. Schulze: Niobium - Proc. Int. Symp., Metallurgical Society of AIME, p163 (1981) 25. S. Berry et. al.: Proc EPAC, p1000 (2004) 26. H. Piel: CERN Accelerator School Proc., p149 (1989) 27. J. Kirchgessner: IEEE Trans. Nuc. Sci., 30, p2901 (1993) 28. V. Palmieri et. al.: Proc. of 4th European Particle Accelerator Conference, p2212 (1994) 29. C. Antoine et. al.: AIP Conf. Proc., p176 (2003) 30. W. Singer et. al.: Materiaux and Techniques, 91, p13 (2003) 31. P. Kneisal et. al.: Proc. of 8th Workshop on RF Superconductivity-1997, p463 (2008) 25

51 32. E. Kako et. al.: Proc. of 8th Workshop on RF Superconductivity-1997, p491 (2008) 33. P. Kneisal et. al.: Proc. of 9th Workshop on RF Superconductivity-1999, p328 (1999) 34. B. Visentin et. al.: Proc. of 1998 European Particle Accelerator Conference, p1885 (1998) 35. G. Ciovati: Proc. of 13th Workshop on RF Superconductivity-2007, p79 (2008) 36. J. Knobloch: Proc. of 9th Workshop on RF Superconductivity-1999, p77 (1999) 37. J. Halbitter: Proc. of 10th Workshop on RF Superconductivity-2001, p292 (2001) 38. R. Graham: AIP Conf. Proc. 927, p21 (2006) 39. G. Myneni et. al.: AIP Conf. Proc. 927, p84 (2006) 26

52 2. Secondary Ion Mass Spectrometry (SIMS) The primary technique used for the surface analysis of the first 60nm of both heat treated and non heat treated large grain Nb in this study was Secondary Ion Mass Spectrometry (SIMS). Other techniques such as Transmission Electron Microscopy (TEM) were also used, but were limited to the characterization of the surface oxide of Nb. In this chapter, the basics of SIMS and characteristics of the SIMS instruments used are reviewed. 2.1 Introduction to SIMS 1 SIMS is a highly surface sensitive analytical technique in material is bombarded by a beam of energetic ions termed primary ions. The transfer of energy from the bombarding primary ions onto the material surface results in sputtering i.e. in the removal of material from the sample surface in the form of both neutral atomic species and also in the form of secondary negative and positive ions. In SIMS analysis, selected polarity secondary ions (either positive or negative ions are chosen) are extracted are extracted into a mass spectrometer by which these secondary ions are separated according to their mass/charge (m/z) ratios using a mass analyzer and then counted. The escape depth of the sputtered species is only a few angstroms, thus providing very high surface sensitivity. During the bombardment by primary ions, which typically have energies in the range of 0.5 to 30 kev, momentum is transferred to the limited area around the point of ion impact in the bombarded material. In addition to sputtering of material from the surface, rearrangement of species in the bombarded material and bond breaking occurs as well as implantation of the bombarding primary ion species into the sample. The ion bombardment induced emission processes include electron emission and the emission of surface particles (atoms or molecules) in both charged and uncharged and often excited states. An important question in SIMS concerns the probability of the formation of the charged state of the emitted atomic or molecular species. This probability of formation of the charged state depends strongly on the chemical environment of the sputtered species. By changing the chemical environment, for example, going from a pure metal to an oxide, the ionization probability of the metal atom may be changed by several orders of magnitude. 27

53 The main advantages of SIMS are its surface sensitivity, detection limit, and depth profiling capability. The technique can often accurately quantify elemental impurities in a sample matrix down to ppm to ppb levels. It is also one of the few techniques with the simultaneous ability to detect both molecules and all elements in the periodic table including hydrogen and has the ability to separate their various isotopes. While SIMS provides the ability to quantitatively determine elemental concentration over 5 orders of magnitude, this is not possible without appropriate standards for quantification. Since SIMS elemental sensitivity is strongly dependent on the sample matrix in which the element is contained, accurate quantification can only occur if a suitable standard of the sample matrix to be analyzed containing a know amount of the impurity in question is available. This allows determination via SIMS analysis of what is termed a relative sensitivity factor (RSF). Quantitative concentration of the selected impurity in an unknown sample can then be determined using this RSF Static and Dynamic SIMS The application of the SIMS technique is divided into two main categories: Static SIMS and Dynamic SIMS. In static SIMS, it is the intent of the analysis to obtain information on the composition of the uppermost layer, without effectively disturbing its composition and molecular structure. To insure that secondary ion information obtained is indicative of the structure and chemistry of the original sample surface, it is necessary that secondary ions originate from a previously un-bombarded portion of the surface. Since a monolayer of Si is 1x10 15 atoms/cm 2 and since this is of a similar order of magnitude with regard to the number of atoms in other materials, the static SIMS limit is set to be less than 1 percent of this number of surface atoms, i.e. the static SIMS bombardment limit is set to be 5x10 12 to 1x10 13 atoms/cm 2.This limit is generally achieved through the use of very low primary ion current densities. For an ion current density of 10-9 A/cm 2, the life time of a monolayer is on the order of some hours. 2 This technique has been applied for characterization of the molecular and elemental constituents of the uppermost mono-layers of metals, semiconductors, plastics, thermally mobile bio-molecules and numerous other applications. 2 28

54 In comparison to Static SIMS, Dynamic SIMS utilizes orders of magnitude higher primary ion current densities, up to some A/cm 2. The resulting fast surface erosion continuously moves the surface being analyzed into the bulk material, thus supplying information on the chemical composition of subsequently deeper layers of the bombarded sample i.e. a depth profile of sample composition. By use of dynamic SIMS, in depth information on elemental concentration can be obtained with high sensitivity, often down to the ppb range. This type of SIMS has been typically applied to metals, semiconductors and integrated circuit analysis The sputtering process Surfaces erode under particle bombardment and this is the phenomenon which is known as sputtering. Other effects also occur as a result if ion bombardment such as the generation of heat, ionization etc. Not all erosion is caused by sputtering. Evaporation, blistering etc. may also be consequences of high bombardment rates. Sigmund 5, in his review of cascade sputtering theory suggested that there are three criteria which define sputtering: 1. It is a class of erosion phenomenon observed on a material surface as a consequence of particle bombardment. 2. It is observable in the limit of small incident particle current. This makes it clear that macroscopic heating and subsequent evaporation by a high intensity beam is not sputtering. 3. It is observable in the limit of small incident particle fluence. This ensures that even a single particle can initiate a sputtering event. There are basically two approaches to modeling the sputtering process. One is to treat the process as a series of hard sphere collisions where classical energy transfer equations can be used. This forms the basis of the collision cascade theory introduced by Sigmund. An alternative approach to modeling this process is the use of one of the computer simulation methods in which the effects of bombardment on something approaching a realistic three dimensional structure is attempted. These latter approaches are still being investigated and research is underway to create 3-D simulations of the sputtering phenomenon. 4 29

55 The Collision Cascade Theory 5 This theory has been typified by the extensive and successful work of Sigmund. When a high energy particle penetrates a solid matrix, there are two major ways it can lose its energy. The first energy loss mechanism is by elastic collisions called nuclear stopping. The second major energy loss mechanism, which only becomes a major source of energy loss at very high projectile velocities via energy transfer to the electrons of the solid, a non collisional inelastic process generating heat, also called electronic stopping. Thus, for ion bombardment at moderate energies, elastic collisions are most important. Keeping the above in mind, three types of sputtering can be distinguished: a) Single knock on or prompt sputtering: occurs at s after primary particle impact. This is a direct impact process between the incident ion and the sputtered particle, involving only surface atoms. b) Slow collisional sputtering: The linear cascade regime which occurs due to the internal flux of moving target atoms intersecting the surface. This is the main process concerned with SIMS and occurs at s after impact. c) Slow thermal sputtering or spike regime: This is also a consequence of the movement of recoil atoms and occurs at s after impact. Unlike b above, the spatial density of the atoms is high, such that the majority atoms within a certain volume are in motion. Slow collisional sputtering is primary contributor to the SIMS secondary ion intensities caused by formation of a collision cascade. Thus, particle surface interactions involve forward and backward effects. Forward effects affect the interior of the bombarded target and include the stopping of the incident primary particle, the deposition of energy and the mixing of target material. Backward effects are those which lead to particle expulsion or sputtering. Clearly, the both effects are involved in sputtering. Fig. 2.1 shows the sputtering process in SIMS. As is seen, the SIMS technique secondary ion emission is as follows. Primary ion beam is directed at the sample surface. These ions penetrate the surface of the material and lose their energy via a series of elastic collisions, forming a collision cascade. Due to the transfer of energy, slow collisional sputtering starts in 30

56 the material and particles are ejected out from the surface. The ejected particles have a low energy (upto 130 ev, with a peak at 5-10 ev for a magnetic sector instrument, see section 2.2.1) and can be positive or negative ions, neutrals or electrons. The ions emitted are hence called secondary ions which are then extracted into and separated in a mass spectrometer, according to their mass/charge ratios. Figure 2.1: SIMS sputtering process Introduction to SIMS Instrumentation 1,4 Various kinds of SIMS instruments exist which differ in complexity, performance and the ways in which they generate, detect and mass separate secondary ions. Fig. 2.2 shows the basic components of a SIMS instrument. The main components are: A source to produce energetic primary ions and a primary column to focus and direct an ion beam at the sample A chamber to mount the sample to be analyzed 31

57 A secondary column to collimate, focus and direct the secondary ions A secondary ion energy analyzer A mass spectrometer to separate the secondary ions according to their mass/charge ratios Detectors to detect these secondary ions and record their intensity or produce an image Vacuum setup The ion source most frequently used for the generation of O + 2, O -, Ar +, or Xe + ion beams is the duoplasmatron, while a surface ionization source is used to produced Cs + ion beam and a Liquid Metal Ion Source (LMIS) is used for Ga + and Bi + 3 ions (See descriptions below). The ion optics used to transport the beam may be able to focus the beam to a diameter as small as 100nm in some cases, but primary ion beam diameters of 10 to 100 μm are more commonly used. 9 The SIMS sample chambers often include elaborate sample changing capabilities or in some case can accept only one sample at a time. The vacuum in the sample chamber can range from 10-7 to (UHV) torr depending on the instrument. In addition, the sample chamber may contain an electron gun to neutralize charge buildup on insulating samples that occurs when bombarded with a primary ion beam. The secondary ion energy analyzer in SIMS is responsible for energy focusing and can vary in design according to the type of instrument. Energy analyzers are important for achieving high mass resolution since the ions extracted from the sample have a range of energies which can impact the mass separation capability of the mass spectrometer. 32

58 Figure 2.1: Block diagram of basic components of a SIMS instrument 9 The mass spectrometer of a SIMS instrument is a critical component which separates the secondary ions according to their mass/charge (m/z) ratios. The mechanism by which mass separation takes place in the mass spectrometer of a SIMS instrument defines the classification scheme of SIMS instrumentation (majorly: magnetic sector, time of flight (TOF) and quadrupole, see succeeding section) Detection of the secondary ions can take place in a counting mode (faraday cup and electron multiplier), or an imaging mode (microchannel plate) (see section 2.3(c)). 2.2 Mass Analyzers used in SIMS Secondary ions sputtered from a sample must be extracted into a mass spectrometer and separated in order to evaluate the impurity levels in a given material. By definition, a mass spectrometer is a device used to separate ions according to their mass/charge ratios (m/z). The techniques used to provide mass filtration differentiates the various types of mass spectrometers. In SIMS, these are classified into the following major types. 33

59 2.2.1 Magnetic Sector Mass Analyzer 19 Magnetic Sector Mass Spectrometers separate ions by application of a magnetic field. When charged particles move through a magnetic field, they experience a force orthogonal to both direction of motion and the magnetic flux lines, resulting in a circular trajectory. The force acting on the particle and the radius of the circular path depend on the velocity of the ion. If ions of all masses are accelerated to a given potential, then the resultant velocity depends on the m/z ratio of each ion. Thus, for an ion of mass m accelerated to a potential V, the kinetic energy is given by : 1/2mv 2 = qv...(2.1) In order for a mass spectrometer to separate a particular mass ion from other ions in an ion beam, the mass to be separated has to fulfill the equation of force imparted to it by the magnetic field and the the force resulting from centrifugal force. The force due to the magnetic field B is given by: Bqv = mv 2 /r, where r is the radius of the circular path or r= mv/bq...(2.2) Inserting the value of v from (2.1) and squaring both sides: r 2 = 2Vm/B 2 q...(2.3) The equation above indicates that for a given acceleration potential (which is fixed for all ions in a particular experiment), the application of a magnetic field B would deflect the beam of ions and would only let the beam of ions with mass m ± Δm to pass through. This is the basis of a magnetic sector mass analyzer. The mass resolution is constant with mass and the mass resolution obtained from this method depends on the characteristics of the magnetic filter and on entrance and exit slit widths. Secondary ions sputtered from a sample have an energy spread ranging up to hundreds of ev and with a peak at 5-10eV. This energy distribution hampers the mass resolution of the magnetic sector, since the beam now possesses ions of different velocities, due to the energy spread. A secondary ion beam with a much narrower energy spread (the ideal is a monoenergetic beam) is achieved using a 90 0 electrostatic sector analyzer prior to insertion of the secondary ion into the magnetic analyzer. In such a device, a radial electrostatic field is 34

60 induced by a pair of cylindrical sector electrodes, kept at a certain potential. Ions of different energies and mass enter the electrostatic analyzer, they are deflected around the sector, where the amount of deflection depends on their kinetic energy. The geometry of the electrostatic sector and the imposed electric field thus reduces the energy spread in the beam via selection of ions of a given narrow kinetic energy spread and thus serves to inject ions of nearly equal energies into the magnetic sector. Figure 2.3: Double focusing mass spectrometer 19 This energy focusing compensates for the energy dispersion of the secondary ions entering mass spectrometer, while the mass dispersion in the electrostatic analyzer is compensated by the mass focusing of the magnetic sector analyzer. The spectrometer is thus called a double focusing mass spectrometer which can achieve high mass resolution (fig2.3). 6 These spectrometers can have a large mass range, high mass resolution (The IMS-6f used in this study easily achieves 4,000 m/ m using the 10% valley definition with a maximum specified resolution of 25,000 m/ m). High energy extraction is used, contributing to the high transmission of about 10-50% of these instruments. The high mass resolution also means that they are very attractive for SIMS analysis in semiconductor research. 7 35

61 2.2.2 Quadrupole Mass Analyzer 20 In a quadrupole mass analyzer, ions are separated according to their mass/charge ratios based on the stability of their trajectories in the oscillating electric fields applied by cylindrical rods kept at opposite potentials. Fig. 2.4 shows the schematic of such an analyzer. Figure 2.4: Schematic of a quadrupole mass spectrometer 20 The system consists of four cylindrical rods, connected together as two opposite pairs. A potential with a constant (DC) component and an oscillating (RF) component is applied to one pair of rods. An equal but opposite voltage is applied to the other pair. The rapid periodic switching of the field sends most ions into unstable oscillations of increasing amplitude until they strike the rods and are not transmitted (non resonant ions). Ions with a certain mass/charge ratio follow a periodic but stable trajectory of limited amplitude and therefore are transmitted (resonant ions). By increasing the DC and RF fields while maintaining a 36

62 constant ratio between them, this resonant condition is satisfied for ions of each ascending mass in turn, allowing the collection of a complete mass spectrum. In instruments employing a quadrupole mass spectrometer, the sample is grounded, unlike in the magnetic sector SIMS instruments, where the sample is held at a specified potential. The polarity of extraction electrodes preceding the quadrupole analyzer determines the polarity of the secondary ions extracted. The mass resolution that can be obtained in these analyzers is dependent on the number of RF cycles that an ion undergoes when it penetrates the region between the rods. For optimum mass separation, the velocity of the ions entering the quadrupole mass spectrometer should be very small (a few ev of energy) and the energy spread should be low. For this purpose, the extraction field is kept low and the energy band pass of the system is narrow (typically 10eV). The drawback to the low extraction fields is low collection efficiency, often resulting in transmissions of only 0.1-1%. Increasing the DC:RF ratio of a quadrupole mass analyzer improves the absolute mass resolution over the entire mass range, while application of a small offset DC voltage improves mass resolution at lower masses. However, the two adjustments also reduce the higher and lower mass transmissions respectively. The quadrupole is usually configured to give constant resolution with acceptable (but decreasing) transmission over the entire mass range. Other ways of increasing the mass resolution and transmission are the use of larger quadrupole (and rod) diameters and higher frequency RF components, along with an electrostatic analyzer (ESA) for energy focusing. 8 Because these analyzers do not employ magnets, mass and thus peak switching for selected ion monitoring can be performed quickly without hysteresis effects, which make them ideal for depth profiling, where one must constantly switch among masses. Also, charge neutralization in these instruments is easy due to the low extraction fields and grounded samples. 9 Quadrupole analyzers are therefore generally characterized by a low ion transmission (1%) and medium mass range (<1000) with a low mass resolution (300 using the 10% valley definition). 37

63 2.2.3 Time of Flight (TOF) Mass Analyzer 10,11,21 When ions are accelerated to a given potential so that they have equal kinetic energies, ions of different m/z ratios will have different velocities. If these ions pass through a field free drift region, they will disperse in time, with the higher mass ions arriving later, thus achieving mass separation according based on time of flight (TOF). This is the basis of the TOF mass analyzer. The kinetic energy of the ion leaving the sample is: 1/2mv 2 = qv...(2.4) where V is the extraction potential v = L/t...(2.5) where L is the length of the flight tube Inserting the value of v from (2.5) in (2.4): t 2 = 2qV/mL 2...(2.6) Since V and L are fixed for a particular experiment and instrument respectively, the mass only depends on t, which is how the ions are separated. Since a TOF mass spectrometer measures the time it takes ions of differing m/z ratios to move from the sample to the detector, it requires that the starting time (the time at which the ions leave the ion source) be well defined. This is typically achieved by pulsing the primary ion beam, which differs from quadrupole and magnetic sector SIMS instruments, where the primary beam operates in a continuous mode. The ions leaving the ion source of a TOF mass spectrometer neither have exactly the same starting times nor exactly the same kinetic energies. Methods for compensating for this energy spread include the use of an electrostatic analyzer prior to entrance into the TOF analyzer, and/or use of an ion optical device called a reflectron, in which ions pass through a mirror and their flight is reversed. A linear field reflectron (fig. 2.5) allows ions with greater kinetic energies to penetrate more deeply into the mirror than ions with lower energies. 38

64 Figure 2.5: TOF mass spectrometer with a reflectron 21 The ions that penetrate more deeply will remain in the reflectron longer and thus take longer to reach the detector, thus offsetting their increased velocity. Using a curved field reflectron instead of a linear one ensures that the ideal detector position for the time of flight mass spectrometer does not vary with the m/z ratio. The reflectron thus helps to improve the mass resolution of the analyzer. TOF-SIMS analyzers have very high transmission of ions (50-100%), which when coupled with a non scanning parallel mass detection provides good sensitivity. TOF analyzers have the highest practical mass range of all SIMS analyzers and can achieve high mass resolution. Charge neutralization is also relatively easy since the extraction field is switched off during the drift time of the ions and low energy electrons can be directed toward the sample surface. The use of these analyzers has been growing rapidly both for Static SIMS and for elemental and molecular depth profiling. All SIMS analyzers have their benefits and drawbacks. Quadrupole and magnetic sector analyzers are used mainly for depth profiling, as dynamic SIMS, while TOF analyzers are often used for surface chemical analysis as a static SIMS technique and for depth profiling. Table 2.1 compares the three mass spectrometers. 39

65 Analyzer Mag sector Quad Table 2.1: Different mass analyzers used in SIMS Mass range Resolution Transmissio n >1000 < Sequential scanning TOF > < Mass det. Disadvantages Primary Applicati on Sequential scanning Parallel scanning Slow mass switching, difficult charge neutralization Low resolution, transmission, narrow energy band pass Limited dynamic range Depth profiling Depth profiling Surface chemical analysis 2.3 CAMECA IMS-6f Magnetic Sector Dynamic SIMS Instrument 29 The primary surface characterization instrument used in this work was a CAMECA IMS-6f magnetic sector double focusing SIMS instrument. Fig.2.6 shows the main components of such an instrument. The primary components of this CAMECA SIMS instrument are: a) Primary Column : Ion beam sources (Duoplasmatron and Cesium Source); Primary Beam Mass Filter (PBMF) ; Primary ion optics b) Sample Chamber : Introductory chamber (Loadlock) ; Analysis chamber c) Secondary Column : Secondary ion optics ; Mass spectrometer ; Direct ion image detectors (Microchannelplate/ flourescent screen); Counting mode detectors (Electron Multiplier (EM) and Faraday Cup (FC)) d) Charge Neutralization : Normal incidence electron gun (NEG) 40

66 Figure 2.6: Schematic of a CAMECA IMS-6f SIMS instrument 29 41

67 2.3.1 Primary Column 21 The primary column is where the ion beam is generated, mass separated, accelerated and focused. The ion sources, a duoplasmatron (for oxygen ions) and the Cesium microbeam source (for cesium ions), generate the ion beam. This ion beam is then mass separated to obtain high intensity Cs + or O ions in the primary beam mass filter. The beam then passes through a series of electrostatic optics consisting of lenses, apertures and deflectors which produce a finely focused beam for a small spot size. Fig 2.7 shows the parts of a primary column. Figure 2.7: Parts of the primary column in a CAMECA IMS-6F instrument Duoplasmatron Ion Source 12 A plasmatron generates an ion beam from a gas source. When an energetic electron interacts with an atom, some of its kinetic energy is transformed into potential energy of the electrons of the atom. If the energy of the bombarding electrons is chosen to be higher than the 42

68 ionization energy of the atom, then efficient ionization is found to occur. In some plasmatrons, the electrons are generated by a cathode which is heated to produce thermionic emission, and accelerated toward an anode to provide them with sufficient energy. The path length of these electrons can be increased by the use of electrostatic and magnetic fields that force the electrons to move in oscillating or spiral trajectories during their flight to the anode. This increases their chance to have ionization interactions with gases even at low pressures (10-3 mbar, in this case, O 2 ). When the number of ions, electrons and neutrals in the source exceeds a critical value, plasma is formed with an equipotential boundary, from which ions are extracted. This is the basic principle of a plasmatron. In the plasmatron used in the CAMECA-IMS 6f, the plasma is created simply due to the high electric field between the cathode and the anode (no thermionic emission) and the ions are extracted through an aperture in a planar anode, An additional electrode known as the intermediate electrode physically compresses the plasma. A magnetic field is also present to concentrate the discharge further, which provides the source its name: duoplasmatron (dual mechanism to compress the plasma). Figure 2.8: Schematic of a duoplasmatron source 23 43

69 Fig. 2.8 shows the schematic of such a source used in the CAMECA IMS-6f to extract oxygen ions. A gas (O 2 ) is introduced at a low pressure to the interior of a hollow cathode through an adjustable leak. Plasma is produced by an arc maintained between the cathode and anode, which is kept at a potential relative to the anode. An intermediate electrode and a magnetic field produced by a coil are used to concentrate the plasma close to the extraction hole on the anode (hole diameter about 400μm). A part of this plasma expands beyond the duoplasmatron due to the pressure difference between the duoplasmatron and the chamber of the gun. The polarity of the extraction potential determines the polarity of the ions to be extracted. To extract negative ions, the axis of the discharge is displaced relative to the positive ion extraction. This is done by decentralizing the intermediate electrode relative to the anode by about 0.8mm. The duoplasmatron is used to generate oxygen ions in the CAMECA SIMS instrument, with both positive and negative polarities. The abundance ratio produced is O + 2 /O + = O - /O - 2 = 10. The advantages of such a source are high brightness due to the dense plasma ( A/m 2 /Sr) and low energy spread (<10eV). 4 With a duoplasmatron, generation of a low energy ion beam required for better depth resolution is difficult due to the high beam density which results in beam spreading due to space charge effects which results in low source brightness when low extraction voltages are used. Moreover, as the impact energy is reduced, the angle of incidence of the beam at impact becomes large with respect to the sample normal, since a lower energy ion beam would experience more deflection away from the sample normal, if the sample is kept at the same positive potential. This makes beam focusing extremely difficult. 44

70 Figure 2.9: Accel/Decel system for the CAMECA IMS-6f 14,15 To provide better source brightness and beam focusing, the CAMECA IMS 6f is equipped with an Accel/Decel system which operates by maintaining an extraction voltage on the duoplasmatron source high enough to avoid space charge effects for the ions exiting the duoplasmatron.after high energy extraction, the primary ions are then decelerated prior to the primary column entrance, providing improved focusing and transmission such that the low energy ions can be more efficiently focused and directed toward the sample. As shown in figure 2.9, the extraction electrode in the duoplasmatron is negatively biased instead of grounded which accelerates the ions. An Einzel lens is added in between the mass filter and duoplasmatron which decelerates the same ions. It has been shown that using this system while lowering the source voltage provides up to a factor of 2.5 increase in the maximum ion beam intensity, with the achievement of the primary beam impact energy as low as 500 ev. 14, Cesium Microbeam Source 16 The Cs microbeam source utilizes surface ionization to produce Cs ions. When a low ionization potential atom (Cs) is adsorbed on a high work function metal (W), the electronic 45

71 distributions of the atom and surface broaden the overlap, allowing movement of electrons between the two materials. If the temperature is raised so that the rate of desorption exceeds the rate of arrival of the Cs atom at the surface, the composite work function of the metal adsorbate increases, resulting in very high increase in the ratio of adsorbed ions to atoms. The element is then desorbed as ions and almost 100% efficiency can be reached in practice. In a Cs microbeam source used on the CAMECA IMS-6f, Cs vapor is generated from a CsCO 3 pellet in a reservoir by heating the pellet at 400 o C. This vapor then comes in contact with a W plate kept at 1100 o C, thus ionizing the vapor to Cs + ions. When an electric field is applied between the plate of W and the extraction electrode forming the ionizer, the Cs + ions are extracted and accelerated. The ionizer and reservoir are independently heated via electron impact. Electrons are accelerated from two heated filaments and electron from these filaments are attracted to the reservoir and ionizer, respectively, by positive voltages ranging from 3kV and 12kV resulting in electron impact heating. This surface ionization source provides both high brightness (>10 6 A/m 2 /Sr) and low energy spread (<1eV). 4 Fig 2.10 shows the schematic of a Cs microbeam source. Figure 2.10 : CAMECA IMS-6F Cs Microbeam source 16 46

72 Primary Beam Mass Filter (PBMF) 17 The PBMF allows selection of the desired source on the IMS-6f and removes unwanted species from the primary ion beam. While the Cs beam generated by the Microbeam Cesium source is quite pure, the beam generated by the duoplasmatron has some unwanted impurities due to the production of unwanted ions during ionization. As the name suggests, the PBMF filters the primary ions so that the desired type of ions (O + 2 /O - or Cs + ) is directed toward the sample. The PBMF in the CAMECA IMS-6F is a magnetic prism designed to be attached to two sources in the instrument. The basic principle of mass separation is the same as for the magnetic sector mass spectrometer. A magnetic field is applied to separate the desired ions from unwanted impurities in the ion beam via their m/z ratios. Spatial dispersion of the unwanted ions allows separation of the desired from the unwanted species. The polarity of this magnetic field is selected depending on the ion polarity or source selection. Mathematically, ions which satisfy the equation below are allowed to pass through: (2Vm/Bq) 1/2 / B = R (2.7) where R is the radius of the magnetic prism. The digital indication of the field value applied is obtained by measuring a voltage across the terminals of a shunt, which is placed across the terminals of a resistor though which the magnetic prism power supply is conducted. This shunt voltage is given by : U shunt = C (mv/q) 1/2 (2.8) where the constant C is determined experimentally Primary Beam Ion Optics 17 After the mass filtering of the primary beam, the selected ions enter the the primary beam ion optics, which consist of the lenses (electrostatic), apertures and deflectors needed to focus, deflect and shape the beam to a fine spot size before it strikes the sample surface (refer to fig. 2.7) There are four electrostatic Einzel lenses in the CAMECA IMS-6f which are named L1, L2, L3 and L4, with L1 present before the PBMF. The remaining three lenses are used to focus the beam. Each lens has one or more associated which are used to center the beam on the optical axes of the respective lenses. A double deflector system, located prior to L4, is used to position the primary beam position on the sample surface and to raster the beam across the 47

73 sample surface., and deflecting the primary beam in the primary column faraday cup (used to measure the beam current). In addition to the deflectors, L3 and L4 are each preceded by an aperture termed the L3 and L4 apertures respectively. These apertures are both composed of a set of four interchangeable and centerable apertures made of molybdenum. The L2 aperture functions as the collimating aperture for the ion source and as the exit slit for the PBMF allowing adjustment of the mass resolving power for the PBMF from 1 to 30. The L4 aperture collimates the primary beam prior to L4 to improve focusing.. A primary beam faraday cup positioned just before the lens L4 provides the ability to measure the primary beam intensity Sample Chamber The CAMECA IMS-6f is equipped with an introductory chamber (Load-lock) and an analysis chamber. The load-lock can hold two sample holders, which is useful in analyzing multiple samples, since one sample can be pumped down to meet the vacuum requirement needed for introduction into the sample chamber while analysis is carried out on the sample inside the analysis chamber. The analysis chamber is kept under ultra high vacuum (UHV) conditions ( torr) to minimize surface contamination. The intro-chamber is also kept under UHV conditions ( 10-8 torr) to facilitate the exchange of samples and minimize waiting time associated with load-lock pump down The Secondary Column The secondary column in the CAMECA-IMS 6f SIMS instrument provides the ability to extract secondary ions from the sample. Extracted secondary ions are then focused, deflected, mass separated and counted.. Mass separation is achieved by a double focusing mass spectrometer and detection and recording the beam intensities is achieved by various types of detectors Secondary Ion Extraction and Optics 18,24 Secondary ions are sputtered from the sample surface by the bombardment of the primary ion beam. These secondary ions are collected using an extraction lens, configured as an immersion lens. In the CAMECA system, the sample is held at a high potential (e.g. ± 4500V) and the immersion lens cover plate is at ground potential. Depending on the polarity 48

74 of the sample, positive or negative ions may be extracted. The polarity of the secondary ions is user selected and independent of the primary beam polarity. In the CAMECA IMS-6f, the immersion lens cover plate (ground potential) is set at a distance of 4.5 mm from the sample. In order to maintain a constant secondary ion beam current, the sample potential must be kept constant. With an insulating sample, this can often be partially achieved by coating the sample surface with a thin layer (<0.02 µm) of a conductor, such as gold. Under these conditions, along with a charge compensating electron gun, only minor changes in sample potential occur and any excess charge buildup is thus neutralized (section 2.3.4). Figure 2.11: Secondary ion extraction and optics 24 After the secondary ions have been extracted from the sample surface by the immersion lens, they are transferred via electrostatic transfer lenses into the mass spectrometer. The transfer lenses form a magnified image of the sample surface at the field aperture position and focus the secondary in beam onto the entrance slit of the mass spectrometer. A contrast aperture, placed between the entrance slit and the transfer lens is used to limit the solid angle of ions entering the mass spectrometer. Smaller contrast apertures therefore result in greater spatial resolution and mass resolution but at the expense of reduced secondary ion intensities. The 49

75 field aperture limits the area of the sample (field of view) from where secondary ions are admitted into the mass spectrometer. Reduced fields of view provide improved mass resolution by limiting ions far from the optical axis from entering the mass spectrometer as well as, if properly selected in conjunction with raster size, eliminate crater edge effects, deleterious to depth profiling depth resolution (section 2.4.3). Again, the above improvements come at the expense of secondary ion intensity reaching the detector. Fig shows the schematic of the secondary ion extraction and optics The Mass Spectrometer 19 The mass spectrometer includes an electrostatic sector analyzer (ESA) to energy filter the incoming secondary ion beam and a magnetic sector (MSA) analyzer which plays the role of a mass dispersing prism. The ESA and MSA are preceded and succeeded by entrance and exit slits respectively, the combination of which is used to adjust the mass resolution. The dispersion of secondary ions produced by a MSA depends upon the momentum of the particles. Thus, in order for the secondary ions to experience a deviation whose angle is a function of mass only, the beam must ideally be mono-energetic and in practice have a sufficiently narrow energy distribution to achieve the desired momentum and thus mass resolution. Since the secondary ions emitted from the sample surface have a range of energies, the use of a magnetic sector mass spectrometer alone leads to a limited mass resolution. Because of the initial energy spread, particles of s given mass will experience an range of dispersions about their centroid deviation. Thus, this equivalent of a chromatic effect must be corrected. To achieve higher mass resolution, the magnetic prism is coupled with an electrostatic sector through an electrostatic lens termed spectrometer lens. To understand the working of the spectrometer, two considerations are made. The first is the path of the ions which leave the surface with velocities directed along the normal path. Figure 2.12 shows the trajectory of such a beam, which has medium, low and high energies (red, yellow and blue color rays in the figure respectively). The red ray (central ray), after a first 90 o deflection in the electrostatic sector goes through the center of the spectrometer lens and experiences a second 90 o deflection in the magnetic sector field. Low and high energy 50

76 ions (yellow and blue rays respectively) comparatively are deflected more and are then bent by the spectrometer lens toward the achromatic point of the magnet. Figure 2.13 shows the beam path for ion energies other than the normal velocities. These trajectories are brought to a common focus on the central ray by the electrostatic prism. This virtual image point produced by the spectrometer lens in turn plays the part of an object point for the magnetic prism and is focused onto a real image point by the prism. Thus, the entire beam is achromatic. Figure 2.12: Normal trajectories of varying energies in the mass spectrometer 19 51

77 Figure 2.13: Non-normal trajectories of the same energies in the mass spectrometer 19 As a consequence of the sputtering process, all ion species are present at relatively low energies, but the energy distribution of secondary ion atomic species is broader then that of polyatomic ions. In order to suppress the molecular and often unwanted (because of mass interferences) species, an energy slit is used which is placed between the electrostatic sector and the magnetic prism. By moving the energy window and providing a voltage offset, the molecular secondary ion species can thus be lowered. However, a more important aspect of the energy slit comes from its utilization to achieve high mass resolution (m/δm > 2400). By narrowing the energy window, ions of nearly the same energies can be made to pass through, thus providng a relatively mono-energetic beam, which increases the resolving power of the mass analyzer Secondary Ion Detectors The CAMECA IMS-6f has multiple detectors for ion imaging and counting mode. The counting mode detectors include the Electron Multiplier (EM) and the Faraday Cup (FC) while the ion imaging detector is the Micro Channel plate (MCP)/ fluorescent screen. The two counting mode detectors provide an extended dynamic range of detection. 52

78 25, 26 a) Electron Multiplier The electron multiplier is the most sensitive detector. If protected from stray ions, neutrals and cosmic rays, then the background count rate is normally less than 0.01 counts per second (c/s). However, the multiplier must also be protected from intense ion beams (>5x10 6 c/s) as they can rapidly lead to degradation of performance. The electron multiplier (EM) used consists of a series of electrodes called dynodes. Each dynode is connected to a resistor chain. The first dynode is at ground potential, so that both positive and negative ions may be detected. The last dynode can be between to V depending on the age and type of multiplier. When a particle (electron, neutral, ion etc.) strikes the first dynode it may produce a few (1, 2 or 3) secondary electrons. These secondary electrons are accelerated to the second dynode that is held at a slightly higher positive potential. On impact, more secondary electrons are generated and a cascade of secondary electrons ensues. Fig shows the working of an Electron Multiplier. The CAMECA IMS-6f EM is equipped with a post acceleration system. When the first dynode is grounded as with earlier CAMECA IMS instruments, the velocity of the secondary particles is fixed by both their mass and the secondary accelerating voltage. Thus, at the same extraction voltage, the sensitivity is lower for higher masses since velocity is inversely related to mass at constant energy. When low extraction voltages have to be used, the instrument sensitivity is reduced because of the decrease of the EM yield due to lower impact energy of the secondary particles onto the first dynode. This loss of electron multiplier performance can be reduced by post-accelerating the secondary particles just before they reach the electron multiplier first dynode. Moreover, at a given secondary extraction voltage, post-acceleration minimizes the mass dependence of the electron multiplier efficiency. The post acceleration system in CAMECA IMS instruments has post acceleration voltage adjustable from -10 to +8 kv For optimum performance, the electron multiplier should operate at sufficiently high voltage so that every ion arrival produces a pulse. This pulse is then amplified and as long as it is above a set threshold, it will be passed to the counting circuit. At a given accuracy, the highest secondary ion intensity which can be measured is limited by the time resolution of 53

79 the pulse counting system. The dead time (τ) of the pulse counting system is the time spent after each event before being able to detect the next one. If N is the number of pulses generated by the electron multiplier and it is assumed that the time between two consecutive pulses is constant (1/N), then the total time to count one pulse becomes 1/N + τ. Thus, the total number of counted pulses will become 1/(1/N+τ), which is the same as N/(1+Nτ). Thus, the dead time correction has to be applied in order to obtain the accurate pulse rate. Figure 2.14: Working of an electron multiplier 26 It is also important to note that when the optical gating capability of the CAMECA IMS-6f is used, vey high instantaneous count rate yielding an overload of the EM can be reached while the apparent count rate measured by the pulse counting system appears much below the limit. 54

80 All of the counts are obtained in the fraction of time that the beam passes over the optically gated detection region. Therefore the electron multiplier count limit may be reached while the apparent count rate measured appears much below the limit of ~3 x 10 6 c/s (EM limit in the CAMECA IMS-6f). As this count rate is reached and exceeded, the accuracy of the ion intensity measurement will be degraded due to the detector dead time. In order to prevent the EM from count rate overload, the software can be used to set up a transition such that the secondary beam is automatically switched into the Faraday cup if count rate reaches the preset value (2 x 10 5 c/s). b) Faraday Cup 27 A Faraday cup detector can detect count rates from 5x10 4 c/s up to about 1E9 c/s. Unlike the electron multiplier it does not discriminate between the type of ion or its energy. It is simple and relatively inexpensive, but its response time is slow. Fig shows the schematic of a Faraday Cup. The detector consists of a hollow conducting electrode connected to ground via a high resistance. The ions hitting the collector cause a flow of electrons from ground through the resistor. The resulting potential drop across the resistor is amplified and sent to the voltage/frequency converter. A plate held at about -80 V in front of the collector, prevents any ejected secondary electrons from escaping and causing an anomalous reading. Figure 2.15: Schematic of a faraday cup detector 27 55

81 c) Micro-channel plate/fluorescent Screen 28 The IMS-6F secondary ion image detector consists of a micro-channel plate device coupled to a fluorescent screen. The micro-channel plates convert secondary ions to electrons and the fluorescent screen converts electrons to photons. It consists of an array of small channels oriented parallel to one another. Each channel is a small hollow glass tube with a conductive inner surface layer. When a secondary ion strikes the inner surface, there is electron emission, these electrons are accelerated and multiplied by collision cascades in the channel. Electrons leaving the channels are further accelerated onto a fluorescent screen resulting in photon emission. The fluorescent screen image is acquired by means of a chargecoupled device (CCD) camera. This system is primarily used in alignment by providing a mass filtered image of the sample surface Charge Neutralization 24 During SIMS analysis, primary ions impinge the sample surface, while secondary ions and/or secondary electrons leave the sample surface. If the ratio of the yield of secondary ions and electrons to primary ions are not equal to 1, a charge imbalance occurs and an excess of charge will accumulate in the sputtered area. If the sample has an intrinsic conductivity, the excess surface charge can be compensated by electrons flowing from conductive sample holder and the potential of the sample surface will remain constant. If the sample is an insulator, the electrical charge will accumulate on the sample surface with deleterious results, such as non-stable or absent secondary ion signal and/or sample high voltage breakdown. When severe charging occurs, the potential of the sample will build up high enough to generate an arc and damage the sample. In either positive mode or negative mode of sample bias, a nonconductive sample always charges positively under positive primary ion bombardment. To neutralize the excess positive charges accumulated on sample surface, a flux of electrons can be introduced to the ion sputtering area. Other methods of charge neutralization include using a conducting layer of coating on the sample and using an O - primary ion beam respectively. The former is not sufficient for very insulating material, while the latter is often impractical due to the very low 56

82 beam current density provided by the O - beam in the CAMECA IMS-6f. Thus the use of an electron beam is the most common method for charge neutralization. In the CAMECA IMS-6F, the normal incidence electron gun (NEG) is used for electron beam charge neutralization. Figure 2.16 is a schematic drawing that illustrates the configuration of the normal incidence electron gun. Electrons are generated from the tungsten filament and the electron beam is directed through the electron column. A magnetic deflector (Bya) deflects the electron beam into the secondary ion optical axis through the immersion lens and towards the sample surface. Figure 2.16: Schematic of the working of the electron gun 24 57

83 Depending on the polarity of the bias of the sample, the neutralization method can be positive mode charging compensation or negative mode charging compensation. The methods described here are used in CAMECA IMS-6f magnetic sector design due to the special geometry of the instrument optics and high electric field which is on the order of kilovolts per mm over the 4.5mm distance between the sample and the immersion lens cover plate (part c of this section) Positive Mode Charge Neutralization using the NEG The sample is positively biased in positive secondary ion mode. Depending on the thickness of the insulating sample, two approaches can be used for charge neutralization. The first approach includes charge neutralization for a thin film insulator on a conductive substrate, and is straightforward. The electron beam impact energy is chosen such that electrons can penetrate the film and render the film conductive via electron beam induced conductivity. The electron beam impacts the sample surface at normal angle of incidence. 38 The second approach is used if the insulator is a thick film or bulk material and the electron beam can not penetrate through the insulator, which makes charge compensation difficult. Pivovarov et al. 39 reported an electron beam based charge neutralization procedure for magnetic sector SIMS analysis of bulk insulators which consists of positioning the electron beam adjacent to or just touching the ion beam raster area. Prior to analysis, the surface of the sample has to be coated with gold which provides a conductive surface layer and which has a high secondary and backscattered electron yield. The advantage of this method is the effective neutralization over a large range of primary ion beam currents, energies and raster sizes providing the ability to cover a wide range of analysis requirements such as differing sputter rates, depth resolutions and detection limits Negative Mode Charge Neutralization using the NEG 40 For negative secondary ion detection with a NEG, both electron source and sample are negatively biased to the same potential. When electrons reach the sample surface they lose their momentum, so a cloud of very low energy electrons is formed above the sample. The availability of low energy electrons in the "cloud" provides self regulated charging compensation. If a region of the sample charges positively, electrons will be attracted to the 58

84 positively charged region from the cloud as needed to compensate for the sample charging. Since a self balancing situation is established where electrons only impact the sample as needed, for maintaining a charge balance, careful regulation of NEG electron beam current is not required. However, an excess of electrons above the sample surface is needed, which requires that primary ion density be kept sufficiently low in order that sufficient electrons are available. 2.4 ION TOF V SIMS Instrument Time of Flight SIMS was used in this study for imaging the grain boundaries of large grain bicrystal Nb, primary due to the high spatial resolution capability of the technique which can provide images with a resolution of 100nm, to study segregation of H, C, O and N impurities on the grain boundaries before and after various heat treatments. For this purpose, an ION TOF V instrument was used, which is equipped with a Cs + and Bi m+ n (n = 1 6, m = 1, 2) primary ion beams for sputtering and analysis respectively. The major differences between the ION TOF V and the CAMECA IMS-6f are the source, mass analyzer and the energy analyzer. Section shows details of the mass and energy analyzers used in a TOF-SIMS instrument, which are similar to the analyzers employed in the ION TOF V. Fig shows the schematic diagram of an ION TOF V SIMS instrument. A pulsed primary ion beam with energy of 1 25 kev is used to bombard the sample surface. The primary ion energy is then transferred to the target atoms and a so-called collision cascade is generated. 31 Part of the energy is transferred back to the surface, allowing the surface atoms and molecular species to overcome the surface binding energy and leading to the sputtering of secondary species from the top surface of the sample. 32,33 These ejected secondary species may include electrons, neutrals, or ions. The majority of the emitted particles are neutrals and only a small amount (4 5%) of the emitted species are positively or negatively charged ions. 30 These ionized particles are then accelerated into a reflectron TOF mass spectrometer. Since the ions all leave the sample at the same time, and are subject to the same accelerating voltage, the lighter ions arrive at the detector before the heavier ones. The "flight" time of an ion is proportional to the square root of its mass. Therefore, all the ions with different masses can be separated during the flight and detected individually. 59

85 Figure 2.17: Schematic of an ION TOF V Time of Flight SIMS instrument 34 Figure 2.18: ION TOF V Time of Flight SIMS instrument at AIF 60

86 Figure 2.18 shows the TOF SIMS instrument present at the Analytical Instrumentation Facility (AIF) in NCSU, with the various sources (C 60 source not used in this study). Hence, the function of the instrument is similar to the CAMECA IMS-6f SIMS, apart from the fact that the incident primary ion beam is pulsed instead of continuous, to aid in the mass separation of the secondary ions according to their time of flight. The TOF SIMS also does not have multiple counting mode detectors as in the CAMECA IMS-6f, which is why the dynamic range of the instrument is lower than that of the CAMECA. The main advantages of this technique over the magnetic sector and the quadrupole instruments are the extremely high transmission, the parallel detection of all masses and a theoretically unlimited mass range. The other important advantage of this instrument is the high spatial resolution achievable due to the Bi beam ion source, which is a Liquid Metal Ion Source (LMIS) Liquid Metal Ion Gun (LMIG) A liquid metal ion gun (LMIG) produces a well-controlled ion beam that can be focused to a small spot size. 36 It is separated into four different parts: Liquid Metal Ion Source (LMIS), Primary focusing unit, Blanker/Chopper, and Buncher. A schematic of the complete system is shown in Figure

87 Emitter Liquid metal ion source Extractor Lens Source Beam Size Aperture Primary focusing unit Lens Magnification Crossover Aperture Blanker Chopper Buncher Blanker/chopper Buncher Primary focusing unit Blanking Aperture Lens Target Sample Figure 2.19: Schematic of a Liquid Metal Ion Gun 35 The Liquid Metal Ion Source generates the beam. Fig shows a schematic of such a source. The source consists of the emitter (tungsten tip, metal or metal alloy reservoir with heating wire), the suppressor, the extractor, and the lens electrode. In the ion source emitter, the reservoir, which contains the source material, is heated to form a liquid metal layer over the tip of a needle. The heating is performed by feeding a well-regulated current through the heating coil surrounding the metal reservoir. When liquid metal is exposed to an electric field, the shape of the liquid starts to deform. As the voltage increases, the effect of the electric field becomes more prominent and this electric field exerts a similar amount of force on the droplet as the surface tension, which leads to the formation of a cone shape, also known as the Taylor cone. The cone has an angle of Due to the extremely high field strength at the Taylor cone, field evaporation of metal ions occurs and ions are emitted from a very small area with a virtual source size on the order of 10 nm in diameter. 35 The ion beam 62

88 kinetic energy is determined by the potential difference between the sample, which is held at ground, and the ion source. The electric field at the tip is mainly generated by the extraction electrode in front of the tip. The emission current is, therefore, controlled by the extraction voltage. The ion source emitter is also surrounded by a suppressor electrode and the suppressor voltage is used for a fine tuning of the emission current. The typical emission current is about 1.5 µa. Reservoir needle Suppressor Lens Source Ion Beam Taylor Cone Extractor Beam Size Aperture Heating Figure 2.20: Schematic of a Liquid Metal Ion Source 35 The extracted ions leave the source through the focusing optics and are pulsed before arriving at the target. The focusing system has three lenses (Lens Source, Lens Magnification and Lens Target) and two apertures, which are used to control the primary ion beam direction and ion current. Normally the ion optical column is used in the crossover mode, meaning that the lens source and lens magnification produce a magnified image at the position of the crossover aperture and the image is then demagnified by the lens target. The overall 63

89 magnification is approximately The final spot size is mainly affected by the chromatic aberrations of both lenses. The primary ion beam emitted from the liquid metal ion source is a continuous ion beam. To be compatible with the time-of-flight (TOF) analyzer, the ion beam must be pulsed. The pulsing system of the liquid metal ion gun consists of a high performance beam blanker/chopper for high speed motionless beam blanking (figure 2.19). It can produce ion pulses with lengths of l to tens of nanoseconds. To further reduce the pulse length, each ion package can be bunched or compressed further along its flight direction in a bunching system. 35 The bunching system works by applying the correct amplitude of the high voltage pulse, which leads to the acceleration of the ions in the rear of the ion packet to catch up to the ones in the front of the packet, allowing all ions to reach the target at the same time. Thus, the bunching system allows an axial compression of the ion packets. Since there are no primary ions lost, secondary ion intensity is usually higher in the bunched mode than in the chopped mode. On the other hand, the energy of the beam is no longer mono-energetic and can have a spread of several hundreds of ev. As a consequence, the sharpness of the beam is reduced and the spatial resolution is increased to ~5 µm, which is not favorable for imaging acquisition. Thus, for image acquisition with better spatial resolution, only the chopped mode is used. The resulting ion beam is then directed toward the sample and the secondary ions are extracted using an extractor. Unlike the CAMECA IMS 6f, the sample in an ION-TOF V is kept at ground potential. The secondary ions extracted are sent to the TOF analyzer, which is coupled with a reflectron or an ion mirror for energy focusing of the ions (section 2.2.3). 64

90 References 1. A Benninghoven et al : Secondary Ion Mass Spectrometry- Basic Concepts, John Wiley and Sons (1987) 2. A Mueller : Thin Film Sol., 12, p439 (1972) 3. L. Wiedmann : Surf. Sc., 41, p483 (1974) 4. J. Vickerman et al : SIMS Principals and Applications, Oxford Science Publications (1988) 5. Sputtering by Particle Bombardment I, Springer Series topics in Applied Physics, 47, p9, Springer Berlin (1981) 6. A. Nier et al : Phys. Rev. 81, p507 (1951) 7. R. Castaing et al : Jour. of Micros., 1, p395 (1962) 8. E_source : Characteristics of Different Mass Analyzers, Fraunhofer Institute for Process Engg. and Packaging : http : // 9. F. Stevie et al : SIMS- A practical handbook for depth profiling and bulk impurity analysis, Wiley Interscience Publications (1989) 10. B. Mamryn et al : Sov. Phys. (English translation), 37, p45 (1973) 11. X. Jang et al : Int. Journ. Mass Spec. Ion. Phys. (1989) 12. CAMECA-IMS 6f user guide : Duoplasmatron Illustration, I-63 (1996) 13. CAMECA-IMS 6f user guide : Cs Microbeam Source, I-65 (1996) 14. M. Schuhmacher et al : Journ. of Vac. Sci. & Tech. B, 18, p529 (2000) 15. CAMECA-IMS 6f Application Note : Acc/Decel system (2000) 65

91 16. CAMECA-IMS 6f user guide : PBMF, I-66 (1996) 17. CAMECA-IMS 6f user guide : Primary Ion Optics, I-77 (1996) 18. CAMECA-IMS 6f user guide : Beam Deflection, I-68 (1996) 19. CAMECA-IMS 6f user guide : Mass Spectrometer, I-16 (1996) 20. E_Source : R. Fleming : SIMS theory tutorial 21. E_Source: The Chemistry Hypermedia Project, E_Source : Atomic and Molecular Mass Spectrometry III, E_Source : Evans Analytical Group website : C. Gu: PHD thesis, Dept. of MSE, NCSU, p21; p68 (2005) 25. CAMECA-IMS 6f Application Note : Post Acc. system (1996) 26. CAMECA-IMS 6f user guide : Electron Multiplier, I-78 (1996) 27. CAMECA-IMS 6f user guide : Faraday Cup, I-77 (1996) 28. CAMECA-IMS 6f user guide : Channel Plate, I-74 (1996) 29. J. Becker: Inorganic Mass Spectrometry, John Wiley & Sons, Weiheim, Germany (2007) 30. J. Vickerman et al.: TOF-SIMS: Surface Analysis by Mass Spectrometry; IM Publications/Surface Spectra: Chichester, UK (2001) 31. M. Pacholski et al.: Chem. Rev., 99, p2977 (1999) 66

92 32. U Oran et al.: Appl. Surf. Sci., 252, p6588 (2006) 33. I. Talian et al.: Surf. Sci., 601, p4158 (2007) 34. E_Source: ION TOF USA Website, TOF-SIMS-TIME-OF-FLIGHT-SURFACE-ANALYSIS.htm (2012) 35. TOF SIMS V - users' guide, ION-TOF GmbH, Münster, Germany. (2012) 36. P. Prewitt et al.: J. Phys. D: Appl. Phys., 13, p1747 (1984) 37. G. Taylor: Proc. R. Soc. London, A, 280, p383 (1964) 38. J. Goldstein et. al: Scanning Electron Microscopy and X-ray Microanalysis, Second Edition, Plenum Press, New York and London p.89 (1992). 39. A. Pivovarov et. al: App. Surf. Sci , p781 (2004) 40. A. Pivovarov et. al: App. Surf. Sci. p786 (2004) 67

93 3. Experimental In the previous two chapters, an introduction to superconducting materials and the motivation and the objectives of this dissertation work were provided. The need for surface characterization of bulk large grain niobium was detailed and the primary method used for impurity analysis (H, C, O and N), SIMS, was described. As indicated in section 1.7, the objective of this work is to characterize the niobium surface to determine the presence and levels of H, C, O and N and to study the effects of various heat treatments on the levels of these impurities. Thus, it is important to understand the sample preparation techniques, heat treatment parameters and SIMS analysis conditions used. 3.1 Sample Preparation All samples were prepared at Jefferson Laboratory. Figure 3.1 shows the sample preparation procedure. High-purity (RRR > 200) large grain Nb samples from CBMM, Brazil 1, having length and width of 5 mm x 7.5 mm and thickness of 2.5-3mm were cut by wire electrodischarge machining from larger disks. Nb (5 x 7.5 x mm) BCP (1:1:1) 130µm removal BCP (1:1:2) 20-30µm removal Nanopolishing BCP (1:1:2) µm removal Heat Treatment (600 o C/ 10hrs) Control No heat treatment after nanopolishing BCP = Buffer Chemical Polishing [HF (49%): HNO 3 (69%): H 3 PO 4 (85%)] Heat Treatment after nanopolishing Figure 3.1: Flowchart for Nb sample preparation for surface analysis 68

94 After cutting to dimensions, the samples were subjected to the following treatments: 1. Etching by BCP (Buffer Chemical Polishing) 1:1:1, removing about 130 μm 2. Etching by BCP 1:1:2, removing about μm 3. Heat treatment at 600 o C for 10 hours in a vacuum furnace 4. Etching by BCP 1:1:2, removing about μm 5. Nanopolishing at Wah Chang, USA (proprietary process) 6. Heat treatment (on select samples) The above sample preparation steps replicate the standard cavity preparation steps 2. Unless otherwise stated, all samples were cut from within a single grain of a large grain Nb disk composed of grains several mm in size, thus making them essentially single crystal (single grain). Following the nanopolishing, some of the samples were heat treated via various regimens of selected temperatures and times, depending on the type of sample and the heat treatment to be studied. Control samples were selected from the samples produced which underwent no further treatment or processing after nanopolishing. For cavity performance experiments, samples were heat treated either in the same furnace as the cavities or with similar heat treatment procedures for consistency. Nanopolishing of the disk samples was performed to provide a surface of the required smoothness required for SIMS and TEM measurements. Cavity samples were not nanopolished subsequent to removal from the cavities. While fine (multi) grained samples were also studied, the main focus of this study was on large grain Nb due the SRF cavity performance advantages provided by the large grain Nb versus polycrystalline Nb (section 1.6) 3.2 Types of Heat Treatments Because of the number of heat treatments performed, a nomenclature was defined for each type of heat treatment which will be followed throughout this dissertation: 1. All 120 o C heat treatments were termed as long term low temperature baking, and will be described as 120 o C/time (hrs), (for example: 120 o C/48hrs) 2. Heat treatments performed at temperatures 600 o C were named high temperature heat treatment and will be described as temperature/time (hrs), (for example: 600 o C/10hrs) 69

95 3. Another type of heat treatment performed involved high temperature heating followed by long term low temperature baking. These will be described as temperature/time (hrs), 120 o C/time (hrs) (for example: 800 o C/3hrs, 120 o C/24hrs) 4. There were some heat treatments performed for very short periods of time in the same furnace and experiment, after high temperature heating which are termed as high temperature heating followed by lower temperature heating. These will be described as: temperature/time (hrs), temperature/time (min) (for example: 800 o C/3hrs, 400 o C/20min). Although all the heat treatments in this category were performed in vacuum 5. An additional type of heat treatment performed involved the introduction of N 2 gas at 10-5 mbar during lower temperature heating step of the samples after high temperature heating. This will be described in the dissertation with N 2 gas pressure in the paranthesis adjacent to the lower temperature heating step [for example: 800 o C/3hrs, 400 o C/15min (10-5 mbar N 2 )] Table 3.1: Table showing the various types of heat treatments performed Type of Heat Treatment Temperature/Time Purpose of Heat Treatment High Temperature Heat Treatments High Temperature Heating followed by Long Term Low Temperature Baking High Temperature Heating followed by Lower Temp. Heating (10-5 torr N 2 ) 600 o C-1400 o C/3-10h 600 o C-1400 o C/3-10h; 120 o C/12-48hrs 800 o C-1000 o C/2-3h; 400 o C-800 o C/10-20min in vacuum or 10-5 torr N 2 Desorption of hydrogen; Dissolution of any interstitial oxygen into bulk Preliminary cavity performance results showed better performance compared to high temperature heat treatments only Nitridation of surface to avoid reabsorption of impurities on cool down Long Term Low Temperature Baking 120 o C/48h Preliminary cavity tests showed consistent cavity performance improvement 70

96 Table 3.1 shows a brief explanation of the reasons behind these heat treatments. Elaborate versions of these explanations would be provided in each chapter (Chapters 3-6), before the discussion of the results. Since some of the heat treatments were performed keeping a particular impurity in mind (H, O, C, or N), it is best to address these particulars before the discussion of characterization results for that impurity in Nb, in designated chapters. Two types of furnaces were used for heat treating the samples. While the standard cavity heat treatment step utilizes an Elnik resistive heating furnace, some of the heat treatments in this work were performed using an induction furnace. The Elnik resistive furnace 3 used for the high-temperature heat treatment of SRF cavities was a high-vacuum furnace with a molybdenum hot-zone; molybdenum (or tungsten) resistive heating elements with cavities/samples heated by radiation from the heating elements. However, one of the challenges was the re-absorption of impurities from the furnace environment during cooldown to room temperature and thus, an induction furnace with a niobium hot zone was used for some of the heat treatments, which was expected to reduce contamination by foreign elements. In fact, such a furnace has been shown to reduce cavity losses by a factor of 2. 3 Figure 3.2: Schematic diagram of the induction heating system used 3 71

97 Table 3.2: Table relating the heat treatments to the respective furnaces Type of Heat Treatment Temperature/Time Furnace Used High Temperature Heat Treatments 600 o C-800 o C/3-10h Resistance Heating Furnace 1000 o C o C/3-6h Induction Furnace 600 o C-1200 o C/3-10h; High Temperature 120 o C/12-48h Heating followed by Resistance Heating Furnace Long Term Low Temperature Baking 1400 o C/3h; 120 o C/12h Induction Furnace High Temperature Heating followed by Lower Temp. Heating 800 o C/3h; 400 o C/10-20min in vacuum or 10-5 torr N 2 (10-5 torr N 2 ) 1000 o C/2h; 800 o C/10min [10-5 torr N 2 ] Resistance Heating Furnace Induction Furnace Long Term Low Temperature Baking 120 o C/48h Resistance Heating Furnace Figure 3.2 shows the components of the induction furnace used for some of the heat treatments, while table 3.2 separates the various heat treatments according to the furnace used. Both the furnaces used were high vacuum furnaces and the samples were heated in vacuum with a total pressure of 10-7 torr. Hydrogen pressures in the furnaces during heat treatment were monitored using a Stanford Research System residual gas analyzer (model RGA100). 3 For all the heat treatments, the furnace was ramped up to the desired temperature and the sample was held at that temperature for the specified amount of time. The furnace was then switched off and the samples were allowed to cool down. For samples undergoing additional low temperature baking and heating steps, the samples were held at 120 o C or 400/800 o C for the desired time and then cooled down to room temperature. 72

98 3.3 Nanopolishing Initial roughness of the Nb samples rendered depth profiling using SIMS very difficult. An important aspect of SIMS analysis is the surface profilometer measurements of SIMS crater depths which provide the analysis depth needed for calculation of the sputter rate of the ion beam (see chp 2). Rough samples with uneven topography make accurate measurement of these craters almost impossible. The rough topography of the large grain Nb unpolished sample in the optical micrograph shown in Fig. 3.3(a) that made it difficult to identify the SIMS craters was thus removed in the nanopolished 23 surface of a large grain sample shown in Fig. 3.3 (b). From this figure, it can be seen that nanopolishing provides a smooth surface allowing accurate SIMS crater measurement. SIMS Craters (a) (b) Figure 3.3: Optical images of niobium surface (a) unpolished and (b) after nano-polishing Note that nanopolishing was not performed for the cavities. However, it was shown via SIMS analysis on polished and unpolished samples that nanopolishing does not appear to affect the impurity content of the Nb surface. Since non-nanopolished cavities and nanopolished samples have the same impurity levels, correlation of impurity levels with cavity is possible. SIMS data from polished and unpolished Nb samples is shown in chapter 4. 73

99 3.4 SIMS: Experimental Parameters Following the heat treatments, samples were analyzed using SIMS for surface impurity levels. To allow direct comparison, control samples which were not heat treated were analyzed using the same SIMS analysis conditions. Before discussing the analysis conditions, it is important to describe the various experimental aspects related to SIMS analysis. The type of primary ion beam to be used, ion yield of the species to be profiled in the particular matrix to be analyzed, and the sputtering rate, are some important ones. Detection limit of the species (H, C, O, N in this case) is the ability to detect the particular secondary ion species to be analyzed. Thus, obtaining quality data is heavily dependent on the experimental conditions used for the analysis Primary Beam 4, 5 The selection of the primary beam species to be used depends primarily on the secondary ion species to be detected as well a on the depth and the lateral resolution needed. The CAMECA IMS-6f used in this work is equipped with O + 2 /O - and Cs + beam primary ion sources (see chapter 2) and all the characterization of Nb was performed using either the Cs + + or O 2 beams. An important criterion for the selection of a primary ion beam (O + 2 or Cs + ) is the enhancement of the ion yield. The primary ion selected is generally determined by whether positive or negative secondary ion yield is favored for the impurity which is to be analyzed. The ionization potential of the element strongly influences the positive secondary ion yield and the electron affinity strongly influences the negative secondary ion yield. Cs + enhances negative secondary ion yields, while O + 2 enhances the positive secondary ion yield. For this reason, Cs + is generally used to detect electronegative elements (H, C, O etc.) while O + 2 is generally used for electropositive elements (Na, K, Mg, Ti etc.). The choice of the ion beam also affects the depth resolution since a more massive primary ion penetrates less deeply into the sample and thus, at a given energy, can provide better depth resolution, as compared to a lighter primary ion. For analysis of multiple species in a sample matrix, it can often be necessary to perform analyses using various analytical conditions including different primary ion species and/or energies to achieve the desired detection limits and/or depth resolution. 74

100 3.4.2 Primary Beam Energy and Angle of Incidence The choice of the primary ion impact energy affects the depth resolution, secondary ion yield and the sputter yield. Sputtering using an ion beam leads to a churning effect, known as ion beam mixing, which causes general mixing of sample species with the result that impurity atoms are pushed deeper into the sample. The mixing depth increases with increasing primary impact energy and increase in mixing depth degrades depth resolution. Thus, lower impact energy is required for better depth resolution. Secondary ion yield and sputter yield are related to impact energy. Figure 3.4 shows the effect of the impact energy on the secondary ion (SI) yield of Si, bombarded at normal incidence with O + 2, Cs +, Ar + and Xe +. As the primary ion beam energy is increased from 2 to 12 kev/atom, the SI yield decreases by a factor of 5 for Cs, and essentially remains constant for O On the other hand, the sputter yield, which is the number of atoms sputtered per incident primary ion, increases with energy for all primary ion species of interest over the energy range 0 to 10keV/atom, as shown in fig.3.5. While sputter yield does decrease at higher energies, this is not the case over the energy range typically used for SIMS analysis. 7 Figure 3.4: The effect of primary ion beam energy on the secondary ion yield 6 75

101 Figure 3.5: Effect of primary beam energy on sputter yield for normal and 60 o angle of incidence from the normal 7 The angle of incidence at which the primary ion beam strikes the sample surface also affects the sputter yield, secondary ion yield and the depth resolution. The angle of incidence is measured from the sample normal and sputter yield increases with increasing angle to a maximum after which the yield decreases. The secondary ion yields decrease as the angle of incidence is increased until grazing incidence is approached, as shown in fig. 3.6 and the depth resolution is seen to improve at higher incidence because the collision cascade in the sputtering process occurs more closely to the surface than at near normal incidence. Relationships for the penetration depth to the primary ion species, beam energy and angle of incidence are given by: O + 2 : R = 2.15 E cos θ (3.1) Cs + : R = 1.84 E 0.68 cos θ (3.2) Where E and θ are the energy and angle of incidence respectively. 8 76

102 The actual angle of incidence depends on the sample potential and the primary beam impact energy. For example a 30 degree nominal angle of incidence (sample normal at 30 degrees to the primary beam), and a sample potential of +4500V, a 10keV positive ion primary beam will experience some deflection away from the sample normal, before striking the sample surface, making the actual angle of incidence to be greater than 30 degrees. If the primary beam energy is reduced, this deflection angle will increase, leading to an even greater actual angle of incidence. Figure 3.6: Effect of angle of incidence on secondary ion yield using an O 2 + and Ar + beam 6 77

103 3.4.3 Sputtering Rate and Detected Area The sputtering rate depends on the mass, energy and the angle of incidence of the bombarding ions, along with the density and bond strength of the sample atoms, with a weak dependence on the crystal orientation and type. Factors including the surface binding energy of the sample and the current density of the ion beam (current/area) also affect the sputtering rate. Sputtering rate and sputter yield are closely related to the mass, energy and angle of incidence of the primary ions. Figure 3.7: The effect of a smaller raster/detected area ratio on the depth profile of As in Si : A denotes the larger raster area (220 x 220 μm 2 ) and B denotes a smaller raster area (80 x 80 μm 2 ), shaded area denotes the detected area (60μm diameter) 4 In SIMS, the primary ion beam current and rastered area are used to control the sputtering rate for depth profiling. If all other factors are kept constant, the sputtering rate increases with the current density of the primary beam. 78

104 Increasing the sputtering rate may decrease the depth resolution because of the greater depth interval between data points, but conversely the detection limit is improved. For elements present as residual gas species (H, C, O, N), additional improvement in detection limit is obtained since contaminants deposited in the analysis crater are a smaller fraction of the atoms removed by sputtering. 9 The SIMS detection area of the sample is of primary importance in terms of depth resolution. In a depth profile, the total rastered or sputtered area cannot be used for detection of secondary ions because of contributions from crater walls. Figure 3.7 shows the effect of a small raster/detected area ratio. For a profile taken with an 80 x 80 μm 2 raster area and a 60 μm diameter detected area, the profile is distorted as compared to a 220 x 220 μm 2 raster area with the same detected area. Thus, depending on the quality of the beam shape, one may be able to obtain usable results from a crater width approximately 2.5 times that of the width or diam diameter of the detected area. 4 Fig. 3.7 shows the effect of too small a ratio of the raster area/ detected area on a depth profile, also known as the crater wall effect SIMS Quantification 4 The purpose of SIMS quantification is to relate the measured secondary ion intensity I A of element A to its concentration C A in the sample. This relationship is given by the expression shown in equation: I A = I p Y P ± f C A...(3.3) where I A is the secondary ion intensity of element A, I p is primary ion current, Y is sputter yield, P± is the ionization probability of A, f is the instrument transmission factor (which includes secondary ion extraction efficiency, mass spectrometer transmission efficiency and detector efficiency for the measured mass to-charge ratio), and C A is the fractional concentration of element A in the surface layer. In this expression, I P can be measured and Y and f are held constant for a particular set of analytical conditions (see importance of this in discussion of equation 3.4 below), while the ionization probability is related not only to the element and the substrate matrix but also to the species of the primary ion in SIMS analysis. The difficulty of elemental quantification in SIMS results from the complexities of ionization. The secondary ion yields vary over six orders of magnitude from element to 79

105 element across the periodic table as shown in figure 3.8 which illustrates the variation of relative sensitivity factor (RSF) in Si matrix under O + 2 and Cs + bombardment. Each element s ion yield is also affected by the matrix in which it is contained (matrix effects). The RSF technique is the most widely used calibration technique for quantitative analysis in SIMS if the impurity to be quantified is less than 1% of the total matrix concentration. Above 1%, care must be taken since the impurity concentration is no longer negligible with respect to the matrix concentration. In this situation, the presence of the impurity may have an impact on both impurity and matrix secondary ion yields. The sensitivity factor is defined according to equation: I m /C m = RSF i (I i /C i )...(3.4) Where I m and C m are the secondary ion intensity and concentration of matrix element and I i and C i are the secondary ion intensity and concentration of element i. RSF i is the relative sensitivity factor of element i. Note that while Y and F are not known, they are constant for both I M and C M and thus they cancel. In trace element analysis, the matrix elemental concentration is assumed to be constant. The matrix concentration can be combined with the elemental RSF i to give a more convenient RSF: RSF = C m RSF i = (I m /I i ) C i or C i = I i /I m (RSF)...(3.5) The RSF is a function of the element of interest, the sample matrix and the SIMS experimental conditions used. If RSF is known for a particular matrix and SIMS experimental conditions, the elemental concentration can be calculated from equation 3.5. An RSF for an impurity i can be calculated using an external standard sample with which has been implanted with a known dose known dose of this impurity. From an ion implanted standard, the RSF is determined from the equation: RSF = (φ I m t)/(d I i ) (3.6) where φ is implanted dose of the impurity, I m is matrix secondary ion intensity in counts/sec, t is the total sputtering time, d is the crater depth and I i is the sum of the detected secondary ions over the depth profile. 80

106 Figure 3.8: RSFs of various elements in Si for O 2 + and Cs + bombardment 4 81

107 The most common and accurate method of obtaining RSFs is via standards produced by ion implantation of the element in the matrix. Figure 3.9 shows a typical quantification and RSF calculation flowchart for measuring O in Nb using an O ion implant in Nb as the standard. Figure 3.9: Quantification procedure followed in SIMS RSF data have been published for a variety of matrices One study showed the instrument dependence (the deviation in RSF within one instrument group) and time dependence (relative standard deviations over five years) can be less than ±50% using the RSF method. 13 Typical precision of quantification with the same instrument and analysis condition is approximately 20% but very precise measurement can reach less than 1%, such as As in Si with NIST standards. 14 Standard samples can be fabricated by ion implantation, using either implantation into a matrix which previously had none of the implanted impurity or a standard addition method in which case a higher level of impurity is implanted into a matrix already 82

108 containing a lower level of the impurity. 15 Implantation standards must meet the following requirements. 1) Sample and standard should match in matrix composition. The ion yield varies if the matrix composition of the sample and the standard are different (matrix effect), which generates errors in the quantification of species in the unknown sample. 2) Matrices used for standards must be homogeneous to ensure the repeatability at different locations on the standard sample. 3) The ion implant dose and energy should be appropriately chosen. The dose should be sufficiently high to obtain good counting statistics but sufficiently low to avoid secondary ion yield changes which generally occur at atomic concentrations higher than approximately 1%. 4 The ion implant energy should be sufficiently high to place the implant peak deeper than the equilibrium depth but sufficiently low to ensure that the implanted ions are in the layer of interest in the case of implanting thin layers SIMS Analysis of Residual Gas Species 9 Since this work involves the characterization of H, C, N, O in Nb, which also are present as residual gases in the SIMS instrument vacuum system, it is important to SIMS analysis techniques used to minimize background contribution and improve detection limits of the instrument. One method to improve the detection limits for H, C, N and O is to improve the vacuum in the analysis chamber. If samples can be inserted into the sample chamber on the day prior to analysis and thus evacuated overnight, a lower residual gas level in the sample chamber can be obtained. Other improvements in SIMS residual gas detection limits can often be obtained by baking of the sample holder and sample before insertion into the vacuum system, and via the use of a cryo-cooled surface close to the sample Another method to reduce background contribution to detected residual gas species in SIMS is increasing the primary ion current density and thus the sputtering rate used for analyses. Studies have shown that detection limits for elements present in the residual vacuum can be improved by increasing the primary beam current density since increased primary current density increases the sample erosion rate compared to the rate of adsorption of gases. 16 To obtain this improvement, the sample erosion rate must be high compared to the adsorption rate of the impurity from the sample chamber vacuum ambient. This condition is met if P < 83

109 10 4 J p, where P is the residual-gas pressure (Pa) and J p is the primary ion beam density in μa/mm A related approach called raster collapse, includes reducing the raster size, which reduces re-deposition from outside the analysis area Raster Reduction 20 A method to estimate residual gas contribution to the detected secondary ion intensities of residual gas species such as H, C, O and N, is the raster reduction technique. In this method, a comparison of the changes in the intensities of the respective matrix and impurity secondary ion intensities resulting from an increase or decrease in raster size while keeping the primary ion current constant, is observed. In the case where the background does not influence the impurity secondary ion intensity, the ratio of the impurity secondary ion intensity to the matrix secondary ion intensity must remain constant regardless of any variation in primary ion current density. For example, if decreasing the raster area by a factor of 2 for the same primary ion current (hence, increasing the beam density and thus the sputter rate by factor of 4) during analysis increases the secondary ion intensities of the impurity element by a factor of four, which is the same as the increase in the matrix secondary ion intensities, then it is confirmed that the ions detected are sample related. Figure 3.10 shows a raster reduction profile (Counts/sec vs Time) of a non heat treated Nb sample. 1E+06 1E+05 Nb- Counts (cts/sec) 1E+04 1E+03 1E x 120 μm 2 60 x 60 μm 2 20nA NbN- 20nA 1E+01 1E Time (s) Figure 3.10: Raster reduction Cts vs Time profile for NbN - in Nb 84

110 As seen in the figure, the impurity species NbN - used to measure N in the sample increases equivalently to the matrix species Nb -, when the raster is reduced from 120 x 120μm 2 to 60 x 60 μm 2, while maintaining the same primary ion beam current. Such a profile shows that there is no measurable background contribution to the NbN - signal and thus the secondary ion intensities are related to the sample. 3.5 SIMS Characterization of impurities in Niobium The above discussion provides the basis for choosing the analysis conditions for the SIMS characterization of H, C, O and N on the Nb surface. As mentioned in sections and 3.4.3, SIMS characterization requires a trade-off between experimental parameters, depending on the information desired from the data. Hence, two types of analyses were performed on the samples for better detection limits and depth resolution respectively. Unless otherwise noted, all SIMS analyses were performed using a CAMECA IMS-6f Dynamic SIMS instrument (details of this instrument provided in chp 2), and analysis conditions were similar for both control and heat treated samples High Energy SIMS analysis For better detection limit and sensitivity, SIMS analyses were performed at a 14.5 kev impact energy (the primary beam was directed at the sample at 10keV energy and the sample was kept at a voltage of -4.5V). Since H, C, O and N are electronegative elements, a Cs + primary ion beam was used which enhances negative secondary ion yield (section 3.4.1), providing higher secondary ion intensity and thus, higher sensitivity. The Cs primary ion current was maintained at 20 na, and the beam was rastered over a 120 x 120 μm 2 area to obtain the desired sputtering rate. A 30μm diameter detection area was used in conjunction with the above sputtered area to avoid crater edge effects which can compromise depth resolution (section 3.4.3). A mass resolution of m/δm = 2200 was used, which was sufficient to avoid mass interferences such as D interference with 2H (required m/δm= 1600) and CH 4 interference with O (required m/δm= 500) To improve vacuum levels and reduce surface contamination, the samples were vacuumpumped overnight and the sample chamber was maintained at about Torr during analysis. Also, liquid N 2 was used in a trap around the primary beam sample interaction 85

111 region to condense residual gas species and minimize the background levels. The raster reduction method (section (3.4.6)) was also employed for some of the samples to estimate background contribution Low Energy SIMS analysis As mentioned in chapter 1, the penetration depth of the magnetic field for Nb is 40-60nm. Thus, it was imperative to characterize impurities on the Nb surface over this depth using low energy SIMS for better depth resolution. Some samples were thus analyzed using a 6keV impact energy Cs + beam (the primary beam energy was chosen to be 5keV while the sample was kept at a voltage of -1V). The primary ion current for this analysis was maintained at 7nA and the primary ion beam was rastered over a 200 x 200 μm 2 area. To avoid crater edge effects which can compromise the desired depth resolution, a 60 μm diameter detection area was used in conjunction with the above sputtered area. The mass resolution and vacuum arrangements were kept the same as in the high energy analysis to minimize mass interferences and residual gas background contribution respectively. 3.6 TOF-SIMS Characterization of Nb Bicrystals Since cavities heat treated with the niobium samples are made of large grain niobium, results from characterization of single grain samples were not fully indicative of the cavities, since grain boundaries were not taken into account. Segregation of impurities at the grain boundaries can degrade cavity performance (see section 1.6 and 4.8), and for this purpose, TOF-SIMS imaging was used to study the grain boundaries of large grain Nb bicrystals for presence of H,C, O and N impurities for heat treated and non heat treated (control) samples Sample Preparation Samples of 7mm x 5mm x 2mm dimensions, similar to the dimension for the single crystal samples, were cut by wire electro-discharge machining at Jefferson Labs, from large grain polycrystalline Nb disc manufactured by CBMM Brazil, in such a way that the grain boundary divided the sample into two parts, with each part side of the sample having a distinct crystallographic orientation. 86

112 Two types of bicrystals were prepared: (a) Bicrystals having different crystallographic orientation combinations: Heat treatment on one of the samples from this set was performed using the Elnik resistive heating furnace for 800 o C/3hrs, 120 o C/24hrs. The other sample was not heat treated and thus, was a control sample. The two samples had differing crystallographic orientation combinations. Electron Back Scatter Diffraction (EBSD) analysis was carried out on both the samples using a TSL/EDAX machine with OIM analysis software to determine the crystallographic orientations of both crystals in each sample. A description and features of this instrument can be found elsewhere. 21, 22 Figure 3.11 shows the EBSD results, the color coded plot shows that the left crystal has (001) orientation while the right crystal has higher orientation indices for both the samples [(213) for control and (215) for heat treated respectively] (a) (b) Figure 3.11: EBSD results for control (a) and heat treated (800 o C/3hrs, 120 o C/24hrs) (b) bicrystal samples. Color coding similar for both samples 87

113 (b) Bicrystals having the same crystallographic orientation combination: Heat treatment on one of the samples from this set was performed using the induction furnace for 1200 o C/6hrs, while the other sample was non heat treated (control). Both samples had the same crystallographic orientation combination TOF-SIMS Imaging: Analysis Conditions Imaging of the bicrystal grain boundaries was performed using an ION-TOF V TOF-SIMS instrument. The details of this instrument are provided in chapter 2 (section 2.4). A 10keV Cs sputter beam was used with a beam current of 20nA. The beam was rastered over a 180 x 180 μm 2 raster area and a Bi 3+ analysis beam with an energy of 25keV was used to analyze the sputter area. The analysis beam current was maintained at 7nA, with a detected area of 100 x 100μm 2. Burst alignment mode (no bunching) was used since this provides better spatial resolution for imaging (see section 2.4.1), with a 100ns pulse width. In the following chapters, the theoretical aspects of the properties of Nb containing H, C, O and N in Nb will be discussed and SIMS and other characterization results are presented for each of the heat treatments based on SIMS analysis using the above analysis conditions. The chapters are divided into the type of impurity characterized. 88

114 References 1. G. Myneni et. al.: AIP Conf. Proc. 927, p84 (2006) 2. G. Ciovati et al: Phys. Rev. Special Topics - Accelerators and Beams, 13, p (2010) 3. P. Dhakal et al: Rev of Sc. Inst., 83, p (2012) 4. F. Stevie et al : SIMS- A practical handbook for depth profiling and bulk impurity analysis, Wiley Interscience Publications (1989) 5. R. Levi-Setti et al : SIMS V, Springer Berlin, p132 (1986) 6. K. Witmaack : App. Surf. Sc., 9, 315 (1981) 7. P.Zalm : Journ. Of App. Phys., 54, p2660 (1981) 8. K. Witmaack : Nuc. Intrum. Meth., 218, p307 (1963) 9. R. Levi-Setti et al : Nuc. Intrum. Meth., 168, p139 (1980) 10. F. Stevie et al : Journ. Vac. Sci. and Tech. A11, p2373 (1993) 11. P. Kahora et al : SIMS VII, Wiley New York, p143 (1988) 12. R. Wilson et al : SIMS VII : Wiley New York, p131 (1988) 13. Y. Homma : SIMS IX, Wiley New York, p135 (1993) 14. NIST Standard Reference Materials, SRM 2137, G. Leta et al : Analyt. Chem., 52, p277 (1980) 16. J. Kaboyashi et. al: Journ. Vac. Sc. Tech., A6, p86 (1988) 89

115 17. J. Dupoy et. al: SIMS VI, Wiley New York, p277 (1988) 18. Y. Homma et. al: SIMS V, Springer, Berlin, p161 (1986) 19. A. Czanderna: Beam Effects: Surface Topography and Depth Profiling in Surface Analysis, New York, p394 (1998) 20. A. Pivovarov et. al: J. Vac. Sci. Technol. A 21(5), p1649 (2003) 21. R. Abart et. al.: Cont. to Min. and Pet, 147, p633 (2004) 22. T. Abe et. al: Trans. of the Japan Soc. of Mech. Engg, A, 69(6), p972 (2003) 23. Samples nanoplished by Wah Chang-USA with proprietary process 90

116 4. Hydrogen Hydrogen has been recognized as being responsible for the degradation of Q o observed in Nb RF cavities, called Q-disease. 1, 2 The characterization of H in Nb, and investigation of the physical and chemical properties of the Nb-H system can provide useful information for understanding of hydrogen penetration with the goal of reducing hydrogen levels in Nb. The Nb-H system has been studied extensively 3, both to understand the corrosion resistance of Nb as well as for hydrogen storage since Nb is among metals able to absorb a large amount of hydrogen even at room temperature (for instance, one can store five times more hydrogen in Nb than in an equivalent volume of liquid hydrogen). Careful investigation can help provide information on the mechanism by which hydrogen enters Nb during the preparation of Nb for superconducting cavities Source of Hydrogen Contamination in Nb Cavities After fabrication, the inner surface of RF Cavities requires processing in order to remove damage layers and any surface contamination. The two most important methods for preparing these surfaces are Buffer Chemical Polishing (BCP) and Electropolishing (EP). BCP is performed using HF, HNO 3 and H 3 PO 4 (1:1:1-1:1:4 in volume). HF provides the F - ions, which complex the Nb 5+ ions and drive them into solution while NO - 3 ions act as an oxidant toward the Nb metal to transform it in to Nb 5+ ions (oxidized). The role of H 3 PO 4 is to control the viscosity and/or etching rate of the mixture. EP involves the etching of the material in a HF-H 2 SO 4 solution (15-85 in volume), and the oxidation of Nb results from anodic polarization of the cavity ( 8-10 volts). F - plays the same complexing role as in BCP and H 2 SO 4 is used as a buffer with high viscosity, known to improve the surface state in many polishing recipes. Both these techniques are known to provoke heavy Q-disease in the cavity and are believed to be the main cause of H contamination. 1,4 H contamination inside Nb is in the form of interstitial individual atoms. Thus, the molecule H 2 is not the only contamination source, and the species present inside aqueous solutions: H 2 O, H +, OH - etc also are also likely candidates. Molecular hydrogen does not dissociate on oxide, although it does readily on bare metals as a result of catalyzation process. While the oxide shields Nb from hydrogen formation in air, it is also known to be amphoteric 5, i.e. it 91

117 can both accept and donate a proton. During chemical etching and electropolishing, which occur in concentrated acid mixtures, the presence of high concentration of H + ions directly affects the H contamination in Nb. Four types of competing reactions are involved during the acidic etching of Nb: 6 (1) Nb + H 3 O + + e - (Nb) H ads Nb + H 2 O (fast) After H ads Nb is formed, which is a fast process, any one of the below reactions can occur, all of which can contribute to H being absorbed in to Nb. (2) H ads Nb + H 3 O + + e - (Nb) Nb + H 2 + H 2 O (slow) (3) H ads Nb + H ads Nb 2M + H 2 (slow) (4) H ads Nb H s Nb (Absorption process, slow, competing with (2) and (3)) H ads represents hydrogen, in the atomic form, adsorbed on the surface of the metal, and H s is the dissolved or absorbed hydrogen under the surface of the metal. Process (4) is probably the main source of H interstitial contamination since molecular H 2 has to dissociate first before diffusing in to Nb, a process which requires additional energy. The composition and the structure of the surface oxide layer, as well as the presence of adsorbed impurities on it also has some influence on the kinetics of reaction (4). Halogen ions like F - and Cl - and/or cathodic polarization of the Nb surface are known to depassivate the oxide layer and favor absorption of H and precipitation of hydrides. 6 Oxidizing conditions result in the dissociation of the hydrides and formation of a new oxide layer. This is the reason that the presence of a strong oxidant like NO - 3 in chemical reduces hydrogen uptake as compared to H uptake in electropolishing. The product of reaction (2) is H 2 O instead of H 2, which explains why molecular hydrogen is not produced during chemical polishing. For instance, the reaction (2) can be rewritten as: (5) H ads Nb + H 3 O + + e - (Nb) + NO - 3 Nb + 5H 2 O + NO 2 Hydrogen contamination has been measured after chemical polishing and is low i.e. a few ppm by weight, maximum in the bulk 7, which is clearly insufficient to generate significant hydride precipitates. However, this hydrogen contamination segregates near the Nb surface, i.e precisely in the penetration depth region of Nb (λ = 40nm) and in this region, the H concentration can reach several %. 8 A well known phenomenon of H in metals is its tendency 92

118 to interact with crystal defects like impurity atoms, dislocations, grain boundaries etc. forming a Cottrell cloud, where H is concentrated and where H levels can reach hydride precipitation limits. 9, Effect of Hydrogen on properties of SRF Nb Interstitial impurities like H are particularly detrimental to SRF cavity performance, since these impurities can affect the mobility of electrons, and have a strong effect on the thermal conductivity of Nb. This performance deterioration manifests in the form of residual losses and overall Q-disease in RF cavities Residual Losses from Hydrides An important residual loss mechanism arises when the hydrogen dissolved in bulk Nb precipitates as a lossy hydride at the RF surface. This residual loss, also known as Q- disease is a subtle effect that depends on many factors, including the hydrogen concentration, the rate of cool down and the amount of other interstitial impurities present on the Nb surface. If the bulk hydrogen concentration in Nb exceeds 2ppm by weight, there is a clear danger of hydride formation at the RF surface during cool-down. The effect can be severe enough to lower the Q o by two orders of magnitude depending on the amount of hydrogen dissolved. 1 As delivered, commercial Nb typically has less than 1wt ppm of dissolved hydrogen because the material in its final form is usually annealed for recrystallization, but as mentioned earlier, the hydrogen concentration can increase during chemical etching of the surface to remove surface damage. According to the phase diagram of the Nb-H system 11, hydrogen precipitation in Nb is not prevalent at room temperature, but as the temperature is lowered (like in the case of cavity operation), the hydrogen concentration needed to form the hydride phases decreases. A cavity can be cooled as slowly as desired at temperatures above 150 K, since the hydrogen concentration to form hydride phases is still high. However, below 150 K, the hydrogen concentration required to form these phases decreases to a dangerously low level, allowing islands of the hydride phase to form even when H concentration is as low as 2wt ppm. The hydride precipitates at favorable nucleation sites, and if these are at the surface, they increase the residual loss. Also, the diffusion rate of hydrogen in Nb below 150 K is still (put value here) and thus significant, so that hydrogen can move to accumulate to 93

119 critical concentrations at nucleation sites. Fig. 4.1 shows the phase diagram of H in Nb and the effect of lower temperatures on H precipitation. Fig. 4.1: Nb-H phase diagram: Hydride phase formation at low temperature and H concentration 12 Thus, when a cavity with a large bulk hydrogen concentration is cooled to liquid helium temperature, the extent of the hydride formation leads to residual losses in the cavity Critical Temperature (T c ) A number of different measurements have been made to investigate the relationship of the hydrogen concentration to the critical temperature of superconducting Nb These are reviewed by Isagawa 12. Figure 4.2 shows data from various research groups on the 94

120 dependence of T c of niobium on hydrogen concentration. The transition temperature was obtained by the dc resistivity method for all measurements. The value ΔT c indicates the temperature width of the transition (full width). In the Isagawa results, (figure 4.2), although T c and ΔT c initially show a slight decrease and increase, respectively, they remain almost constant for the hydrogen concentrations between 2.5 and 24 atomic %. These results agree with those reported in other reference which indicated that there were no detectable changes in T c with increasing hydrogen concentration. Fig. 4.2 : T c vs H conc. observed by various groups using the DC resistivity method In the literature, however, contradictory results have been reported which indicated that T c decreased monotonically with the increase of hydrogen, as seen in figure 4.2 from Wiseman and Horn et al According to Isagawa, these contradictions can be ascribed to the 95

121 difference of sample form, purities of starting materials, and measurement methods. Samples which were reported to have shown the T c decrease were fine powders 16,18 or fine strips. 17 Thus, conflicting dependence of T c on the hydrogen concentration has been reported, but if the decrease of T c with hydrogen concentration is assumed to be correct, then H detrimentally affects the cooling requirements for the superconducting transition to occur in Nb, thereby increasing energy and thus operating costs Electrical Resistance Hydrogen is known to increase the resistance of Nb both in the normally conductive and in the superconducting state. Isagawa 12 measured the Nb resistance loaded with upto 20% atomic H, with an unloaded sample measured as the standard, at 285 K (normal state) and 9.5 K (near the superconducting transition temperature). Measurements were taken using a 4 point DC resistance measurement method, and the ratio of the resistance increase (ΔR) to the unloaded sample resistance (R) were plotted vs. the H concentration for both temperatures. Fig. 4.3 (a) and (b) show the results of these measurements for 285 K and 9.5 K respectively. Comparison of these results was also made with another study by Westlake 19, and the results were seen to be similar at 285 K. (a) (b) Fig. 4.3: Increase in resistance of Nb with increasing H conc. at 285K (a); Similar for (b) but at 9.5 K. 12,19 96

122 The excess resistivity is believed to be due to the scattering of the conduction electrons of Nb by interstitial hydrogen atoms which disturb the periodic variation of potentials within the Nb lattice, an effect which can severely hamper the performance of Nb SRF cavities Q disease As mentioned in section 4.2.1, hydrides can precipitate at temperatures below 150 K in Nb even with low H bulk concentrations. In fact, the H diffusion is fairly rapid (10-7 cm 2 /sec) between 150 K and 60 K, after which the diffusion is not as rapid. Thus, the length of time for which a cavity is kept between K determines the extent of hydride formation, which leads to residual losses and adversely affects the Q o of the SRF Nb cavity. Fig. 4.4 shows the effect of cooling rate on the Q o of SRF cavities. As seen, a fast cooling rate leads to higher Q o for a given electric field, while a slow cooling rate of 3-70 hours can lead to lower Q o values, which can be two orders of magnitude lower. This degradation of Q o is termed as Q- disease. 20 Fig. 4.4: Faster cooling rates can avoid Q o degradation by surpassing hydride precipitation temperatures 20 97

123 Hydrogen is thus seen to be a major factor in Q disease because of the high mobility of H in Nb even at 100 K temperature and its ability to lead to phase transitions between 150 K and 60 K while cavity cool down to cryogenic temperatures. The introduction of H in the material during chemical and electropolishing in high concentrations thus makes it imperative to study H levels in Nb and come up with ways to eliminate this hydrogen contamination. 4.3 Solubility, Diffusion and Site Occupancy of Hydrogen in Nb Solubility The solubility of any gas in a metal depends on the temperature, pressure and the type of lattice. For Nb, which has a BCC lattice structure, the tetrahedral and octahedral sites provide interstitial spaces for the H atoms to reside. The solubility depends on the pressure and temperature as: 21 S= s o.p 1/n. exp (±Q s /(nrt))...(4.1) Where n is the no. of gas atoms in one molecule, p is the partial gas pressure, s o is a constant and Q s is the activation energy of dissolution. At constant temperature, the solubility of a gas is determined by the Sievert s law : S α p 1/2.(4.2) for a diatomic gas in thermodynamic equilibrium. 21 The ± for the exponential term in equation 4.1 is an indication of whether the gas solubility increases or decreases with temperature respectively. The equilibrium hydrogen pressure p and concentration C H of hydrogen dissolved in niobium in dilute solutions as a function of temperature follow Sievert s law with the following equation: 39 C H (at.%) = 2.2 x 10-2 (p(h 2 )) 1/2 exp(2000/t)..(4.3) T is in K and p is in pa. This equation is valid from 275 < T < 2275 K and C H < 5 at%. Thus, H has the unique behavior in Nb such that at low pressures and high temperatures, the H dissolution in Nb has a negative heat of solution and thus is exothermic. This implies that, unlike in other metals, H solubility decreases with temperature at a constant low pressure. An estimate of the H solubility in the α phase of Nb above 400 o C and a constant pressure (below 10 4 pa) is given by: 22 S (H 2 /Nb) = exp (4240/T) mol/m 3. pa 1/2 (4.4) 98

124 The positive exponential factor again indicates that the H solubility decreases with increasing temperature. On the other hand, as seen in the phase diagram for Nb-H, below 77 o C (350 K), with H in a saturated gas phase or in air, the hydrogen solubility increases with temperature in the α phase, which is the terminal solid solution of H in Nb: 12 C H at% = exp ( /T) mol/m 3.pa 1/2.(4.5) with T in o K and constant pressure. Early important studies of the solubility of H in Nb were performed by Albrecht et al, for a temperature range of o C. In this work, it was found that at a constant pressure, the isobars of Nb-H system had a deceasing trend with an increase in temperature. Fig. 4.5 shows the different solubility curves of H in Nb at 10, 100 and 1000mm Hg. The invariance in the middle of the curves was believed to be due to the formation of a second phase, along with the α phase. It was also observed that the Sievert s law for H in Nb was only valid for dilute solutions (H/Nb= 0.055) 23, as seen in equation 4.3. Note that the first half of each of the curves shows the primary α-phase solid solubility of H in Nb. Figure 4.6 shows the pressure- composition relationship for the Nb-H system, studied by Veleckis et.al 55, A reference of this type of a diagram not only indicates the equilibrium pressure of hydrogen at different temperatures, but also provides information on the solubility of hydrogen at different temperatures and pressures. For example, the solubility is seen to be very small at 1000 o C, particularly under high vacuum conditions (0.001% atomic under 10-3 Pa). 99

125 Fig. 4.5: Increasing solubility of H in Nb with decreasing temperatures at various pressures (mm Hg) 23 Fig.4.6: P-composition isotherms for different concentrations of H in Nb

126 At low temperatures however, there is a possibility of pick up of hydrogen because of the solubility increasing with decreasing temperatures. It has been observed in literature that below 600 o C, oxide like sorption layers are formed on the Nb surface and these act as barriers for exchange of hydrogen between gas and metal Diffusion Hydrogen diffusion in Nb is extremely fast. Diffusivity and solubility are oppositely influenced by temperature in Nb, as is the case for palladium and the other two group V b metals, vanadium and tantalum. Volkl et al have summarized six measurements of diffusivity of H in Nb over the temperature range -50 to 600 C, from which an average diffusion coefficient equation can be deduced as: 24 D = 5 x 10-4 exp (-1230/T) cm 2 /sec (4.6) where T is in K. Fig. 4.7 shows a plot of the diffusion coefficients of H in various metals vs temperature. The temperature dependence of the diffusion coefficient is evident as it increases with increasing temperature for Nb-H and Nb-D. Note that the plot is of ln(d) vs 1000/T but the values of ln(d) have been replaced by the values of D. The activation energy of diffusion of H in Nb has been experimentally determined to have a mean value of 10,100 cal/mol (0.45 ev/atom). 43 An sense of how fast hydrogen diffuses in Nb can be gained by comparing its diffusion coefficient to that of H in Si. Table 4.1 shows such a comparison, where hydrogen is seen to diffuse about 22 orders of magnitude faster than in Si. 101

127 Fig. 4.7: Nb-H and Nb-D diffusion coeff. vs Temperature 24 Table 4.1: Comparison of H and D diffusion coefficients in Nb and Si respectively at 300 o K 24,25 Matrix H Diff. Coeff. (cm 2 /sec) D Diff. Coeff. (cm 2 /sec) Si 1 x x Nb 2 x x 10-6 Hydrogen has a large mobility, not only in Nb, but in other metals as well. 57 Thus, a brief review on the various mechanisms of diffusion of H in metals is essential Since hydrogen has a small mass compared to other interstitials, quantum effects in diffusion are likely to be observed for hydrogen. Fig.4.8 shows the diffusion process of hydrogen at 102

128 various temperatures. 60 At the lowest temperatures, hydrogen is delocalized in the form of a band state, unless it is trapped by lattice defects. The propagation in the band state is limited by the scattering on thermal phonons or lattice defects. At some higher temperature, the H interstitial will be localized at, or about, a specific interstitial site. The elementary step of the diffusion process is now a thermally activated jump from one to another interstitial site. The H particle might execute the jump by tunneling from one to another interstitial site or by hopping over the potential barrier. In the first case, thermal activation is necessary to bring the energy levels of both sites to the same height. In the second case, a higher activation energy is required to overcome the barrier. Hence the occurrence of movement is more frequent at higher temperatures. Finally, at the very high temperatures the hydrogen interstitial will mainly be in states above the potential barriers and diffusion of H becomes similar to that occurring in a dense gas or a liquid, where many collisions occur. Collisions in this case are with the thermally fluctuating host lattice. Fig. 4.8: Diffusion processes of hydrogen at different temperatures 60 This regime is thus called fluidlike diffusion. Interstitial diffusion is characterized by the jump rate of the interstitial and an accepted view of interstitial diffusion involvesa jump process that is short compared to the average time between two consecutive jumps. In fluidlike diffusion, hydrogen atoms perform so many jumps per second, especially in bcc 103

129 metals, that the average time between two consecutive jumps is short compared to the duration of a jump and states above the barriers contribute increasingly to the diffusion process thus making the diffusion is extremely fast. The boundaries separating the different diffusion regimes are not sharply defined and one can consider subdivisions and modifications within one regime. Also, it is not certain that all regimes actually occur in a given system. It can however be inferred that, even at the lowest temperatures, hydrogen can diffuse in Nb via tunneling Such a high mobility of H in Nb poses problems for SIMS characterization of H on and near the Nb surface. As mentioned in chapter 2, for accurate determination of H concentration, SIMS analysis requires the determination of the Relative Sensitivity Factors (RSF) of H in Nb, which is determined by noting the H intensities obtained from the SIMS depth profile an ion implant of a know energy and dose H implant in Nb. The integration of the H intensity obtained from the SIMS depth profile of the implant peak is thus essential for obtaining the relative sensitivity of H in the Nb matrix. The problem arises when the H distribution of H in Nb cannot be obtained due to the extremely high mobility of H in this material. Since SIMS measurements require a certain area of the sample to be sputtered, this sputtering may create an H concentration gradient, due to the sputtered area having reduced hydrogen. Such a concentration gradient along with the creation of vacancies due to sputtering induced surface damage could lead to a very fast movement of H atoms in Nb in which, H from surrounding areas of the sample could diffuse extremely rapidly into the sputtered area. Since this is the area from which the SIMS H secondary ion intensity is extracted, this rapid diffusion into this region could result in detection of constant hydrogen intensity. A non heat treated large grain sample of Nb and a single crystal sample of Si were implanted with D. Figures 4.9 and figure 4.10 show the results of the SIMS depth profile analyses of these samples. As seen in the figures, analysis of D in Si (Figure 4.9) shows the expected implant distribution of D resulting from ion implantation. The SIMS depth profile of D in Nb does not show any indication of the expected implantation profile. Whether due to the inherent high diffusion coefficient of D and H in Nb or whether the diffusion rate is further 104

130 enhanced by sputter enhanced diffusion, this result clearly supports the hypothesis that diffusion of D (and thus also H) in Nb is very rapid (Table 4.1). In order to be able to perform useful SIMS depth profiles of H in Nb, either the diffusion coefficient of hydrogen in Nb has to be lowered (analysis at lower temperatures) or the sputtering rate of analysis has to be more rapid than the diffusion rate of H in this material. Such sputter rates are not possible using the CAMECA IMS-6f SIMS instrument and the 70K sample temperature possible using the cold stage on the ION-TOF 5, TOF SIMS is not sufficiently cold to sufficiently reduce the diffusion rate of H in Nb. Thus it was not possible to obtain H in Nb in depth concentration distributions. This issue does not arise with C, O, N analysis using SIMS, since the diffusion of these species is about 10 orders of magnitude lower than H in Nb (see Chapters 4 and 5) Si Implant 1E+20 28Si Counts 1E+08 Concentration (atom/cm3) 1E+19 1E+18 1E+17 1E+16 1H Counts 2D conc 1E+07 1E+06 1E+05 1E+04 1E+03 1E+02 1E+01 Counts (cts/sec) 1E Depth (um) Fig. 4.9 : D in Si implant shows an implant peak 1E

131 Nb implant (C ts/sec) 1.E+08 1.E+07 1.E+06 1.E+05 1.E+04 1.E+03 1.E+02 1.E+01 1.E+00 93Nb Counts 1H Counts 2D Counts Depth (um) Fig : D in Nb implant shows a constant D signal Site Occupancy As depicted in section 1.3.3, the BCC structure of Niobium allows the possibility of 12 tetrahedral interstitial sites and 6 octahedral interstitial sites (Fig. 1.7). The size of the octahedral site is given by the relation, R oct = 0.155r Nb and that of the tetrahedral site is given by: R tet = 0.291r Nb. Thus, the size of the tetrahedral site in Nb is larger than the octahedral site. Hydrogen occupies the tetrahedral sites in Nb. 26 To understand why this is so, the sizes of the tetrahedral and octahedral sites must be calculated and compared to the size of a hydrogen atom. From table 1.2, the size of the tetrahedral site is: R tet = x r Nb = Å The size of the octahedral site is: R oct = x r Nb = Å 106

132 Since the size of the H atom is determined by its covalent radius (0.31 Å), it is clear that only the tetrahedral sites provide sufficient space for the H atoms to fit. Although the H atoms are able to occupy the tetrahedral sites of Nb, there has been some lattice expansion observed in single crystal Nb samples at low hydrogen concentrations. This low concentration phase is the α phase (see phase diagram, figure 4.1), which is essentially the pure Nb body centered cubic (bcc) phase interstitially alloyed with H atoms randomly distributed over tetrahedral sites in the crystal lattice. Fig shows this lattice expansion in the form of Δa/a. 27 Fig. 4.11: Lattice expansion in Nb due to H SIMS Characterization of H in Nb. Hydrogen has been recognized as a strong contributor to the degradation of Q o observed in Nb RF cavities, called Q-disease. The characterization of H in Nb along with the investigation of the physical and chemical properties of the Nb-H system may provide information helpful in preventing hydrogen penetration and perhaps preventing any 107

133 meaningful H incorporation into the Nb surface. An empirical treatment which reduces Q- disease, discovered in 1998, utilizes a low-temperature ( o C, 48 h) baking of the cavities in ultrahigh vacuum. 28 In essence, it is believed that this baking is responsible removal of hydrogen and other impurities from the Nb surface (60 nm, penetration depth) and thus restoring the lost Q o. Since the treatment does not fully remedy the loss of Q o, further study of the effect of heat treatment as a purification step, including heat treatments at higher temperatures ( o C), is indicated. Several observations have been found in the literature which may link removal of hydrogen to the baking effect : (i) Thermal desorption studies on Nb foils, where the natural oxide layer was removed by cycles of thermal treatments and Ar ion sputtering showed hydrogen desorption peaks at 130 o C and 198 o C which were interpreted as hydrogen desorption from surface and subsurface sites. 29 Several experimental and theoretical studies 33,34 have concluded that Nb surfaces exhibit a strongly bound subsurface state of hydrogen. (ii) Measurements by positron annihilation spectroscopy (PAS) show that the defect density (vacancies) increases with hydrogen concentration in Nb samples. 35 This region of high defect density, which leads to areas of normally conducting state of Nb, has been termed as a hot spot due to its propensity for joule heating. Hydrogen entering Nb during chemical etching or mechanical polishing may possibly be the source of the high defect density in the hot spot samples measured by EBSD. 36 (iii) Hydrogen affects the magnetic behavior of Nb by lowering the magnetic susceptibility for increasing H concentration. 37 Measurements of hydrogen concentration profiles in Nb are difficult due to the high diffusivity of H in Nb. The fact that surface contamination even in high vacuum can produce false signals requires low H partial pressure in the vacuum to provide any prospect of detection of H on the first 40nm of the Nb surface. SIMS is a powerful technique for H surface analysis due to the high depth resolution (10-20nm and 1-2nm at low energies) of the technique and its ability to detect hydrogen while minimizing background effects. The latter can often be confirmed even without the use of standards, by a specific technique known as raster reduction (see section 3.4.6)

134 H near surface levels were characterized and compared for both non heat treated and heat treated Nb (120 o C 1400 o C) using SIMS. The effects of these heat treatments on the apparent relative surface concentrations of hydrogen in Nb were thus obtained and correlated to the Nb SRF cavity performance parameters, essentially the Q o. Preliminary SIMS mass spectra of a non heat treated sample showed intense hydride peaks as shown in figure Some of these NbH x hydride peaks were not observed in a heat treated sample (800 o C/3hrs, 400 o C/20min). Figure 4.13 shows the mass spectra of the heat treated sample, showing both a significant decrease in the intensity of the hydride peaks and the disappearance of some of the peaks, indicating hydrogen desorption after heat treatment. Thus, initial SIMS data confirmed the feasibility of high temperature heat treatments for H removal. Counts/sec 1E+07 1E+06 1E+05 1E+04 1E+03 1E+02 - ṈbH NbH - 2 Nb NbH 3 - NbH 4 - NbH 5-1E+01 1E Mass (a.m.u.) Fig. 4.12: Intense niobium hydride peaks in a non heat treated sample 109

135 1E+07 1E+06 1E+05 Nb - Counts/sec) 1E+04 1E+03 NbH - 1E+02 NbH 2-1E+01 1E Mass (a.m.u.) Fig. 4.13: Less intense hydride peaks in a 800 o C/3hrs, 400 o C/20min heat treated Nb sample as compared to control, higher hydrides are not seen after heat treatment Heat Treatments Chapter 3 describes the sample preparation, heat treatment parameters and SIMS analysis conditions used. In brief, four types of heat treatments were studied: 1. Low temperature baking (120 o C/48hrs) 2. High Temperature Heat Treatments (600 o C-1400 o C/3-10h) 3. High Temperature Heating followed by Long Term Low Temperature Baking (600 o C o C/3-10h; 120 o C/12-48hrs) 4. High Temperature Heating followed by Lower Temperature Heating (800 o C-1000 o C/2-3h; 400 o C-800 o C/10-20min) in vacuum or 10-5 torr N 2 Although a brief discussion of the rationale behind these heat treatments is given in section 3.2, it is important to discuss this rationale with the perspective of hydrogen as an impurity. As seen in section 4.4, an increase in the cavity performance has been reported for the long term low temperature baking of the SRF cavities in literature. 28 It is believed that this baking is responsible for the removal of hydrogen and other impurities from the Nb surface (60 nm, penetration depth) and thus the restoration of the lost Q o. Moreover, the solubility of hydrogen in Nb decreases with increasing temperatures, especially above 600 o C in vacuum 110

136 conditions (section 4.3.1), which is why it was important to observe any changes in hydrogen after high temperature heat treatment (heat treatment type 2) of bulk Nb at various elevated temperatures (600 o C-1400 o C). Heat treatment type 3, which involved high temperature heating followed by long term low temperature baking, was performed since preliminary results from Ciovati et.al 36 showed that, after low temperature long term baking of some of the large grain SRF cavities previously subjected to a high temperature heat treatment (specifically 600 o C/10hrs and 800 o C/3hrs), the cavity performance was further improved beyond the initial efficiency increase. 36 Heat treatment type 4 involved the introduction of N 2 during lower temperature heating of the samples after high temperature heating was done. The intent of this heat treatment was the nitridation of the Nb surface to prevent surface contamination of the sample on cool down thus preventing the adsorption and possible incorporation of H and other impurities (C, O) into the Nb surface once the furnace reached room temperature (see chapter 6) Results Raster Reduction SIMS Analyses Before the SIMS results can be discussed, it is important to ascertain the background contribution to hydrogen levels measured by SIMS. Raster reduction (section 3.4.6) was performed on some of the control and heat treated samples to check for background contribution. Figure 4.14 (a) shows the raster reduction profile for H - and Nb - for a control sample, where the raster was reduced to 60 x 60 μm 2 area from 120 x 120μm 2 area while maintaining a primary ion beam current of 20nA. 111

137 Counts (cts/sec) 1E+07 1E+06 1E+05 H- 60 x 60 μm x 120 μm 2 20nA 20nA Nb- 1E Time (s) Fig (a): Raster reduction for H - and Nb - for a control sample As seen in the figure, the hydrogen secondary ion intensities increase proportionally to the matrix (Nb - ) secondary ion intensities, indicating no background contribution to the hydrogen levels. Thus, it may be inferred that hydrogen in a control sample is above the detection limit of the instrument and thus that H intensity is proportional to that measured at the sample surface. Similar raster reduction results were observed for all the control samples analyzed in this study. However, SIMS results obtained from a treated sample that was high temperature heat treated to 600 o C/10hrs (figure 4.14(b)),showed that the H intensities do not increase proportionally to the increase in Nb intensities after raster reduction, indicating that there is some background contribution in the hydrogen intensities. This indicates that there is a contribution to the hydrogen intensities observed from contamination from the vacuum, and thus the hydrogen level in the sample is lower than that measured by SIMS. This was observed for all the heat treated samples for which raster reduction was performed. 112

138 Since the raster reduction profiles for all the heat treated samples were similar, it can be assumed that the background contribution to the hydrogen intensities was the same for all heat treated samples and thus, a comparison between the H levels of heat treated samples would be valid. 1E+06 Counts (cts/sec) 1E+05 1E x 120 μm 2 60 x 60 μm 2 20nA 20nA H- Nb- 1E Time (s) Fig (b): Raster reduction for H - and Nb - for a heat treated sample High Temperature Heat Treatments Figure 4.15 shows the H - /Nb - ratios vs depth for 600 o C-1400 o C/time heat treatments as compared to a control sample over a sputtered depth of about 1 μm. All the heat treatments show a significant decrease in the H levels as compared to a control sample. Specifically for the 1200 o C/6hrs heat treatment, the H levels are seen to decrease by a factor of 100 as compared to the control sample, while the 600, 1000 and 1400 o C heat treated samples show higher H levels in comparison to the 1200 o C heat treated sample. Nevertheless, these heat treatments also show significantly lower H levels as compared to the control sample. 113

139 Control 600C/10hrs 1000C/6hrs 1200C/6hrs 1400C/3hrs 100 Control H/Nb Ratio C/6hrs 1400C/3hrs 600C/10hrs C/6hrs Depth (um) Fig. 4.15: SIMS H - /Nb - vs Depth data for 600 o C-1400 o C heat treated Nb samples vs contro This difference in hydrogen levels obtained via the various heat treatments cannot be explained in terms of the time for which the heat treatment was performed, since the diffusion of hydrogen is sufficiently fast in Nb that the amount of hydrogen desorption should not vary significantly with the time of heating (in hours). Table 4.2 shows the average H 2 pressure in the furnace during heat treatment. A direct correlation is observed between the H 2 pressure and the H/Nb ratio of the sample after heat treatment, indicating that a higher H 2 pressure in the furnace leads to less desorption of hydrogen from the sample during heat treatment producing the variation in the hydrogen level observed after various heat treatments. 114

140 Table 4.2: Heat treatments and H 2 pressure in the furnace Heat Treatment (Furnace) H/Nb Ratio Average H 2 pressure (torr) 600 o C/10hrs (Resistive) x o C/6hrs (Induction) x o C/6hrs (Induction) x o C/3hrs (Induction) x High Temperature Heating followed by Long Term Low Temperature Baking Figure 4.16 shows the SIMS results after the high temperature heating followed by long term low temperature baking heat treatments: o C/2-10h, 120 o C/12-48h. As seen in the figure, the H - /Nb - ratio for these heat treatments is up to a factor of 100 lower than the control sample. The lower hydrogen levels obtained after these heat treatments as compared to the single high temperature heat treatments (section ) indicate that the long term, low temperature baking significantly contributes to the lowering of H in Nb. Table 4.3 shows the average H 2 pressure in the furnace for the respective heat treatments. Similar to the high temperature heat treatments (previous section), a direct correlation is observed between the H 2 pressure and the hydrogen level in the sample after the heat treatment indicating that the higher pressure may cause less desorption of hydrogen for that particular heat treatment. 115

141 100.0 Control 600C/3hrs,120C/48hrs 800C/3hrs,120C/12hrs 1000/2hrs,120C/12hrs 1200/2hrs, 120C/12hrs 1400C/3hrs,120C/12hrs Control 10.0 H/Nb Ratio /2h,120C/12h 1400C/3h,120C/12h 600C/3h,120C/48h 800C/3h,120C/12h 1200/2h 120C/12h Depth (um) Fig SIMS H - /Nb - vs Depth data for different high temperature heated-low temperature baked Nb samples vs control Table 4.3: Heat treatments and H 2 pressure in the furnace Heat Treatment (Furnace) H/Nb Ratio Average H 2 pressure (torr) 600 o C/10hrs, 120 o C/48hrs (Resistive) x o C/3hrs, 120 o C/12hrs (Resistive) x o C/2hrs, 120 o C/12hrs (Resistive) x o C/2hrs, 120 o C/12hrs (Resistive) x o C/3hrs, 120 o C/12hrs (Induction) x

142 High Temperature Heating followed by Lower Temperature Heating There were two types of heat treatments performed in this step. One was the 800 o C/3hrs- 400 o C/20min heat treatment performed under vacuum conditions through out. The other heat treatment involved high temperature heating, followed by a lower temperature heating during which N 2 gas was introduced into the chamber for a short period of time (given below) at 10-5 mbar pressure during this heat treatment process. Two samples were heat treated according to the latter method of heat treatment with one being further baked at 120 o C for 6hrs. N 2 was introduced to induce nitridation of the Nb surface to prevent surface contamination of the sample on cool down with the goal of preventing the adherence of H and other impurities (C, O) to the Nb surface as the furnace returned to room temperature (see chapter 6). Figure 4.17 shows the H/Nb ratios for these heat treatments, where the H levels are again seen to be up to a factor of 100 lower than a control sample Control 800C/3h,400C/20min N2,120C/6h 800/3h,400/20min 1000C/2h,800C/10min N2 Control 10.0 H/Nb Ratio C/2h, 800C/10min N2 800/3h,400/20min 800C/3h,400C/20min N2,120C/6h Depth (um) Fig SIMS H - /Nb - vs Depth data for high temperature heating/low temperature heated Nb samples vs control 117

143 The lowest H level is seen in the sample that was additionally baked at 120 o C for 6hrs. Table 4.4 shows the average H 2 pressure in the furnace during heat treatment. Similar to all the other heat treatments, higher the hydrogen pressure in the furnace, higher is the hydrogen level observed in the sample after the respective heat treatment. It can be thus inferred that a higher average hydrogen pressure in the furnace is leading to lesser hydrogen desorption during heat treatment. Table 4.4: Heat treatments and H 2 pressure in the furnace Heat Treatment (Furnace) H/Nb Ratio Average H 2 pressure (torr) 800 o C/3hrs, 400 o C/20min (Resistive) x o C/2hrs, 800 o C/10min N 2 (10-5 torr) (Induction) x o C/3hrs, 400 o C/20min N 2 (10-5 torr), 120 o C/6hrs (Resistive) x Discussion Table 4.5 shows the H - /Nb - ratios obtained from all the above heat treatments. The H/Nb ratios were used to characterize the H levels in the samples since the extremely fast diffusivity of H in Nb precluded the ability to obtain absolute concentration as discussed in section It can be inferred from the data that hydrogen is desorbed from Nb at high temperatures. As is seen from the solubility and diffusion data, above 600 o C and below 10 4 pa pressure, hydrogen diffusion in Nb is extremely fast while solubility decreases considerably. Thus, at these temperatures, H is expected to desorb from Nb provided the vacuum levels are met although the cool down from these temperatures might result is some re-absorption of hydrogen from the furnace atmosphere. 118

144 Table 4.5: Average H - /Nb - ratios obtained from SIMS data after various heat treatments Heat Treatment H - /Nb - (H - /Nb - None High Temperature Heat Treatment ) Control /(H - /Nb - ) HT 600 o C/10 hrs o C/6 hrs o C/6 hrs (new furnace) o C/3 hrs (new furnace) High Temperature Heat Treatment with Low Temperature Baking 600 o C/10 hrs, 120 o C/48 hrs o C/3 hrs, 120 o C/12 hrs o C/2 hrs, 120 o C/12 hrs o C/2 hrs, 120 o C/12 hrs o C/3 hrs; separate exp. : 120 o C/12 hrs (new furnace) High Temperature Heat Treatment with Lower Temperature Heating in Vacuum or N o C/3 hrs, 400 o C/20 min o C/3 hrs, 400 o C/20 min ( N 2 at 10-5 mbar), 120 o C/6hrs 1000 o C/2 hrs, 800 o C/10 min ( N 2 at 10-5 mbar) (new furnace)

145 It can be inferred from table 4.5 that a high temperature heat treatment of the sample desorbs less hydrogen from Nb compared to both high temperature heating followed by long term low temperature baking and high temperature heating followed by a lower temperature heating of the samples. Also, at all of the high temperature baking temperatures employed, the surface oxide layer of Nb dissociates thus facilitating desorption and diffusion of H from Nb, unlike in the control sample. The high temperature heat treatments drive out the hydrogen from Nb due to the decreased solubility of H in Nb at high temperatures and low pressures. As seen in the figure 4.15, higher H levels are seen for some heat treatments. This can be attributed to the average pressure of H 2 in the furnace during heat treatment since pressure is directly proportional to temperature even for a non ideal gas. The higher pressure may lead to less desorption of hydrogen from Nb, which is why more hydrogen is seen in Nb after 1000 o C and 1400 o C heat treatments as compared to 600 o C and 1200 o C (new heat treatments respectively. For heat treatments involving a high temperature heating followed by baking (120 o C) for long periods of time in the same experiment (fig.4.16), the H levels seen using SIMS are lower than those obtained using high temperature heating alone. No definitive explanation for this decrease can be provided at this point, since the fast diffusion of hydrogen should remove all the hydrogen at high temperatures and additional baking of the samples should not show any more hydrogen desorption. A possible explanation may be that the high temperature heat treated samples experience some re-absorption of hydrogen on cool down, while the additional low temperature baking allows formation some kind of an oxide or other barrier layer on the surface which does not reduced the rate hydrogen re-absorption on cool down after baking. The same may hold true for the samples heated to high temperatures followed by a lower temperature firing, since although the surface oxide dissociates at about 250 o C, some monolayers of oxygen may be present even at temperatures above 250 o C (sections and 4.3.1). Figure 4.18 shows the dependence of the the H - /Nb - ratios measured using SIMS after the various heat treatments on the average partial pressure of H 2 measured in the furnace during the various heat treatments. to. It is seen in the figure that the best fit of the data shows a 120

146 linear relationship between (p(h 2 )) 1/2 and H - /Nb -. Thus, it appears likely that a higher partial pressure of hydrogen leads to less desorption of hydrogen from Nb during heat treatment H2 pressure (p^1/2)/ (Torr)^1/ H-/Nb- ratio Fig. 4.18: Hydrogen average partial pressure in the furnace vs H - /Nb - ratio measured by SIMS after heat treatment Role of the Nb Surface Oxide Layer Bare metals easily react with hydrogen, but Nb owes its corrosion resistance to 5-10nm thick oxide layer on the surface, which can form in air at room temperature and and which passivates the surface to the absorption/desorption of hydrogen. 3 It has been shown in literature that this oxide layer is mainly composed of Nb 2 O 5-x and dissociates on heating Nb above 250 o C in high vacuum conditions. 41,42 The high temperature (> 250 o C) desorption of hydrogen seen in the above data is attributed to the dissociation of the surface oxide film of Nb at these temperatures. As the Nb is heated to higher temperatures (> 250 o C), e.g. high temperature heat treatments 600 o C, this oxide 121

147 layer is removed from the surface and the hydrogen is free to diffuse into or out of the material. The oxide reforms upon cooling of the sample to room temperature at which point hydrogen from the furnace atmosphere is not able to penetrate the Nb surface. In order to test this hypothesis, the following experiment was performed: 1. Since the Nb surface oxide can be removed using HF, a heat treated Nb sample (800 o C/3hrs, 400 o C/20min) was etched in full strength (49%) HF by immersing it in the acid for 15 min, essentially removing the surface oxide layer 2. The etched sample was then immediately rinsed with water. Since the the oxide layer had been removed, it was expected that hydrogen would be absorbed by Nb. 3. The rinsed sample was then dried and analyzed using SIMS. The expectation of the experiment was that hydrogen would be re-introduced into Nb and would show high levels. Fig shows the SIMS analysis of the heat treated sample both before and after etching. The increased H levels seen in the sample after etching thus corroborate the above mentioned hypothesis. 1E+07 1E+07 Counts (cts/sec) 1E+06 1E+05 1E+04 H Cts C o u n ts ( cts/sec) 1E+06 1E+05 1E+04 H Cts 1E Time (s) (a) (b) Fig. 4.19: SIMS analysis (cts vs time) of 800 o C/3hrs, 400 o C/20 min heat treated sample: (a) before and (b) after etching in HF for 15 min and rinsing with water, H levels seen to increase to control sample H levels in (b) 1E Time (s) 122

148 Further corroboration of these results were obtained by performing SIMS analysis on a heat treated sample (1200 o C/2hrs, 120 o C/12hrs) which had been rinsed with a high pressure water jet after heat treatment. Figure 4.20 shows the SIMS H - /Nb - vs depth data for such a sample. Since no etching of the Nb 2 O 5 was performed on the sample post heat treatment, the surface oxide layer on the Nb is intact, and it was expected that hydrogen would not penetrate into Nb. This hypothesis is verified by the data shown in figure Control H-/Nb- Ratio C/2hrs,120C/12 hrs + Water Rinse Depth (um) Fig.4.20: H - /Nb - levels remain low even after high pressure water rinsing of a heat treated sample owing to the surface oxide layer 123

149 4.4.5 Conclusion from Results It is seen from the above data that in order to remove hydrogen from Nb, high temperature heat treatments have to be performed, preferably above 400 o C. The fast diffusion of H in Nb poses a problem in terms of hydrogen quantification using SIMS, but relative hydrogen levels can be determined by comparison of the H - /Nb - ratios obtained from control and heat treated samples. The variations of these ratios observed between different heat treated samples can be attributed to the average H 2 pressure in the furnace during heat treatment. It is also observed that since the oxide layer acts as a passivation layer on the Nb surface that does not allow H to penetrate. The effect of this oxide on the surface hydrogen levels can be investigated using low energy SIMS analysis, which provides better depth resolution. 4.5 Low Energy SIMS Characterization of H in Nb Since the Nb surface oxide layer acts a barrier for the hydrogen absorption/desorption as seen from the previous section, it can be inferred that hydrogen does not rapidly diffuse in the oxide. Thus, it is important to perform low energy SIMS analysis of the Nb surface to characterize hydrogen in the oxide. Low energy SIMS analysis was performed to provide improved depth resolution on various non heat treated (control) and heat treated large grain Nb samples for all types of heat treatments. Chapter 3 describes the analysis conditions used for low energy SIMSanalysis Results and Discussion Long Term Low Temperature Baking (120 o C/48hrs) Figure 4.21 shows the H - /Nb - ratio vs depth (up to 100nm) comparison of the low temperature baked (120 o C/48hrs) sample to a non heat treated sample (control). As is seen from the figure, there is a small decrease in the H levels after the baking. This small decrease could be responsible for the reduction in Q- disease after baking a cavity at these temperatures for long periods of time, seen in literature

150 1000 Control H-/Nb- Ratio C/48hrs Depth (um) Fig. 4.21: H - /Nb - Ratio vs Depth for control vs 120 o C/48hrs baked large grain Nb samples. This small decrease in the H levels on the surface could be attributed to the surface oxide layer of Nb, which does not dissociate at these temperatures, but might be sufficiently altered to allow small amounts of hydrogen to desorb from Nb in high vacuum conditions (maintained in the furnace). Another important aspect seen in this data is that the 2-6nm depth region is devoid of any hydrogen detectable under these analysis conditions for both control and heat treated samples, confirming that the surface oxide layer acts as a passivation film and there is no hydrogen present in this layer, indicating that hydrogen does not diffuse in the oxide High Temperature Heat Treatment Figure 4.22 compares the H - /Nb - ratio (over a depth of up to 100nm) of the high temperature heat treated (1200 o C/6hrs) sample to a non heat treated sample (control). As is seen from the figure, there is a significant decrease in the H levels after this heat treatment, which also seen in the high energy data shown in section This decrease in the hydrogen levels can be 125

151 explained by the dissociation of the surface oxide layer above 250 o C and the subsequent desorption of H from the Nb surface, due to the decrease in solubility of hydrogen in Nb at low pressures and high temperatures H -/N b - ratio Control 1200C/6hrs Depth (um) Fig. 4.22: H - /Nb - Ratio vs Depth for control vs 1200 o C/6hrs heat treated large grain Nb samples. It is also seen from this data that is that the 2-6nm depth region is devoid of any detectable hydrogen for both control and heat treated samples, similar to the result seen for the low temperature baked Nb sample High Temperature Heating with Long term Low Temperature Baking Figure 4.23 shows the H - /Nb - ratio (over a depth of up to 100nm) comparison of the high temperature heat treated and long term low temperature baked (1200 o C-1400 o C, 120 o C/12hrs) samples to a non heat treated sample (control). As is seen from the figure, there is a significant decrease in the H levels after this heat treatment, which is also seen in the high energy data shown in section The difference in the H levels between the heat treated samples in the figure can be attributed to the H 2 pressures in the furnace (shown in 126

152 section ). Again, this decrease in the hydrogen levels vs a control sample can be explained by the dissociation of the surface oxide layer above 250 o C and the subsequent desorption of H from the Nb surface Control H-/Nb- Ratio C/3h,120C/12h 1200C/2h,120C/12h Depth (um) Fig. 4.23: H - /Nb - Ratio vs Depth for control and 1200 o C-1400 o C, 120 o C/12hrs heat treated large grain Nb samples. No H in the 2-6nm depth region In this data as well, 2-6nm depth region is devoid of any detectable hydrogen for both control and heat treated samples, similar to the result seen for the long term low temperature baked Nb sample (120 o C/48hrs) and the high temperature heat treated (1200 o C/6hrs) Nb sample (figure 4.25 and 4.26 respectively) High Temperature Heating with Lower Temperature Heating Figure 4.24 shows the H - /Nb - ratio (over a depth of up to 100nm) comparison of the high temperature heat treated and low temperature heated (800 o C/3hrs, 400 o C/15min at 10-5 mbar N 2, 120 o C/6hrs) sample to a non heat treated sample (control). As is seen from the figure, there is a significant decrease in the H levels after this heat treatment, which also seen in the high energy data shown in section Again, this decrease in the hydrogen levels can be 127

153 explained by the dissociation of the surface oxide layer above 250 o C and the subsequent desorption of H from the Nb surface. In this heat treatment as well, the surface region of Nb is seen to be devoid of hydrogen, but the profiles are not similar to the other heat treatments, possibly due to the nitrogen atmosphere heating, as opposed to high vacuum heating in the second step H-/Nb- Ratio Control 1 800C/3hrs,400C/15 min N2,120C/6hrs Depth (um) Fig. 4.24: H - /Nb - Ratio vs Depth for control vs 800 o C/3hrs, 400 o C/15min in 10-5 mbar N 2, 120 o C/6hrs heat treated large grain Nb samples. No H on the Nb surface Conclusion Low energy SIMS analysis further confirms desorption of hydrogen from the Nb surface after heat treatment. Significant decrease in the H levels is seen, especially after high temperature heat treatments. This analysis condition seems to complement high energy analysis data and the role of the surface oxide layer as a passivating film on the surface of large grain bulk Nb. Quantification of H in Nb is seen to be difficult for both low and high energy analysis conditions, but the low energy data provides improved characterization of the 128

154 first 2-6nm of the Nb surface, which shows that hydrogen is not present in the surface oxide layer and diffusion of hydrogen in this layer does not take place. This provides an opportunity to quantify hydrogen in Nb, using the surface oxide as the matrix and thus, the characterization of hydrogen in niobium oxide is important and can provide an estimate of the hydrogen concentration in Nb using SIMS. 4.6 SIMS Characterization of Hydrogen in Niobium Oxide It is seen from the previous data that H does not rapidly diffuse in the surface oxide of Nb, thus making it important to characterize H in Niobium Oxide. Since H diffuses extremely fast in Nb, an estimate of the concentration of H on the Nb surface was difficult due to the non availability of a quantification factor owing to the constant hydrogen signal obtained from the H in Nb implant standard and absence of a H implant peak in Nb. The immobility of H in niobium oxide provides an opportunity to quantify it in this oxide, and correlate this quantification to the actual niobium matrix using SIMS Experimental In order to improve SIMS characterization of H in Nb oxide and to provide sufficient oxide thickness to allow utilization of ion implantation of known impurities, a thicker oxide was needed. Non heat treated Nb samples were anodized to obtain a 150 nm oxide thickness on the surface. Prior studies have shown that the anodized layer is expected to have properties similar to that of the native oxide. 44,45 The anodized layer had sufficient thickness to alow low energy implantation to locate elements like H within the oxide. Note that both H and D can be implanted and the peak of the D can be placed at a different depth than that of the H. This capability facilitates the extraction of information from the layer and from the interface between layer and substrate. 46 The energy of the H implant was chosen to put the peak of the implant (projected range) in the oxide, and the energy for the D implant to put the peak at the interface between the oxide and Nb substrate. H was implanted at 5 kev to a dose of 1x10 16 atoms/cm 2 and D at 7 kev to a dose of 2x10 15 atoms/cm 2. Low energy SIMS analyses were performed on the anodized H and D implants using the analysis conditions mentioned in chapter 3 (section 3.5.2) 129

155 4.6.2 Results and Discussion Figure 4.25 shows the implant profile via SIMS analysis. The 150nm surface oxide is characterized by the O - secondary ion intensities. The implantation energy provides the implanted H to be well within the oxide. H shows an implant peak in the oxide, unlike analysis performed on implants of H in Nb (fig. 4.7), which showed a constant signal of hydrogen and the absence of an implant peak. The low levels of H in Niobium Oxide and the fact that hydrogen does not diffuse in the oxide facilitate this result. Oxide Substrate 1E+23 1E+08 Concentration (atom/cm3) 1E+22 1E+21 1E+20 1E+19 1E+18 1E+17 D Conc. O Counts H Conc. Nb Counts Depth (um) 1E+07 1E+06 1E+05 1E+04 1E+03 1E+02 Counts (cts/sec) Fig. 4.25: H and D implant peaks seen in niobium oxide; insignificant change in the Nb - intensities from in the oxide to in the substrate; D implanted to show peak at the interface It is also noted that the D was implanted to provide the D peak right at the niobium oxideniobium interface. The dotted line indicates the D profile shape if the D implanted into the Nb metal were not mobile. The H is clearly seen to be mobile in the Nb substrate. Well defined H and D implant shapes were obtained in the oxide, but only constant levels were 130

156 measured in the Nb. Clearly, the implanted D and H are more mobile in the metallic phase than in the oxide and this supports the hypothesis that this oxide film is a barrier to H absorption and desorption. It seems clear that niobium oxide provides a good barrier for movement of H. As mentioned earlier, a continuous oxide less than 10nm thick could be an effective barrier to H absorption or desorption. 1E+23 Concentration (atom/cm3) 1E+22 1E+21 1E+20 H Conc. 2E22 at./cm 3 1E Depth (um) Fig. 4.26: H concentration estimate for H in Nb in a control niobium sample, the atomic concentration is found to be 2x10 22 atoms/cm 3, which when divided by the Nb atomic density (5.44x10 22 atoms/cm 3 ) provides the concentration estimate to be 37 atomic % H in Nb For this ion implanted sample, the H peak in the niobium oxide can be used to calculate the H quantification factor (RSF) in niobium oxide. Since there is very little difference in the Nb matrix signal observed in the oxide compared with that in the substrate, the RSF obtained for H in the oxide can be used to estimate the concentration of H in Nb. 47 Using this method, the H concentration on the Nb surface was estimated for a non heat treated (control) Nb sample. Figure 4.26 shows the quantified profile for H in Nb for such a sample. The H concentration on the Nb surface is found to be 2x10 22 atoms/cm 3 which, when divided by the Nb atomic 131

157 density (5.44x10 22 atoms/cm 3 ) provides the estimate of H in a control Nb sample surface to be 37 at.% Conclusion H is less mobile in niobium oxide, as seen from the SIMS characterization of implanted H and D in an anodized non heat treated large grain Nb sample. This mobility decrease can be used to provide a quantification estimate of hydrogen in niobium using the oxide as the matrix, since the Nb - secondary ion intensities show very little change from in the oxide to in the niobium substrate. Using the RSF obtained from the oxide, the concentration estimate of H in Nb is found to be 37 atomic %, indicating that there is a high level of hydrogen on the niobium surface. 4.7 SIMS Characterization of Non-Nanoplished Nb Chapter 3-Section 3.1 shows the various steps for sample preparation for the large grain bulk Nb samples. Although this procedure is the standard cavity preparation procedure, the cavities were not nanopolished, while the samples were, for SIMS analysis. Thus, it is important to determine whether the H results obtained from nanopolished samples can be used to infer H levels in the cavity surfaces. For this purpose, SIMS analysis of nonnanopolished samples was performed and compared to their nanopolished counterparts. Figure 4.27 shows the H/Nb ratio of nanopolished and non-nanopolished control samples. The non-polished control sample went through the same sample preparation procedure (section 3.1) except the final nanopolishing step. SIMS analysis conditions were the same as for low energy SIMS characterization (Chp 3, section 3.5.2). The H levels in both the samples are observed to be similar, indicating that nanopolishing does not affect the surface H levels and nanopolished samples are indicative of the cavities tested. 132

158 Control (Nanopolished) Control (Non-Nanopolished) 1000 Control (Nanopolished) H/Nb Ratio 100 Control (Non- Nanopolished) Time (sec) Figure 4.27: H/Nb vs time (sec) for nanopolished and non-nanopolished control samples 4.8 TOF-SIMS Imaging of Hydrogen in Large Grain Nb Bicrystals It is believed that the grain boundaries of large grain Nb contribute to a reduction in the performance of the cavities, due to impurities that nucleate on the grain boundaries. 48 Grain boundaries are "weak" areas in a niobium surface that may segregate impurities. They may form "weak links" which are considered to be a prime candidate for causing the degradation in the performance. Other effects include reducing thermal conductivity, and enhancing the penetration of external magnetic fluxes. 48 It is thus important to study the grain boundaries of large grain Nb, to investigate H segregation at the grain boundaries, before and after heat treatment, since this segregation might lead to cavity performance deterioration. 133

159 4.8.1 Results and Discussion As mentioned in chapter 3, section 3.6.1, two types of bicrystal sample sets were prepared: bicrystals having different crystallographic orientation combinations, where the orientation of the individual crystals were determined using EBSD, and bicrystals having the same crystallographic orientation combination Bicrystals having different crystallographic orientation combinations One of the samples in this set was heat treated for 800 o C/3hrs, 120 o C/24hrs using the Elnik resistive heating furnace (section 3.6.1). Section also provides EBSD data for the control and heat treated bicrystals having different crystallographic orientation combinations. Notably, the upper crystal had a (001) orientation for both the control and heat treated bicrystals, while the lower crystal orientation was (213) and (215) for control and heat treated samples respectively. H-/Nb- intensity Lower crystal H-/Nb- control sample 2.3 A/s Upper crystal (001) 6.65 A/s Depth (um) H - /N b - in t e n s it y H-/Nb- for 800C/3hrs,120C/24hrs HT sample Upper Crystal (001) 4.75 A/s Lower crystal 6.5 A/s Depth (um) Fig. 4.28: Dynamic SIMS results for control (left) and heat treated (800 o C/3hrs, 120 o C/24hrs) (right) bicrystal samples. The sputtering rate differences may be attributed to the different crystal orientations. As seen in previous data, lower H levels observed in the heat treated sample 134

160 Dynamic SIMS was used to characterize the surface of the near grain boundary regions for both samples. The analysis conditions used were the same as shown in chpt 3, section H levels were significantly different both within the bicrystal samples and between heat treated and control samples. Fig shows the H - /Nb - intensity extracted from the SIMS measurements. As seen in the figure, H - /Nb - intensity varies within the same sample for both samples. This may be attributed to the difference in the crystallographic orientations in the two crystals, which further causes a difference in the sputtering rates. Notably, the sputtering rate difference between the two crystals of the control sample is almost a factor of three. Also, as seen in previous sections, the control sample tends to harbour more H on the surface than the heat treated sample. The upper crystal mentioned here in the figure has a (001) orientation, while the lower crystal has (213) and (215) orientation for control and heat treated samples respectively. For imaging across the grain boundaries, TOF-SIMS imaging was conducted to a total depth of 0.25μm. Fig shows the ion images for the control (a) and the heat treated (b) samples. No segregation of hydrogen was seen for both control and heat treated samples. Hydrogen is also seen to be concentrated near the grain boundary region for the heat treated sample upper crystal. Grain Boundary (001) (001) (213) (215) (a) (b) Fig. 4.29: TOF SIMS ion images for control (a) and heat treated (800 o C/3hrs, 120 o C/24hrs) (b) bicrystal sample grain boundaries. No segregation of H seen at the grain boundary for both samples 135

161 The intensity difference across the interface for both samples could be an effect of the sputtering rate difference owing to the different crystal orientations. The important result obtained from this analysis was that hydrogen is not seen to segregate on the large grain Nb bicrystal grain boundary both before and after heat treatment. This result corroborates the findings of Rudd et. al. 49, which shows that hydrogen does not permeate into Nb along the grain boundaries, but through the crystal lattice Bicrystals having the same crystallographic orientation combination In order to account for the effect of crystal orientation on the secondary ion intensities and obtain a better understanding of the heat treatment effect on the grain boundaries of large grain Nb bicrystal Nb samples, a separate set of samples were characterized using TOF- SIMS imaging with the same conditions mentioned in section However, to eliminate any orientation effects, a large sample was sectioned from the same grain boundary area and cut into half. One part of this large sample was heat treated to 1200 o C/6hrs in the induction furnace after BCP (1:1:2, 7μm removal) while the other part was not heat treated after the same BCP etching (control). (hkl) H - Grain Boundary (hkl) H - (h k l ) (h k l ) (a) (b) Fig. 4.30: TOF SIMS ion images for control (a) and heat treated (1200 o C/6hrs) (b) bicrystal sample grain boundaries with similar crystal orientation of the upper crystal for both samples and lower crystal for both samples. 136

162 Fig shows the TOF SIMS images of the control (a) and heat treated (b) Nb samples. The upper crystal in both the samples is similarly oriented, while the lower crystal also has the same orientation in both the samples. Again, H is not seen to be segregated at the grain boundary, similar to the results obtained after heating the bicrystal in the resistive furnace (previous section) Conclusion TOF SIMS analysis of the grain boundary region of large grain bicrystal Nb shows that hydrogen is not segregated at the grain boundaries for both non heat treated and heat treated samples. Initial analysis suggested that the difference in the H - intensities may be related to the different crystallographic orientation of the crystals in the samples. However, subsequent analysis performed on the grain boundary area of similarly oriented Nb bicrystals show similar results. 4.9 Effect of Hydrogen on SIMS Nb - Secondary Ion Intensities SIMS secondary ion intensities should remain the same for the matrix species in different samples to obtain accurate quantification estimate of the impurity element in the matrix, since the impurity species secondary ion intensities are normalized to the matrix intensities to obtain the RSF for that impurity in the matrix (in this case, H in Nb). However, for SIMS characterization of impurity elements in Nb, it was seen that the Nb - intensities changed between the high temperature heat treated and control samples. Figure 4.31 shows a comparison of the Nb - secondary ion intensities for a heat treated sample (HT = 800 o C/3hrs, 400 o C/20min) with a control sample. It is seen in the figure that the Nb - intensities are a factor of three higher for the control sample as compared to the heat treated sample. 137

163 1.E+06 Nb Cts vs Depth 1.E+05 Cts/sec 1.E+04 Nb Cts Control Nb Cts (HT) 1.E Depth(um) Fig. 4.31: Nb - intensities a factor of 3 higher for a control sample vs a heat treated sample (HT= 800 o C/3hrs, 400 o /20min) This change in intensity was seen in all high temperature heat treatments. Figure 4.32 shows the I - (Control)/I - (Heat Treated) intensity ratios for both Nb and H of all heat treatments performed up to 1400 o C. It can be seen from this figure that the ratio is about the same for all higher temperature heat treatments, while it is similar to the control sample for only the 120 o C/48hrs baked sample, which showed similar hydrogen levels as the control sample respectively. It can thus be inferred that for high temperature heat treated large grain single crystal Nb, which shows a large decrease in hydrogen levels after heat treatment, the Nb - intensities are also lower than in the non heat treated Nb, while the Nb - intensities do not show a change between samples having similar amounts of hydrogen, as detected by SIMS. Thus, it can be concluded that hydrogen is the main cause of change in the Nb - intensities between the high temperature heat treated and control Nb samples. 138

164 I- (C on trol)/ I- (H eat Treated) Temp (deg.c) H Nb Fig.4.32: Both Nb and H follow similar trends in the control to heat treatment intensity ratios of secondary ions vs heat treatment,. H intensity ratios have been scaled for comparison This change in the Nb - intensities provided some difficulty in obtaining accurate quantification factors (RSFs) for impurity profiling in heat treated samples, since the RSFs obtained from non heat treated, ion implanted standards, could not be used for impurity quantification in heat treated Nb. To obtain precise quantification factors, a heat treated large grain (single crystal) Nb sample (1000 o C/2hrs, 120 o C/12hrs) was implanted with one of the impurity elements (O), with the same dose and energy as for the non heat treated standard (5e15 at./cm 2 and 180 kev respectively) and SIMS analysis was conducted using the CAMECA-IMS 6f, with the same analysis conditions mentioned in chapter 3, which were also the analysis conditions for depth profiling O in Nb (chp 5). Fig shows a comparison between the heat treated and non neat treated O in Nb ion implant standards. It was observed that the RSF for the heat treated standard was a factor of 3 139

165 lower than that of the non heat treated standard (control), which was equivalent to the decrease in the Nb - intensities. Thus, this new RSF was used for the quantification of O in Nb, and a correction factor (factor of 3 decrease) was used for other impurity elements (C, N). Concentration (atoms/cm^3) 1E+22 1E+21 1E+20 1E+19 RSF (Control): 4.5e19 at./cm 3 RSF (HT): 1.5e19 at./cm 3 O - (Control) O - (1000 o C/2hrs) 120 o C/12hrs ) 1E Depth (um) Fig. 4.33: HT= 1000 o C/2hrs,120 o C/12hrs; Comparison of the non heat treated (control) and HT oxygen implants; O RSF seen to decrease by a factor of 3 for the heat treated implant It was only possible to speculate with respect to the reason for the observed change in the Nb - intensity. 53 Figure 4.34 shows a comparison of the mass spectra of one of the heat treated samples (800 o C/3hrs, 400 o C/20min) to a control sample. As expected, intense hydride peaks are seen in the control sample. The region below the peaks in the control sample consists of ions having fractional masses, which are believed to be metastable ions

166 From literature, it is known that Nb is used as a catalyst in MgH 2 to facilitate the desorption of H XRD experiments have shown that Nb forms a metastable phase (NbH 0.6 ) with hydrogen in MgH 2, which assists in faster diffusion of hydrogen through the lattice. This phase exists in the α + β region of the Nb-H phase diagram, which is the same region observed by the hydrogen concentration estimate of a control Nb sample, using SIMS (section 4.6.2). It can be speculated that this phase might be responsible for the metastable ion contribution to the mass spectra of the control sample in fig.4.40 and during depth profiling, these metastable ions might be disintegrating in the mass spectrometer, causing the increase in the Nb - intensities. Since heat treated Nb has very low hydrogen levels, compared to the control sample, the metastable phase might not exist in heat treated Nb, as shown in fig Counts/sec 1E+07 1E+06 1E+05 1E+04 1E+03 - NbH - 2 NbH - Nb NbH 3 - NbH 4 - NbH 5 - Counts/sec) 1E+07 1E+06 1E+05 1E+04 1E+03 Nb - NbH - 1E+02 1E+02 NbH 2-1E+01 Metastables 1E Mass (a.m.u.) 1E+01 1E Mass (a.m.u.) Figure 4.34: Mass spectra of Control (left) and 800 o C/3hrs, 400 o C/20min heat treated Nb (right) samples Thus, during dynamic SIMS analysis, where sputtering produces secondary negative and positive ions along with electrons and neutrals, the extracted negative metastable ions from the metastable phase NbH 0.6, may be disintegrating in the mass spectrometer according to the reactions portrayed in figure Note that these reactions have been depicted for negative secondary ions only. 141

167 Negatives only: Nb - + H (neutral) β-nbh NbH 0.6 NbH H - + Nb (neutral) Fig. 4.35: Possible secondary ion reactions relating to increased Nb - intensities in SIMS anlysis of a control Nb sample 142

168 References 1. B. Bonin et al : Proc. Of 5 th Workshop on RF Superconductivity-1991, p210 (1991) 2. B. Aune et al : Proc. Of Linear Acc. Conference -1990, p253 (1991) 3. C. Antoine et al : AIP Conf. Proc. 671, p176 (2002) 4. E. Kako et al : Proc. Of 9th Workshop on RF Superconductivity-1999, p179 (1999) 5. P. Pascal : Nouv. Tr. Chim. Min., Paris, Masson (1958) 6. G. Khaldiv et al : Russ. Chem. Rev., 56(7), p605 (1987) 7. C. Antoine et al : Proc. Of 5 th Workshop on RF Superconductivity-1991, p616 (1991) 8. B. Roux et al : Vacuum, 46 (7), p929 (1995) 9. J. Rodriguez et al : Script. Metall., 17, p159 (1983) 10. W. Staiger et al : Zeits. Phys. Chem., 164, p991 (1989) 11. J. Smith: Bull. All. Ph. Diag., 4 (1), p39 (1983) 12. S. Isagawa: J. Appl. Phys, 51(8), p4460 (1980) 13. W. Albrecht et al : J. Electrochem. Soc., 106, p981 (1959) 14. G. Rauch et al: J. Less-Common Met., 8, p99 (1965) 15. J. Welter et. al: Z. Phys. B, 27, p227 (1977). 16. H. Horn et al: J. Am. Chern. Soc. 69, p2762 (1947). 17. C. Wiseman: J. Appl. Phys. 37, p3599 (1966). 18. K. Ralls et al: J. Less-Common Metals, 2, p127 (1966). 143

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170 36. G. Ciovati et al: Phys. Rev. Special Topics - Accelerators and Beams, 13, p (2010) 37. U. Kobler et al: J. Less-Common Met., 84, p225 (1982). 38. F. Stevie et al: SIMS- A practical handbook for depth profiling and bulk impurity analysis, Wiley Interscience Publications (1989) 39. H. Behrens et al: Gases and Carbon in Metals: Thermodynamics, Kinetics, and Properties. Pt. VIII, Group VA Metals (2), Niobium, Fachinformationszentrum Energie Physik Mathematik GmbH, Karlsruhe, Germany (1976) 40. G. Khaldeev et al: Russ. Chem. Rev., 56(7), p605 (1987) 41. R. Brewer et al : Phys. Rev. Lett., 62(15), p1760 (1989) 42. F. Palmer et al: Journ. Electr. Spec. Rel. Phen., 52, p511 (1990) 43. T. Ogutarni: Metall. Trans., 2, p3035 (1971) 44. A. Dacca' et al: App. Surf. Sci. 126, p219 (1998) 45. M. Grundner et al: J. Appl. Phys. 51, p397 (1980) 46. F. Stevie et al: J. Vac. Sci. Technol., B18, p483 (2000) 47. F. Stevie et al: SIMS- A practical handbook for depth profiling and bulk impurity analysis, Wiley Interscience Publications (1989) 48. R. Graham: AIP Conf. Proc. 927, p21 (2006) 49. D. Rudd et al.: Journ. Phys. Chem., 66, p351 (1962) 50. B. Wunderlich: Thermal Analysis, Academic Press, New York, p137 (1999) 145

171 51. E. Pungor: A Practical Guide to Instrumental Analysis, Boca Raton, Florida, p181 (1995). 52. J. Bulak et al: Journ. Phase Equ. and Diff., 28, 5, p422 (2007) 53. J. Veleckis et. al: J. Phys. Chem, 73 (3), p683 (1969) 54. P. Williams and D.Griffis: private communication 55. M. Rose: Mass Spectrometry for Chemists and Biochemists; Cambridge University Press (1996) 56. J. Pelletier et. al: Phys. review B, 63, p52103 (2000) 57. J. Volk et. al: Hydrogen in Metals 1, Springer-Verlag, Berlin (1976) 58. J. Sussmann: Ann. Phys., 6, p135 (1971) 59. A. Stoneham: Ber. Bunsenges. Physik. Chem. 76, p816 (1972) 60. E. Maksimov et.al: Soy. Phys. Uspekhi, 18, p481 (1976) 61. A. Stoneham: Collective Phenomena 2, p9 (1975) 62. C. Gupta et. al: Extractive Metallurgy of Niobium, CRC Press (1994) 146

172 5. Oxygen Besides hydrogen, Nb tends to harbor other interstitial impurities which may be detrimental to the RF cavity performance. One of these interstitial impurities is oxygen. It is known that Nb has a tendency to form a surface oxide layer, even in air, about 5-10nm thick which acts as a passivating film for the absorption/desorption of hydrogen (see chapter 4). Also, in terms of dissolved impurities, oxygen is also important due to the high affinity of Nb for oxygen above 200 C. 1 Thus, it is important to characterize the surface of Nb for oxygen and understand the source of this contamination and the effect it has on the RF cavity performance. 5.1 Source of Oxygen Contamination in Nb Cavities A major source of the surface oxygen contamination on Nb is the natural surface oxide layer. Although beneficial as a passivating film through its ability to block the absorption of H, there has been several speculations that this surface oxide is an important contributor to residual resistance in Nb SRF cavities. 2-5 A source of dissolved oxygen contamination is the fabrication of the cavity. Steps taken during formation of the cavity structures can lead to interstitial oxygen being dissolved into the bulk. For example, care has to be taken while trimming the formed Nb cavity half shell due to the high reactivity of oxygen to Nb during lathe machine trimming. 1 While electron beam welding the cavity half shells together, total pressures above 10-5 torr in the chamber can lead to oxygen contamination in the cavity. 1,6 Although discussion of the details of the sources of oxygen contamination are not as prevalent in literature as hydrogen, it is generally accepted that the surface oxide layer, fabrication steps and water rinsing of the cavities after etching treatments 7 appear to have a role in interstitial oxygen contamination in Nb. 5.2 Effect of Oxygen on the properties of SRF Nb Like hydrogen, oxygen affects the residual resistance of superconducting Nb in the form of residual losses. 1 Oxygen is also known to decrease the critical temperature (T c ) of Nb which directly affects the cavity performance. 8 Thus, it is important to review these effects in detail. 147

173 5.2.1 Residual Losses due to Oxides It has been suggested that the natural oxide layer of niobium, typically 5-10 nm thick, is an important contributor to residual resistance. 2-5 Literature suggests that the outermost oxide on Nb, irrespective of sample preparation, is Nb 2 O The loss tangent of Nb 2 O 5 has been estimated by Kniesal et.al. to be tan(δ) < 10-6 from a Nb cavity with a 40nm oxide layer produced by electrolytic anodization. 13 Systematic experiments by Palmer have shown that the lower limit of residual losses via oxide contribution is about 1-2nΩ. 14 In this study by Palmer, Nb cavities were fired at 1400 C to dissolve the oxide layer into the bulk, which was confirmed by Auger studies on similarly treated Nb samples. RF tests on the oxide free cavities showed a residual loss of 5-10nΩ, which is comparable to a Nb cavity with its oxide layer intact. When this oxide layer was regrown by controlled exposure to pure oxygen (0.1 torr for 2-16 hrs), the increase in the residual loss was about 1-2nΩ Oxide as a Source for Magnetic Flux Pinning Sites A well understood and controllable source of residual loss is trapped DC magnetic flux from insufficient shielding from the earth s magnetic field, or other DC magnetic fields in the vicinity of the cavity. To obtain the highest Q o, a superconducting cavity must be well shielded from the earth s field. Under ideal conditions, if the external magnetic field is less than the critical magnetic field for Nb (200mT), the DC flux will be expelled from the bulk of a superconductor due to the Meissner effect. However, if there are lattice defects or other inhomogeneities in the material, the flux lines may be pinned and trapped within the material contributing to a performance loss. 1 It is suspected that the oxide layer on the surface of Nb serves as a source of flux pinning sites. 15 Experimental studies show that when Nb cavities are cooled in the presence of a small DC magnetic field, all the flux within the volume of the cavity is trapped Effect on Critical Temperature (T c ) and Resistivity (10 K) DeSorbo studied the effect of oxygen concentration in the bulk on the critical temperature (T c ) of superconducting Nb. 16 The findings show a linear relationship between the oxygen atomic % vs T c of Nb, with the T c decreasing by a value of 0.93 K/atomic% of oxygen. Table 148

174 5.1 shows the results obtained from these measurements. As seen in the table, T c decreases linearly reaching a minimum value of 5.83 K for 3.84% atomic O, which is near the solubility limit of O in Nb at 1000 C. 17 It was however seen that as the concentration of oxygen was increased to above the solubility limit of O, the T c increased, reaching 9 K for 6.43% atomic O. Figure 5.1 shows this trend for decreasing T c (characterized by the ratio of pure Nb T c (T c0 ) to T c of oxygen loaded Nb) upto 4% atomic O in Nb. 16 Table 5.1 also shows the resistivity dependence on O concentration in Nb at 10 K. The resistivity was seen to increase monotonically with oxygen concentration. 16 More information on the apparatus and technique used in this study can be found in ref.16. Table 5.1: Decreasing critical temperature with increasing O concentration in Nb. Resistivity increase in Nb with increase in the O concentration 16 Material Critical Temp. (T c ) K Resisitvity (Ω-cm) Nb x 10-8 Nb % at. O Nb % at. O x 10-6 Nb + 1.8% at. O x 10-6 Nb + 2.6% at. O x 10-6 Nb % at. O x 10-6 Nb % at. O Figure 5.1: Dependence of T c on solute (O) concentration; O seen to decrease T c linearly with increasing concentration

175 Both the decrease in T c and resistivity (10 K) can detrimentally affect cavity performance, thus showing that increased O levels in Nb can negatively affect the Q o of the RF cavities. Besides these effects, O being an interstitial impurity, also reduces the superconducting energy gap (see chapter 1), thus making it easier for the Cooper pairs to dissociate to normal conducting electrons on the surface and decrease the surface superconductivity Solubility and Diffusion of Oxygen in Nb Solubility Figure 5.2 shows the phase diagram of the Nb-O system, up to 80% atomic O concentration. 18 Figure 5.2: Nb-O phase diagram

176 The diagram shows the maximum solid solubility limit of O in Nb of about 8% atomic at a temperature of 1860 C. At room temperature however, oxygen is less soluble in Nb. 18 There have been many studies conducted on the solid solubility of O in Nb, all showing a limit of about 4-8% atomic O in Nb near temperatures of 1500 C. Based on sintered alloys heated in the temperature range 1600 C-1700 C, Brauer 19 concluded that the solid-solubility limit of oxygen in niobium lay below 4.7 at.%, whereas Seybolt 20, using a Sievert s apparatus to meter the amount of oxygen added, found values ranging between 1.4 and 5.5 at.% oxygen in the temperature range 775 C-1100 C. According to Elliot 21, the solid-solubility of oxygen rises from 1.4 at.% at 500 C to about 3.9 at.% at 1800 C, whereas Bryant 22, using hardness as a criterion of oxygen content, found that the concentration ranged from 0.7 at.% oxygen at 700 C to 5.5 at.% at 1550 C. Finally, Gebhardta et. al. 23 using a combination of X-ray diffraction, micro-hardness, electrical resistivity measurements and micrographic studies found that the concentration varied from 1.1at.% oxygen at 750 C to 5.5 at.% oxygen at 1540 C. Bryant 22 studied the oxygen solid solubility in Nb at a temperature range of 700 C-1550 C using an oxygen absorption apparatus, with a pressure of 1atm O 2. The findings showed that oxygen solubility in Nb followed an Arrhenius plot and increased with temperature in this temperature range, as shown in the phase diagram in figure 5.2: -Ln (N) = 8600/(RT) (5.1) where N is the atomic fraction of O in Nb. Figure 5.3 shows the Arrhenius relationship of O solubility in Nb by Bryant. It can thus be inferred from literature that the solid solubility of O in Nb increases with temperature in air from 700 C-1550 C, while the room temperature solid solubility is low. 151

177 Figure 5.3: Oxygen solubility in Nb follows an Arrhenius plot Oxidation of Nb In addition to oxygen being taken into interstitial solid solution within the b.c.c. niobium structure, numerous oxides and sub-oxides can be formed depending on the temperature and the oxygen pressure. 22 Abruzov et. al. 24 studied the oxidation of Nb in air. The oxidation was monitored by continuous weighing in the range C by steps of 50 C, and the oxidation products were examined via X-ray diffraction, which showed that the oxide formed at C in 5-9 hrs was orthorhombic and termed α-nb 2 O 5, while the oxide that formed at C also showed an orthorhombic structure, but with different lattice parameters (termed β-nb 2 O 5 ) The results indicated that there are two stages in the oxidation of Nb at any temperature. The first (slow) stage is due mainly to solution of oxygen in the metal; the second (faster) stage is the formation of oxide. It was considered that oxygen dissolves in the metal in both stages but that its relative importance in the oxidation alters

178 The following oxidation process is derived from data obtained from various sources in the literature: Finely divided α-nb 2 O 5 containing excess oxygen 25 is formed at the surface of the saturated solution of oxygen in niobium at C. The ratio of volumes of oxide to metal is large, so the oxide formed at this temperature is gradually shed, which allows the oxygen to reach the surface; the unprotected metal is then oxidized. A layer of NbO starts to form under the α-nb 2 O 5 at C 26, the oxygen reaching the surface gives rise to NbO as well as α-nb 2 O 5, and the layer adheres firmly to the surface. The oxygen concentration in the α-nb 2 O 5 therefore falls, and the compound changes to stoichiometric composition. The volume ratio for Nb 2 O 5 to NbO is smaller than the volume ratio for Nb 2 O 5 to Nb, so the Nb 2 O 5 formed near the surface of the NbO adheres firmly to the metal. The oxygen then has to diffuse through the Nb 2 O 5, the NbO, and the solid solution. This implies a reduced oxidation rate in this temperature range (Figure 5.4). The thickness of the NbO layer is increased at C and it tends to have vacant cation and anion sites 27,28, so oxygen ions diffuse through it to the metal, and niobium ions diffuse to the NbO-Nb 2 O 5 interface, which can give rise to Nb 2 O 5 deficient in oxygen. 25 This tends to facilitate the oxidation, as is observed in Figure 5.4. Formation of the high-temperature β-nb 2 O 5 affects the oxidation at C and the rate falls (Figure 5.4). This phase appears after 5-7 hrs at 880 C and a characteristic feature is that whiskers (monocrystals) of β-nb 2 O 5 grow on the surface of the oxide. 29 These appear to be produced by evaporation of the high-temperature phase, followed by deposition on defect sites; the vapor of β-nb 2 O 5 can condense on micropores and cracks, which tends to seal these and thus help to protect the metal. This healing process accelerates between o C and the oxidation rate falls (Figure 4.4), while the lattice of β-nb 2 O 5 becomes more nearly perfect, which should also hinder the diffusion of oxygen and hence reduce the oxidation rate

179 Figure 5.4: Kinetics of oxidation of Nb, 3hr periods; q (mg/cm 2 ) is the oxidation characteristic (ratio of the gain in weight to the initial surface area) 24 The above oxidation process arises from heating in air. However, for Nb cavities and Nb samples baked ( o C) in vacuum, surface analysis studies 5,30-32 have shown that a decomposition of the natural oxide Nb 2 O 5 layer into metallic suboxides NbO and NbO 2 and oxygen diffusion towards the bulk are initiated Pressure-Temperature relationships in the Nb-O system Figure 5.5 shows the pressure-temperature diagram for the Nb-O system determined by Taylor et. al. 33, as log 10 (po 2 ) vs 1/T. The NbO and NbO 2 curves mark the boundaries for the various phases in the system. Any phases under the conditions of temperature and pressure to the left of the NbO curve consist essentially of α-nb solid solution while those immediately to the right of the NbO line are found to be two-phase, consisting of α-nb+nbo. Beyond the curve for NbO 2, three phases, α-nb+nbo+nbo 2, exist. 154

180 Figure 5.5: Pressure-Temperature diagram for Nb-O system 33 Taylor et.al 33 estimated the relationship between pressure and temperature for the Nb-O system using thermodynamic calculations from Worrell s data. 34 Worrell had estimated the free energy of formation ΔF of NbO 2 by measuring the CO equilibrium pressure from the reaction: Nb 2 O 5 (solid) + C (graphite) NbO 2 (solid) +CO (gas); using solid pellets of NbO 2 +Nb 2 O 5 +C in a CO atmosphere at o K, and had come up with the relation:- ΔF = -185, *T cal/mole (5.2) where T is the temperature in K. This value was then used by Taylor et.al 33 to calculate the equilibrium dissociation pressure of NbO 2 for the reaction: NbO 2 Nb+ O 2, using the relation: Log 10 p(o 2 ) = ΔF/RT = -40,340/T (5.3) 155

181 This calculation suggested that in order to outgas oxygen from NbO 2, a pressure of 1.48 x Torr of oxygen was required at 1500 C and judging by figure 5.5, the pressure required to remove gaseous oxygen from the NbO would be even lower. 33 Thus, Taylor et al. postulated that only at oxygen pressures substantially below torr and T= 1500 C is it possible to remove gaseous oxygen from the interstitial α-nb primary solid solution, and the computed pressure corresponds to the pressure of oxygen, at equilibrium, when the primary α-nb solid solution is saturated and the monoxide phase NbO is about to form. Such a process should be reversible, indicating that with solid niobium, there is no way to remove all traces of oxygen from the metal using conventional vacuum of Torr and unlike hydrogen, Nb can oxidize at high temperatures, even at pressures below torr. 33 Thus, as shown in figure 5.5, any removal of oxygen from the α-nb solid solution for pressure-temperature values (clearly higher than Torr) along the NbO curve, would be in the form of NbO vapor rather than gaseous O Diffusion of Oxygen in Nb Diffusion of oxygen in Nb is found to be many orders of magnitude slower than hydrogen in Nb. Ang 35 studied the diffusion of oxygen in Nb using a torsional pendulum technique. Figure 5.6 shows the diffusion coefficient results from these findings. Like hydrogen, the diffusion rate of oxygen in Nb increases with temperature, given by the relation: D = 1.5 x 10-2 exp (-27600/RT)..(5.4), 300< T< 2300 K The activation energy was found to be 27,600 cal/mole 35 From figure 5.6 and figure 3.5 (section 3.3.2), a comparison of the diffusion coefficients of oxygen in Nb and hydrogen in Nb can be made. Oxygen in Nb: 5 x cm 2 /sec (300 K) Hydrogen in Nb: 2 x 10-5 cm 2 /sec (300 K) It is seen that hydrogen diffuses times faster in Nb than oxygen at room temperature. 156

182 Figure 5.6: Diffusion Coeff. vs Temp ( o K) for oxygen in Nb 35 The one dimensional approximate diffusion length of oxygen is given by 47 : L= 2*(D*t) 1/2 cm..(5.5) Where D is the diffusion coefficient of oxygen in Nb at a given temperature (from above figure) and t is the time in seconds. 5.4 Oxygen Pollution Model A recent model which attempts to explain the deterioration of cavity performance due to the surface oxide layer is the oxygen pollution model. 4 According to this model, cavity performance degradation is caused by magnetic vortices being pushed in the Nb surface at a magnetic field value lower than the critical magnetic field, due to regions of high oxygen concentration, such as grain boundaries, which result in a local reduction of the surface 157

183 barrier of Nb, from the value of about 200mT (critical magnetic field) down to about 100 mt. 4 A refinement of this model by Ciovati 3 included the effects of both oxygen diffusion and oxide decomposition in determining the oxygen concentration at the metal/oxide interface, after low-temperature baking of the cavity. The calculation showed a minimum oxygen concentration in the temperature range of 120 C 150 C bakeout for 48 hours on the Nb surface, in good agreement with the baking parameters which gave the highest improvement of the cavity performance. 3 However, some contradictions of this model have been seen via experimental results: (i) Cavity performance degradation was not re-established in a previously baked cavity, after additional baking at 120 C/48 hrs in 1 atmosphere of pure oxygen although measurements on samples by SIMS confirmed the presence of a higher oxygen concentration at the Nb surface after this baking. 36 (ii) Cavity performance did not improve after baking a cavity in situ at 400 C/2 hrs, although the interstitial oxygen should have diffused deeper into the bulk by baking at such high temperatures, according to the diffusion equation (5.5). 37 (iii) Measurements of the subsurface oxygen profile by diffuse x-ray scattering on a single crystal Nb sample revealed a diffusion length of only about 2 nm (figure 5.7) after baking at 145 C/5hrs, in spite of the 40 nm diffusion length, as calculated using the diffusion equation (5.5). 38 Figure 5.7: X-Ray diffuse scattering results: only 2nm diffusion of oxygen seen after 145 C/48hrs baking

184 All these observations point to the importance of a systematic study of the effect of various heat treatments on the oxygen concentration on the Nb surface, via surface analytical methods. 5.5 SIMS Characterization of Oxygen in Nb The detrimental effects of both the surface oxide layer and dissolved oxygen in Nb make the characterization of this impurity on the Nb surface important. Various contradictions seen in the oxygen pollution model (section 5.4) make it necessary to perform a systematic study on the effects of both low temperature baking and high temperature heat treatments on the surface oxygen concentration of Nb. SIMS can be a powerful technique to study the surface of Nb before and after these heat treatments, owing to the excellent depth resolution and detection limits achievable by this technique. SIMS analysis was performed using a CAMECA IMS-6f on various heat treated bulk Nb samples and compared to a control (non heat treated) sample. Typical analysis conditions are mentioned in chapter 3, section Heat Treatments As mentioned in chapter 3, the following heat treatments were performed on the samples: 1. Low temperature baking (120 o C/48hrs) 2. High Temperature Heat Treatments (600 o C-1400 o C/3-10h) 3. High Temperature Heating followed by Long Term Low Temperature Baking (600 o C-1400 o C/3-10h; 120 o C/12-48hrs) 4. High Temperature Heating followed by Lower Temperature Heating (800 o C o C/2-3h; 400 o C-800 o C/10-20min) in vacuum or 10-5 torr N 2 For different types of heat treatments, the idea was to dissolve more oxygen from the near surface regions to the bulk, hence cleaning the first 40nm of the niobium surface. Since the diffusion of oxygen is faster at higher temperatures, it was expected that higher temperatures would diffuse the oxygen more into the bulk Results and Discussion Figure 5.8 shows the concentration profile of the O implanted Nb standard. O was implanted into the Nb after heat treatment to 1000 C/2hrs, 120 C/24hrs. The O implant peak is clearly 159

185 seen in the figure, unlike the hydrogen in Nb implant. This indicates that O does not have a high diffusion coefficient in Nb and the Gaussian shape shows the SIMS analysis does not affect the O distribution. 1E+22 Concentration (atoms/cm^3) 1E+21 1E+20 1E+19 O Conc. 1E Depth (um) Figure 5.8: SIMS Depth profile of O in Nb implant RSF =1.5e19 at/cm 3, extracted from the implant peak for quantification Before the SIMS results can be discussed it is important to ascertain the background contribution to oxygen levels measured by SIMS. Raster reduction (section 3.4.6) was performed on some of the control and heat treated samples to check for background contribution. 160

186 Figure 5.9(a) shows the raster reduction profile for O - and Nb - for a control sample, where the raster was reduced to 60 x 60 μm 2 area from 120 x 120μm 2 area while maintaining a primary ion beam current of 20nA. As seen in the figure, the oxygen secondary ion intensities do not increase proportionally to the matrix (Nb - ) secondary ion intensities, indicating some background contribution to the oxygen levels. Thus, it may be inferred that oxygen is near detection limit of the instrument and the actual amount of oxygen might be even lower in a control sample. This was observed for all the control samples analyzed in this study. 1E+06 Counts (cts/sec) 1E+05 1E+04 1E x 120 μm nA Nb- O- 60 x 60 μm nA. 1E Time (s) Figure 5.9 (a): Raster reduction for O - and Nb - for a control sample, some background contribution seen indicating that oxygen is near detection limit of the instrument. 161

187 Counts (cts/sec) 1E+06 1E+05 1E+04 1E x 120 μm nA 60 x 60 μm 2 20nA Nb- O- 1E Time (s) Figure 5.9 (b): Raster reduction for O - and Nb - for a heat treated sample, showing oxygen above detection limit and the measured oxygen levels have no background contribution However for a heat treated sample, high temperature heat treated to 600 o C/10hrs, the raster reduction profile in figure 5.9(b) shows the same increase in the oxygen intensities as the matrix (Nb - ) signal after raster reduction indicating that the oxygen intensities observed are real and are related to the sample, with no background contribution. This was observed for all the heat treated samples for which raster reduction was performed (800 o C/3hrs, 120 o C/12hrs and 600 o C/10hrs, 120 o C/48hrs respectively) High Temperature Heat Treatments Figure 5.10 shows the O concentration vs depth for 600 C-1400 o C heat treatments as compared to a control sample, vs depth o C heat treatments show dissolution of oxygen in to the material, with the baseline oxygen concentration highest (almost an order of magnitude compared to a control sample) for the 1200 o C heat treatment, indicating diffusion of oxygen into the bulk. This is expected, since from equation 5.5, the diffusion length for oxygen at 1200 o C is in mm range and oxygen should diffuse into the bulk and show such a concentration profile. 162

188 Control 600C/10hrs 1000C/6hrs 1200C/6hrs 1400C/3hrs 1E+22 O Concentration (at./cm^3) 1E+21 1E+20 1E C/3hrs 600C/10hrs 1200C/6hrs 1E+18 Control Depth Figure 5.10: Oxygen concentration for heat treated samples (600 o C-1400 o C) and control However, the 1400 o C heat treated sample shows an unexpected diffusion profile. At 1400 o C/3hrs, the diffusion length of oxygen calculated from eq. 5 5 should be about 4.5mm. However, instead of oxygen diffusing into the bulk and baseline concentration increase, a steady decrease in the oxygen concentration is observed upto 2.5 μm depth. This diffusion profile is possibly due to the gettering of oxygen by dissolved titanium from the furnace at these temperatures, which will be discussed in section 5.6. The small peak at 2.5µm in the O profile is due to outgassing event from ion pump High Temperature Heating followed by Long Term Low Temperature Baking Figure 5.11 shows the O concentration vs depth for 600 C-1400 o C/120 o C heat treated samples in comparison to a control (non heat treated) sample. 163

189 Control 600C/10hrs,120C/48hrs 800C/3hrs,120C/12hrs 1E C/2hrs,120C/12hrs 1200C/2hrs,120C/12hrs 1400C/3hrs,120C/12hrs O Concentration (at./cm^3) 1E+21 1E+20 1E C/3h,120C/12h 1E C/2h,120C/12h 800C/3h,120C/12h Depth (um) Figure 5.11: Oxygen concentration for various high temperature heated/low temperature baked heat treated samples vs. non heat treated sample (control) As seen in the figure, higher the temperature of heat treatment, more prominent is the diffusion profile, with almost all heat treated samples showing nearly the same baseline level of oxygen concentration. However, these baseline levels are achieved at a much deeper depth as the temperature of heat treatment is increased. Since all the samples were baked at 120 o C for the same amount of time (12hrs) after high temperature heating, time can be ruled out for the difference in diffusion profiles. This observation confirms that at higher temperatures of heat treatment, there is more diffusion of oxygen into the material. Both the 1200 o C and 1400 o C/120 o C heat treated samples show increased oxygen diffusion away from the sample surface, as compared to the lower temperature heat treatments, with the 1400 o C/120 o C heat 164

190 treated sample showing a step in the oxygen concentration within the first 100nm and more diffusion, as compared to the 1200 o C/120 o C heat treated sample. This step in the oxygen profile cannot be explained. One explanation can be that since Ti contamination on this sample shows the same diffusion profile as oxygen (see section 5.6), more oxygen might be diffusing into the material in the form of titanium oxide after the 120 o C bake. However, this is rather unlikely since the Ti profiles for the same heat treatment do not show this step High Temperature Heat Treatment with Lower Temperature Heating There were two types of heat treatments performed in this these experiments. One was the 800 o C/3hrs-400 o C/20min heat treatment performed in vacuum conditions throughout. The other type of heat treatment involved high temperature heating, followed by a lower temperature heating, where N 2 gas was introduced in the chamber for about 10-20min at 10-5 mbar pressure during the experiment. Two of the samples were heat treated according to the latter method of heat treatment, with one of them being further baked at 120 o C/6hrs as part of the process. The introduction of N 2 was made in an attempt to achieve nitridation of the Nb surface to prevent surface contamination of the sample on cool down with the goal of preventing the adherence of impurities (H,C,O) to the Nb surface once the furnace reached room temperature (see chapter 6). Figure 5.12 shows the O concentration for the 800 o C/3hrs, 400 o C/20min heat treated sample as compared to a non heat treated (control) sample. It can be inferred from the figure that this heat treatment does not lead to any significant change in the baseline oxygen concentration as compared to a control sample. 165

191 Control 800C/3h,400C/20min 1E+22 O Concentration (atoms/cm^3) 1E+21 1E+20 1E+19 1E+18 Control 800C/3h,400C/20min 1E Depth (um) Figure 5.12: Oxygen concentration for 800 o C/3hrs, 400 o C/20min heat treated sample as compared to a control Nb sample A more conclusive result is seen for the 800 o C/3hrs, 400 o C/20min (N 2 at 10-5 mbar), 120 o C/6hrs heat treated sample, as shown in figure There seems to be no change in the O concentration before (control) and after this heat treatment. However, heat treating the sample at a higher temperature 1000 o C/2hrs and holding it at 800 o C/10min at a N 2 pressure of 10-5 mbar induces some oxygen diffusion away from the surface, into the material. This result is rather surprising, since the 1200 o C/6hrs heat treated sample did not show this profile, even though it was heated to a higher temperature in the same new furnace. This indicates that holding the sample at 800 o C for 10min is responsible for this diffusion. 166

192 Control 800C/3h,400C/20min N2,120C/6h 1000C/2h,800C/10min N2 1E+22 O Concentration (atoms/cm^3) 1E+21 1E+20 1E+19 1E+18 1E+17 Control (black profile below the green profile) 1000C/2h, 800C/10min N2 800C/3h,400C/20min N2,120C/6h Depth (um) Figure 5.13: Oxygen concentration for 800 o C/3hrs, 400 o C/20min (N 2 at 10-5 mbar), 120 o C/6hrs; and the 1000 o C/2hrs, 800 o C/10min (N 2 at 10-5 mbar) heat treatments as compared to a control sample Conclusion High temperature single heat treatments ( o C) are seen to dissolve oxygen into the Nb bulk, as seen from SIMS analysis of these samples as compared to a control sample, which show an increase in the baseline oxygen concentration. The amount of oxygen dissolved depends on the temperature of heat treatment. For the high temperature heat treated and baked (120 o C) samples, there appears to be an oxygen diffusion profile from the sample surface into the bulk, which is more prominent for temperatures above 1000 o C. Thus, some diffusion of oxygen is seen in the first 100nm of the Nb surface after high temperature heating followed by a low temperature bake of the samples,. For verification of this diffusion, low energy SIMS analysis for better depth resolution is essential (See section 5.7). For heat treatments above 1000 o C, especially o C, the apparent surface contamination of oxygen is seen, and it is suspected that oxygen 167

193 dissolution into the bulk is prevented due to the gettering of oxygen by substitutional metallic impurities, such as titanium. 5.6 SIMS Characterization of Titanium in Niobium The importance of the 1400 o C/3hrs heat treatment has been established by the fact that the Q o improvement in the cavity undergoing this same heat treatment produced a record breaking 200% (see chapter 7). In section , where data for hydrogen levels is shown for this heat treatment, it is seen that the H - /Nb - ratio was at the similar levels after this heat treatment as was seen for other high temperature heat treatments. However, oxygen was seen to contaminate the Nb surface after this heat treatment, and it was suspected that a metallic gettering agent was responsible for preventing long distance diffusion of oxygen (mm range for o C, eq.5.5) for both the 1200 o C and 1400 o C heat treatments Experimental Mass spectra of the 1400 o C/3hrs heat treated Nb sample were taken using an ION-TOF V TOF-SIMS instrument, with a 7nA Bi + 3 analysis beam at 25keV energy. The spectra showed high levels of titanium on the surface as compared to other metallic impurities which showed extremely low intensities. A Ti in Nb standard was prepared via implantation of Ti into an Nb control sample to allow quantification of Ti in the 1400 o C/3hrs and 1200 o C/2hrs, 120 o C/12hrs heat treated samples. The CAMECA-IMS 6f was used for quantitative analysis of Ti in Nb, using a 5.5keV impact energy O + 2 primary ion beam, to enhance the positive secondary ion yield of electropositive Ti to and to optimize Ti secondary ion yield and thus sensitivity. A 200μm raster with a 60μm detection area and a 120nA primary beam current was used for this analysis. Mass resolution of M/ΔM=2200 was chosen so as to minimize any mass interferences, while the analysis chamber was also kept under ultra high vacuum (UHV) conditions ( Torr) to minimize contamination. In addition, to obtain high depth resolution concentration profiles of Ti, low energy (1.25 kev impact energy) 20nA O + 2 primary ion beam was used. The raster size chosen was 220 x 220 μm 2, with a 60μm diameter detected area, again, to avoid crater wall effects. The mass resolution for this analysis was also m/δm = High resolution analysis was only 168

194 performed for the 1400 o C/3hrs, 120 o C/12hrs heat treated sample to investigate the high Ti levels seen on the surface in the high energy SIMS depth profiles Results and Discussion Figure 5.14 shows the TOF-SIMS mass spectrum for the 1400 o C sample over the Ti mass range and clearly identifies the presence of Ti on the sample surface. Figure 5.15 shows the depth profile of the ion implant standard of Ti implanted in Nb. The implant peak, which can be clearly seen, was used to extract the relative sensitivity factor (RSF) required for quantitative analysis of Ti in the Nb samples. Intensity (counts) 3 x Ti Mass (u) Figure 5.14: TOF-SIMS mass spectrum of 1400 o C /3hrs heat treated sample. High Ti levels detected on the surface 169

195 Concentration (atom s/cm ^ 3) 1E+20 1E+19 1E+18 1E+17 1E+16 Ti Conc Depth (um) Figure 5.15: Concentration profile of Ti in Nb implant standard; implant peak clearly seen Figure 5.16 shows the Ti concentration versus depth profile in Nb for the 1200 o C/2hrs, 120 o C/12hrs, 1400 o C/3hrs heat treated samples, and a non heat treated sample. No titanium is seen on the surface of the control sample, while the both the heat treated samples show what appears to be a Ti diffusion profile from the surface into the material. The 1400 o C/3hrs heat treated sample showed diffusion of Ti to a depth of about 3um. These Ti levels are surprising, since Ti was not intentionally introduced into the Nb. It is believed that the source of this contamination is the induction furnace used for heat treatment since the susceptor of the furnace is made of an Nb-Ti alloy. More diffusion of Ti was observed in the 1400 o C/3hrs heat treatment probably as a result of the vapor pressure of Ti being significantly higher at this temperature ( 10-5 torr) 39, apparently leading to its increased dissolution in the Nb heat treated at this higher temperature. 170

196 1E+23 Ti concentration (atoms/cm^3) 1E+22 1E+21 1E+20 1E C/2h,120C/12h 1400C/3hrs Control 1E Depth (um) Figure 5.16: Concentration profile of Ti in Nb for 1400 o /3hrs and 1200 o C/2hrs, 120 o C/12 hrs heat treated samples in comparison to a control sample. An interesting aspect of this analysis can be seen in figure 5.17, which shows the oxygen concentration for the same samples mentioned in figure Based on this data, it can be inferred that the oxygen concentration profiles match the profiles seen for titanium in figure 5.16, supporting the notion that oxygen is gettered by titanium in the material thus preventing oxygen diffusion into the bulk of the material. This conclusion is aided by the knowledge that Ti has a higher affinity for oxygen than Nb. 40 This gettering of oxygen by Ti was also seen by Padamsee 41 after heat treating Nb to 1400 o C, similar to the temperatures used in the heat treatment of the sample above. 171

197 1E+22 O Concentration (atoms/cm^3) 1E+21 1E+20 1E C/3h 1200C/2h,120C/12h 1E+18 Control Depth (um) Figure 5.17: Concentration profile of O in Nb for 1400 o /3hrs and 1200 o C/2hrs, 120 o C/12 hrs heat treated samples. O concentration seems to match the Ti conc. profiles seen in figure This diffusion profile of oxygen was not observed in other heat treatments and it is known that Ti increases the solubility of oxygen in Nb at high temperatures. 42 To obtain the higher SIMS depth resolution needed to provide more information about the surface region, analyses were also conducted at lower impact energy (1.25keV O + 2 ) (see section for experimental details). Figure 5.18 shows SIMS low energy analysis of the 1400 o C/3hrs, 120 o C/12hrs heat treated sample. Note that the first 5nm of the surface shows a very high concentration of Ti in Nb, possibly an anomaly resulting from the additional oxygen in the surface oxide layer which may provide additional positive secondary ion yield enhancement for the Ti. 43 For depths below 5nm i.e. below the surface Nb oxide, the Ti concentration can be accurately measured. At this depth and below, Ti concentration was measured to be about 1 atomic %, up to 40nm of the surface, indicating a high level of Ti. 172

198 1E+24 Ti Concentration (atom s/cm ^3) 1E+23 1E+22 1E+21 1E C/3h,120C/12h Depth (um) Figure 5.18: Concentration profile of Ti in Nb for 1400 o /3hrs, 120 o C/12 hrs heat treated sample, using low energy O 2 + primary ion beam (1.25keV). Ti concentration estimated to be about 1% atomic fir 5-50nm Conclusion High levels of Ti were seen in the 1400 o C/3hrs and 1200 o C/2hrs, 120 o C/12hrs heat treated samples. The source of this Ti contamination is believed to be from the furnace used for heat treatments. The diffusion of oxygen into the material observed in these samples can be explained by the fact that Ti getters oxygen, preventing movement of oxygen into the bulk. Low energy analysis of the 1400 o C/3hrs, 120 o C/12hrs heat treated sample shows a high level of Ti in the 5-50nm depth of Nb surface ( 1 atomic %) supporting the notion that the Ti was introduced to the Nb through its surface via a source within the oven used for heat treatment. 5.7 Low Energy SIMS Characterization of Oxygen in Nb Low energy SIMS analysis was performed to provide improved depth resolution and to study the oxygen concentration in the first 40nm of large grain (single crystal) Nb sample surfaces, since this is the depth which is of prime importance in terms of RF cavity performance. To obtain improved depth resolution, analyses were performed on the same samples analyzed in section 5.5 but in this case, the SIMS instrument was configured to provide a low energy Cs + 173

199 primary ion beam (6keV impact energy), with the analysis conditions mentioned in chapter 3, section Results and Discussion As mentioned in section 5.4 and in chapter 4, the long term low temperature baking has been seen to improve the SRF cavity performance. Section 5.4 shows that according to the oxygen pollution model, this might be explained by the diffusion of oxygen into the bulk from regions of high concentrations of interstitial oxygen near the surface. Thus, SIMS analysis of samples after this baking process can provide useful information of the long term baking on the oxygen levels in the first 40nm of the Nb surface. Besides this baking, other types of heat treated samples shown in section 5.5 were also analyzed using low energy SIMS, to characterize oxygen levels in the first 40nm of the Nb surface and to study the effects of these heat treatments on the Nb surface oxide layer Long term Low Temperature Baking (120 o C/48hrs) Figure 5.19 shows the high depth resolution oxygen concentration profile for the 120 o C/48hrs baked sample in comparison to a non heat treated (control) sample over the first 100nm depth. The high concentration of oxygen on the surface for both samples confirms the presence of the surface oxide layer. There appears to be some diffusion of oxygen observed for this baking, but the depth over which the length of the diffusion is not observed to be as great as predicted by the diffusion length equation (25nm for 120 o C/48hrs from equation 5.5). However, the small amount of diffusion seen (1-2nm), corroborates the results from x- ray diffuse scattering (figure 5.7). 38 This small diffusion of oxygen as compared to that predicted by equation 4.5 can be explained by the speculation that oxygen diffusion might be compensated by the segregation of oxygen back to the surface on cool down. 4,38 174

200 1E+24 Concentration (atoms/cm3) 1E+23 1E+22 1E+21 1E+20 1E+19 Control 120C/48hrs 1E Depth (um) Figure 5.19: Concentration profile of O in Nb for 120 o C/48 hrs baked sample in comparison to a control sample. Small diffusion of oxygen (1-2nm) seen after baking (Low energy SIMS, 6keV Cs + ) High Temperature Heat Treatment Figure 5.20 shows the oxygen concentration profile for the 1200 o C/6hrs heat treated sample in comparison to a non heat treated (control) sample over the first 100nm depth. As seen in the figure, the first 5nm of the surface has similar, and high oxygen concentration indicating the presence of the surface oxide layer. However, the 20-60nm region of the surface is observed to have a slightly higher oxygen level for the heat treated sample, as compared to the control sample. As seen from the high energy data in section , oxygen ahs dissolved via diffusion into the material at this temperature, leading to higher oxygen in the 20-60nm region of the surface after 1200 o C heat treatment. 175

201 1E+24 O Concentration (atoms/cm^3) 1E+23 1E+22 1E+21 1E+20 1E C/6h Control 1E Depth(um) Figure 5.20: Concentration profile of O in Nb for 1200 o C/6 hrs heat treated sample in comparison to a control sample (Low energy SIMS, 6keV Cs + ) High Temperature Heat Treatment with Long term Low Temperature Baking Figure 5.21 shows the oxygen concentration vs. depth (for the first 100nm) comparison of the 800 o C/3hrs, 120 o C/12hrs and 1200 o C/2hrs, 120 o C/12hrs samples as compared to a non heat treated sample (control). Similar data is shown for the 1400 o C/3hrs, 120 o C/2hrs heat treated sample vs. control in figure Higher levels of oxygen can be clearly observed for the 1200 o C/2hrs, 120 o C/12hrs sample in figure 5.21, possibly due to the titanium (see figure 5.18 in section 5.6) diffusion at these high temperatures. The concentration in the 10-40nm region is about atomic % oxygen for this sample. The diffusion profile observed in the first 100nm for the baked 800 o C/3hrs, 120 o C/12hrs in section , using high energy SIMS, is also confirmed using low energy SIMS, in figure This diffusion could be due to the additional baking step performed after high temperature heat treatment of the sample and as described in the higher energy data, it is also observed here that higher the temperature of initial heat treatment, more is the oxygen diffusion seen as compared to the control sample. 176

202 1E+24 O C oncentration (atom s/cm ^3) 1E+23 1E+22 1E+21 1E+20 1E+19 1E C/3h,120C/12h 1200C/2h,120C/12h Control Depth(um) Figure 5.21: Concentration profile of O in Nb for 800 o C/3hrs, 120 o C/12hrs and 1200 o C/2hrs, 120 o C/12hrs heat treated samples in comparison to a control sample (Low energy SIMS, 6keV Cs + ) 1E+24 O Concentration (atoms/cm^3) 1E+23 1E+22 1E+21 1E+20 1E+19 1E C/3h,120C/12h Control Depth(um) Figure 5.22: Concentration profile of O in Nb for 1400 o C/3hrs, 120 o C/12hrs heat treated sample in comparison to a control sample (Low energy SIMS, 6keV Cs + ) 177

203 In figure 5.22, the step of oxygen observed in the high energy data is confirmed for the 1400 o C/3hrs, 120 o C/12hrs heat treated sample. A much better explanation of this step oxygen profile can be provided at this point. While the 1400 o C/3hrs heat treatment incorporates titanium oxide onto the surface of Nb (section 5.6), the additional 120 o C bake might be pushing the oxygen (and not Ti) from this oxide further into the material and causing the step profile, as seen in the figure High Temperature Heat Treatment followed by Lower Temperature Heating Figure 5.23 shows the oxygen concentration vs. depth (for the first 100nm) of the 800 o C/3hrs, 400 o C/15min at 10-5 mbar N 2, 120 o C/6hrs sample in comparison to a non heat treated sample (control). The profiles are seen to match indicating no change in the oxygen levels after this heat treatment, similar to the resuls obtained from high energy SIMS analysis (section ) 1E+24 Control 800C/3h,400C/15min N2, 120C/6h O C oncentration (atom s/cm ^3) 1E+23 1E+22 1E+21 1E+20 1E C/3h,400C/15min N2, 120C/6h 1E+18 Control Depth(um) Figure 5.23: Concentration profile of O in Nb for (800 o C/3hrs, 400 o C/15min at 10-5 mbar N 2, 120 o C/6hrs) heat treated sample in comparison to a control sample (Low energy SIMS, 6keV Cs + ) 178

204 5.7.2 Conclusions It is observed from low energy SIMS analysis of the first 40nm of the Nb surface that there is no significant difference in the oxygen concentration in the first 5nm of the surface, compared to a control sample indicating that the surface oxide remains intact even after all types of heat treatments performed on the samples. In the nm region however, the oxygen concentration is seen to be higher for higher temperature heat treated samples. The 1400 o C/3hrs, 120 o /12hrs heat treated sample shows the highest concentration of oxygen in the 10-60nm region of the surface, possibly due to the gettering of oxygen by Ti. The 1200 o C/2hrs, 120 o C/12hrs heat treated sample also shows high levels of oxygen in the 10-60nm region for similar reasons, as also observed in section 5.6. In contradiction to the calculation from equation 5.5, the 120 o C/48hrs long term baking is not observed to diffuse oxygen into the material by a diffusion length of 25nm, possibly due to oxygen segregation back on the surface during cool down which compensates for any kind of diffusion of oxygen into the material. 5.8 TEM Characterization of the Surface Oxide of Niobium The similar surface oxygen (5 nm depth) observed in the heat treated samples using SIMS, and the fact that the surface oxide acts as a passivation layer on the surface of niobium, as seen in chapter 4, makes it important to perform characterization of the surface oxide and to study the effects of heat treatment on the surface oxide thickness. In this section, Transmission Electron Microscopy is used to provide further information of surface oxide on Nb Experimental A FEI Quanta 200 3D Focused Ion Beam (FIB) system with a Ga + ion beam was used to prepare TEM cross sections of both the heat treated and control samples (treatment of samples similar to the procedure used for SIMS characterization, see section 3.1). More information on the FIB technique can be found in a summary publication. 44 Bright field imaging of the samples, using Transmission Electron Microscopy (TEM) was performed with a Hitachi HF 2000 microscope. 45,46 using a 200keV electron beam. 179

205 5.8.2 Results and Discussion Figure 5.24 shows the TEM image of a non heat treated sample (control). The image shows a continuous surface oxide with a sharp Nb-Oxide interface. The oxide thickness was found to be 6nm. There were no sub-oxides observed near the surface oxide. Figure 5.24: TEM image of a control Nb sample (a) (b) Figure 5.25: TEM image of a control (a) and heat treated (800 o C/3hrs, 140 o C/3hrs) Nb (b) sample 180

206 Figure 5.25 shows a comparison of the oxide thickness of a control sample (a) to a heat treated (b) sample (800 o C/3hrs, 140 o C/3hrs). The oxide thickness is observed to be less than 10nm for both the samples. From figure 5.24, it can be inferred that the oxide thickness varies between control samples, but a general measurement of <10nm oxide is seen for both the heat treated and control samples. Similar results were seen by Tian 46 on TEM analysis of baked Nb (120 o C) samples, where the oxide thickness was observed to be about 8nm for the control sample and 7nm for the baked sample, indicating no significant effect of baking on the oxide thickness, possibly due to the segregation of oxygen back on the surface during cool down. This result corroborates the results obtained from low energy SIMS analysis of a similary heat treated sample (800 o C/3hrs, 120 o C/12hrs, see figure 5.21), where there was no significant difference seen for the oxygen concentration on the first 5nm of the surface, as compared to the control sample Conclusion In the samples analyzed by TEM, no significant effect of heat treatment on the surface oxide thickness was observed. The surface oxide is observed to be continuous, with a sharp Nb- Oxide interface and less than 10nm for both heat treated and control samples. No suboxides were observed via TEM, which has been found to be similar to some observations by other researchers 46. These results support the SIMS characterization data obtained for oxygen in the first 5nm of the Nb surface (section 5.7) 5.9 TOF-SIMS Imaging of Oxygen in Large Grain Nb Bicrystals The oxygen pollution model (section 5.4) identifies grain boundaries as a region for high oxygen concentration which can deteriorate cavity performance. Thus, it is important to study the grain boundaries of large grain Nb for any oxygen segregation both before and after heat treatment. As described in chapter 3, section 3.6.1, two types of Nb bicrystals were prepared for surface imaging at the grain boundary: bicrystals having different crystallographic orientation combinations and bicrystals having the same crystallographic orientation combinations. Details of sample preparation and crystallographic orientation determination are provided in 181

207 section Surface imaging was performed at the grain boundary using an ION-TOF V TOF SIMS instrument (chp 2) with analysis conditions mentioned in chapter 3, section Results and Discussion Bicrystals having different crystallographic orientation combinations Figure 5.26 shows the ion images for the control (a) and the heat treated (b) (800 o C/3hrs, 120 o C/24hrs) samples. No segregation of oxygen was seen for either the control or the heat treated samples. It should be noted that in the Figure 5.26, the upper crystal is a (001) oriented crystal while the lower crystal has higher miller indices for both control and heat treated Nb (EBSD results from chapter 3, section 3.6.1) with the control sample having a (213) orientation and the heat treated sample having a (215) orientation for the lower crystal respectively. Although the intensity differences across the grain boundary interface are insignificant, these differences could be the result of differing sputtering or primary beam implantation conditions resulting from the differing crystallographic orientations of the crystals across the grain boundary. The important result obtained from this analysis is that that there is no segregation of oxygen at the large grain Nb bicrystal grain boundary for either the heat treated or control samples. Grain Boundary (001) (001) (213) (215) (a) (b) Figure 5.26: TOF SIMS ion images of control (a) and heat treated (800 o C/3hrs, 120 o C/24hrs) (b) bicrystal samples. 182

208 Bicrystals having the same crystallographic orientation combination In order to account for the effect of crystal orientation on the secondary ion intensities and obtain a better understanding of the heat treatment effect on the grain boundaries of large grain Nb bicrystal Nb samples, a separate set of samples was characterized using TOF-SIMS imaging using the same instrument settings described above. However, to eliminate any orientation effects, a large sample was sectioned from the same grain boundary area and cut into half. One part of this large sample was heat treated to 1200 o C/6hrs in the new furnace after BCP (1:1:2, 7μm removal) while the other part was not heat treated after the same BCP etching (control). Figure 5.27 shows the TOF SIMS images of the control (a) and heat treated (b) Nb samples. The upper crystal in both the samples is similarly oriented, while the lower crystal also has the same orientation in both the samples. Again, oxygen is not seen to be segregated at the grain boundary for both the control and heat treated (1200 o C/6hrs) samples. (hkl) Grain Boundary (hkl) (h k l ) (h k l ) (a) (b) Figure 5.27: TOF SIMS ion images for control (a) and heat treated (1200 o C/6hrs) (b) bicrystal sample grain boundaries with similar crystal orientation combination for both samples. 183

209 5.9.2 Conclusion TOF SIMS analysis of the grain boundary region of large grain bicrystal Nb shows that oxygen does not segregate at the grain boundaries for both non heat treated and heat treated samples. Initial analysis suggested that the small differences in the O - intensities observed in the different crystals may be related to differing sputtering or primary beam implantation conditions resulting from their differing crystallographic orientations. However, subsequent analysis performed on the grain boundary area of similarly oriented Nb bicrystals show that there is indeed no segregation of oxygen at the grain boundaries of the bicrystals. 184

210 References 1. H. Padamsee et al : RF Superconductivity for Accelerators, 2 nd edition, Wiley New York (2008) 2. H. Padamsee: Supercond. Sci. Technol., 14, pr28 (2001) 3. G. Ciovati: App. Phys.Lett., 89, p22507 (2006) 4. G. Ciovati et. al: Phys. Rev. Special Topics: Acc. and Beams, 13, p22002 (2010) 5. G. Ciovati: Journ. App. Phys., 96(3), p1591 (2004) 6. W. Singer et. al: Physica C, 386, p379 (2003) 7. C. Antoine et al : AIP Conf. Proc. 671, p176 (2002) 8. A. Septier: Proc. Workshop on RF Superconductivity, KfK,53, p3019 (1980) 9. M. Grunder et al.: J. Appl. Phys., 51 (1), p397 (1980) 10. M. Grunder et al.: Surf. Sc., 136, p144 (1984) 11. J. Halbritter: Kernforschungszentrum Karlsruhe, Primarbericht, , P10A. 12. T. Asano et al: KEK Report 88 (2) (1988) 13. P. Kniesal et al.: Proc. of 1972 App. Superconductivity Conf., Annapolis, p657 (1972) 14. F. Palmer et. al: Proc. of 3rd Workshop on RF Superconductivity, Argonne, IL, p309 (1988) 15. C. Valet et al.: Proc. of 1992 European Part. Acc. Conf., Editions Frontieres, p1295 (1992) 185

211 16. W. DeSarbo: Phys. Rev., 132 (1), p107, J. Halbritter: Appl. Phys. A, 43, p1 (1987) 18. E_Source: ASM phase diagrams center, G. Brauer: Anorg. Allgem. Chem,1, p248 (1941) 20. A. Seybolt: Trans. AIME, 200, p770 (1954) 21. R. Elliott: Trans. Am. Soc. Metals, 52, p990 (1960) 22. R. Bryant: Journ. Less Comm. Met., 4, p62 (1962) 23. E. Gebhardtand et. al: Z.Metallk,54, p443 (1963) 24. M. Abrizov et.al.: Izvestiya Vuz. Fizika, 2, p. 129 (1965) 25. T. Hurlen: J. Metals, 89, 8, p (1961) 26. P. Kofstad et. al: Transac. Met. Soc. AIME, 2, p221 (1961) 27. P. Phelps et. al.: J. Inst. Metals, 88, p301 ( ) 28. G. Anderson et.al: Acta Chem. Scand., 11, p1065 (1957) 29. W. Hicks: Trans. Met. Soc. AIME, 221(2), p191, (1961) 30. A. Dacca et.al: Appl. Surf. Sci., 126, p219 (1998) 31. A. King et.al: Thin Solid Films, 192, p351 (1990) 32. Q. Ma et.al.: J. Appl. Phys. 96, p7675 (2004) 33. A. Taylor et.al: Journ. Less Comm. Met., 13, p313 (1967) 186

212 34. W. Worrell: J.Phys. Chem.,68, p952 (1964) 35. C. Ang: Act. Mett.,1, p123 (1953) 36. G. Ciovati et.al: Phys. Rev. ST Accel. Beams, 10, p (2007). 37. G. Eremeev et. al: Proceedings of the 13 th Workshop on RF Superconductivity, Beijing, p356 (2007) 38. M. Delheusy: Ph.D. thesis, University of Paris-Sud IX and Stuttgart University (2008) 39. P. Desai: Int. J. of Thermophys., 8 (6), p781 (1987) 40. R. Kirchheim: Act. Metall., 27, p869 ( H. Padamsee: IEEE Trans. of Mag., Mag-21, 2, p1007 (1985) 42. K. Bryant, Journ. Less Comm Met.,4, p62 (1961) 43. F. Stevie et al : SIMS- A practical handbook for depth profiling and bulk impurity analysis, Wiley Interscience Publications (1989) 44. F. Stevie et. al: Introduction to Focused Ion Beams, Springer, Berlin (2005) 45. E_Source: H. Tian: Surface Oxide Study On Solid Niobium For Superconducting RF Accelerators presentation at SRF Workshop 2007, Fermi lab (2007) 47. D. Porter et. al: Phase Transformations in Metals and Alloys, CRC Press, Boca Raton, Florida, USA (1992) 187

213 6. Carbon and Nitrogen Carbon and nitrogen are interstitial impurities in SRF Nb which have not been studied as extensively as oxygen and hydrogen. This is probably due to their low solubility in Nb even at elevated temperatures as compared with oxygen and their less deleterious effect on superconductivity as compared with hydrogen. 1 However, high concentrations of these impurities can adversely affect SRF cavity performance making it important to characterize C and N in Nb, especially in the first 40nm of the Nb surface. Incorporation of these impurities occurs primarily during the SRF cavity fabrication process. Recrystallization heat treatment of ingot Nb and electron beam welding of the cavity cell half structures can induce carbon and nitrogen into the material Effect of Carbon and Nitrogen on properties of SRF Nb Like hydrogen and oxygen, carbon and nitrogen form interstitial solutions with Nb, and thus are detrimental to cavity performance since their presence affects the mobility of electrons and decreases the superconducting energy gap (see chapter 1). As a result, it is easier for the Cooper pairs to dissociate to normally conducting electrons near the surface and decrease the surface superconductivity. 1 Limited data is available for the effects of interstitial carbon on the superconducting properties of Nb, which is the reason that only the effects of nitrogen on the resistivity and critical temperature of Nb will be discussed in this section. 2 Desorbo 2 studied the effect of nitrogen concentration on the critical temperature and resistivity in the non superconducting state of Nb up to and beyond the solubility limit of N by dissolving nitrogen in Nb at high temperatures and quenching the alloy to retain the solid solution or phase formation, depending on the temperature. The results showed that nitrogen present in niobium in amounts equal to 0.33 at.% depressed the transition temperature to 9.1 K from 9.4 K measured for pure Nb, representing a slight decrease in critical temperature. This concentration was the solubility limit of nitrogen in niobium quenched from 1200 o C. Increasing the cooling rate (for example, cooling the specimen in helium exchange gas instead of in vacuum) after introducing the gas at an elevated temperature, resulted in only a slight change in the critical temperature. At a 188

214 nitrogen concentration exceeding the solubility limit as in Nb with 1.64 at.% N, T c rose to a value of 9.24 K, but was still found to be lower than pure Nb. For resistivity in the normal state 2, it was found that nitrogen present interstitially in the niobium lattice enhanced the resistivity. In Nb with 0.33 at.% N (the solubility limit), the resistivity was found to be 1.9x10-6 Ω-cm for a specimen cooled in vacuum after introducing the N at 1200 o C. For somewhat more rapid cooling, the resistivity increased to approximately 2.4x10-6 Ω-cm. For Nb containing 1.64 at.% N (N concentration in excess of the solid solubility limit) and cooled in vacuum, the resistivity was found to be equal to 1.8x10-6 Ω-cm. Thus, this represented little or no change with the similarly heat treated niobium containing the smaller nominal nitrogen concentration. Figure 6.1 shows the relationship of resistivity in the normal state for various compositions of Nb+N. Figure 6.1: N conc. dependence on resistivity of Nb 2 ' 189

215 6.2 Solubility, Diffusion of Carbon and Nitrogen in Nb Solubility of C in Nb The solubility of carbon in Nb is low, compared with oxygen (see chapter 4) and nitrogen (see following section). Figure 6.2 shows the phase diagram for the Nb-C system. Several investigators 6-10 have measured the solubility of C in Nb in the temperature range C. The solvus in fig. 6.2 was investigated by Horz et.al 8 on the basis of the resistometric measurements between 1500 and 2200 C and may be described by the relation: c = 4190 exp( /T) 8, for 1500 o C <T < 2200 o C...(6.1) where c is in at.% C. Internal friction studies of low C alloys quenched and aged at 310 o C for different times place the solubility limit at that temperature at less than at.% C. 11 Figure 6.2: Nb-C phase diagram

216 6.2.2 Solubility of Nitrogen in Nb Figure 6.3 shows the partial phase diagram of the N-Nb system, down to 50% atomic Nb. Although the solid solubility of nitrogen is seen to be higher than C, it is less than oxygen for elevated temperatures. 13 Studies by Ang et.al 14 have placed the solubility limit of N in Nb at 0.33% atomic N at a temperature of 1200 o C. The solubility equation was given as: c = 1.7 exp (-4600/RT) for 300 o C < T < 1200 o C...(6.2) where c is in atomic % nitrogen Cost et. al 15 provided an equation for the terminal solubility above 1150 o C as: c = 720 exp (-20,000/RT) for 1150 o C < T < 2230 o C...(6.3) where c is in atomic % nitrogen. Figure 6.3: Phase diagram of the N-Nb system

217 6.2.3 Pressure-Temperature Relationships in the Nb-N system Nitrogen absorption in Nb is reversible and follows Sievert s law (see section 3.3) up to a concentration of about 2 % atomic nitrogen. 16 Figure 6.4 shows the equilibrium plot of the Nb-N system relating N concentration and pressure at various temperatures. The plot indicates that for a temperature of over 1000 o C, the nitrogen pressure should be kept below 10-5 mbar so as not to exceed the solubility limit of nitrogen in Nb. 16 The solubility of nitrogen in Nb, in terms of pressure and temperature can be estimated by the relation: 16 Log (C N ) = ½ log p(n 2 ) /T 3.76, for 1770 K < T < 2470 K..(6.4) in the primary solid solution of nitrogen in Nb (α phase shown in fig. 6.4) Figure 6.4: Equilibrium pressure-temperature-concentration diagram of the Nb-N system

218 6.2.4 Diffusion of Carbon and Nitrogen in Nb (a) Nitrogen Diffusion of nitrogen in Nb is found to be many orders of magnitude slower than hydrogen in Nb. Ang 17 studied the diffusion of nitrogen in Nb using a torsional pendulum technique. Figure 6.5 shows the diffusion coefficient results from these findings. Similar to hydrogen and oxygen, the diffusion rate of nitrogen in Nb increases with temperature, given by the relation: D = 9.8 x 10-2 exp (-38,600/RT) for 300< T< 2300 K..(6.5), The activation energy was found to be 38,600 cal/mole 17 Figure 6.5: Diffusion Coeff. vs Temp (K) for nitrogen in Nb

219 (b) Carbon Powers et. al 18 studied the diffusion of C in Nb using an elastic relaxation technique. The diffusion equation was given to be: D = 3.8 x 10-2 exp (-33,000/RT) for 300< T< 650 K..(6.6), with the activation energy: 33,000cal/mol 18 From figures 4.5 (section 4.3.2), 5.5 (section 5.3.4), and equation 6.6, a comparison of the diffusion coefficients of oxygen in Nb, hydrogen in Nb and nitrogen in Nb can be made: Nitrogen in Nb: 1 x cm 2 /sec (300 K); Carbon in Nb: 3 x cm 2 /sec (300 K) Oxygen in Nb: 5 x cm 2 /sec (300 K) ; Hydrogen in Nb: 2 x 10-5 cm 2 /sec (300 K) Thus, for comparison purposes, it can be seen that hydrogen diffuses and times faster in Nb than oxygen and nitrogen respectively, and carbon diffusion is faster than nitrogen in Nb. The one dimensional approximate diffusion length of nitrogen and carbon is given by 19 : L= 2*(D*t) 1/2 cm..(6.7) Where D is the diffusion coefficient of nitrogen in Nb at a given temperature (from above figure and equation) and t is the time in seconds. 6.3 SIMS Characterization of Carbon and Nitrogen in Nb Carbon and Nitrogen contamination from the fabrication steps of a large grain Nb SRF cavity can be detrimental to cavity performance, as a result of the effect of these impurities on the superconducting properties of Nb, as seen in section 6.1. The utilization of various heat treatments after cavity surface treatments have been seen to be promising for the removal of carbon and nitrogen interstitials from the first 40nm of the Nb surface. 20 However, a systematic study on the effect of heat treatments on the surface carbon and nitrogen levels in Nb needs to be conducted to understand the role of these impurities in cavity degradation. SIMS can be a powerful technique to study the surface of Nb before and after these heat treatments, owing to the excellent depth resolution and detection limits achievable by this technique. 194

220 6.3.1 Results Chapter 3 discusses the sample preparation and SIMS analysis conditions used for the surface characterization of niobium. Due to furnace limitations, partial pressures of gaseous species involving N and C (e.g. CO and N 2 ) could not be monitored for the heat treatments. One of the non heat treated (control) samples was ion implanted with C and N at a dose of 1E15 atoms/cm 2, and implantation energy of 140 and 160 kev respectively, to obtain a SIMS quantification standard for both impurities for concentration measurements. Since N - ions have practically no ion yield in SIMS, NbN - secondary ions were used to monitor the nitrogen intensities. Figure 6.6 shows the concentration profiles of the C and N implanted standard. The C and N implant peaks are clearly seen in the figure, unlike the hydrogen in Nb implant. This indicates that both C and N do not have a high diffusion coefficient in Nb and the Gaussian shape shows the SIMS analysis does not affect the C and N distribution. Gaussian distributions for both impurity implant species are clearly seen in the figure, unlike the hydrogen in Nb implant. 1E+22 Concentration (atom/cm3) 1E+21 1E+20 1E+19 1E+18 NbN C 1E Depth (um) Figure 6.6: Depth profile of C and N in Nb implant. 195

221 As mentioned in section 4.9, the Nb - intensities decrease by a factor of 3 for high temperature heat treated samples. Since the RSF for oxygen implanted in Nb for a heat treated sample was also found to be a factor of 3 lower, it was assumed that the Nb- secondary ion would exhibit similar behavior for the C and N implants in a heat treated Nb sample. Thus, a correction factor of 1/3 for both C and N RSFs was applied to improve concentration accuracy for C and N in high temperature heat treated Nb samples. No correction was necessary for control samples. For the samples, various types of heat treatments were performed. The nomenclature provided in chapter 3, section 3.2 will be followed in this chapter as well. The main types of heat treatments performed were: (i) Long term low temperature baking; (ii) High temperature heat treatment; (iii) High temperature heating followed by low temperature heating and (iv) High temperature heating followed by lower temperature heating (heating in vacuum or in 10-5 mbar N 2 atmosphere). Before the SIMS results can be discussed, it is important to ascertain the background contribution to C and N levels measured by SIMS. Raster reduction (chapter 3 section 3.4.6) was performed on some of the control and heat treated samples to check for background contributions. Figure 6.7 (a) shows the raster reduction profile for C - and NbN - for a control sample, where the raster was 196

222 Counts (cts/sec) 1E+06 1E+05 1E+04 1E+03 1E x 120 μm 2 20nA 60 x 60 μm 2 20nA C- Nb- NbN- 1E Time (s) Figure 6.7(a): Raster reduction for C - and NbN - for a control sample reduced to 60 x 60 μm 2 area from 120 x 120μm 2 area while maintaining a primary ion beam current of 20nA. As seen in the figure, both C and N secondary ion intensities show a proportional increase to the matrix (Nb - ) secondary ion intensities, indicating no significant background contribution to the C and N levels. Thus, it may be inferred that both C and N in a control sample are above the detection limit of the instrument and are sample related. This was observed for all the control samples analyzed in this study. Similar observations can be made for both C and N secondary ion intensities for a heat treated sample (800 o C/3hrs, 120 o C/12hrs), fig. 6.7 (b), where both impurity secondary ion intensities are seen to increase by factor 4 (equivalent to the sputter rate and the matrix Nb - secondary ion intensity increase) indicating that in the heat treated samples as well, the C and NbN levels reflect actual C and N concentrations and are both are above the detection limit of the instrument. This was observed for all the heat treated samples for which raster reduction was performed (600 o C/3hrs, 120 o C/48hrs and 600 o C/10hrs respectively) 197

223 Counts (cts/sec) 1E+05 1E+04 1E+03 1E x 120 μm 2 20nA C- Nb- NbN- 20nA 60 x 60 μm 2 1E Time (s) Figure 6.7(b): Raster reduction for C - and NbN - for a control sample High Temperature Heat Treatments (a) Carbon Figure 6.8 shows the C concentration versus depth for 600 o o C heat treatments as compared to a control sample. The baseline C concentration in all the heat treated samples is seen to be higher than a control sample, indicating that C is dissolving into the material as a result of the heat treatments. For the 1200 o C/6h heat treatment, the C concentration is a factor of six higher than the control sample. Since the diffusion length of C for these heat treatments (600 o C-1400 o C) is μm depending on the temperature and time (equation 6.7), it can be speculated that any carbon deposited on the surface will diffuse into the bulk and increase the baseline C concentration on heat treatment, as can be seen in figure 6.8. However, the source of this deposited C cannot be identified. One explanation is that C rich gaseous species (e.g. CO) in the furnace, might be responsible for contaminating the Nb surface and hence the contamination could be attributed to the furnace. 198

224 1E+21 Control 600C/10h 1000C/6h 1200C/6h 1400C/3h C Concentration (atoms/cm^3) 1E+20 1E+19 1E C/6h 1400C/3h 600C/10h Control Depth (um) Figure 6.8: Carbon concentration for heat treated samples (600 o C/-1400 o C) versus non heat treated sample (control) (b) Nitrogen SIMS data is shown for the nitrogen concentration in figure 6.9. The N levels in the material seem to be variable after heat treatment and show no particular trend. While the 1000 o C heat treated sample shows a higher N concentration, the 600 and 1200 o C show a factor of three lower N level as compared to a control sample. Also, the 1400 o C heat treatment shows a very different N profile as compared to the other heat treatments, since the amount of nitrogen on the surface is observed to be the same as the control sample, but this level continuously decreases down to 1.5 μm depth where the level reaches a baseline concentration similar to the 600 and 1200 o C heat treatments. One explanation of this observation is that nitrogen may 199

225 have outgassed from the bulk of the sample and may have re-deposited on the surface upon cool down, unlike in the 600 and 1200 o C heat treatments where re-deposition of nitrogen on cool down might not have taken place. No plausible explanation can be made for the increase in the N levels observed after the 1000 o C heat treatment, without the knowledge of the N 2 partial pressure in the furnace during heat treatment (a limitation of the apparatus used for these heat treatments) 1E+20 Control 600C/10h 1000C/6h 1200C/6h 1400C/3h NbN Concentration (atoms/cm^3) 1E+19 1E C/6h 600C/10h 1000C/6h Control 1400C/3h 1E Depth (um) Figure 6.9: NbN concentration for heat treated samples ( o C) versus non heat treated sample (control) 200

226 High Temperature Heat Treatment and Long Term Low Temperature Baking (a) Carbon Fig shows the C concentration versus depth for 600 o C-1400 o C/120 o C heat treated followed by baking samples. The 800 o C/120 o C heat treatment shows a C peak at about 0.3μm depth, indicating diffusion of carbon from the surface to this depth. The carbon concentration is seen to be the highest for this sample and almost an order of magnitude higher than the control sample. 1E+21 Control 600C/10h,120C/48h 800C/3h,120C/24h 1000C/2h,120C/12h 1200C/2h,120C/12h 1400C/3h,120C/12h C Concentration (atoms/cm^3) 1E+20 1E C/3h,120C/24h 1200C/2h,120C/12h 1400C/3h,120C/12h 600C/10h,120C/48h Control 1E Depth (um) Figure 6.10: Carbon concentration for various high temperature heated/low temperature baked heat treated samples versus non heat treated sample (control) 201

227 The 1200 and 1400 o C/120 o C heat treated samples only show a factor of two increase in the C concentration as compared to a control sample, while the 600 and 1000 o C/120 o C heat treatments have similar baseline C concentrations as compared to a control sample. For all the heat treatments except the 800 o C/120 o C heat treatment, the surface C level is high (the first data point) as compared to a control sample indicating C contamination from the furnace. When compared with the single heat treatments (previous section), C seems to show a small diffusion profile until 0.3μm for the extra baking step- heat treatments. This clearly indicates that C is diffusing into the material as a result of the 120 o C/12-48hrs baking and contaminating the surface more as compared to the high temperature heat treatments, with no baking. (b) Nitrogen Figure 6.11 shows the nitrogen concentration versus depth for 600 o C-1000 o C/120 o C heat treatments. 1E+20 Control 600C/10h,120C/48h 800C/3h,120C/24h 1000C/2h,120C/12h NbN Concentration (atoms/cm^3) 1E+19 1E C/2h,120C/12h 600C/10h,120C/48h Control 800C/3h,120C/24h 1E Depth (um) Figure 6.11: NbN concentration for various high temperature heated/low temperature baked heat treated samples versus non heat treated sample (control) 202

228 The nitrogen levels are seen to decrease by a factor of three after each heat treatment, indicating desorption of N from the material on heat treatment. Figure 6.12 shows the same data for o C/120 o C heat treated samples as compared to a control sample. The 1400 o C/120 o C heat treatment shows the same N profile as the 1400 o C single heat treatment (section ). However, the 1200 o C/120 o C heat treatment seems to incorporate N on the Nb surface (upto 0.5 μm). The lower baseline N concentration in this sample as compared to a control sample indicates that the surface N contamination is a result of the 120 o C baking. 1E+20 Control 1200C/2h,120C/12h 1400C/3h,120C/12h NbN Concentration (atoms/cm^3) 1E+19 1E C/2h,120C/12h Control 1400C/3h,120C/12h 1E Depth (um) Figure 6.12: NbN concentration for 1200 o C/2hrs, 120 o C/12hrs; and the 1400 o C/3hrs, 120 o C/12hrs heat treatments as compared to a control sample High Temperature Heat Treatment with Lower Temperature Heating Two types of heat treatments were performed involving high temperature heat treatment with lower temperature heating. One was the 800 o C/3hrs-400 o C/20min heat treatment performed in vacuum conditions throughout. The other heat treatment involved high temperature heating, followed by a lower temperature baking during which N 2 gas was introduced in the 203

229 chamber for about 10-20min at 10-5 mbar pressure. Two of the samples were heat treated according to the latter method of heat treatment, with one of them being further baked at 120 o C/6hrs in the same experiment. N 2 was introduced with the goal of nitridation of the Nb surface to prevent the adherence of impurities (H,C,O) to the Nb surface once the furnace reached room temperature (see following section). (a) Carbon Figure 6.13 shows the C concentration for the 800 o C/3hrs, 400 o C/20min heat treated sample in comparison to a non heat treated (control) sample. There seems to be a factor of three increase in the C level after this heat treatment. This increase in the C concentration is less in comparison to the 800 o C/120 o C heat treated sample (section ) indicating that C is dissolving into the material less during this heat treatment, which is expected, since the longer baking step is not involved in this heat treatment. Figure 6.14 shows the C concentration comparison for the 800 o C/3hrs, 400 o C/15min N 2 (10-5 mbar), 120 o C/6hrs and the 1000 o C/2h, 800 o C/10min N 2 (10-5 mbar) heat treated samples to a non heat treated (control) sample. The 1000 o C/2h, 800 o C/10min N 2 (10-5 mbar) heat treated sample seems to have helped in reducing the C contamination in the material as seen in the figure, where a slight decrease in the C concentration is observed as compared to a control sample. However, the 120 o C extra bake in the latter heat treatment diffuses the carbon into the material, up to about 0.4 μm of the surface. Even though the introduction of N 2 might have helped in reducing surface contamination, the extra baking step appears to have incorporated C from the furnace atmosphere and diffused it into the material. 204

230 1E+20 C C oncentration (atoms/cm^3) 1E+19 1E+18 Control 800C/3h,400C/20min 1E Depth (um) Figure 6.13: C concentration for 800 o C/3hrs, 400 o C/20min heat treated sample as compared to a control Nb sample 1E+22 C Concentration (atoms/cm^3) 1E+21 1E+20 1E+19 1E+18 1E C/2h, 800C/10min N2 800C/3h, 400C/15min N2,120C/6h Control Depth (um) Figure 6.14: C concentration for 800 o C/3hrs, 400 o C/15min (N 2 at 10-5 mbar), 120 o C/6hrs; and the 1000 o C/2hrs, 800 o C/10min (N 2 at 10-5 mbar) heat treatments as compared to a control sample 205

231 (b) Nitrogen Figure 6.15 shows the N concentration comparison for the 800 o C/3hrs, 400 o C/20min heat treated sample to a non heat treated (control) sample, over a depth of about 2μm. The nitrogen levels are seen to decrease after this heat treatment, which is similar to the result observed in figure It can be only speculated that this decrease in nitrogen levels, especially for the extra baking step (section ) and extra heating step respectively, can be due to a surface layer formation which does not allow nitrogen to adsorb on the surface and reenter the material on furnace cool down. As mentioned in the introductory part of this section and in preceding chapters for hydrogen and oxygen, the introduction of N 2 in the furnace was utilized for nitridation of the Nb surface and subsequent reduction of impurity re-deposition from the furnace atmosphere on cool down. Thus, it is important to examine the specifications of these nitridation experiments 1E+20 N bn C oncentration (atom s/cm ^3) 1E+19 1E+18 Control 800C/3h,400C/20min 1E Depth (um) Figure 6.15: N concentration for 800 o C/3hrs, 400 o C/20min heat treated sample as compared to a control Nb sample 206

232 6.3.2 Nitridation Experiments One potential problem with the utilization of heat treatments for the removal of residual gas impurities from the Nb surface is that these gases may be reabsorbed by Nb upon cooling of the sample and subsequent exposure to air and water. A potential remedy for this problem had been proposed based on a theoretical calculation. This calculation indicated that a thin (10 nm) nitride layer would form on the Nb surface at some intermediate temperature (i.e. 400 o C) as a result of thermal diffusion of nitrogen into the Nb during the furnace cooldown. 21 The expectation was that this nitride layer would passivate the Nb surface and prevent hydrogen and oxygen absorption from the atmosphere. The nitride layer must be thin enough so as to not to change the superconducting properties of Nb significantly. This calculation, based on the solution of and diffusion of nitrogen into niobium, showed that it should be possible to grow a nitride layer about 10 nm thick on the Nb surface by admitting N 2 in the furnace at a partial pressure of about 5 x 10-6 Torr at 400 o C for 15 min. 20,21 In order to test the above hypothesis, two large grain (single crystal) Nb samples were heat treated in an attempt to provide nitridation of the surface. One of the samples was heat treated for 800 o C/3hrs and then cooled to 400 o C, at which point, N 2 gas was introduced into the furnace at a pressure of 10-5 mbar, and the sample was held at 400 o C for 15min and then cooled down. The other sample was heat treated using the same procedure, but was additionally baked for 120 o C/6hrs after the 400 o C/15 min hold in nitrogen atmosphere. SIMS analyses were performed on both these same samples via low energy Cs + primary ion beam (6keV impact energy) for improved depth resolution, using analysis conditions mentioned in chapter 3 (section 3.5.2). Figure 6.16 shows the N concentration over 500nm depth for both the 800 o C/3hrs with 400 o C/15min (N 2 at 10-5 mbar) and 800 o C/3hrs with 400 o C/15min (N 2 at 10-5 mbar) and 120 o C/6hrs heat treated samples along with data obtained from a control sample. 207

233 Control 800C/3h,400C/15min N2 800C/3h,400C/15min N2,120C/6h NbN Concentration (atoms/cm^3) 1E+20 1E+19 1E+18 1E Depth (um) Figure 6.16: NbN concentration for 800 o C/3hrs, 400 o C/15min (N 2 at 10-5 mbar) and 800 o C/3hrs, 400 o C/15min (N 2 at 10-5 mbar), 120 o C/6hrs heat treatments, as compared to a control sample; (SIMS 6keV Cs + ) As seen in the figure, nitrogen levels near the surface are not elevated, indicating a very low surface concentration of nitrogen. The results show conclusively that a nitride layer was not formed on the surface. It was thus speculated that a higher temperature may be needed for nitridation of the surface at 10-5 mbar. Subsequently, another large grain (single crystal) Nb sample was heated to a temperature of 1000 o C for 2hrs, held at 800 o C for 10min in 10-5 N 2 partial pressure and cooled down to room temperature. SIMS analysis was performed on this sample using the same conditions as mentioned in chapter 3, section (high energy SIMS) Figure 6.17 shows the results for this heat treatment as compared to a control sample over 600nm depth. This heat treatment confirmed the presence of a high nitrogen concentration on the surface (1% atomic), indicating that it was indeed possible form a nitrogen rich layer on the Nb surface at a temperature of about 800 o C. However, as seen in sections 4.4, 5.5 and fig. 208

234 6.14, the introduction of this nitride layer did not lead to any significant decrease in the impurity contamination (H, C, O) on the Nb surface. 1E+21 NbN Concentration (atoms/cm^3) 1E+20 1E+19 1E+18 Control 1000C/2h, 800C/10min N2 1E Depth (um) Figure 6.17: N concentration for 1000 o C/2hrs, 800 o C/10min (N 2 at 10-5 mbar) as compared to a control sample Conclusions SIMS analyses of C in Nb before and after various heat treatments suggest that there is some carbon contamination on the surface of some heat treated Nb samples, which may originate from the heat treatment furnace. The baseline C concentration is seen to increase after almost all types of heat treatments. SIMS analyses of samples processed to produce a nitrogen rich surface show that nitrogen can be incorporated into the Nb surface at 800 o C. However, this surface nitrogen did not reduce impurity contamination during cool down after heat treatment. 209

235 6.4 Low Energy SIMS Characterization of Carbon and Nitrogen in Niobium Low energy SIMS analysis was carried out to improve depth resolution to allow study the C and N concentrations in the first 40nm of large grain (single crystal) Nb sample surface in order to obtain more detailed information for these species over the depth of prime importance in terms of RF cavity performance. Analysis conditions were typical for low energy analysis, as shown in chapter 3, section Results and Discussion The following results are divided into the type of heat treatment performed on the samples which were: low temperature baking (120 o C/48hrs); high temperature heat treatments (1200 o C/6hrs), high temperature heat treatment with low temperature baking ( o C/3hrs, and 120 o C/12hrs), and high temperature heat treatment and low temperature heating in N 2 (800 o C/3hrs and 400 o C/15min N 2 at 10-5 Torr, 120 o C/6hrs) Long Term Low Temperature Baking (120 o C/48hrs) (a) Carbon Figure 6.18 shows the carbon concentration profile for the 120 o C/48hrs baked sample in comparison to a non heat treated (control) vs depth. As seen in the figure, the carbon concentration in the baked samples is up to two orders of magnitude higher for the 40-60nm region. It is clear that there is carbon contamination present at the near surface after baking. It is possible that this carbon impacts the cavity performance after this heat treatment (see chapter 7). The non heat treated sample itself has a high carbon level in the first 10nm, but C drops off rapidly to a significantly lower level as compared to the heat treated sample. 210

236 1E+21 C Concentration (atoms/cm^2) 1E+20 1E+19 1E C/48h Control 1E Depth (um) Figure 6.18: Concentration profile of C in Nb for 120 o C/48 hrs baked sample in comparison to a control sample. High C levels seen after baking (Low energy SIMS, 6keV Cs + ) 1E+20 NbN Concentration (atoms/cm^3) 1E+19 1E+18 Control 120C/48h 1E Depth (um) Figure 6.19: Concentration profile of NbN for 120 o C/48 hrs baked sample in comparison to a control sample. No difference observed after baking (Low energy SIMS, 6keV Cs + ) 211

237 (b) Nitrogen Figure 6.19 shows the N concentration profile for the 120 o C/48hrs baked sample in comparison to a non heat treated (control) sample over the first 100nm depth. As seen in the figure, there is no change in the nitrogen levels before and after baking, while the N concentration measured in both the samples is observed to be very low High Temperature Heat Treatment (a) Carbon Figure 6.20 shows the C concentration profile vs depth for the 1200 o C/6hrs heat treated sample as compared to a control sample. As seen in the figure, the C concentration is higher as compared to the control sample, possibly due to the contamination from the furnace. The surface carbon deposition from the furnace might have diffused into the material at this elevated temperature. (b) Nitrogen For figure 6.21, the nitrogen levels are observed to be higher over the first 20nm in the heat treated sample indicating that there is some incorporation of nitrogen during this heat treatment. Similar to the C contamination, nitrogen from the furnace atmosphere might have diffused into the material at this high temperature. This behavior is not observed in the high energy SIMS data (section ), due to the relatively poor depth resolution and the consequent limitation to study the first 20nm of the surface, due to less data points. 212

238 1E+21 C Concentration (atoms/cm^3) 1E+20 1E+19 1E+18 Control 1200C/6h 1E Depth (um) Figure 6.20: Concentration profile of C in Nb for 1200 o C/6 hrs heat treated sample compared with a control sample. High C levels seen after baking (Low energy SIMS, 6keV Cs + ) 1E+23 NbN Concentration (atoms/cm^3) 1E+22 1E+21 1E+20 1E+19 1E C/6h Control 1E Depth (um) Figure 6.21: Concentration profile of N for the 1200 o C/6 hrs heat treated sample compared with a control sample. (Low energy SIMS, 6keV Cs + ) 213

239 High Temperature Heat Treatment with Long Term Low Temperature Baking (a) Carbon Figure 6.22 shows the C concentration versus depth for the high temperature heat treated and low temperature baked 1200 o C/2hrs with 120 o C/12hrs and the 1400 o C/3hrs with 120 o C/2hrs samples along with C in a non heat treated sample (control). Again, it can be observed that higher the temperature of heat treatment, higher the carbon concentration near the surface possible, as a result of enhanced diffusivity of the surface- deposited carbon with temperature. 1E+21 Control 1200C/2h,120C/12h 1400C/3h,120C/12h C C oncentration (atoms/cm^3) 1E+20 1E+19 1E C/3h,120C/12h 1200C/2h,120C/12h Control 1E Depth (um) Figure 6.22: Concentration profile of C in Nb for 1200 o C/2hrs with 120 o C/12hrs and 1400 o C/3hrs with 120 o C/12hrs heat treated samples in comparison to a control sample (Low energy SIMS, 6keV Cs + ) (b) Nitrogen Figure 6.23 shows N levels in the 1200 o C/2hrs with 120 o C/12hrs and 1400 o C/3hrs, with 120 o C/2hrs heat treated samples compared to a non heat treated sample (control). Although the former heat treatment shows only slightly higher nitrogen levels as compared to the 214

240 control sample, the latter heat treatment shows a significant increase of nitrogen in the first 10nm of the sample surface. Furthermore, the shape of the N concentration profile matches the shape of the oxygen concentration profile observed for the same heat treatment (see fig.5.22), with a step in the profile at about 60nm depth, indicating that both oxygen and nitrogen are diffusing into the material during the 120 o C/12hrs baking, since this step is not seen for only the 1400 o C/3hrs heat treatment in both impurity profiles (see chapter 5) 1E+22 Control 1200C/2h,120C/12h 1400C/3h,120C/12h NbN Concentration (atoms/cm^3) 1E+21 1E+20 1E+19 1E C/3h,120C/12h 1200C/2h,120C/12h Control 1E Depth (um) Figure 6.23: Concentration profile of NbN for 1200 o C/2hrs with 120 o C/12hrs and 1400 o C/3hrs with 120 o C/12hrs heat treated samples in comparison to a control sample (Low energy SIMS, 6keV Cs + ) High Temperature Heat Treatment with Lower Temperature Heating (a) Carbon Low energy results for N for similar heat treatments have already been shown in section Figure 6.24 shows the C concentration profiles vs depth for the 800 o C/3hrs with 400 o C/15min (N 2 at 10-5 mbar) and 800 o C/3hrs with 400 o C/15min (N 2 at 10-5 mbar) and with 120 o C/6hrs heat treated samples along with the C concentration obtained from a control 215

241 sample. As seen in the figure, the 800 o C/3hrs with 400 o C/15min (N 2 at 10-5 mbar) heat treatment shows a very high amount of carbon (5-10% atomic) in the first 100nm of the surface, as compared to the control sample. This very high concentration of C has not been seen in any of the other heat treatments. The source of this heavy contamination is unclear. It is possible that heavy contamination from the furnace might be forming a carbide on the Nb surface, since this amount of carbon can only be present in the form of a carbide (see Nb-C phase diagram, fig. 6.2), at 800 o C temperature. 1E+22 Control 800C/3h,400C/15min,120C/6h 800C/3h,400C/15min C Concentration (atoms/cm^3) 1E+21 1E+20 1E+19 1E C/3h,400C/15min N2,120C/6h Control 800C/3h, 400C/15min N2 1E Depth (um) Figure 6.24: C concentration for 800 o C/3hrs with 400 o C/15min (N 2 at 10-5 mbar) and 800 o C/3hrs with 400 o C/15min (N 2 at 10-5 mbar) and with 120 o C/6hrs heat treatments, as compared to a control sample; (SIMS 6keV Cs + ) Conclusion It has been observed using low energy SIMS analysis of C and N in Nb, before and after various heat treatments that almost all the heat treatments lead to an increase in the carbon concentration on the Nb surface. High levels of contamination are seen for the 120 o C baking, 216

242 o C heat treatments, and even higher C contamination is seen for the 800 o C/3hrs with 400 o C/15min (N 2 at 10-5 mbar) heat treatment. A high concentration of N is observed for the 1400 o C/3hrs with 120 o C/12hrs heat treatment, with a matching profile to the oxygen concentration for the same heat treatment. This is believed to be caused by the extra baking step, since the 1400 o C/3hrs heat treatment alone, does not show a step in the concentration levels for both impurities. 6.5 TOF-SIMS Imaging of Carbon in Large Grain Nb Bicrystals Grain boundaries in fine grained Nb often contain high levels of impurities. 20 Even though the total volume of grain boundaries are smaller in large grain Nb, interstitial impurities may localize at the grain boundary region which can have a deleterious impact on cavity performance. Thus, it is important to study the grain boundaries of large grain Nb for any C segregation both before and after heat treatment. As described in chapter 3, section 3.6.1, two types of Nb bicrystals were prepared for surface imaging at the grain boundary, bicrystals having different crystallographic orientation combinations and bicrystals having the same crystallographic orientation combinations. Details of sample preparation and crystallographic orientation determination are provided in section Surface imaging was performed at the grain boundary using an ION-TOF V TOF SIMS instrument (chp 2) with analysis conditions mentioned in chapter 3, section Results and Discussion Bicrystals having different crystallographic orientation combinations One of the samples in this set was heat treated for 800 o C/3hrs, 120 o C/24hrs using the Elnik resistive heating furnace (section 3.6.1). Section shows the EBSD data for the control and heat treated bicrystals having different crystallographic orientation combinations. Notably, the upper crystal had a (001) orientation for both the control and heat treated bicrystals, while the lower crystal orientation was (213) and (215) for control and heat treated samples respectively. Fig shows the C - secondary ion images for the control (a) and the heat treated (b) samples. It can be seen in the control sample image that there is no segregation of C at the grain boundary. However, the results for the heat treated sample clearly show a high C level 217

243 at the grain boundary indicating segregation of carbon in the region. The regions around the grain boundary area are depleted of C, indicating that C appears to have diffused to the grain boundary during heat treatment with the result that C is depleted in the near grain boundary regions. It should be noted that the small intensity difference across the interface for both samples may be the result of sputtering rate differences due to the different crystal orientations. (001) Grain Boundary (001) (213) (215) (a) (b) Figure 6.25: TOF SIMS ion images of C - for control (a) and heat treated (800 o C/3hrs with 120 o C/24hrs) (b) bicrystal sample grain boundaries Bicrystals having the same crystallographic orientation combination In order to account for the effect of crystal orientation on the secondary ion intensities and obtain a better understanding of the heat treatment effect on the grain boundaries of large grain Nb bicrystal Nb samples, a separate set of samples was characterized using TOF-SIMS imaging under the same experimental conditions. However, to eliminate any orientation effects, a large sample was sectioned from the same grain boundary area and cut in half. One part of this large sample was heat treated to 1200 o C/6hrs in the new furnace after BCP (1:1:2, 7μm removal) while the other part was not heat treated after the same BCP etching (control). Fig shows the TOF SIMS images of the control (a) and heat treated (b) Nb samples. While the upper and lower crystal orientations in the samples differ, the two samples have 218

244 identical crystal orientations. In the figure, the grain boundary region is not clearly visible for both samples, possibly due to no measurable change in the intensities of the impurity at the grain boundary, which indicates no segregation of C both before and after heat treatment. This heat treatment was carried out in the induction furnace obtained by Jefferson Labs, and it is possible that there was considerably less C contamination during heat treatment, as compared to the previous heat treatment, shown in figure (hkl) Grain Boundary (hkl) (h k l ) (h k l ) (a) (b) Figure 6.26: TOF SIMS ion images of C - for control (a) and heat treated (1200 o C/6hrs) (b) bicrystal sample grain boundaries with similar crystal orientation of the upper crystal for both samples and lower crystal for both samples Conclusion TOF SIMS analysis of the grain boundary region of large grain bicrystal Nb shows that C levels are higher and C is segregated at the grain boundaries for the 800 o C/3hrs with 120 o C/24hrs heat treated samples, possibly due to contamination from the furnace. However, subsequent analysis performed on the grain boundary area of a sample with similarly oriented Nb bicrystals shows that there is no segregation of C for both control and the 1200 o C/6hrs heat treated samples. This result is possibly due to the heat treatment being performed in a new furnace and with less carbon contamination on heat treatment. 219

245 References 1. H. Padamsee et al : RF Superconductivity for Accelerators, 2 nd edition, Wiley New York (2008) 2. W. DeSarbo: Phys. Rev., 132 (1), p107, G.Hardy et.al: Phys. Rev., 93, p1004 (1954). 4. A. Giorgi et.al: Phys. Rev, 125, p857 (1952) 5. B. Matthias et.al: Phys. Rev., 87, p99 (1952). 6. M.L. Pochon et. al: Reactive Metals, 2, Interscience Publishers, New York, p327 (1959) 7. R. Elliott et.al: Trans. ASM, 53, p13. (1961) 8. G. Horz et.al.: Journ. Less-Comm. Met., 35, p97 (1974). 9. H. Kimura et.al: Trans. Japan Inst. Met., 2, p98 (1961) 10. E. Gebhardt et. al.: Z. Metallkunde., 57, p682 (1966) 11. R.Powers et. al.: J. Metals Trans., 9, p1285 (1957) 12. J. Smith et. al.: Journ. Nuc. Met., 148, p1 (1987) 13. H. Okamoto: ASM International, Binary Alloy Phase Diagrams, 2 nd ed., 3, Materials Park, Ohio, U.S.A, p2690 (1990) 14. C. Ang et. al: Journ. Met. AIME, 197, p1035 (1953) 15. J. Cost et. al: Act. Mett., 11, p231 (1963) 16. C. Gupta et. al: Extractive Metallurgy of Niobium, CRC Press (1994) 220

246 17. C. Ang: Act. Mett.,1, p123 (1953) 18. R. Powers et. al: Journ. App. Phys., 30, p514 (1959) 19. D. Porter et. al: Phase Transformations in Metals and Alloys, CRC Press, Boca Raton, Florida, USA (1981) 20. G. Ciovati et. al: Phys. Rev. Special Topics: Acc. and Beams, 13, p22002 (2010) 21. P. B. Wilson: SLAC Technical Note No. SLAC-TN-71-7 (unpublished) (1971) 221

247 7. Cavity Performance Results, Conclusions and Future Work 7.1 Cavity Performance An important aspect of this study is to correlate the SRF cavity performance for various heat treatments with the results obtained using SIMS. SIMS provides impurity characterization in Nb in the first 40nm depth region of the surface (see preceding chapters) Technique for Cavity Performance Measurements 1 The cavity testing process was performed at Jefferson Laboratories. The most important figure of merit in evaluating the RF performance of the cavity is its intrinsic quality factor (Q o ) and it is important to understand the procedure to obtain this quality factor. 1 To excite the resonant mode of a cavity, it needs to be connected to an RF source. In practice, a setup such as the one shown in figure 7.1 might be used. Figure 7.1 : Cavity with RF probes 1 222

248 The power from the RF source is carried to the cavity via a coaxial cable. This center conductor protrudes into the beam tube of the cavity, which forms the outer conductor. The strength of this input coupler is adjusted by changing the penetration of the center conductor. The fixed output coupler is also called a transmitted power probe since it picks up power transmitted through the cavity. After the RF source is turned off, the total power being lost is the power dissipated in the cavity walls (ohmic losses) and the power that leaks out from each coupler: P tot = P c + P e + P t..(7.1) Where P c is the cavity wall power dissipation, P e is the power leaking out of the input coupler, and P t is the transmitted power. Since, Q tot can be analogous to the Q o (see chp 1): Q tot = ωu/p tot..(7.2) where U is the stored energy, which satisfies the relation: du/dt = -P tot = -ωu/q tot...(7.3) If there is no field emission or other anomalous losses, the solution to this differential equation is simple: U = U o exp(-ωu/q tot )...(7.4) The energy in the cavity thus decays exponentially with a time constant: τ tot = Q tot /ω..(7.5) and by measuring the decay time, the total or loaded quality factor of the cavity can be obtained. And equation 6.1 can be rewritten: ωu/q tot = ωu/q o + ωu/q e + ωu/q t Or: 1/Q tot = 1/Q o + 1/Q e + 1/Q t.(7.6) And thus, the Q o of the cavity can be obtained. 223

249 7.1.2 Cavity Performance Results Table 7.1: Cavity performance results for cavities heat treated with the Nb samples; next to the Q o values (in parentheses) is the magnetic field for the values reported. All measurements taken at 2 K temperature. 224

250 Table 7.1: Continued Table 7.1 shows the cavity performance values (Q o ), with some typical operating magnetic fields used at Jefferson Labs, in parentheses, for cavities which were heat treated with the large grain Nb samples shown in previous sections of this work Hydrogen: Primary Factor in Cavity Performance For all high temperature heat treatments (600 o C-1400 o C), it is observed via SIMS analysis that there is considerable decrease in the hydrogen levels (up to a factor of 100), as seen in chapter 4. For the first 40nm of the surface, this large reduction of hydrogen is believed to be the primary cause for the performance improvement of the cavities. Specifically, the 77% increase in the cavity performance for the 800 o C/3 hrs, 400 o C/20 min heat treated cavity can be attributed to the very low hydrogen levels seen in the samples heat treated with the cavity, using SIMS, since the other impurities (C, O, N) show insignificant differences, as measured by SIMS. Similar observations can be made for almost all the high temperature heat treated cavities, barring some heat treatments, which along with hydrogen also show some differences in C and O in the first 40nm of the Nb surface. For example the record 200% increase in the cavity efficiency for the 1400 o /3hrs heat treatment cannot be explained by the 225