Microstructural and morphological aspects of plutonium hydride
A thesis submitted to the University of Manchester for the degree of Nuclear EngD in the Faculty of Engineering and Physical Sciences
2015
Martin Brierley
School of Materials
Contents
Contents ...... 2
Table of figures ...... 10
Word count ...... 16
Glossary ...... 17
Abstract ...... 18
About the Author ...... 19
Declaration ...... 19
Copyright statement ...... 20
Personal Statement ...... 21
Acknowledgements ...... 22
Chapter 1 Introduction ...... 23
1.1 Aims of the research ...... 24
1.2 Introduction to plutonium ...... 25
1.2.1 Origins of plutonium ...... 25
1.2.2 Isotopes and radioactivity ...... 26
1.2.3 Fission ...... 27
1.2.4 Allotropes of plutonium ...... 28
1.2.4.1 Alpha plutonium ...... 29
1.2.4.2 Delta plutonium ...... 30
1.2.5 Transition between phases ...... 32
1.2.6 Chemical reactivity ...... 33
1.2.7 Uses of plutonium ...... 35
1.2.8 Why is the corrosion of plutonium important? ...... 35
1.2.9 Health ...... 36
1.2.10 Pyrophoricity ...... 38
1.2.11 Environment ...... 39
1.3 Introduction to plutonium oxide ...... 40
1.3.1 Rate of reaction with plutonium metal ...... 43
1.3.2 Stability ...... 43
1.4 Plutonium hydride ...... 44
1.4.1 Rate of reaction with plutonium metal ...... 48
1.4.2 Stability of the hydride phases ...... 49
1.4.3 Sources of hydrogen ...... 49
Chapter 2 Literature review ...... 50
2.1 Historical research into plutonium hydride ...... 51
2.1.1 Initial studies ...... 51
2.1.2 Solid solubility of hydrogen in plutonium and plutonium hydride 52
2.1.3 Hydriding kinetics ...... 52
2.1.4 Activation energy ...... 53
2.1.5 Pyrophoricity ...... 53
2.1.6 Hydride growth behaviour ...... 54
2.1.7 Hydride nucleation ...... 55
2.1.8 Continued growth ...... 56
2.2 Hydriding work undertaken on other metal-hydride systems ...... 56
2.2.1 5f elements (actinides) ...... 57
2.2.1.1 Uranium ...... 57
2.2.2 4f elements (lanthanides)...... 60
2.2.2.1 Cerium ...... 60
2.2.2.2 Gadolinium ...... 60
2.2.2.3 Holmium...... 61
2.2.3 Transition metals ...... 61
2.2.3.1 Zirconium and zirconium alloys ...... 61
2.2.3.2 Titanium ...... 62
2.3 Models of the hydriding process ...... 63
2.3.1 Nucleation models ...... 63
2.3.2 Growth models ...... 65
2.3.3 General models ...... 66
2.3.4 Comparison of models ...... 67
Chapter 3 Experimental ...... 68
3.1 Background ...... 69
3.2 Cerium experiments ...... 69
3.2.1 Material details ...... 69
3.2.2 Pre-hydriding sample preparation ...... 70
3.2.2.1 Sawing ...... 70
3.2.2.2 Cleaning ...... 71
3.2.2.3 Surface preparation ...... 71
3.2.3 Exposure to hydrogen ...... 71
3.2.3.1 Gas handling rig ...... 71
3.2.3.2 Reaction cell ...... 72
3.2.3.3 The hydriding reaction ...... 72
3.2.4 Post hydriding sample preparation ...... 73
3.2.5 Light microscopy ...... 74
3.2.6 Secondary Ion Mass Spectrometry ...... 74
3.2.7 Scanning Electron Microscopy ...... 76
3.2.7.1 Secondary Electron Imaging ...... 77
3.2.7.2 Backscattered Electron Imaging...... 78
3.2.7.3 Energy Dispersive Spectroscopy ...... 78
3.2.8 Sample coating ...... 79
3.2.9 Ion milling ...... 79
3.2.10 Vickers hardness ...... 80
3.2.11 Nanoindentation ...... 81
3.2.12 Atomic Force Microscopy...... 86
3.3 Plutonium experiments ...... 88
3.3.1 Preparation for hydriding reaction ...... 88
3.3.2 Surface analysis preparation ...... 90
3.3.3 Cross sectional analysis preparation ...... 90
3.3.4 Sample coating ...... 91
3.3.5 Ex-situ hydriding ...... 91
3.3.6 Scanning electron microscopy for plutonium ...... 92
3.3.6.1 Analysis modes ...... 94
3.3.6.2 Sample transfer to the FEGSEM ...... 95
3.3.6.3 Ion milling ...... 95
3.3.6.4 In-situ hydriding ...... 96
Chapter 4 Philosophy for the research ...... 99
4.1 Overall aims ...... 100
4.1.1 Does the overall shape of the hydride evolve isotropically? .... 100
4.1.2 Does the microstructure provide evidence of growth characteristics? ...... 100
4.1.3 Is the hydride morphology directed by strain introduced into material surrounding a dilated hydride? ...... 100
4.2 Experimental philosophy ...... 101
4.3 Experimentation on cerium ...... 104
4.3.1 Aims ...... 104
4.3.2 Initial work ...... 105
4.3.3 Microstructure of the cerium hydride reaction sites ...... 106
4.3.4 Nanoindentation of the cerium/cerium hydride interface ...... 109
4.4 Experimentation on plutonium ...... 109
4.4.1 Active commissioning of the FEGSEM ...... 109
4.4.2 The morphology and anisotropic growth of plutonium hydride reaction sites 112
4.4.3 The reaction between hydrogen and electro-refined plutonium observed by in situ electron microscopy ...... 113
4.4.4 In situ hydriding of mixed phase α/δ-Pu ...... 115
Chapter 5 The morphology and anisotropic growth kinetics of cerium hydride reaction sites 118
5.1 Cover Page ...... 119
Chapter 6 The microstructure of cerium hydride growth sites ...... 126
6.1 Cover Page ...... 127
Chapter 7 Probing the cerium/cerium hydride interface using nanoindentation 136
7.1 Cover Page ...... 137
Chapter 8 The anisotropic growth morphology and microstructure of plutonium hydride sites ...... 142
8.1 Cover Page ...... 143
8.2 Title ...... 144
8.3 Author names and affiliations ...... 144
8.4 Abstract ...... 144
8.5 Keywords ...... 144
8.6 Introduction ...... 145
8.7 Method ...... 146
8.7.1 Hydriding reaction ...... 146
8.7.2 Scanning Electron Microscopy ...... 147
8.8 Results ...... 147
8.8.1 Hydriding experiment ...... 147
8.8.2 Surface Characterisation ...... 147
8.8.3 Cross-sectional analysis ...... 149
8.9 Discussion ...... 151
8.9.1 Accommodation of hydride dilation ...... 151
8.9.2 Growth anisotropy ...... 152
8.9.3 Hydride/metal interface ...... 153
8.9.4 Mixed phase hydride ...... 153
8.9.5 Grain boundaries ...... 154
8.9.6 Crystallite orientation ...... 154
8.9.7 Triangular Projections ...... 154
8.9.8 Discontinuous interface mechanism ...... 155
8.10 Conclusions ...... 156
8.11 Acknowledgements ...... 157
8.12 References ...... 157
Chapter 9 The reaction between hydrogen and electro-refined plutonium observed by in situ electron microscopy ...... 165
9.1 Cover Page ...... 166
9.2 Title ...... 167
9.3 Author names and affiliations ...... 167
9.4 Abstract ...... 167
9.5 Introduction ...... 167
9.6 Experimental ...... 168
9.7 Results and discussion ...... 170
9.7.1 In situ surface examination ...... 170
9.7.2 Post experimental cross section ...... 172
9.8 Conclusions ...... 176
9.9 Acknowledgements ...... 177
9.10 References (see main references section) ...... 177
Chapter 10 In situ growth of hydride reaction sites on an alpha/delta mixed phase plutonium alloy ...... 178
10.1 Cover Page ...... 179
10.2 Title ...... 180
10.3 Author names and affiliations ...... 180
10.4 Abstract ...... 180
10.5 Introduction ...... 180
10.6 Experimental ...... 181
10.6.1 In-FEGSEM hydriding reaction ...... 182
10.6.2 Material ...... 182
10.6.3 Sample preparation ...... 182
10.7 Results ...... 183
10.7.1 Sample state before hydrogen exposure ...... 183
10.7.2 In situ hydriding ...... 185
10.7.3 Post hydriding cross sections ...... 190
10.8 Conclusions ...... 193
10.9 Acknowledgements ...... 193
10.10 References ...... 194
Chapter 11 Overall discussion ...... 196
11.1 Does the overall shape of the hydride evolve isotropically? ...... 198
11.2 Does the microstructure provide evidence of growth characteristics? 200
11.2.1 Cerium ...... 200
11.2.2 Electro-refined plutonium ...... 201
11.2.3 Delta plutonium ...... 203
11.2.4 Mixed α/δ phase plutonium ...... 204
11.3 Is the hydride morphology directed by strain introduced into material surrounding a dilated hydride? ...... 206
11.4 Hydride reaction site growth mechanism ...... 210
11.5 General observations ...... 211
11.5.1 On passivation ...... 211
11.5.2 Sites of nucleation ...... 214
11.6 Future work ...... 217
Chapter 12 Conclusions ...... 218
References ...... 221
Table of figures
Figure 1 The alpha plutonium lattice. The unit cell is shown by the box. Created using Mercury 3.6 with data from the Crystallography Open Database [26]...... 30
Figure 2 The low alloy Pu-Ga phase diagram [35] ...... 31
Figure 3 The face-centred cubic delta plutonium unit cell. Created using Mercury 3.6 using data from the Crystallography Open Database [26]...... 32
Figure 4 Dilatometry data for pure plutonium indicating the large changes in volume compared to that of aluminium over the same temperature range [8]...... 33
Figure 5 Overlapping radial distribution functions of the 7s and the 5f orbitals for plutonium [25]. 34
Figure 6 The effect of plutonium in the body – fission tracks which have occurred due to the radioactive decay of plutonium at the bone surface are evident as short black lines [58] 36
Figure 7 A video still recorded during exposure of the vigorous reaction that occurs when a plutonium sample covered in hydride formed at room temperature to oxygen. 39
Figure 8 Schematic of how the surface of plutonium oxidises in an O2 environment, and how an oxidised plutonium surface auto converts from PuO2 to Pu2O3 in UHV conditions [69]...... 40
Figure 9 The equilibrium phase diagram for Pu-O [72] ...... 41
Figure 10 The CaF2 structure of PuO2 with the plutonium atoms as the larger blue spheres. For the type C Pu2O3 structure, 25 % of the red spheres would be missing. Created using Mercury 3.6 using data from the Crystallography Open Database [26]...... 42
Figure 11 Standard Chemical Picture of Plutonium Oxidation in Dry Air [80] ...... 43
Figure 12 PCT traces of plutonium hydrided to H/Pu ratios above the PuH2 stoichiometry [86] . 45
Figure 13 Top: The Pu-H diagram of Ellinger et al. [23] . Bottom: The equilibrium phase diagram for Pu-H [63]. Phase I is a mixed phase with a PuH2 and hydrogen in solution in the plutonium metal, Phase II is the PuH2 phase, which adopts the fluorite structure and
appears to retain some solubility for hydrogen around stoichiometric PuH2. Phase III is a mixed phase between PuH2 and PuH3, which arises as the tetrahedral fcc lattice sites become filled. Phase IV occurs at around H/Pu of 2.75 which corresponds to complete filling of the fcc lattice, which given sufficient temperature, forces a change to a hexagonal structure...... 46
Figure 14 Crystal structure of PuH2, large spheres are plutonium, the black spheres are occupied tetrahedral sites [95] ...... 48
Figure 15 The cubic crystal structure of PuH3 , white spheres are occupied octahedral sites and black spheres are occupied tetrahedral sites [95] ...... 48
Figure 16 A typical rate vs. time curve for plutonium hydriding. The four regions identified refer to different steps in the hydride reaction process: (1) induction, (2) nucleation or acceleration, (3) bulk hydriding, and (4) termination [90]...... 54
Figure 17 Microstructure of the cerium metal used in the studies...... 70
Figure 18 A Process & Instrumentation Diagram for the cerium hydriding rig...... 72
Figure 19 Diagram of a double focussing magnetic sector analyser...... 75
Figure 20 The interaction volume of electrons in a bulk sample and generation of signal in an SEM. 77
Figure 21 Indentations made in cerium metal (left) and cerium hydride (right) ...... 81
Figure 22 Top: diagram indicating three possible lengths that could be used to calculate the spring constant of the cantilever; the overall length (A), the length to the tip of the indenter (B), the freely deflecting length (C). Bottom: measurement of the freely deflecting portion of the cantilever...... 84
Figure 23 A 10 x 10 grid of increasing nanoindenter load (r to l) from 0.5 to 2.75 V at steps of 0.25 V. The load increased right to left...... 85
Figure 24 Power fit of the data plotted for the indentation depths and apparent hardness. 86
Figure 25 Distortion from the punching process extends a distance of around 200 µm into the surface...... 89
Figure 26 A high resolution XRD trace of the region that contains the main PuO2,
Pu2O3 and δ-Pu peaks. Also visible in the as-polished trace is an α-Pu peak. A number of
thermal treatments were subsequently carried out in situ to assess the stability of the various phases present. 90
Figure 27 A schematic diagram of the plutonium hydriding rig. The majority of the system is housed within a nitrogen filled glove box described by the orange rectangle...... 92
Figure 28 Top: The plutonium capable FEGSEM. Bottom: A simplified three dimensional representation of the FEGSEM...... 93
Figure 29 The Process & Instrumentation Diagram for the FEGSEM gas dosing system 97
Figure 30 The pressure-temperature phase diagram for cerium [164] ...... 103
Figure 31 Top: an optical image of a hydride reaction site cross section the extent of the hydride is the yellow area and would have also been the gray area. Bottom: an EDS oxygen map (green) of the same hydride reaction site overlaid onto the backscattered electron image. The sample was re-polished between the two images...... 108
Figure 32 Commissioning results of the ion mill on the FEGSEM used on plutonium. Top: a pillar masked by a particle on the surface allowed the etched depth to be measured after 130 min. Bottom: etched crater with grain detail revealed after 370 min. .... 111
Figure 33 Cracking and disintegration of etch pillars following a high temperature oxidation at 445 °C...... 112
Figure 34 An area of the α-Pu matrix with a high density of micro-cracks...... 114
Figure 35 The α/δ microstructure of the as cast 0.3 wt% alloy...... 116
Figure 36 a) A top-down view of a plutonium hydride reaction site. b) The structure of the central core of each site consisted of multiple platelets aligned towards the centre (indicated by cross). c) Radial fractures were found within the core, in all cases this followed the radially oriented platelet boundaries...... 159
Figure 37 a) The surface oxide layer immediately surrounding the hydride reaction site was observed to have been cracked by the dilation of the hydride site. b) The cracks are oriented in the radial and circumferential directions and extend beyond the edge of the hydride site. 160
Figure 38 a) Subsurface structure of a plutonium hydride reaction site revealed by ion beam milling. b) A magnified section (red box in a) of the hydride-metal interface...... 161
Figure 39 Cross section of a typical plutonium hydride reaction site imaged using backscattered electrons. The sites were found to inhabit an oblate shape in cross section. Each site had a discontinuous interface with the parent plutonium metal...... 161
Figure 40 a) Growth of hydride into the plutonium bulk following grain boundaries (identified in black), b) infill of the grains occurs by successive growth of aligned linear hydride features, c) preferred growth into one grain compared to its neighbour was observed. 162
Figure 41 A number of triangular projections were found along the interface between plutonium and plutonium hydride have a high aspect ratio and consist of an inner region of aligned platelets surrounded by a tapered sheath of non aligned material. Grain boundaries (GB) and a triple point junction (TP) are identified in black...... 163
Figure 42 A suggested schematic of the hydriding process: (1) Incubation, (2) Early growth, (3) Oxide failure, (4) Continued growth...... 163
Figure 43 A schematic of the growth processes: (1) an oblate hydride reaction site surrounded by plutonium crystallites, (2) grain boundary precipitation, (3) plane slip and aligned hydride growth along slip lines and (4) stress induced triangular projections...... 164
Figure 44 Microstructure of the electro-refined plutonium ...... 168
Figure 45 A three dimensional representation of the field emission gun scanning electron microscope arrangement...... 169
Figure 46 Secondary electron images of the reacted sample showing (a) surface cracks, spalled regions and extensive hydride coverage, (b) fine detail of the hydride product in a spalled region. 171
Figure 47 Secondary electron image of a typical hydride protrusion associated with a micro crack in the plutonium...... 172
Figure 48 low concentration PCT curves from unalloyed Pu. The lower four traces are from the δ-Pu temperature range [87]...... 174
Figure 49 Cracking was visible throughout the hydride product with some cracks in close proximity to the hydride/metal interface...... 175
Figure 50 Backscatter electron image of a cross section through a surface protrusion reveals hydride had formed within the micro cracks intersecting the surface. .... 176
Figure 51 (a) Back scatter electron image showing the microstructure of the mixed phase alloy, the denser α-Pu domains have a brighter contrast. (b) EPMA WDS map of Ga distribution in the sample (different area to a) showing distinct Ga segregation between the grain centres and grain boundaries...... 184
Figure 52 SEI micrographs showing (a) a typical thin sheet grain boundary precipitates ca. 100 nm thick at a α/α grain boundaries and (b) a typical cubic precipitate found in δ-domains. 185
Figure 53 Secondary electron image of a plutonium hydride particle which shows the material is comprised of agglomerated needles of plutonium hydride ca 160 nm wide with a common orientation...... 187
Figure 54 Secondary electron image showing hydride formation within an alpha region and brittle fracture of the alpha phase close to the interface...... 188
Figure 55 Backscatter electron images showing the distortion of δ-domains at the hydride/metal interface ...... 189
Figure 56 Backscatter electron images showing a) a region of mixed phase surface adjacent to the hydride interface, b) a detailed view of the black ringed area in (a) which shows distortion in the δ-domains and fracture in the α-domain...... 190
Figure 57 Secondary electron images of the hydride/metal interface at the metal surface. 191
Figure 58 Backscatter electron images of cross sections through the same site, (a) highlighting the oblate morphology as the radius (r) is greater than the depth (d) and (b) detailing of the reaction interface...... 192
Figure 59 Advanced precipitation of hydride at a sheet-like α/α grain boundary precipitate, here labelled GBP...... 193
Figure 60 Top: an unusually large spherical hydride which appears to have formed below a large, flatter surface hydride. Bottom: a typical oblate cerium hydride in cross section showing a two layered oxide/hydride structure of an air-exposed hydride...... 199
Figure 61 An edge-on view of a spalled piece of surface oxide showing a change in morphology at around 2.2 µm depth...... 202
Figure 62 Transformation of δ-Pu to α’-Pu beneath polishing scratches near to grain boundaries with lower gallium concentration in a partly homogenised sample...... 207
Figure 63 a) The radial compressive stresses (depicted in red) and the tangential tensile stresses (green) that surround the hydride at the surface. b) The extent of the stressed zone is anisotropic because of the ability of the forces normal to the surface to push the hydride outwards reduces the resultant force...... 208
Figure 64 Laths surrounding a cerium hydride reaction site...... 209
Figure 65 Three hydride site cross sections of different sizes; site a) ca. 400 µm diameter, site b) ca. 304 µm diameter, site c) ca. 141 µm diameter; each with a crust extending ca. 50% of the depth of the hydride...... 213
Figure 66 Cracks observed in the cerium metal often had hydride reactions sites nucleated along them ...... 215
Figure 67. Backscattered electron image of an iron rich inclusion seen in the cross section of an oxidised cerium hydride reaction site. The inclusion was analysed at higher magnification using EDS (red overlay bounded by the black rectangle), the optical image (bottom right) shows the remaining hydride as a yellow tinted area, indicated on the main image as the purple overlay, the green tint is the EDS oxygen map...... 216
Word count
Word count chapters 1-4, 8-12, references: 52907
Word count chapter 5: 4383
Word count chapter 6: 5984
Word count chapter 7: 2719
Total Word count: 65993
Glossary
AFM Atomic Force Microscopy AWE Atomic Weapons Establishment bcc Body centred cubic BEI Backscattered Electron Imaging CCD Charge Coupled Device CF Conflat DHC Delayed Hydride Cracking EBSD Electron Back-Scatter Diffraction EDS Energy Dispersive Spectroscopy EngD Engineering Doctorate EPMA Electron Probe Micro Analysis fcc Face centred cubic FEGSEM Field Emission Gun Scanning Electron Microscope FIB Focussed Ion Beam hcp Hexagonally close packed HYDEC Hydride/Dehydride/Casting process HYDOX Hydride/Oxidation reaction IAEA International Atomic Energy Agency IR Infra-red NTP Normal Temperature and Pressure PCT Pressure-Composition-Temperature PECS Precision Coating and Etching System SEI Secondary Electron Imaging SEM Scanning Electron Microscopy SiC Silicon Carbide SIMS Secondary Ion Mass Spectrometry SPL Surface Passivating Layer STV Sample Transfer Vessel UHV Ultra High Vacuum WDS Wavelength Dispersive Spectroscopy XRD X-ray diffraction
Abstract
The University of Manchester Martin Brierley Nuclear EngD Microstructural and morphological aspects of plutonium hydride September 2015
Plutonium is a hazardous radioactive material; the α-particles that are emitted are particularly damaging to health should contamination be inhaled or ingested into the body. During long term storage a number of conditions have been observed which can cause plutonium to corrode, which liberates particles from the surface. It is imperative to understand the processes involved in the corrosion of plutonium during long term storage to predict the likely state that metallic pieces may be found should subsequent handling be required. The growth mechanisms of plutonium hydride beyond the nucleation stage are not well understood. Detailed characterisation of the microstructural features associated with hydride reaction sites is required to develop a mechanistic understanding of the growth stage of hydrogen corrosion. Suitable processes and analysis methods were developed using cerium as an analogous material to δ-plutonium; during this stage, the knowledge of the corrosion of cerium by hydrogen was significantly improved using in situ gas dosing equipment, metallographic preparation, light microscopy, scanning electron microscopy (SEM), ion milling, secondary ion mass spectrometry (SIMS), atomic force microscopy (AFM) and vacuum nanoindentation. SEM and ion milling methods developed on cerium were subsequently used on the analysis on pre-formed and passivated hydride reaction sites on δ-plutonium .In situ exposure of electro-refined plutonium and a Pu 0.3 wt% Ga alloy were investigated without prior exposure to oxygen, revealing the as-formed microstructure of the hydride reaction product to be analysed. Subsequent metallographic preparation was used to confirm findings from the in situ analysis. The highest resolution analysis of the hydride product formed on cerium, delta plutonium and electro-refined plutonium has been obtained to date. Hydride reaction sites formed on cerium and δ-Pu were observed to be oblate, confirming growth anisotropy. A mechanism for the anisotropic growth was proposed where the stress fields introduced into the metal surrounding a lower density hydride play a significant role in further development of a hydride reaction site, causing failure of the surface oxide diffusion barrier surrounding a hydride reaction site.
About the Author
The author graduated from the University of Liverpool in 1998 with a 2.2 in Physics before briefly working at Didcot power station in the drawing office. A move in careers followed, becoming a development scientist at Printable Field Emitters Ltd (PFE), based at the Rutherford Appleton Laboratories. During the time at PFE, the author was involved in materials I-V testing and analysis, specialising in Scanning Electron Microscopy (SEM) and Atomic Force Microscopy (AFM). In 2002, the author commenced a part time Master’s degree In Advanced Materials at the University of Surrey, choosing to study the absorption of hydrogen by pressed pellets of palladium powder. A merit was obtained for the Master’s degree in 2008. During the Master’s, in 2004, the author moved to work at the Atomic Weapons Establishment (AWE) as a plutonium metallurgist and quickly became recognised as an expert in SEM and AFM techniques. The opportunity arose to undertake a Nuclear EngD in 2011 while at AWE. The remainder of the experience has been in the pursuit of this EngD.
Declaration
No portion of the work referred to in the thesis has been submitted in support of an application for another degree or qualification of this or any other university or other institute of learning.
Copyright statement
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Personal Statement
I dedicate this to my children, Rowan and Jasper. They gave me focus when I needed it and despite being only eight and six whilst I was writing up the thesis, they were able to understand that sometimes I needed to be given enough space to collect my thoughts. They are two very inquisitive children and have always been interested in the pictures they would see on screen while I was writing papers for the EngD. I hope that our discussions have inspired them to be the intelligent people that I know they are, and to achieve whatever they set out to do.
I love you both dearly.
I began to take an interest in amateur astronomy at a similar time to starting the EngD and throughout this course I have often mused about the fact that, quite accidentally, I have ended up studying the two elements (cerium and plutonium) that were named after planets which have subsequently lost planet status These planets were named Ceres and Pluto after characters from Greek mythology. Ceres was downgraded to an asteroid a long time ago and both it and Pluto were reclassified as dwarf planets in 2006 by the International Astronomical Union (IAU).
Interestingly I also studied the reaction of hydrogen with palladium for my Master’s thesis, named after the astronomical body Pallas, which also used to be classed as a planet but remains unlikely to be reclassified as a dwarf planet.
The last few months have been doubly interesting because I as I write up my research on how hydrogen interacts with the surfaces of these elements in fine microstructural detail; space probes are on their way to Ceres and Pluto, which have at the time of writing started sending detailed the most detailed images back to earth.
Martin Brierley
Acknowledgements
Funding This work was funded by AWE plc. AWE do an excellent job of developing their staff to the highest levels, long may this ethos continue. Many thanks go to Dave Geeson for initially suggesting the idea and helping me to set up the initial area of research for the EngD. Thanks also to Gordon McGillivray for supporting my proposal to gain funding. I hope the results of the EngD will prove to be valuable to the programme.
Discussions/guidance I must thank John Knowles, my Industrial Supervisor. There are not many people who would choose to study plutonium with all its unstable behaviour for a doctorate, let alone having to work in such a highly regulated environment to be able to do so. John brought a lot of experience about hydrogen dosing of difficult elements such as plutonium, uranium and cerium and helped me forge a way through the process.
Michael Preuss, My academic supervisor. Having been in the nuclear industry for so long before starting the EngD, it was helpful to have the fresh eyes of a knowledgeable academic to be able to discuss the intricate nature of plutonium and how best to approach resolving various problems that were encountered.
I had plenty of discussions with colleagues at AWE during the EngD. Sometimes it was helpful advice and others it was a fresh perspective. Some of the ideas made it through, some faltered, but the process of discussing various ideas helped me gather my thoughts.
Technical support A special mention must go to the Technical Support Staff, without whom, the plutonium work would have been very difficult to achieve. This included moving the samples from glovebox to glovebox as needed, careful sample preparation and loading and unloading the FEGSEM. I thank you Duane, Lester, Jean, Hannah, Liz, Bill and Stuart for all the technical support and nothing ever being too much trouble.
Chapter 1 INTRODUCTION
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1.1 Aims of the research
This EngD was commissioned by AWE at the University of Manchester to investigate corrosion mechanisms of plutonium by hydrogen with the aim of understanding how microstructure of the metal interacts with the growing hydride phase and to understand how hydride forms the morphologies observed.
The aim of this research programme is to develop a deeper understanding of the formation of plutonium hydride. The central theme to this thesis is a metallographic analysis of the hydride reaction sites and the exploration of relationships between hydride and parent material microstructures. In particular, changes to the parent material surrounding the hydride reaction sites are of specific interest. One important question to answer is regarding hydride growth morphology. Currently at the commencement of the EngD, the morphology of hydride sites was not well understood, which has implications on predicting quantity of hydride formed during storage. When plutonium hydride is exposed to oxygen a vigorous exothermic reaction takes place, which liberates and scatters particulate from the hydride reaction sites. It is possible that if enough hydride reacted with oxygen that the temperature of the plutonium metal could be raised to the extent that ignition could take place. Knowledge of the hydride morphology is important in calculating the likely amount of hydride present during storage.
The transformation of plutonium metal into plutonium hydride and associated breakup of the reaction product into a powder has been used in the recovery and re- use of plutonium as metal or oxide using the HYDEC and HYDOX processes respectively [1,2]. These processes work because plutonium hydride forms as a powder from the metal, which can be recovered and refined for subsequent use.
A model to predict how much corrosion is likely to be present on a metal plutonium component following storage would be valuable to allow predictions of the amount of corrosion which could be expected. Mathematical models exist for uranium corrosion by hydrogen [3,4] describing the time for a hydride to precipitate (the incubation time) but none exist which are specific to plutonium corrosion owing to the scarcity of data on the hydriding mechanisms during storage. The terms ‘initiation’, ‘incubation’ and ‘induction’ are used by different authors to describe the
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same period. Currently, the uranium model considers the rate determining step for production of hydride reaction sites to be the incubation time for a local region of plutonium to exceed the terminal solubility for hydrogen. This is based on the native oxide present on the surface acting as a diffusion barrier to hydrogen. Once terminal solubility has been exceeded then a hydride reaction site can precipitate. The continued growth of hydride reaction sites following initial precipitation is not well understood.
1.2 Introduction to plutonium
1.2.1 Origins of plutonium
Plutonium is a predominantly man-made element, and at the time of writing it is estimated that there is over 1500 tonnes of plutonium in the world [5], of which very nearly all has been manufactured. Heavy elements are produced by neutron capture processes when stars reach the end of their lives and supernovae occur [6]. Most isotopes of plutonium are highly radioactive with the most stable isotope (244Pu) having a half life of 76 million years. When compared to the age of the planet Earth (~4 billion years) any plutonium which existed when the Earth formed would have decayed into other species. Small amounts of 244Pu have been found in bastnasite rocks [7] where 85 kg of the rock was found to contain sufficient 244Pu to be detectable once processed. This is an extremely low relative abundance of -20 plutonium at 9.5 x10 gPu/grock and is a negligible composition in the earth’s crust.
Plutonium was first created and detected in 1941 by Seaborg, Kennedy, and Wahl [8] using a cyclotron at the University of California at Berkley and was the second new man made element after neptunium. The atoms of plutonium were created by impacting deuterons into targets of 238U, giving rise to 238Pu. The high activity of 238Pu was sufficient to detect the alpha decay of the few atoms produced.
Industrially the production of plutonium occurs via capture of neutrons from the high neutron flux within the core of a reactor. Under neutron bombardment the nuclei of the 238U isotope can absorb one neutron to become 239U (half life of 23.5 min) and decays via beta particle emission to form 239Np (half life of 2.33 days), which then decays via another beta particle emission to 239Pu (Equation 1). The half
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life of 239Pu is 24,000 years, which is sufficient to be considered stable. The time spent in the reactor core determines the relative proportions of isotopes of plutonium produced, as more time allows the 239Pu nucleus to absorb further neutrons, creating 240Pu, or higher isotopes.
Equation 1 The plutonium production process [9]
1.2.2 Isotopes and radioactivity
Plutonium has 17 known isotopes which are all radioactive (refer to Table 1). The various isotopes decay by a range of radioactive decay paths including alpha emission (α), beta emission (β), electron capture and spontaneous fission. The radioactive half life of plutonium varies considerably according to the isotope in question, from 7x107 y for 244Pu to as low as 20 min for 233Pu. There are also seven isotopes which undergo spontaneous fission but to a lesser proportion than decay via other modes. 239Pu and 241Pu are fissile and can be induced to fission following capture of neutron; 240Pu is fertile and can capture an incident neutron to become fissile as 241Pu. Reactor produced plutonium only exists in any meaningful amount as 238Pu, 239Pu, 240Pu, 241Pu, and 242Pu. The short half life of 243Pu (ca 5 h) hinders higher isotopes from being produced in a reactor. One of the aspects of radioactive materials is that due to the decay process of radioactive isotopes, an amount in- growth of the daughter products occurs [10], and dependent on the half-life of the daughter products, further isotopes in the decay chain can in-grow. This in-growth alters the composition of the plutonium over a period of years to the parts-per-million (ppm) level. In aged material, this can affect the metallurgical microstructure by precipitation of second phase particles such as uranium compounds.
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Pu Half life Decay Mode (%) Fissile Isotope
α β Electron Spontaneous Capture Fission 228Pu 4 ms 100 232Pu 34.1 min 23.6 77.6 233Pu 20 min 0.125 99.885 234Pu 9.0 h 6 94 235Pu 26 min 0.003 99.997 236Pu 2.85 y 100 1.37E-77 237Pu 45.6 d 0.00424 99.996 238Pu 86.4 y 100 1.85E-74 239Pu 24360 y 100 3.0E-108 Fissile 240Pu 6580 y 100 5.75E-65 Fertile 241Pu 14.4 y 0.00245 99.998 2.4E-14 Fissile 242Pu 3.79x105 y 100 5.54E-46 243Pu 4.95 h 100 244Pu 7.6x107 y 99.879 0.121 245Pu 10.5 h 100 246Pu 10.85 d 100 247Pu 2.27 d 100 Table 1 Isotopes of plutonium [11], [12], [13]
1.2.3 Fission
There are three aspects to fission that need to be considered in the case of plutonium, and all three are involved with a number of plutonium isotopes. The three aspects are: whether an isotope can undergo Spontaneous Fission; whether a nucleus is Fissile; whether a nucleus is Fertile.
Spontaneous fission is the disintegration of a nucleus into two daughter products without any external influences. During this process, a number of neutrons are often released with some kinetic energy. Six of the plutonium isotopes are known to undergo spontaneous fission. The daughter products produced depend on the isotope which undergoes fission, giving rise to a large, but finite number of possibilities of daughter products for each isotope, for instance it is known that as a result of fission, uranium forms over 80 different fission products [14].
Fissile means that a nuclear material is capable of undergoing nuclear fission upon capture of a neutron, causing a splitting of the nucleus in the same manner as that seen in spontaneous fission [15]. Fertile means that the nucleus of a nuclear
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material can be converted to fissile material by capturing a neutron to become one isotope higher in the elemental species, i.e. 240Pu can absorb a thermal neutron and change into 241Pu rather than undergoing a fission event [15].
The neutrons emitted during spontaneous fission can induce fission in other plutonium nuclei, when sufficient mass of 239Pu exists in close proximity; the neutrons generated in fission events are captured to ensure that a chain reaction propagates. This is the concept of criticality and the minimum amount of mass required to make a sphere of plutonium undergo a self-sustaining nuclear reaction is known as the critical mass [16].
1.2.4 Allotropes of plutonium
Plutonium is an unusual element in many ways: one of which is that it exhibits six allotropes between room temperature and has an anomalously low melting point of 641 °C [17]. No other elements exhibit as many allotropes by variation of temperature. The nature of the metallic bonding depends on the valence electrons which determine the structure of the allotropes. The valence electrons in plutonium are the 5f electrons, which have a narrow valence band. As the actinides increase in atomic number, the f-band starts larger than the other bands and the light actinides exhibit d-band behaviour, arising from the itinerant 5f electrons, but contraction of the 5f orbital renders the electrons localised from americium onwards. The transition between d-like and f-like behaviour occurs at Pu [18].
A seventh allotrope, the ζ-phase can be formed at pressures between ca 2 and 32 kBar (3.2 x 109 Pa) and between temperatures of ca 280 °C and 520 °C. Very little is understood about the properties of the ζ-phase since this phase only exists at elevated temperature and pressure and is difficult to work with.
The room temperature stable phase of plutonium is the α-phase, which is very hard and brittle meaning that machining of α-Pu is difficult. At 112 °C α-Pu transforms to the β-phase, causing large volumetric changes. These two factors make α-Pu an undesirable phase for manufacture.
As with other metallic systems that form multiple allotropes it is possible to stabilise plutonium phases beyond their pure limits by the addition of small amounts of other elements. The more ductile δ-phase can be stabilised to improve machining
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properties and it is most common to find examples of the δ-phase stabilised using small amounts of aluminium or gallium.
The extent of solubility of an element in a host matrix of another element is described by the Hume Rothery rules [19]. If one or more of the rules cannot be achieved then it is unlikely that they will form a miscible solid solution and other phases or inter-metallic compounds can be formed. The most significant Hume Rothery rules for plutonium alloying are [20]
Size factor – alloys only form extensive solid solubility with atomic radii within 15% of each other.
Electronegativity – alloys with elements differing largely in Pauling electronegativity are likely to form compounds, not solid solutions.
Crystal structures – occur at characteristic electrons per atom ratio. Elements with similar ratios are more likely to have higher solubility in each other i.e. if similar crystal structures exist, more likely to form solid solutions not compounds.
Plutonium is electropositive, with a value of 1.28 Pauling units [21], an atomic radius of 1.53 Å [22]. Plutonium does not have complete miscibility with any other element, however neptunium, and uranium exhibit extended solubility in the α, β and ε-phases and americium has a considerable solubility in the δ-phase. Plutonium tends to form intermetallic compounds and mixed phases [23] rather than solid solutions.
This project is principally concerned with the formation of PuH2+x during long term storage of α and δ-Pu-Ga alloys because these are the only phases that can be stabilised at room temperature. A description of relevant properties of the α and δ phases is given.
1.2.4.1 Alpha plutonium The α-phase of plutonium is a simple monoclinic lattice that has 16 atoms in the unit cell, the cell dimensions are a = 6.183 Å, b = 4.822 Å, c = 10.963 Å, the cell angles are: β = 101.79° [24,25]. The density of alpha plutonium is 19.86 g cm-3. The monoclinic structure allows for a high density by employing different bond lengths between atoms, and is a modification of the hexagonally close packed (hcp) lattice
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structure. The unit cell is manifested as two inverted layers of 8 individual atoms (see Figure 1).
Figure 1 The alpha plutonium lattice. The unit cell is shown by the box. Created using Mercury 3.6 with data from the Crystallography Open Database [26].
The α-phase has a relatively high strength, with a Vickers hardness of around
0.82 GPa (260 HV10) [27] and a bulk modulus of 54.4 GPa [28]. The phase is very brittle, giving less than 0.1% elongation during tensile testing [29]. The thermal conductivity of alpha plutonium is the second lowest of the metals at 6.74 W m-1 K-1.
1.2.4.2 Delta plutonium The δ-phase is a face-centred cubic lattice (fcc) (Figure 3), with cell dimensions of a = 4.6371 Å [24,30]. This gives a unit cell volume of 99.704 ų.giving rise to a density of 15.92 g cm-3. The δ-plutonium allotrope is unusual pure elemental metals in that in pure plutonium it exhibits a negative coefficient of thermal expansion of -8.6x10-6. This is unusual in pure elements – only silicon exhibits a negative thermal expansion at cryogenic temperatures. Negative thermal expansion is more commonly found in certain compounds such as water or zirconium tungstate. The high number of slip planes in the fcc lattice means that the δ-phase is soft and easily machined. Hardness values range between 0.28 GPa and 040 GPa for newly
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produced Pu-Ga δ-stabilised alloys [31,32]. The bulk modulus is ca. 30 GPa [33]. The thermal conductivity of gallium-stabilised delta plutonium is around 9.2 W m-1 K- 1. The aim to get a stable, easily-worked material with minimum alloying composition means that the phase of interest must be the δ-phase.
The δ-phase field boundaries can be stabilised to greater extents by the addition of a number of trivalent alloying constituents, such as cerium, americium, gallium, aluminium, scandium [23], and thorium, zinc, zirconium [34]. It is possible to stabilise the δ-phase to room temperature using a few atomic per cent of aluminium, gallium, cerium or americium. An expanded version of the metastable low alloy Pu- Ga phase diagram is shown in Figure 2a, an alternative form of the Pu-Ga phase diagram was measured by Russian researchers (Figure 2b), which indicates that given sufficient time there is the possibility that a eutectoid decomposition would take place. Since the timescale for such a decomposition at the temperatures (and the associated rate of diffusion) are extremely long (ca 11,000 years), the equilibrium (Russian) phase diagram does not represent the likely structure that would be encountered in a real world situation.
Figure 2 The low alloy Pu-Ga phase diagram [35]
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The available data for scandium [23] suggests it may not stabilise the δ-phase to room temperature, whereas americium stabilises plutonium in the δ-phase to room temperature over a wide composition range, between ~5 at% and 75 at% but is not used in the stabilisation because of the high gamma ray activity of americium, leading to high dose rates to nearby staff. Americium is an in-growth product of the β-radioactive decay of 241Pu.
The δ-phase has been observed to transform to the α-allotrope via a martensitic reaction under an applied stress [36], or at low temperatures [37]. Since no diffusional processes take place during this transformation, gallium remains trapped in the α-matrix and the phase is referred to as the α’-phase.
Figure 3 The face-centred cubic delta plutonium unit cell. Created using Mercury 3.6 using data from the Crystallography Open Database [26].
1.2.5 Transition between phases
There are large changes in volume during the heating of plutonium as the crystal allotropes change from one structure to another (Figure 4). There is some disagreement in the published literature about the actual onset temperature and enthalpies of the transitions [38–40]. It has been suggested that differences in experimental procedure may give rise to the wide range of differences found [41,42]. There is also disagreement over the boiling point of plutonium, with values ranging from 3230 °C to 3327 °C [43–45]. The generally agreed transition temperatures are shown in Table 2.
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Table 2 Phase transition details
Transition Temperature (°C) Volume Change (%) Enthalpy Change (J mol-1) (a) (b) (c) αβ 112 + 10 3706 βγ 215 + 3.5 478 γδ 310 + 7 713 δ δ’ 451 - 0.5 84 δ’ε 480 - 3 1841 εLiquid 641 - 3 2824 LiquidGas 3230 (d) N/A 345200 (a) [46] (b) Calculated from [47] (c) [40] (d) [48]
Figure 4 Dilatometry data for pure plutonium indicating the large changes in volume compared to that of aluminium over the same temperature range [8].
1.2.6 Chemical reactivity
Plutonium exhibits complex chemistry, having five possible oxidation states [49]: Pu (III), Pu (IV), Pu (V), Pu (VI) and Pu (VII), defined by the behaviour of the
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valence electrons in the 5f energy levels. The high availability of electrons and similar energies between the 5f and 6d orbitals gives plutonium a wide range of compounds that can be formed, with up to four of the five states able to exist at the same time.
The lanthanide series is defined by the 4f electrons and the chemistry is remarkably similar along the series, due to the fact that the 4f energy levels sit within the 5s energy levels, and the valence electrons are therefore partially shielded from the external environment. In the case of the actinides, the 5f electrons sit mainly within the 7s electrons, but have a longer tail to the distribution, meaning that there is a small probability that part of the 5f orbital exists outside the 7s orbital. This means that the 5f electrons have an influence on the chemical nature of the actinides (Figure 5) and as the actinide series progresses; the 5f orbital contracts, bringing it fully within the 7s. Once this occurs, the bonding behaviour changes from being d- like (transition metal) to being f-like (rare earth) [8], making the heavy actinides less reactive. This changeover in the actinide series occurs at plutonium.
Figure 5 Overlapping radial distribution functions of the 7s and the 5f orbitals for plutonium [25].
The number of possible oxidation states complicates prediction of the spread of plutonium in the environment. The Pu (IV), Pu (V) and Pu (VI) states are most important for water based dispersal, as these oxidation states are stable at the pH of natural environmental waters [50]. Analysis of groundwater at the Nevada Test Site
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show that plutonium in the water table could be traced back to a particular test, some 1.3 km north of the water sampling site, a distance consistent with the estimation of the ground water migration velocity [51].
1.2.7 Uses of plutonium
The uses for plutonium are fairly limited. As a material there are significant issues with the use of plutonium, it is not commercially available and significant hazards are associated with it. The difficulty in working with a material which changes phase so easily limits its usefulness as a structural material. The nuclear fission which occurs in plutonium provides its main technological use in mixed oxide fuel [52] and defence applications [53–55]. The high α particles emission activity of 238Pu causes disturbance to the crystal lattice which raises the temperature. This self-heating effect is used in radioisotope thermoelectric generators to power small space craft [56].
1.2.8 Why is the corrosion of plutonium important?
The corrosion of plutonium by hydrogen during storage is important in a number of ways, all of which stem from the properties of the metal and hydride. As will be discussed later, the surface hydride is very friable and the stress induced by the volume difference between metal and hydride compound is known to generate respirable particles [57]. On exposure to oxygen, the hydride converts to the oxide in a very exothermic reaction, the vigorous nature of this reaction can disperse particles around a localised area. This presents a major hazard for the any operation involving plutonium which has been stored for both the operators undertaking the task and for the potential to dispersal into the wider environment.
Metallic plutonium is expensive to produce therefore once produced is required to last for a considerable period during storage. It is important to understand all possible corrosion mechanisms which can degrade the material. The corrosion of plutonium metal by hydrogen has implications for the long term storage of components manufactured from metallic plutonium.
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1.2.9 Health
The isotope of plutonium most commonly used (239Pu) is mainly an α-emitter, plutonium with its relatively high activity is considered to be a health hazard when ingestion or injection occurs. Externally to the body, α-emission is relatively benign because the large α-particle is stopped by the dead layer of skin on the outside of the body [58]. Within the body however, α-emission is particularly undesirable since α-particles do not travel far in matter and deposit their energy in a short space, causing a large amount of damage and ionisation in that space. When in the body, plutonium is both biologically and chemically attracted to bone, and aggregates on the surface of bone. These aggregates then deliver a concentrated dose of α- particles to the tissue surrounding the bone (see Figure 6), causing significant and prolonged damage to these tissues, often giving rise to cancers and deformations [14].
Figure 6 The effect of plutonium in the body – fission tracks which have occurred due to the radioactive decay of plutonium at the bone surface are evident as short black lines [58]
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The greatest risk of damage to health by plutonium in today’s situation is via dispersal of fine particulate of a plutonium species, and the subsequent inhalation or ingestion of the particulate. The plutonium currently in the environment is there as a result of atmospheric testing of nuclear weapons (>5 tons) and much smaller quantities from accidents at nuclear power stations such as Chernobyl and Fukushima [59,60]. The most likely scenarios for members of the population to become contaminated and inhale, ingest or inject plutonium are: via a reactor core meltdown and subsequent environmental dispersal in a manner similar to Chernobyl or Fukushima; leaching into groundwater from geological storage of nuclear waste; atmospheric dispersal of loosely-bound corrosion product from stockpiled plutonium metal. The most likely set of people who are likely to become contaminated with plutonium are those persons working directly with the element. In this case, strict controls are in place to ensure safety of plutonium workers. Plutonium affects health of workers, and others who might come into contact with it in mainly two ways, firstly by exposure to the external radiation arising from the radioactive decays of the plutonium isotopes and other impurities in the material, and secondly by introducing plutonium contamination into the body, usually by inhaling, ingesting or injection.
Understanding the corrosion process of plutonium is important in the long term storage of metal since corrosion products have the propensity to form plutonium containing particles, loosely bound to the surface [43], which could be easily scattered, thereby spreading radioactive contamination.
The Fundamental Safety Principles of safety when working with nuclear and radioactive materials are set out under the auspices of the International Atomic Energy Agency (IAEA) and the Safety Objective is stated as: “The fundamental safety objective is to protect people and the environment from harmful effects of ionizing radiation.” [61]. In order to meet this objective, the IAEA state that measures have to be taken:
(a) To control the radiation exposure of people and the release of radioactive material to the environment;
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(b) To restrict the likelihood of events that might lead to a loss of control over a nuclear reactor core, nuclear chain reaction, radioactive source or any other source of radiation;
(c) To mitigate the consequences of such events if they were to occur.
Both (a) and (b) apply to the case of metallic plutonium in that the potential of loss of control over a radioactive source and potential release of radioactive material to the environment are increased when dealing with particulates containing plutonium. As a result prevention of the spread of radioactive contamination is normally done by having a staged set of barriers to the spreading of contaminated material. Work on plutonium is accordingly undertaken in a glove box environment in a dry nitrogen atmosphere.
Containers being used for long term storage of plutonium during the 1960s in the USA were found to have pressurised, due to radiolysis of adsorbed vapours or of the plastic bags used as a contamination barrier [62]. In some cases, storage containers had failed due to expansion of the solid plutonium [43]. The plutonium was stored in plastic bags within metal tins. The failure of the containment resulted from radiolysis of adsorbed vapours and the plastic storage bags by the α particles emitted from the plutonium liberated hydrogen. The hydrogen reacted with the plutonium, causing an overall volume increase of the stored Pu, which caused perforation in the containers. Upon unpacking of the failed storage containers even in a low oxygen glove box, the package was found to be hot to touch and had expanded significantly, indicating that the exposure to small levels of oxygen had caused a reaction with plutonium hydride.
1.2.10 Pyrophoricity
The pyrophoric nature of plutonium hydride in air means that the presence of hydrides is a fire hazard and can potentially spread Pu contamination outside of suitable containment It has been suggested that the method by which the hydride is formed determines the ignition temperature of the hydride when exposed to air [63]. Confirming the work of previous experiments into the stability of plutonium hydride in contact with air [64], Haschke et al. described the difference in form of the hydride
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produced at low temperatures (less than 100 °C) to be of a finely divided, loose powder, compared to that produced at higher temperatures (greater than 270 °C) being of a more coarse material. The coarse material produced at higher temperature is non-pyrophoric; whereas the fine powdered material formed at low temperature is very pyrophoric and ignites spontaneously on contact with air (Figure 7). The temperature at which the hydriding reaction occurs is usually due to self- heating arising from the exothermic hydriding reaction. The rate of reaction governs the heat generated, and is a function of the amount of hydrogen available to the plutonium.
Figure 7 A video still recorded during exposure of the vigorous reaction that occurs when a plutonium sample covered in hydride formed at room temperature to oxygen.
Metallic plutonium is also pyrophoric, as the reaction with oxygen is exothermic [65]. Bulk plutonium does not ignite burn in air below 500 °C, but metallic fines can become pyrophoric at temperatures as low as 150 °C. Clearly these values are above the storage temperature of the plutonium metal, however there is concern that hydride present, being pyrophoric at room temperature, could ignite the plutonium (either bulk or fines) potentially leading to a serious spread of contamination. Therefore understanding the mechanism of formation of hydride is of great importance for continued safe operations involving metallic plutonium.
1.2.11 Environment
The affects of plutonium on the environment are due to the potential for health effects in people, flora and fauna that inhabit that environment. With a half-life of ca 24000 years, plutonium would contaminate the environment for an extended period. Contamination of the environment by a nuclear establishment causes significant
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social and political uproar as was the case when a radioactive particle was found on Dounreay beach [66–68], with the potential for significant loss of reputation, sanctions, regulatory improvement notices and ultimately closure of the establishment.
1.3 Introduction to plutonium oxide
Plutonium reacts with any available oxygen, rapidly forming a passivating layer. The oxidation rate on unalloyed plutonium has been measured as high as 0.21g Pu cm-2 min-1 at 900 °C. The nature of the oxide on plutonium is complex; the outer surface forms PuO2, whereas the metal/oxide interface forms Pu2O3. The reaction of oxygen with plutonium is a complex mechanism. The most chemically stable oxide of plutonium in air is plutonium dioxide (PuO2). XPS measurements have been used to demonstrate that the oxide can transform from PuO2 to Pu2O3 in the absence of an excess of O2 [69], a model was proposed that the nature of the oxide on plutonium is not a single species of oxide, but a multi layered system with the sesquioxide forming at the plutonium/oxygen interface and then plutonium dioxide forming towards the outer surface (Figure 8).
Figure 8 Schematic of how the surface of plutonium oxidises in an O2 environment, and
how an oxidised plutonium surface auto converts from PuO2 to Pu2O3 in UHV conditions [69].
The reaction of plutonium is similar to that of cerium in that it forms rapidly in air and forms an adherent layer [70]. The rate of oxidation is higher than that for uranium, being (in dry air) around 5.6x10-7 mg m-2 min near room temperature, rising to 1.25x10-1 mg m-2 min at 350 °C, this rate is much higher in moist air [71].
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Two different stoichiometries exist for plutonium oxide, namely sesquioxide, in the form Pu2O3 and dioxide, in the form PuO2. The Pu2O3 phase forms whenever there is oxygen in contact with the metal, and only when a critical abundance of oxygen is in contact with the plutonium can the PuO2 stoichiometry form. Other concentrations may be found as defected forms of the stoichiometric phases.
Figure 9 The equilibrium phase diagram for Pu-O [72]
The PuO2 form has only one allotropic form, the CaF2 structure which is based on the fcc lattice, where the plutonium atoms occupy the face centred positions with the oxygen atoms at interstices [73]. The unit cell for PuO2 is fcc and has a lattice parameter a of 5.47 Å [74], giving a volume of 163.667 ų. Comparison with the fcc δ-phase Pu lattice shows that the PuO2 lattice structure has increased in volume by a factor of 1.64 over that of the metal. The pure Pu2O3 form is dimorphic; having the ability to form a hcp or a bcc structure dependent on the temperature at which it was formed [75]. The α-Pu2O3 has a composition of PuO1.52, and is only stable below 300 °C, and β-Pu2O3 forms when oxidation occurs between the temperatures of 300-2033 °C. Once formed, the β-phase Pu2O3 is stable to room
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temperature, and transformation to the α-Pu2O3 does not readily occur. The crystal structure of α-Pu2O3 is bcc; the β-Pu2O3 form is hexagonal, which has the same lattice parameters as hexagonal La2O3 [76], The oxide over-layer in contact with plutonium metal is somewhat different from pure oxides, existing as a type C sesquioxide (Figure 10), which occupies a CaF2 fluorite structure but with 25 % oxygen deficiencies, similar to that observed in Ce2O3 [77].
Figure 10 The CaF2 structure of PuO2 with the plutonium atoms as the larger blue
spheres. For the type C Pu2O3 structure, 25 % of the red spheres would be missing. Created using Mercury 3.6 using data from the Crystallography Open Database [26].
PuO2 has been reported as having a hardness of around 666 Knoop (2.22
GPa), taken from extrapolations of measurements made in conjunction with UO2 and
PuO2 -40% UO2 [78]. The hardness of the β-Pu2O3 phase of around 2.20 GPa has been reported for β-Pu2O3 in a matrix of PuO and β-Pu2O3. α-Pu2O3 has been measured with a hardness of 5.34 GPa in a matrix of uranium molybdenum.
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1.3.1 Rate of reaction with plutonium metal
Oxygen reacts quickly with plutonium metal, tarnishing in minutes [8].Work using ellipsometry on cleaned surfaces of plutonium oxidised at low pressures of oxygen of 10-2 Torr (around 13 Pa) in vacuum show that whilst initial oxidation is rapid, the oxide provides a semi-protective barrier to further oxidation, and the rate slows considerably, following a parabolic oxidation rate [79], indicative of a diffusion controlled reaction. Other models of oxide growth [80] suggest that following the slowing of the parabolic rate, steady state corrosion is achieved (see Figure 11), governed by the spallation of PuO2 from the oxide surface once sufficient thickness has been achieved.
Figure 11 Standard Chemical Picture of Plutonium Oxidation in Dry Air [80]
The rate of oxidation is considerably affected by the moisture content of the air to which it is exposed, in dry air the rate can be as low at 20 nm per year, increasing by a factor of 200 in moist air, and at temperatures of around 100 °C can be as high as 104 times the rate found in dry air [81].
1.3.2 Stability
The reaction Pu(cr) + O2(g) ⇌ PuO2(cr) gives a Gibbs energy (ΔG) of -998.1 kJ mol-1 [82], which is a spontaneous reaction because ΔG < 0. In a situation where an excess of oxygen is available, PuO2 should readily form. The Gibbs energy of -1 formation of Pu2O3 is -1580.4 kJ mol [82], which indicates that the Pu2O3 oxide
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should form via the reaction 4Pu(cr) + 3O2(g) ⇌ 2Pu2O3 (cr), however this requires 4 moles of plutonium to create 2 moles of Pu2O3, giving a 2:1 ratio of Pu:Pu2O3 compared to that of the 1:1 ratio of Pu:PuO2, therefore the energy per mole of Pu -1 required to create Pu2O3 is -790.2 kJ mol , which is less negative than that of PuO2.
In this situation ΔG(PuO2) < ΔG(Pu2O3) < 0 indicating that in a situation where an excess of monatomic plutonium is available, such as at the oxide/metal interface the
Pu2O3 reaction is favourable, but in stochiometric condition Pu +O2 ⇌PuO2 or where an excess of oxygen exists, the PuO2 reaction is favoured.
The most chemically stable oxide of plutonium is PuO2; it is unreactive against most chemicals and is slow to dissolve in strong acids such as H3PO4 and HNO3 [73]. Studies have shown that in an ultra high vacuum environment, plutonium will undergo auto reduction and the outer PuO2 layer will lose oxygen and reduce to the type C Pu2O3 [69].
1.4 Plutonium hydride
Early work involving Pressure-Composition-Temperature (PCT) investigations [83], determined that the stoichiometry of plutonium hydride is either in the form of plutonium dihydride (PuH2), or plutonium trihydride (PuH3). PuH2 forms between H/Pu ratios of ca 0 to 2.0, and at increased pressures, a slightly sub-stoichiometric H/Pu ratio of 2.9 was achieved. Plutonium has been found to react with hydrogen to form the compound PuH2 from room temperature up to 800 °C [84], from pressures of 10 mbar (1 x 103 Pa) upwards [85], hydrogen begins occupying the octahedral sites in the PuH2 fluorite lattice growing into the PuH2 by filling the remaining spaces in PuH2 but at concentrations above H/Pu of 2.7 a change in crystal structure is evident in the PuH3 from the small plateaus in Figure 12.
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Figure 12 PCT traces of plutonium hydrided to H/Pu ratios above the PuH2 stoichiometry [86] .
Higher fidelity analysis was undertaken on the Pu-H system below the H/Pu ratio of 2 [84], which suggested that above the temperature of 500 °C there was appreciable solubility of hydrogen in the plutonium up to a maximum H/Pu ratio of about 0.2. The miscibility gap between H/Pu of 0.2 and 1.95 indicates that two phases exist which are the PuH2 phase in the H saturated Pu metal. This study also highlighted that the upward deviation of the isotherms from horizontal occurs at a value of H/Pu less than 2.0. It was suggested in the paper that this occurs because a degree of solubility of Pu in PuH2 exists or that PuH2 is non-stoichiometric, having a range of compositions.
The phase diagram of Pu-H (Figure 13) indicates that there is a degree of solubility of hydrogen at temperatures in excess of ~400 °C, in pure plutonium this would incorporate some of the δ, δ’, ε and liquid phases. The PCT curves obtained
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during early work on hydrogen absorption measurements appeared to suggest there was no solubility of hydrogen in plutonium below this temperature [84], however later work concentrating on the solution region indicated that up to 1-2 at% hydrogen was soluble in plutonium to temperatures as low as 350 °C [87]. The solubility at room temperature is vanishingly small. Extrapolation from the recent study of Richmond et al. suggests the solubility in the δ-phase is only 1.1x10-5 H/Pu.
Figure 13 Top: The Pu-H diagram of Ellinger et al. [23] . Bottom: The equilibrium phase
diagram for Pu-H [63]. Phase I is a mixed phase with a PuH2 and hydrogen in
solution in the plutonium metal, Phase II is the PuH2 phase, which adopts the fluorite structure and appears to retain some solubility for hydrogen around
stoichiometric PuH2. Phase III is a mixed phase between PuH2 and PuH3, which arises as the tetrahedral fcc lattice sites become filled. Phase IV occurs at around H/Pu of 2.75 which corresponds to complete filling of the fcc lattice, which given sufficient temperature, forces a change to a hexagonal structure.
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The interaction of plutonium and hydrogen is completely reversible and dissociation can occur with a reduction in pressure to around 3 µmHg (0.4 Pa) at temperatures in excess of 400 ºC [57]. Removal of the hydrogen causes a complete dehydriding of the plutonium. Upon dehydriding, the plutonium hydride turns into a plutonium metal powder.
The reaction rate with hydrogen is extremely rapid, a fact that is often made use of in the HYDEC and HYDOX processes for reclamation of plutonium from alloys into pure metal for re-use. However, the reaction between hydrogen and plutonium does not begin immediately and an initiation time exists where no apparent reaction can be observed taking place either by weight gain, reduction in pressure, or visual growth of hydride reaction [64,83,88–91]. This induction time has been ascribed to the time required for hydrogen to pass through the semi-protected plutonium dioxide layer and mirrors the effect found on other materials which form hydride reaction sites such as cerium [92] and uranium [93].
The PuH2 form of the hydride is fcc, having the CaF2 structure [73] (Figure
14), and occurs between a composition of PuH2.0 and PuH2.7. The lattice parameter for PuH2 is 5.359 Å at a composition of PuH2.0, giving a unit cell volume of 153.9 ų and a calculated density of 10.40 g/cm³. In the PuH2 compound, hydrogen exists as interstitial atoms the tetrahedral interstices of the lattice. The hardness of plutonium hydride has been reported as between 1.138 GPa and 1.295 GPa using Vickers micro indentation on hydride inclusions found in that material [94], which is harder than the α-Pu and δ-Pu metal phases.
PuH3 retains the cubic fluorite form (Figure 15) until a H/Pu ratio of 2.7 is exceeded, with a mixture of the cubic and hexagonal phases existing between the ratios of 2.4 to 2.8 [63]. Above this ratio, the hydride exists as a hexagonal structure
(refer to phase IV in Figure 13). The hexagonal PuH3 has a unit cell volume of 250.95 ų and a density of 9.61 g cm-3 [73].
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Figure 14 Crystal structure of PuH2, large spheres are plutonium, the black spheres are occupied tetrahedral sites [95]
Figure 15 The cubic crystal structure of PuH3 , white spheres are occupied octahedral sites and black spheres are occupied tetrahedral sites [95]
1.4.1 Rate of reaction with plutonium metal
-2 -1 In dry H2 the rate of reaction can be as high as 7.0 mmol H2 cm hr [96].and the reaction interface can move at a speed of 20 cm h-1 [8]. The reaction of hydrogen with plutonium is very exothermic, and has on occasion been noted to cause the reaction products to glow red; in some cases, enough heat is produced to sinter the particles together [57]. This exothermic reaction can heat the plutonium and affect the kinetics of the reaction.
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1.4.2 Stability of the hydride phases
The reaction Pu(cr) + H2(g) ⇌ PuH2(cr) gives a Gibbs energy (ΔG) of -135 kJ mol-1 [97], which is a spontaneous reaction because ΔG < 0. In a situation where an excess of hydrogen is available, PuH2 should readily form. PuH3 can form from PuH2 by further filling of the fcc lattice via the reaction 2Pu(cr) + 3H2(g) ⇌ 2PuH3 (cr). The -1 Gibbs energy of formation of PuH3 is -156 kJ mol [97], which requires one mole of plutonium atoms per mole of PuH3 formed. The energy of formation of PuH3 is marginally more negative than that of PuH2, ΔG(PuH3) < ΔG(PuH2 ) < 0 which suggests that there is a much reduced driving force to form PuH3 from PuH2 than there is to form PuH2 or PuH3 from Pu metal. As more hydrogen is required to form
PuH3 than PuH2 then the reaction path should be Pu + H2 ⇌ PuH2 then PuH2 + H2 ⇌
PuH3. Where hydrogen is available to plutonium the PuH2 phase will readily form from plutonium metal. Where an excess of hydrogen is available a small driving force -1 (ΔG = -21 kJ mol ) exists to form PuH3 from PuH2.
1.4.3 Sources of hydrogen
As a result of investigations into the rupturing of storage containers containing cast metal plutonium, [43] sources of hydrogen that could subsequently be involved with corrosion of plutonium were identified as the product of radiolysis of hydrocarbon based residues left on the surface of the castings following handling and manufacturing processes. The casting was also stored within a sealed plastic bag to act as a containment barrier against migration of radioactive particulates; within an unsealed tin. The plastic bag prevented corrosion gas from dispersing and also acted as a source of hydrogen.
In a similar case involving uranium where a water source was present the water reacted with the uranium forming uranium oxide, hydride and hydrogen. The hydrogen rich atmosphere in the drum exploded on opening with the hydride providing the ignition source [98].
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Chapter 2 LITERATURE REVIEW
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2.1 Historical research into plutonium hydride
2.1.1 Initial studies
Early work into the reaction of hydrogen with plutonium was conducted by I.B. Johns in 1944 [83]; the amount of hydrogen absorbed into the plutonium (99.9 % pure) was measured and found to be absorbed up to a ca 2.9 H/Pu ratio. A small change in composition was observed as a result of a large change in pressure, which suggested that the hydriding reaction of plutonium is not simply the pressure dependent Pu + H2 ↔ PuHx equilibrium reaction. During the reaction, an initial incubation time of about 8 minutes was observed and the hydriding reaction completed in a further 7 minutes. The reaction was observed to be exothermic.
Johns undertook pressure-composition-temperature (PCT) work in order to map out the absorption/reaction with temperature and pressure as a function of the H/Pu ratio. For compositions 2 ≤ H/Pu ≤ 3 a parabolic relationship was suggested. At H/Pu ratio < 2, a plateau pressure was discovered, suggesting that a mixed phase existed. Sorption isotherms showed a degree of solubility of hydrogen in plutonium.
The work of Johns was improved and extended [84] supposedly using Pu of a higher purity, although Mulford and Sturdy quoted their purity at 99.6 %, whereas Johns quoted the purity at 99.9 %. In this study, the temperature range was extended beyond the melting point of Pu and ranged from 400-800 °C. Deviation of the PCT isotherms from the horizontal (constant pressure) portion was noted at a value lower than the H/Pu ratio of 2, suggesting, in the view of the authors of the paper, that either there is a solubility for Pu in PuH2 at increased pressures; or that that PuH2 is slightly sub-stoichiometric and exists at a range of values of H/Pu. In this work, Mulford & Sturdy suggested that the initial curve found on the isotherms show limited solubility for H in Pu above 500 °C. No effects on the isothermal curves during the decomposition of hydride were noted due to the solid state phase transitions δ → δ’ and δ’ → ε. The authors applied the van’t Hoff equation to -1 determine heats of evolution for PuH2 and PuD2 as 37.4 & 35.5 kJ mol respectively.
Mulford and Sturdy then further suggested from that the unpublished work of a colleague (Ellinger) that PuH2 adopts the CaF2 (C1) structure with a lattice
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parameter of a = 5.359 Å with a density of 10.4 g cm-3. Haschke and Stakebake [99] confirmed the earlier work of Mulford and Sturdy [84] that the equilibrium hydride composition occurs at PuH1.95, and reported a similar the enthalpy of formation.
More recent work by Richmond et al. [87] focussed on absorption of hydrogen into plutonium to a concentration below the solubility limit at elevated temperatures. A difference in the hydriding behaviour of pure Pu was observed compared to that of Pu 2 at% Ga. The terminal solubility at temperatures 320-625 °C was suggested to occur near 1-2 at %.The enthalpy of solution was reported.
2.1.2 Solid solubility of hydrogen in plutonium and plutonium hydride
The Pu-H phase diagram (Figure 13b) suggests that there is no solubility for hydrogen in plutonium below about 350-400 ºC. Upon initial exposure to hydrogen at temperatures lower than this, the plutonium will react with the hydrogen to form small regions of PuH2 within a Pu matrix. The low composition region where hydrogen is in solution is known [87] and the phase boundary between solution and saturated solution was found to be between 1-2 at.% H over a temperature range of 350-625 °C. With further exposure to hydrogen imparted via increasing the pressure, the
PuH2 formed initially can transform into PuH3 phase. The phase diagram of Pu-H shows that there is a small mixed phase field between PuH2 and PuH3.[86]
2.1.3 Hydriding kinetics
The hydriding kinetics of plutonium powder with a surface area of 0.1-0.2 m2 g-1 was measured gravimetrically under constant hydrogen pressure [100] and reported the activation energy was 33.5 kJ mol-1. The plutonium powder was prepared from a 1 wt% gallium alloy bulk plutonium sample by a cycle of hydriding and then dehydriding under vacuum. The dependency on pressure was close to P1/2 which suggests the reaction is dissociation controlled.
Near isothermal kinetic experiments were also undertaken gravimetrically on bulk samples by Ogden et al. [101] to determine the hydriding kinetics at 1, 10 & 20 atm between 60 and 360 °C. The isothermal aspect of the experiment was achieved by bonding the plutonium samples to beryllium heat sinks so that heat generated by the hydriding process would not influence the rates observed. A change in the rate of
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hydriding and the morphology of the product formed was reported at 181 °C, which was ascribed to a reordering of the hydride into a super lattice in a manner similar to that observed in some lanthanide hydrides.
2.1.4 Activation energy
The results of rates determined barometrically for plutonium by Haschke and Allen [102] suggested a negative activation energy of -6.7 kJ mol-1 for the plutonium hydrogen reaction. Negative activation energy was not found in similar work. The work of Stakebake [100] suggested a figure of 35 J mol-1 for the activation energy. In this case, the low activation energy suggests that the hydriding rate is relatively independent of temperature. Activation energy was also measured gravimetrically by Bowersox using deuterium [89], and found to be around 2.2 kcal.mol-1 (9.2 kJ.mol-1). More recently work by Kenney and Harker on the specific hydriding rates on δ- -2 -1 stabilised plutonium found the rate increased from 10.3 to 28.0 μmol H2 cm s over the temperature range 26 to 153 °C and reported a positive value for the activation energy of 5 kJ mol-1, which compares favourably with previous work by Bowersox [103] and Ogden [101].
2.1.5 Pyrophoricity
Considerable work has been undertaken concerning the pyrophoricity of plutonium assure the safety of plutonium operations [64]. Metallic plutonium is not pyrophoric below 150 °C but a review into the pyrophoricity was conducted for the Rocky Flats Plant [65], collating information on the process of spontaneous ignition, which reported the specific surface area of the plutonium is critical for the ignition to take place at temperatures below 150 °C. Plutonium fines will spontaneously ignite in air above this temperature. Powdered plutonium hydride in the fcc form, which covers PuH2 and PuH2+x formed at low temperatures, is pyrophoric at room temperature; but the hexagonal PuH3 form required heating to 270 °C for ignition to occur [97].
The ignition of plutonium was investigated by heating either cubes or foils in either air or oxygen until ignition occurred using the burning curve and shielded ignition methods. It was found that as the specific area approached 11 cm-2 g-1, the ignition temperature decreased to sub 300 °C [104].
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2.1.6 Hydride growth behaviour
The basic steps observed to occur when gaseous hydrogen comes into contact with a plutonium surface are laid out by McGillivray et al. [90]. These follow the Bloch and Mintz mechanism for other hydride-forming metals and are stated as such in the paper. The steps are shown in Figure 16, and are described as: (1) Induction, where the hydrogen has been introduced to the plutonium, but no consumption of gas or formation of hydride occurs; (2) Acceleration, where hydride begins to precipitate at nucleation sites in the metal and visible spots of hydride are formed and subsequently grow on the plutonium surface; (3) Bulk Hydriding, where the surface is completely covered by hydride and the reaction only takes place at the Pu/Pu-H interface; (4) Termination, where either the supply of plutonium or hydrogen gas has been exhausted.
( ( ( ( 1) 2) 3) 4)
Figure 16 A typical rate vs. time curve for plutonium hydriding. The four regions identified refer to different steps in the hydride reaction process: (1) induction, (2) nucleation or acceleration, (3) bulk hydriding, and (4) termination [90].
During early investigations into the reaction of hydrogen gas with plutonium, it was noted that no reaction reaction was observed to take place for a while [83]. This period has been termed variously: the incubation, induction or initiation period. Initiation time of hydriding on nearly pure plutonium was originally recorded as 8 minutes on a 99.9 % alloy [83], subsequent work by Stiffler and Curtis [57] has shown that the variability of the initiation time could range from a few minutes to as
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much as several hours. Brown et al. [64], found that reaction only occurred upon heating the sample to a temperature between 100-200 °C, and in this case, the initiation time was of the order of 20-30 minutes. It has been suggested [62] that this was an indication that the hydrogen reaction is controlled by a surface passivating layer (SPL).
McGillivray et al. [85,90] used ultra-pure hydrogen in a reaction cell at a number of different pressures to determine the effect of pressure on the initiation time of hydriding. This was measured using an image capture system through the window in the reaction cell, and the initiation time was determined by the time taken from introduction of the gas to the first visible hydride spot on the surface of the sample.
The oxide formed on the surface of the plutonium is considered in a number of papers to be a semi-protective film on the surface of the metal, leading to the incubation time that is found during experiments on other aspects of the hydriding of plutonium. The incubation time on oxide-free plutonium is absent when exposed to hydrogen unlike plutonium with a native oxide [100].
Recent work by Dinh et al. suggests that the Pu2O3 layer on the surface of the metal significantly affects the hydriding behaviour of plutonium [105], with the suggestion that Pu2O3 acts as a catalyst to the formation of plutonium hydrides. This suggestion was made following examination of plutonium samples where the surface had been modified either by scratching the surface with a diamond scribe or by heating the sample to 110 °C. Both sets of experiments reported increased nucleation and growth kinetics over that of the control sample. The presence and thickness of Pu2O3 in the experiments was not measured directly, but inferred from previous literature. The work of Morrall et al. on auto reduction of PuO2 to Pu2O3 [69] suggests that a PuO2 film remains on the outer surface of the oxide caused by re- oxidation of the surface by background species even in UHV conditions. This film may however exist in the form of islands of PuO2 on the Pu2O3 film.
2.1.7 Hydride nucleation
There appears to be very little firm evidence about the nucleation process of
PuH2 either on the surface or within the bulk of the plutonium metal.
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2.1.8 Continued growth
A study by Saw et al. measured the temperature rise of single hydride reaction sites using infra-red (IR) pyrometry which gave results 85 % lower than predicted by calculations of the temperature rise expected from growth of hemispherical hydride reaction sites [106]. The difference was ascribed to a lower heat concentration arising from hydride that preferentially formed along grain boundaries In their study, the sample was either mounted vertically or at 45º to the vertical for imaging purposes using the IR camera, and it became evident that the spalled reaction product accumulated above reaction sites on the inclined samples, but fell away from the vertical samples. It was considered by the authors of the paper that this spalled product consisted of grains of plutonium, resulting from hydrogen diffusing rapidly along grain boundaries and then forming hydride around each grain, the inter-granular stress causing grains to be wholly separated from the bulk.
More recently the morphology of hydride sites have been investigated by Dinh et al. [107] using cross sectional polishing and focussed ion beam (FIB) analysis. Their work suggested that plutonium hydride reaction sites could grow as either hemispheres or oblate hemispheres. In addition, they observed a mixed phase hydride ahead of the main hydride/metal interface.
2.2 Hydriding work undertaken on other metal- hydride systems
Carrying out research on plutonium is difficult owing to the hazardous nature of the material. Strict limits and restrictions are in place to prevent criticality and spread of radioactive contamination from occurring. It is often helpful to plutonium research to draw comparisons with other less harmful materials, usually termed surrogates or analogues, to develop a greater understanding of the mechanisms involved and how they might be applicable to plutonium. Unfortunately, plutonium does not have any identical surrogate materials similar to the comparison of 238U to 235U, as the only relatively stable plutonium isotope (244Pu) exists in very small quantities. Indeed, the early Manhattan project work used uranium as a surrogate for plutonium before sufficient plutonium was available. It transpired that this was not
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necessarily always a useful comparison [108] but often specific aspects of other systems can be compared to plutonium in the attempt to form a greater understanding.
The materials with greatest significance are discussed first, starting with the other actinide materials, where the bonding is determined by the 5f valence electrons.
2.2.1 5f elements (actinides)
Of the 5f elements only thorium and uranium have isotopes which are safer to handle than plutonium, and consequently most of the research into oxidation and hydriding of the 5f elements and hydriding has been carried out on these materials. Of these, the hydriding behaviour of uranium has by far received more attention due to the use of uranium metal in UK nuclear reactor fuel (Magnox, AGR). Investigations into the hydriding behaviour of thorium appear to be limited to measurement of the hydriding rate [109], and measurement of the properties of thorium hydride. With increasing interest in thorium reactors for power generation, it is expected that the amount of reported work in literature will increase accordingly in the near future.
2.2.1.1 Uranium Uranium is often considered as a surrogate material for plutonium in that it has a relatively high rate of oxidation in dry air of around 5x10-7 mg cm-2 min at 40 °C, rising to around 1.25x10-1 mg cm-2 min at 350 °C [71]. In terms of hydriding behaviour however, uranium only forms the UH3 structure, compared to the compounds of PuH2 and PuH3 available to plutonium, and so may not pose an identical match to that of plutonium. In addition, plutonium is usually stabilised in the fcc δ-phase for ease of working, but α-uranium has an orthorhombic structure. These differences mean that the mechanisms of the plutonium system cannot necessarily be inferred from work undertaken on uranium.
The early stages of the reaction between uranium and hydrogen were investigated by Owen and Scudamore using an in situ hot stage microscope [110]. They found that there were two types of attack on the uranium surface – large areas and a fine dispersion of smaller spots. The large area attack was associated with the presence of uranium carbide inclusions whereas the finely dispersed spots had no
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obvious initiation site. Owen and Scudamore also found that rate of formation (nucleation rate) of hydride sites decreased with increasing thickness of air-formed oxide layers. The authors of that report discussed the factors that cause local attack, termed “spotwise reaction”.
Investigation into the nucleation of hydride sites on uranium has been carried out in order to gain a statistical understanding of any preferred locations of nucleation on the uranium surface [111]. The conclusions drawn from this work were that the time for nucleation of hydrides measured by XRD was consistent with predictions made from the literature. The nucleation of fresh sites slowed by about 3 min, and investigation by SEM showed that initial nucleation had formed heterogeneously, and then coalescence of hydride sites subsequently occurred, forming (in some cases) large grain-like areas which had been preferentially corroded. The corrosion appeared as smooth bumps, presumably subsurface.
The dependency of the gross hydriding kinetics on uranium was conducted in terms of temperature and pressure [93], The results were obtained at both constant pressure and constant temperature, and the findings were that the linear reaction speed was in the region of 6.7 µm min-1 to 8.4 µm min-1, and the apparent activation energy was of the order of 7 kcal mol-1. The conclusions made were that the hydride reaction product presents no significant barrier to the diffusion process, and that the rate of reaction is governed by the precipitation process of the hydride on the metal surface.
Measurement of the diffusion and solubility of hydrogen in the dioxide of uranium was carried out using thermal desorption methods by Wheeler [112]. This work is important because the dioxide species is usually a semi protective surface film on reactive metals. Understanding the solubility and diffusion rate through the dioxide allows calculations to be made regarding the likely induction time that may exist before hydride reaction sites nucleate. Wheeler found that hydrogen had a low -1 -4 -5 2 -1 solubility in UO2 of 0.4 µg g , but a high diffusion rate between 10 and 10 cm s at temperatures between 1250 K and 770 K respectively.
Modern analysis techniques have allowed researchers to investigate the probable mechanisms which determine hydride nucleation and growth. Morrall et al. used time of flight SIMS (ToF-SIMS) to investigate hydrides grown on uranium using
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deuterium to remove any ambiguity about pre dissolved or background levels of hydrogen [113]. This work revealed that the hydride reaction sites formed as UD3.
A micro-textural investigation into the role of twinning and deformation mechanisms on the location of hydride sites on α-uranium was undertaken by Bingert et al. [114]. The work was significant because it was able to establish a statistical link between the location hydride sites and misorientation boundaries using EBSD. This challenges the often-held belief that relative oxide thickness is responsible for the location of the hydride sites. The work was of interest for the case of plutonium because it investigated the effect of low symmetry crystal allotropes (α- uranium is orthorhombic), where deformation occurs mainly by twinning. Later work by Jones et al. confirmed these findings [115].
Recent work on the effect of a hydride on the structure of the surrounding uranium metal has been reported [116]. In their study, Jones et al. first exposed a sample to hydrogen to grow hydride reaction sites and then used a combination of focussed ion beam (FIB) milling and electron backscatter diffraction (EBSD) to identify the extent of deformation to the surrounding uranium microstructure. The results of that work revealed that ductile deformation and crystallographic twinning of the microstructure occurred in the grains abutting the hydride. They concluded that the damage induced during deformation processes increased hydrogen diffusion pathways and made continued growth of hydride reaction sites more likely than nucleation of fresh hydride reaction sites.
SIMS was used to demonstrate that uranium carbo-nitride inclusions in uranium metal act as sites of preferential nucleation on mechanically polished uranium, but were not themselves consumed during formation of hydride [117]. Not all the carbide inclusions had hydride growth associated with them, leading the authors to suggest that the polishing method may have disturbed the protective oxide around a proportion of the carbide inclusions present. Their work also showed that no preference for carbide inclusions occurred on electropolished samples. This is important in a metallurgical approach to plutonium hydride research because plutonium requires electropolishing to remove the Beilby layer introduced during mechanical polishing [118], however the presence of plutonium carbide inclusions in plutonium metal is not common.
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2.2.2 4f elements (lanthanides)
2.2.2.1 Cerium Cerium is often considered a good analogue material for δ-plutonium for a number of aspects for direct comparison in understanding the physical metallurgy of plutonium [18,25,119–125], and alloying studies [126]; to study oxidation [70,75,127– 129]; has been used in both oxide and metal form to investigate the immobilisation of plutonium for long term storage[130,131].
The rate of gaseous reaction of hydrogen with cerium at room temperature was investigated by Gayer and Bos [92], and subsequently at temperatures between 120 °C and 625 °C by Gayer and Melotik [132]. The reaction rate followed sigmoidal function during the testing and proceeded via the via the four stage process (Figure 16) consisting of induction, acceleration, bulk hydriding and termination.
Later work by Sarussi et al. [133] on the kinetics and mechanism of cerium hydride formation observed that no initiation time occurred following a heat treatment to 330 °C in vacuum prior to the hydriding of cerium. The hydriding kinetics were measured barometrically. The cerium was exposed to a range of pressures between 1-32 atm (1x105-3.2x106 Pa) and temperatures in the range of 273-363 K. The outcome of this work was that the reaction front velocity (i.e. the rate at which the hydride layer advanced into the metal) was measured to a reproducibility within 10 % and found to be more dependent on temperature than on pressure, with hydriding rates ranging from 2.7 µm min-1 at 1 atm and 273 K to 72 µm min-1 at 32 atm and 363 K. Activation energies were calculated from the hydriding rates and found to be in the range of 0.14 to 0.25 eV. The hydride in this study formed as a near continuous film on the surface of the cerium sample and the reaction front velocities were discussed to be roughly three times the rate found for other metal hydride reactions.
2.2.2.2 Gadolinium The initiation of hydride forming materials is known to be enhanced by vacuum pre-treatment. An investigation was carried out to assess the effect of vacuum pre-treating of gadolinium at temperature of 470 K [134]. The effects of three different parameters were investigated, namely microstrain, the activation procedure
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used, and change of any barrier to hydrogen diffusion. The effect of microstrain due to the polishing process on initial precipitation of hydrides was considered and investigated by polishing to different surface finishes, thereby altering the degree of microstrain introduced. The degree of microstrain was determined by X-ray line broadening methods. The effect of the activation procedure was measured by comparing the hydriding results for activated samples against deactivated samples. The effect of the hydrogen diffusion barrier was considered in two parts: first were adsorbed functional groups, and second was the role of the oxide.
The effect of the thermal pre-treatment on the gadolinium was to alter the samples in four ways: 1) the mechanically induced stress from polishing was relaxed, 2) the surface oxide partially diffused into the bulk of the metal, 3) carbon species were desorbed from the surface, and 4) chemisorbed hydroxyl groups were also desorbed. All four effects could lead to an increased activation of the gadolinium surface.
The kinetics of precipitation of hydrides and the growth velocity of subsequent growing patches of hydride were investigated using in-situ hot stage optical microscopy [135]. The effect of a low pressure pre-exposure to hydrogen was found to increase the initiation time over that of plain gadolinium.
2.2.2.3 Holmium Samples of holmium have been exposed to gaseous hydrogen at 500 °C and the type and morphology of hydrides formed were characterised [136]. The solubility of hydrogen in holmium at 500 °C is about 0.4 H/Ho, but drops rapidly on cooling. The hydrides formed were observed both at the surface and within the bulk of the holmium samples. The bulk hydrides in the holmium formed in two morphologies, a lamellar form within metal grains, and a patch-type form at grain boundaries.
2.2.3 Transition metals
2.2.3.1 Zirconium and zirconium alloys Significant work has been undertaken into the hydriding behaviour of zirconium, since zirconium is often used in nuclear reactors as the primary containment in the form of cladding for the nuclear fuel, and for other applications
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requiring low neutron capture cross section with good mechanical characteristics [137].
Hydrogen in the zirconium matrix has been shown to cause embrittlement in zirconium alloys, but by far the biggest source of concern has been that of Delayed Hydride Cracking (DHC) [138]. The process of DHC means that sub critical cracks in zirconium can propagate through the material until the critical crack length is reached by cycling hydriding at the strained region near the crack tip, causing embrittlement and extension to the crack. The process repeats, and DHC can remain a problem for a considerable period of time once the fuel components have cooled.
Surface hydrides have been observed in zirconium following hydriding at 800 °C. The hydrides took the form of small clusters on the surface following the electropolishing process. The clusters tended to form preferentially at grain boundaries and at dislocations in the zirconium crystal [139]. The inference was made that these surface hydrides formed during the electropolishing process, i.e. at low temperatures, and form at concentrations as low as 10 ppm. At higher concentrations of around 00 ppm, the zirconium hydride appeared upon cooling as platelets aligned to the 101 0 } direction in the zirconium. Nucleation of hydrides in the case of zirconium showed a preference for grain boundaries, but in all cases in this study, all hydrides were orientated to the 101 0} direction. In later work, the formation of hydrides resulting from hydrogen segregation to the surface and grain boundaries is discussed [140] and tended to be strongest at the Zr(0001) and the Zr( ) surfaces.
2.2.3.2 Titanium The precipitation of titanium hydride has been investigated using techniques such as in-situ TEM, where orientation relationships between the parent titanium lattice structure and the hydride have been found [141]. The habit of the titanium hydrides were found to be {0110}, and at high hydrogen concentrations {0225}. Other studies involving both TEM and XRD have found orientation relationships between the direction and the planes, with hydride nucleation occurring preferentially at interfaces [142]. The α-phase has a lower solubility for hydrogen, and so precipitation occurs first and at low concentrations in this phase; only at much higher concentrations does nucleation
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occur in the β-phase. The morphology of titanium hydrides takes the form of lamella within laths between laths of the parent alpha grain.
2.3 Models of the hydriding process
The models which have been developed to describe the mechanism of hydriding of plutonium have been reviewed and the main points are described below, with a comparison between different models.
2.3.1 Nucleation models
The hydrogen corrosion of plutonium is believed to follow the Glascott model [85,90]. The Glascott model is concerned only with the initiation time before hydride precipitation takes place. It is this calculation that allows the expected amount of hydride formed at some time in the future to be calculated, provided that the nature of the material and the storage conditions are sufficiently well understood. Once precipitation of a hydride has taken place then the continued growth is fitted to empirical data
In the Glascott model [3], the hydriding of plutonium is considered to be analogous to that of uranium. The reasons for this are that both metals are reactive with oxygen and rapidly form a non-protective oxide layer, the oxide species of both plutonium and uranium (PuO2 and UO2) have the CaF2 structure [97]. In addition the observed hydriding reactions appear to follow similar patterns. There are more data on the hydriding of uranium to compare a model to because depleted uranium 238U is much easier to work with than 239Pu, being a β-particle emitter and non fissile, whereas plutonium is an α-particle emitter with a relatively high activity and the fissile nature poses strict limitations on working practices to prevent criticality occurring.
The Glascott model only considers the PuO2 oxide present on the surface of the plutonium because the diffusion of rate hydrogen through Pu2O3 is much quicker than PuO2.
The initiation of the hydriding reaction to be controlled by the oxide limiting diffusion of hydrogen to the plutonium metal [85], this is similar to the situation found for uranium, and helps support the assertion of Glascott that the method for plutonium is analogous to that of uranium. Following from this, the thickness of oxide
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on the plutonium is considered to be the controlling factor in the speed at which hydrogen reaches the metallic surface, and therefore controls the initiation time. This has been observed by McGillivray et al. [90]. The initiation time in the model is the time at which visible hydride spots can be first observed. This time coincides with the point at which hydrogen begins to be consumed in the reaction. The initiation time was also observed by McGillivray et al. to have an inverse relationship to the pressure of hydrogen in the reaction. Glascott suggests that there may be a critical pressure of hydrogen, below which no hydride attack will take place.
The Nucleation rate is considered to be a vital parameter in the hydriding process, where a higher nucleation rate causes the hydrogen to be consumed quicker for a given linear reaction speed. The nucleation rate increases with hydrogen pressure. This means that at high pressure, more hydride sites nucleate than at low pressure, and different distributions occur. Increasing the pressure of hydrogen causes the concentration of hydrogen at the gas/oxide surface of the oxide to increase. In the Glascott model, the variation in nucleation rate with temperature is considered to be the same as that of the variation of hydrogen diffusion in the dioxide with temperature. Analysis of the size of the interstitial spaces in PuO2 suggests that molecular hydrogen is unlikely to be the species diffusing through the oxide due to the molecule being too large, but that diffusion of atomic hydrogen is possible in sub-stoichiometric PuO2-x because the tetrahedral sites in the CaF2 structure which are normally occupied by oxygen atoms allow migration between octahedral sites. At stoichiometric PuO2 and above, the tetrahedral sites are occupied however and the diffusion rate should reduce significantly, Haschke and Stakebake propose that superionic conduction can occur, arising from the high mobility of the oxygen atoms in the CaF2 structure [97]. Defects in the oxide are thought to increase the diffusion rate through the oxide as they can provide short- cuts to the usual route of lattice diffusion through the oxide.
The starting point of hydriding in the Glascott model is the point at which hydrogen is exposed to the outer oxide surface and considers the initiation time to be the time delay until the time at which the terminal solid solution is exceeded and hydride precipitates. Once each hydride site has site has precipitated, the large volume expansion exhibited when Pu becomes PuH2 causes a rupture in the oxide diffusion barrier above the reaction site, allowing ready access to the hydride and
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metal underneath leading to unconstrained growth of the reaction site. The growth of the hydride is then thought to be isotropic.
2.3.2 Growth models
In 1971, a review of the work done into the corrosion properties of Pu was conducted by Stakebake [62]. In it he also dealt with the pyrophoricity and radiation effects. The basis of the review was to help determine the safe storage conditions and safe storage period of plutonium, its alloys and oxides. A more recent review by Haschke and Stakebake [97] has incorporated more recent work into a review of the hydriding process. A brief summary of the main points arising from the reviews are that two main forms of corrosion exist – oxidation and hydriding and the process is governed by adsorption, radiolysis and oxidation mechanisms. Dry air or oxygen gives a limited oxidation rate, whereas moisture in the atmosphere increases the rate of oxidation. Oxidation is quicker in moist inert gases than moist O2 or air, suggesting that a protective oxide forms in moist air, but remains non protective in the inert environment. Plutonium oxide exists on the metal surface consists of two layers,
Pu2O3 in contact with the plutonium and PuO2 at the outer surface. Transport of oxygen to the metal surface occurs via micro-cracks/grain boundaries in the oxide.
Plutonium hydriding exhibits an initiation time, suggesting that the oxide poses a barrier to the diffusion of hydrogen to the plutonium metal. Stakebake suggested that hydrogen can be generated in sealed storage containers as a result of radiolysis, catalysis or thermal decomposition of contaminants from the manufacturing and handling processes or from plastic bags used as for storage.
Disintegration of metal occurs when exposed to moist air, which does not occur in dry air, suggesting that a hydride step is involved. Alloying with some elements reduces the oxidation rate in air; while others increase the oxidation rate, for instance Plutonium alloyed with 1 wt% Ga oxidises more slowly than pure Pu, but hydrides more quickly.
Dinh, et al. conducted a study to evaluate the role of Pu2O3 on the hydriding process. In the paper existing corrosion products, especially PuO2, on the surface of plutonium are considered to strongly influence the initiation. A suggestion was made that the hydride reaction either occurs beneath the Pu2O3 layer via a catalysed
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reaction once penetration of the PuO2 layer occurred via thermal or mechanical means; or that the oxide layer acts as a diffusional barrier to the absorption of hydrogen into the plutonium, and that nucleation occurs as an autocatalytic process at the metal-oxide interface [105]. The reduction of the diffusion barrier provided by the PuO2 layer was discussed to potentially occur via local defects such as micro- cracks, impurities or oxygen vacancies.
The Condon model for uranium [143], and the later Kirkpatrick and Condon model [144] may also adequately explain the continued growth phase of plutonium hydride. They stated that when sufficient hydride product has formed, a compressive strain builds up leading to spalling of the hydride product, this spallation ensures that hydrogen can subsequently gain access to the sample surface, giving rise to the high rate of corrosion of the surface.
2.3.3 General models
Hydride reaction sites are referred to in a number of Bloch and Mintz papers as ‘Growth Centres’ [145], a terminology used by other authors across a number of different hydriding metals. Initially, Bloch and Mintz proposed two models for the hydriding reaction of uranium [93], the first considers hydrogen diffusion through a protective product hydride layer; the second considers hydride growth at the hydride- metal interface. The case of uranium is generally considered applicable to that of plutonium [105]. In the first model, the hydride product formed presents a barrier that reduces the rate at which hydrogen can reach fresh un-reacted metal. The model also suggests that the formation of a constant thickness hydride layer could explain a number of the changes in the kinetics of the hydriding reaction – the constant thickness occurring due to hydride flaking off from the surface once a certain thickness had been reached. The second model considered the hydride to be an easily permeable barrier to hydrogen diffusion and that the growth process is governed by the reaction speed at the metal surface.
A review of the kinetics and mechanisms of surface hydride growth on metals in general was carried out by Bloch and Mintz [146] which further led them to suggest a number of possible mechanisms for the hydriding reaction: interface controlled, i.e. determined by the rate of change from metal to hydride; diffusion controlled, governed by the rate at which hydrogen can diffuse through the Surface
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Passivating Layer (SPL) to the surface; surface controlled, the development of sufficient hydrogen in the correct form for hydriding to occur via process such as chemisorption.
Hydride precipitates once the hydrogen level reaches saturation in the metal. The controlling effect on the initial precipitation is governed by transport of hydrogen to a favourable site and by the mechanical properties of the site. Precipitation occurs preferentially at the surface because the surface has the highest concentration of hydrogen and lowest activation energy. Under certain conditions, precipitation also occurs at other discontinuities in the bulk, such as grain boundaries. The nucleation of surface hydrides is discussed, and the hydride nucleates below the SPL, the volume change induces a strain near the surface, which fractures the SPL. Hydrides tend to favour nucleation at preferred sites, namely grain boundaries or inclusions. Growth of the nuclei is generally spherical, but often some impingement to the spherical growth occurs laterally, leading to the suggestion that certain metallographic features can affect the rate of growth.
2.3.4 Comparison of models
Considerable similarities exist between the models. Most discuss the initiation period prior to initial precipitation of hydride. The models all consider the incubation period as the time taken for sufficient hydrogen to diffuse through a protective surface species such as an oxide layer on the surface of the metal. When the terminal solution for hydrogen in a metal has been exceeded, then the hydride will nucleate and growth of the hydride will take place. The models describe nucleation sites for the relevant metal hydrides as being the result of defects in the oxide. In general, the microstructure of the metal is considered. None of the models discuss factors in the metal which may affect subsequent growth of precipitated hydride reaction sites.
All the models for plutonium hydriding are derived from models covering the hydriding process for uranium or other metals; when the models have been applied to plutonium, they draw heavily on knowledge gained derived from hydriding experiments of uranium and other metals.
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Chapter 3 EXPERIMENTAL
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3.1 Background
The criticality hazards posed when working with plutonium increase the complexity of experimentation, placing restrictions on the amount of mass and the amount of neutron moderating or reflecting materials allowed in the same location. The radioactive and toxic hazards require workstations to be fully enclosed to prevent the spread of contamination. With these restrictions and the additional constraint that limited amounts of material are available for experimentation; the experimental methodology was developed using cerium as an analogous material before applying them to plutonium. Additional work was carried out on the cerium samples once hydride had been formed because insight to the hydriding process could be gained.
This chapter will be limited to the describing the techniques used on the work of the EngD and how any information is generated using that technique. Other techniques are available, but the reason for selection of these particular techniques will be explained in Chapter 4: Philosophy for the Research.
To take full advantage of the low toxicity and lack of radioactive decay of cerium, the work was conducted using different equipment than for plutonium, the equipment was similar although some differences existed. The plutonium work was carried out in full a containment glove box system in an atmosphere of dry nitrogen or under high vacuum. The cerium work was carried out in an open laboratory, under high vacuum or an argon filled glove box as necessary to maintain sample integrity.
3.2 Cerium experiments
3.2.1 Material details
The cerium metal was supplied by Goodfellows as a 12.7 mm diameter 99.9 % pure (rare earth content) rod in an as-cast condition. The main impurity in the cerium was oxygen, which had precipitated as discrete dendrites of CeO2 [147] within the matrix of the metal, observed as black dendritic features in Figure 17. The size and position of the dendrites were closely matched to that of the grains, usually found in the interior of a grain, suggesting that some oxidation had taken place while
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the cerium was liquid during the casting process and that solidification of the metal had nucleated and grown around the oxide dendrites. The cerium had an average grain size of 12.6 ± 4.5 µm based on 102 grains; each grain was measured along the long axis and the short axis.
CeO2 inclusions
Figure 17 Microstructure of the cerium metal used in the studies.
3.2.2 Pre-hydriding sample preparation
3.2.2.1 Sawing A Buehler Isomet 5000 slitting saw was filled with lapping oil as a coolant and the cerium rod was cut into slices 1 mm thick. Water based coolants were not used because the fines were thought to react vigorously with water. The samples when cut were stored in lapping oil to prevent undue oxidation between processing steps. The samples were prepared as they were required to maintain as high a quality surface as possible for the gas dosing experiments.
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3.2.2.2 Cleaning After storage or preparation using lapping oil it was necessary to remove the oil by sonication in ethanol. The lapping oil was not soluble in ethanol; rather this process de-wetted it from the sample and caused it to sink to the bottom of the ethanol in the beaker. The samples were then stored in dry ethanol until they could be loaded into the gas reaction chamber.
3.2.2.3 Surface preparation The samples for chapter 5 were prepared by hand grinding samples with 600 grit SiC papers under lapping oil. The sample was then cleaned as described above. A rough surface was retained for the test because experience has shown that greater contrast exists between the hydride reaction sites and the metal surface when a rough surface is retained. This preparation method also most closely matches that used previously for plutonium work [90].
Hand polishing cerium with diamond suspensions in lapping oil gave an unsatisfactory finish; therefore a water based polishing procedure was developed. The samples used in chapters 6 and 7 were hand ground to 4000 grit using SiC papers using water as a coolant The samples were then polished to 1 µm diamond using water based Buehler MetaDi diamond suspension, which gave a good finish.
3.2.3 Exposure to hydrogen
The hydride reaction rig consisted of two parts, the gas handling rig and the reaction cell.
3.2.3.1 Gas handling rig A simple diagram for the gas handling rig is shown in Figure 18. Gas was supplied by a Parker Balston hydrogen generator, at a purity of >99.9999 % which was used to fill LaN5 beds. This process purified the hydrogen further. A storage volume was then charged from the LaNi5 beds and then isolated. All dosing of the reaction cell and sample was made from the storage volume, which allowed measurement of the gas consumed. Gas admittance to the reaction cell was made using an MKS 248 control system which measured the pressure from a 722A MKS Mini Baratron and controlled the pressure using an MKS 248 control valve. Intertek Orchestrator measured the gas pressure and the temperature within the reaction cell
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and hydriding rig. The temperature was measured using a K-type thermocouple. All other actuation of valves on the rig was under manual control.
Figure 18 A Process & Instrumentation Diagram for the cerium hydriding rig.
3.2.3.2 Reaction cell The cerium was exposed to hydrogen in a visual reaction cell. The reaction cell was constructed from standard conflat (CF) fittings, consisting of a CF40 double nipple, a CF40 window which allowed the reaction to be viewed and a CF40 flange with a thermocouple feed through and a diaphragm valve to a Swagelok VCO connector. The VCO connector was a Viton elastomeric seal, which allowed the cell to be readily dismounted from the gas handling rig to be loaded and unloaded in a fume cupboard or glove box as necessary. Use of the visual cell allowed direct imaging for the growth and allowed a decision to be made about termination of the experiment. The entire cell could be heated. Loading the reaction cell was carried out in laboratory air, a process which took 30 min. This ensured that the samples all received the same air exposure.
3.2.3.3 The hydriding reaction The reaction cell was loaded and connected to the gas handling rig and entire rig was evacuated using turbomolecular pumps backed by rotary vane pumps, the pressure was measured at 10-8 mbar (10-6 Pa) at the vacuum pump. The CCD
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camera was focussed onto the sample surface and the reaction cell was heated to 30 °C and was allowed to pump overnight.
The basis of the hydriding procedure for cerium was to deliver very pure hydrogen to cerium within a reaction vessel with a good vacuum at a defined temperature until hydride sites could be observed growing. The pressure of hydrogen was fixed per experiment, but varied over the course of a number of experiments from 10 mbar (1 x 103 Pa) to 1000 mbar (1 x 105 Pa). The sample used in chapter 5 was exposed to hydrogen at a pressure of 100 mbar (1 x 104 Pa), the samples used in chapters 6 and 7 were exposed to 10, 50, 300 mbar (1 x 103, 5 x 103, 3 x 104 Pa). The hydrogen was kept at constant pressure (within 5 %) during the test using the pressure controller. The samples were viewed using a CCD camera, with the sample nearly filling the field of view, and images were acquired using Image Pro Plus as the reaction proceeded.
3.2.4 Post hydriding sample preparation
Following completion of the gas dosing experiments, the reaction cell was taken to an argon filled glove box to be opened. There was a concern that the hydrides would react vigorously on contact with air, an inevitable step for these samples during loading and handling procedures. The surfaces of the samples were swabbed in the glove box using a cotton bud in an attempt to remove loose hydride material. The samples were stored in an argon filled glass vessel until they could be analysed. No changes were observed while the samples remained in the glove box or the airtight vessel. As it was necessary to transfer the sample through air into each instrument for analysis, one of the samples was removed from the argon and placed in laboratory air for 60 s before being returned to the argon. A slow reaction occurred; the sample showed signs of degradation over the next 20 min, with a visible flakes spalling from the hydride reaction sites.
A number of analyses were carried out on the sample in an ‘as-is’ state without any further preparation, these were: optical microscopy, SIMS and SEM. Ion beam milling was used on a number of hydride reaction sites in conjunction with SEM to remove the surface and investigate the sub surface structure.
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The remaining work on the samples namely: optical microscopy, SEM and EDS, Vickers hardness indentation and nanoindentation required cross sectional polishing. To do this, the samples were mounted edge-on in epoxy resin and ground using 2500 grit SiC papers until the required hydride site was revealed. This process was laborious, because the hydride sites were small (50 to 100 µm) and sparsely populated on the surface of the samples, and involved an iterative process of grinding followed by polishing to 1 µm diamond. Once the hydride had been revealed, the sample was polished using Buehler MetaDi 6 µm, 3 µm, and 1 µm diamond suspensions before a final polish using Struers OP-S 40 nm colloidal silica suspension.
3.2.5 Light microscopy
Conventional bright field light microscopy was used on the cerium because the hydrides exhibited strong contrast against the oxidised cerium surface. The hydride reaction sites on the surface of the samples were imaged with no other preparation. Cross sections were also analysed. Following surface preparation, there was a time window of around 10 min to acquire any data before growth of fresh oxide had obscured any features of interest. Bright field light microscopy involves illuminating a surface of interest using light and projecting the reflected light through a lens onto a detector such as a CCD detector to record an image. The microscope used as a Zeiss AxioImager M1m with lenses of magnification 2.5, 10, 20, 50 and 100.
3.2.6 Secondary Ion Mass Spectrometry
Secondary Ion Mass Spectrometry (SIMS) was used to confirm that small spots found on the surface following the hydriding procedure were associated hydrogen, confirming early stage formation of hydrides. This was especially important for smaller hydride reaction sites grown on cerium which were too small to be observed in situ. The SIMS investigation confirmed that the small sites observed by post experimental microscopy was cerium hydride.
SIMS directs a primary beam of ions (either positively or negatively charged) onto the surface of the sample of interest. This study used oxygen ions (O2-) to sputter sample craters. The primary ions impact the surface and cause a damage
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cascade. Secondary ions from the cascade are liberated from the sample surface by a sputtering event. The secondary ions are then collected and passed through a mass spectrometer, and the mass to charge ratio is used to identify the atoms or molecules liberated from the sample surface. The mass spectrometry in this work used a Cameca IMS 3f double focussing magnetic sector spectrometer, giving high- resolution serial-acquisition. The magnetic sector is a 90° arc within a magnetic field which the secondary ions are passed through (refer to Figure 19 for a diagram). Charged particles follow circular paths in a magnetic field and for each mass/charge ratio with a given kinetic energy a particular radius is obtained. Using this radius, very tightly defined m/z ratios can be measured. Varying the magnetic field strength selects a different mass/charge ratio and acquisition of an entire spectrum can be built up in a serial manner. The SIMS instrument used also had an electrostatic sector which reduces the energy spread in the ions which occurs when they are removed from the sample. This works because particles in an electrostatic field travel in parabolic paths, so particles with greater energy travel deeper into the electrostatic field than less energetic particles, and the energy spread is then minimised upon refocusing. As the particles liberated are often large in size, the surface sensitivity of the SIMS process is very high, emerging from about the first three atomic layers into the bulk of the sample [148]. This technique was not available to use on this EngD for plutonium.
Figure 19 Diagram of a double focussing magnetic sector analyser.
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3.2.7 Scanning Electron Microscopy
The cerium samples were imaged using a JEOL JSM7000F Field Emission Gun Scanning Electron Microscope (FEGSEM) to obtain high resolution images of the hydride reaction sites. SEM is primarily an imaging technique which accelerates electrons to energy between 1 and 30 keV and subsequently focussing them onto the sample surface. When electrons are aimed at a sample they interact with the matter in a small interaction volume [149] below the beam spot. The signal of interest emerges from different amounts of the interaction volume, depending on the energy of the signal and the sample that it must travel through to emerge. Secondary electrons emerge only from a small volume near the beam spot on the surface, but backscattered electrons and X-rays can emerge from a considerable distance within the interaction volume (Figure 20). This distance is dependent on a number of factors, most significantly the energy of the primary electrons, and the atomic mass of the sample and the energy of the emerging photon. The signal emerging from the sample is scanned in a raster fashion and the intensity generated at any point is collected and re-constructed on the display screen. For this work, the accelerating voltage on the FEGSEM was set to 15 kV for cerium and plutonium. The diamond nanoindenter tip was imaged at 5 kV to minimise charging.
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Figure 20 The interaction volume of electrons in a bulk sample and generation of signal in an SEM.
3.2.7.1 Secondary Electron Imaging Secondary Electron Imaging (SEI) is formed using secondary electrons, and is the highest resolution technique within the SEM. The secondary electrons are produced when the primary electrons from the beam scatter inelastically from the outer electrons of an atom in the matrix, transferring a random proportion of the incident energy to the atomic electrons. In some cases the energy transferred is greater than the binding energy, sufficient to ionise an atom, the electrons which escape the atoms are called the secondary electrons. The secondary electrons with enough energy to overcome the work function of the material then exit the sample. The energies of the secondary electrons are therefore considerably lower than the primary exciting beam and are given by . The low energy means that only the electrons near to the surface emerge (Figure 20), giving the small volume important in high resolution imaging. Because secondary electrons energies
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are low, they are accelerated towards the detector with an electric field to increase the brightness of signal. Image contrast is formed as a function of the number of electrons which emerge from the sample preferentially towards the detector because they have a lower probability of escaping the electric field. When the surface impinged by the beam is locally orientated towards the detector, this gives higher secondary electron yields and they are usually used to image surface topography in high resolution for this reason.
3.2.7.2 Backscattered Electron Imaging Backscattered Electron Imaging (BEI) is formed using the primary electrons scattered elastically by the nucleus of the atoms in the sample. Should an electron in the primary beam pass sufficiently close to the nucleus, then the positive charge of the nucleus can cause them to be backscattered to high angles (up to 180º). The backscattered electrons are collected by an annular detector around the lens aperture. As the scattering interaction in BEI is elastic, the primary electrons emerge from the sample with the same energy as they entered, except where additional scattering events have occurred. Backscattered electrons can therefore emerge from deeper within the sample than that of secondary electrons, and there is a reduction in resolution as a result (Figure 20). The image contrast is mainly formed as a function of the average nuclear electronic charge for a region, as a higher charge gives a higher probability of scattering an electron towards the detector. It is particularly appropriate to examine hydride material because the hydride has a lower density than the cerium or plutonium matrix that contains it, giving strong contrast. Other contrast mechanisms exist, particularly that of channelling contrast [149], where atomic planes aligned with the electron beam allow deeper penetration of the electrons, meaning that on average they return fewer backscattered electrons than planes of atoms at a greater angle to the beam. The increased diameter of the interaction volume that backscattered electrons emerge from reduces the effect of topography on signal and degrades the spatial resolution. BEI is usually used to image materials where different phases of average atomic number exist.
3.2.7.3 Energy Dispersive Spectroscopy Energy Dispersive Spectroscopy (EDS) is a technique used commonly on SEM to detect and measure characteristic X-rays generated by the interaction of the
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electrons with the material within the interaction volume. Energy is transferred from the electrons in the primary beam to the electrons in the atoms in the bulk of the sample. The energy transferred to the atomic electrons is sufficient to promote the electrons to higher energy levels within the atom, which subsequently decay to lower energy levels via the release of an X-ray, equal to the difference between the energy levels. Since the energy levels for each atom are characteristic of the atomic species, this technique allows quantitative speciation of the elements in the sample to the 0.1 wt% level [149]. The limitations of this technique are that the X-rays of elements below beryllium (which includes hydrogen) do not have sufficient energy to excite current generation of EDS detectors. The peak width of EDS is relatively wide, meaning that some overlaps in characteristic X-rays can obscure trace elements; the generation of X-rays occurs within the entire interaction volume, meaning that the spatial and depth resolution of EDS is much lower than either SEI or BEI.
3.2.8 Sample coating
Electron bombardment in an SEM causes a build up of negative charge in the surface of any poorly conducting material. Samples mounted in resin for polishing are insulating and build up of charge in an area under investigation, which can affect the positioning of the electron beam, and would not be an accurate reproduction of the sample surface particularly at low accelerating voltages. Cerium samples were coated using carbon evaporation in an Emitech K 750 turbo pumped coater. Thin films of evaporated carbon do not heavily attenuate X-rays, which allows detection of light elements using EDS.
3.2.9 Ion milling
The ion milling of cerium hydride reaction sites was carried out on a Gatan Precision Etching and Coating System (PECS) using Ar ions. The PECS uses a penning gun within the processing chamber to ionise the Ar. The argon ions are accelerated towards the sample and impact the sample, transferring kinetic energy to the sample surface which causes material to be sputtered away from the surface. Material sputters at different rates for different species and crystal orientations, which can help reveal subsurface microstructural features.
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The PECS had a base pressure in the range of 10-3 Pa. Laboratory grade argon gas was admitted into the chamber, raising it to the operating pressure of 10-2 Pa. Etching of the sample in chapter 5 was carried out at normal incidence, 300 µA, 8.6 keV in steps of 30 min. The sample was transferred between from the PECS to the FEGSEM for imaging between etching steps. Each air transfer was completed in ca. 30 s. The process took around one week to complete, and the experiment was terminated when the penning etching gun on the PECS began to short out due to build up of cerium whiskers internally.
3.2.10 Vickers hardness
The hardness of cerium and cerium hydride was initially measured for the paper in chapter 5 using conventional Vickers diamond pyramid hardness indentation. Vickers indentation typically leaves a large footprint on soft materials such as cerium and for this work the load selected was the minimum setting on a LECO V-100-A2 Vickers hardness indenter at 1 kg force. This measurement was only possible using the polished cross section of a sample where a very large hydride site (ca. 0.6 mm diameter) had been allowed to grow until sufficient hydride for three indentations was available.
Vickers hardness indentation uses a pyramidal diamond with an angle of 136° between opposite faces, or a face angle of 68° from the normal to the sample. The load was applied over a time of 10 s to allow for creep and the resulting indentation was measured across the diagonals. The load is applied by the area of the sample contacted by the diamond indenter and as such a stress is applied to the sample. The hardness of a sample is given by Hardness = Load/Contact Area. The contact area and the diagonal measurements of the indent are directly proportional and a conversion factor related to the geometry gives the Vickers hardness number, but since hardness is the stress at which no further penetration occurs into the sample it is now more usually quoted in MPa or GPa.
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Figure 21 Indentations made in cerium metal (left) and cerium hydride (right)
3.2.11 Nanoindentation
Nanoindentation carried out on the interface between cerium and cerium hydride to determine the relative hardness between the two phases and to map any strain fields around the hydride. In nanoindentation, a hard, sharp tip of well defined geometry is forced into the surface under investigation at a given load. The depth of the indentation is a measure of the contact area between the tip and the sample. The contact area and the load apply a pressure or stress to the sample (force per unit area). The final depth occurs when the load does not cause further indentation to the sample.
At the very low loads used in nanoindentation (typically mN), the shape of the tip is critical in obtaining sensible results. A pyramid is difficult to manufacture so that all faces meet at a single point, so a three sided pyramid is used instead. Generally, there are two types of nanoindenter tip, the Berkovich and the cube corner. The Berkovich has a face angle of 65.27° and is designed to have the same depth to contact area relationship as the Vickers indenter tip [150]. The relationship between 2 depth and area is given by A=24.5hc . The Berkovich indenter tip is the most commonly used nanoindenter.
A cube corner indenter was used for this work however for a number of reasons. Berkovich indenters indent to a shallow depth but the contact area is high. The aim of the nanoindentation was to investigate the difference between hydride and metal and any interfacial region. The Berkovich footprint would have been too large to provide any reasonable fidelity of data. The cube corner has a face angle of
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35.26°, making much smaller contact area for the same amount of indentation depth. 2 The relationship between depth (hc) and area (A) is given by A=2.60hc [150] The small contact area allowed a greater number of indentations in the region of interest.
The rapid oxidation of cerium and the low sampling depth of nanoindentation would have rendered the surface unsuitable for study by nanoindentation, because each successive indentation would see further oxidation of the sample, giving differing results each time. This meant that within a short time after final surface preparation, the nanoindentation would measure the oxide and not the metal or hydride. Cube corner nanoindentation was available in close proximity to the polishing equipment used to cross section the cerium samples using a Veeco EnviroscopeTM Atomic Force Microscope (AFM). The EnviroscopeTM had an integral vacuum chamber to maintain the condition of the sample surface capable of an operating pressure in the order of 10-6 mbar (10-4 Pa). This was achieved using turbomolecular pumping backed by a scroll pump. The sample was transferred into the AFM directly following the colloidal silica polish and the AFM vacuum chamber and evacuation commenced within 60 s of completion of polishing; even so, some oxidation will have occurred in that time.
Some drawbacks exist when using an AFM style nanoindenter. The load applied to the tip cannot be measured directly, but is instead calculated from the spring constant and deflection of the cantilever using the formula F=-kx. Each nanoindentation cantilever is supplied with a nominal value of spring constant (k), however it was sensible to confirm this value. This was done by measuring the dimensions of the cantilever using a JEOL JSM7000F FEGSEM. The spring constant for a cantilever is given by the formula k= (Ewt3) / (4L3) [151]
The table below shows the nominal values for the cantilever vs. measured values. The width and thickness measurements were accurate; however there was a significant discrepancy in the length of the cantilever being 315.4 µm instead of the nominal 350 µm.
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Nominal Measured Young's Modulus for Stainless steel E (N m-2) [152] 195 GPa 195GPa Length, L (µm) 350 315.4 Width, w (µm) 96.9 96.9 Thickness, t (µm) 12.9 12.9 Spring constant, k (N m-1) 237 323
A value of 323 N m-1was taken as the actual spring constant based on the full length of the cantilever as is conventional. Careful consideration of the physics involved in the use of a cantilever style nanoindenter to measure the force leads one to consider the following: the nominal spring constant is taken from the full length of the cantilever, measured at 315.4 µm (A in Figure 22, top), however the diamond is adhered to the cantilever so that the force is not acting at the full length but at the tip of the diamond (B in Figure 22, top), which in this case is at a length of 286.2 µm. The situation is yet more complex however; the extra thickness from a tetrahedral diamond with a higher young's modulus (1220 GPa) would heavily affect the actual spring constant over the portion to which the diamond is adhered. This would all serve to make the cantilever stiffer than the nominal value over this portion. The part of the cantilever that did not have the diamond adhered to it, i.e. the part free to deflect, was 233.3 µm in length (C in Figure 22, top and Figure 22, bottom. The spring constants are shown for each measurement in the table below.
Measurement Description Length (µm) Spring Constant (N m-1) A Total length 315.4 323 B Length to tip 286.2 433 C Free length 233.3 799
The tip of the diamond was inspected in the FEGSEM following completion of the work. No evidence of blunting was observed after the measurements had been completed.
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Figure 22 Top: diagram indicating three possible lengths that could be used to calculate the spring constant of the cantilever; the overall length (A), the length to the tip of the indenter (B), the freely deflecting length (C). Bottom: measurement of the freely deflecting portion of the cantilever.
A grid of increasing loads from 0.5 V deflection (32 nN) to 2.75 V deflection (371 nN) was made on cerium metal in steps of 0.25 V (16 nN) to select an appropriate load to use on the hydride interface (Figure 23). The 2.75 V deflection indentations gave indents that were of an appropriate size to fit a number across the hydride interface without overlap and retain a size that would allow reliable measurement on an uneven surface and so it was rounded to 3 V deflection (405
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nN) for the first set of indentations on the cerium hydride/cerium interface. A second set of indentations was made at 10 V deflection (1.35 mN) because the 3 V deflection set were visible in the cerium, but difficult to observe in the hydride.
Figure 23 A 10 x 10 grid of increasing nanoindenter load (r to l) from 0.5 to 2.75 V at steps of 0.25 V. The load increased right to left.
The average column values for indent depth (i.e. hardness measurement) from the grid in Figure 23 and the average values for the indentations made in cerium metal at 405 nN and 1.35 mN were plotted against the apparent hardness, calculated using the load applied gave a non linear response, meaning that the depth of indentation was not proportional to the load applied, the apparent hardness had an increasingly steep trend toward lower indentation depths a graph was plotted (Figure 24). This has been observed previously for other metals and is referred to as an indentation size effect [153] where the hardness increases with decreasing indent depth. The apparent hardness is a sum of the actual hardness and the surface
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effects. This is most probably occurs due to oxidation occurring during the approximate 60 s air transfer or the use of water based polishing media and work hardening from the polishing. Fitting a power series fit to this graph and extrapolating suggested that, the indent depth on this material would need to be in the order of 1– 1.6 µm to ensure that the contribution from the work hardened material plus oxide was negligible, i.e. comparable with earlier nanoindentation work on cerium [154]. Obviously, this degree of extrapolation is excessive, but the instrument was not capable of delivering micron sized indentations and, as such, the absolute values of hardness were not used, instead favouring relative measurements for each phase. This technique was not available for the EngD project to use on plutonium.
1000000
100000
10000
1000
100
y = 64700x-1.699 Apparent Hardness Apparent Hardness (GPa) 10 R² = 0.9335
1
0.1 0 50 100 150 200 250 300 350 400 450 500 Indent Depth (nm)
Figure 24 Power fit of the data plotted for the indentation depths and apparent hardness.
3.2.12 Atomic Force Microscopy
Prior to and following nanoindentation the sample was imaged in the AFM using the nanoindenter tip. This is not ideal as an imaging tool because the stiffness of the cantilever makes it less sensitive than a standard cantilever and the mass of the diamond again reduces the sensitivity. Standard AFM cantilevers are usually considerably sharper than a nanoindenter, giving better spatial resolution. Using the
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nanoindenter tip for subsequent measurement means that the tip is as large as the indent it is used to measure, potentially not measuring the full depth of indent; however this was mitigated by using a slow scan rate of 0.5 Hz. The resonant frequency was sufficiently high that a number of tip interactions should occur over the deepest part of each indentation. It was necessary to use the nanoindenter tip for imaging because opening the vacuum chamber to air to change cantilevers would have allowed the cerium sample to oxidise further, potentially obscuring the nanoindentations, and re acquiring the indentations using a second cantilever is difficult.
The AFM used in the investigation of the material surrounding the hydride site was a Veeco EnviroscopeTM. This AFM is unusual in that it is housed within a medium to high vacuum chamber. This allowed significant benefits when working with freshly prepared surfaces of cerium in that the oxidation of the surface can be slowed to an extent where investigations can take place of the metal without oxidation obscuring the details and altering the results.
The AFM is capable of most of the usual modes of operation, however for this application TappingMode AFM was used to image the sample. TappingMode imaging uses a cantilever vibrating near its resonant frequency as the surface sensing device. As the cantilever is brought close to the surface, the interaction of the tip with the surface causes a decrease in the amplitude of vibration. Feedback control on the AFM maintains the amplitude decrease at a constant value following the surface shape as it is scanned in the x and y directions by moving the cantilever in the z direction as appropriate. The movement in the z direction is plotted as the tip is scanned in x and y across the sample to give the height map of the sample. Because the tip interacts with the sample, each interaction loses a small amount of energy to the sample matrix. The amount of energy lost per interaction is a function of material stiffness and is exhibited as a change in the phase angle between the oscillation of the cantilever in free space and the cantilever oscillating in contact with the surface. A map of the phase angle lag gives a map of differing material properties [155]. This technique was not available for the EngD project to use on plutonium.
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3.3 Plutonium experiments
The techniques used to investigate the hydriding of plutonium are limited to the facilities that are specifically installed for plutonium analysis and are made available to this EngD project. All work on plutonium was carried out in a full containment, nitrogen purged glove box system.
Four samples were made available for this project: two δ-stabilised Pu-Ga alloy samples with ca. 50 µm grain size; one electro-refined unalloyed Pu sample with a grain size ranging from 20 µm to over 100 µm with torturous grain boundaries; one mixed phase Pu-Ga alloy, with a 32/68 % α-Pu to δ-Pu mixture with non- equiaxed grains ca. 30 µm across by >100 µm in length.
3.3.1 Preparation for hydriding reaction
12 x 12 x 2 mm tiles of heat-treated δ-stabilized Pu-Ga alloy were punched from plutonium sheet using a 12 x 12 mm square punch in a fly press. Plastic deformation was found in cross section (Figure 25) to extend a distance into the bulk of around 200 µm and was exhibited as heavily distorted grains. The surfaces were ground to 400 grit using a Buehler Minimet and dry ethanol. This coarse surface preparation was carried out to ensure a low reflectivity to aid image capture and determination of the interface position between hydride material and the ground surface. There may be a degree of damage imparted to the subsurface metal of the sample from the grinding process, which could take the form of plastic distortion, recrystallisation, reversion to α-Pu, and/or the formation of a Beilby layer. XRD showed an amount of α’ present following grinding (Figure 26), which may in turn affect the structure of any oxide formed and potentially how hydride reaction sites could nucleate. The samples were cleaned using ethanol, loaded into the reaction cell and evacuated to 10-8 mbar (10-6 Pa).
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Figure 25 Distortion from the punching process extends a distance of around 200 µm into the surface.
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4500 As polished PuO2 At 180°C 4000 Pu2O3 Room T Room T after 24h 3500 At 290°C PuO2 Room T 3000 δ-Pu Room T after 2 days Pu2O3 2500 δ-Pu α-Pu 2000
1500 Intensity (a.u.) Intensity 1000
500
0 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 2q (°)
Figure 26 A high resolution XRD trace of the region that contains the main PuO2, Pu2O3 and δ-Pu peaks. Also visible in the as-polished trace is an α-Pu peak. A number of thermal treatments were subsequently carried out in situ to assess the stability of the various phases present.
3.3.2 Surface analysis preparation
Tiles were cut from discs of cast material using a hacksaw before filing to size 10 x 10 x 3 mm. The samples were prepared by a standard routine, which was: mounting in epoxy resin, grinding until plane using diamond grinding discs and lapping oil as a lubricant before polishing to 1 µm using diamond pastes and lapping oil. The samples were subsequently broken out of the epoxy resin and electropolished to reveal grain detail. The electropolish used 10 % nitric acid in ethylene glycol as the liquid on a Struers LectroPol. Electropolishing has been shown to remove the contributions of subsurface sample damage arising from mechanical polishing [118], however, small amounts of reaction product between the plutonium and electropolish chemicals may still be present after this step.
3.3.3 Cross sectional analysis preparation
For cross sectional analysis, the samples were prepared using the same routine as above but were not broken out of the mount before electropolishing.
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These samples were then coated using evaporated aluminium to provide a conductive coating and sent for analysis in the FEGSEM.
3.3.4 Sample coating
For plutonium, the samples were coated with evaporated aluminium using an Edwards 306 Evaporator which is incorporated into a glove box. Aluminium is used because it is difficult to use and maintain a carbon braid coating system in a glove box fitted with thick leaded gloves. This does have drawbacks because X-rays are attenuated by the aluminium coating, particularly those of lighter elements such as oxygen, as a result, EDS of the hydride cross sections in chapter 8 could not show the extent of oxidation which occurred during passivation. The plutonium oxide formed is insulating, therefore sample coating is necessary for conductivity of the surface and across the resin mounting medium and minimises surface oxidation in the glove box and sample transfer.
3.3.5 Ex-situ hydriding
The design of the hydriding rig for plutonium (Figure 27) was similar to that used on cerium (see Figure 18) although it is housed within a dry nitrogen glovebox and has a nitrogen dilution tank to avoid the potential build up of an explosive atmosphere in the exhaust system. The plutonium hydriding reaction cell is permanently connected to the hydriding rig within the glove box, with sample loading via a dedicated loading port. The sample is mounted vertically on a heater with the CCD camera above it. Gas admittance to the reaction cell was made using an MKS 248 control system which measured the pressure from a 722A MKS Mini Baratron and controlled the pressure using an MKS 248 control valve. Intertek Orchestrator measured the gas pressure and the temperature within the reaction cell and hydriding rig. The temperature was measured using a K-type thermocouple. In addition, the valves can be operated remotely using pneumatic actuators via Programmable Logic Controllers.
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Figure 27 A schematic diagram of the plutonium hydriding rig. The majority of the system is housed within a nitrogen filled glove box described by the orange rectangle.
3.3.6 Scanning electron microscopy for plutonium
A JEOL JSM7000F Field Emission Gun Scanning Electron Microscope (FEGSEM) was used for analysis of the plutonium samples. The FEGSEM is unique to AWE and has been designed specifically for analysis and experimentation on plutonium. The FEGSEM comprises three chambers (Figure 28). The first chamber is the analysis chamber of the SEM where the imaging and analysis take place. The second chamber is a transfer chamber used to load samples from the glove box and move samples between the first and third chambers. The third chamber is a preparation chamber which houses an Ar ion gun, sample heating and cooling stage and a gas dosing capability. The adjoining chambers are designed so that the plutonium sample is kept under vacuum between experimentation and analysis, with movement by magnetically coupled transfer rods, to prevent oxidation of the hydride sites between hydriding and analysis.
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Figure 28 Top: The plutonium capable FEGSEM. Bottom: A simplified three dimensional representation of the FEGSEM.
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3.3.6.1 Analysis modes The imaging modes used on the plutonium samples were SEI and BEI for high resolution imaging and material contrast imaging respectively. The principle of operation for both analysis modes were discussed earlier. EDS was not used for the EngD project because hydrogen is not detected by this technique, and the aluminium coating prevented acquisition of oxygen X-rays from the sample.
Electron Back-Scatter Diffraction (EBSD) uses backscattered electrons generated near the surface which are subsequently diffracted by the near surface material into patterns of Kikuchi lines. The Kikuchi patterns are representative of the orientation that the crystal presents to the surface, and maps of the crystallographic orientations can be built up by rastering across the surface. EBSD has a spatial resolution of around 20 nm for heavier elements in an elliptical spot due to tilt and with a depth resolution of around 10 nm [149], is very surface sensitive and requires the sample to be orientated towards the detector.
Numerous attempts were made to characterise the orientation of crystal grains in the plutonium samples used during this investigation using EBSD prior to dosing with gas, but to date this has been unsuccessful. Delays in commissioning the FEGSEM for radioactive samples meant that there was insufficient time to optimise this technique prior to carrying out the work for the EngD. There are a number of possible variables which need further investigation, such as the depth of the oxide layer on the sample prior to ion milling. Previous attempts by others [121,156] have obtained Kikuchi patterns on δ-stabilised Pu, but in their work they had the advantage of a Auger spectrometer which allowed them to measure definitively when the oxide layer had been removed by Ar ion milling. The failure to obtain Kikuchi patterns could be because the polished surface may consist of distorted material in a Beilby layer [157], however plutonium oxidises readily under any conditions where even trace amounts of oxygen are available, including high vacuum with a partial pressure of 1 x 10-7 torr (1.33 x10-5 Pa) of oxygen [158]. An oxide with an amorphous or crystalline structure substantially smaller than the beam interaction volume would prevent Kikuchi patterns being obtained.
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3.3.6.2 Sample transfer to the FEGSEM To ensure that the samples were not unduly exposed to air between glovebox processing and operators are not exposed to plutonium contamination, the movement of plutonium from the glovebox to the FEGSEM employs a double door Sample Transfer Vessel (STV). This method maintains the glove box atmosphere from undocking of the STV from the glove box until docking to the FEGSEM. While every effort is taken to ensure that contamination does not escape from the STV during the docking and undocking process a possibility remains that it could occur. To mitigate the effects should this occur the loading/unloading procedure is a protected operation; participating staff are required to wear full face respirators with particulate filters, hoods, coveralls, overshoes and two pairs of gloves with the first pair taped to the coveralls to carry out the process. The author was involved in these loading and unloading operations by physically carrying out the process and directing the loading and unloading aspects involving the FEGSEM. The docking and undocking process from the glove box system was under the direction of the technical team supervisor. The entire operation was supervised by nuclear material controllers and health physics surveyors. Once the STV is been docked onto the FEGSEM the double door system can be opened and the transfer port and STV are evacuated to rough vacuum and the sample is then inserted via a gate valve by a magnetically coupled transfer rod. Removal of the STV requires filling the STV to glove box pressure with N2 before returning to the docking port on the glove box.
3.3.6.3 Ion milling Ion milling uses a Vacuum Generators EX05 Ar ion gun mounted on the preparation chamber. The ion gun is evacuated via a differential pumping system to ensure that the pressure in the ion gun is ca. 10-7 mbar (10-5 Pa), Ar gas is then bled into the ion gun where it is ionised by bombardment of electrons from a hot wire filament. The argon ions are then accelerated towards the sample and focussed using electrostatic lenses.
The conditions used for ion milling were 3 kV accelerating potential, 15 mA filament current, chamber base pressure 10-5 Pa and argon pressure in the ion gun 10-2 Pa. The beam current cannot be measured directly because the stage has piping attached to it to convey the cooling nitrogen, which is earthed. The beam was
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defocused to an approximate 1 mm diameter spot on the sample surface, and the sample normal was tilted to 72° to the incident ion beam.
3.3.6.4 In-situ hydriding The potential risk of a hydrogen explosion within a vacuum chamber contaminated with radioactive material places stringent controls on the quantity and availability of hydrogen gas for experimentation purposes. The maximum quantity allowed to be available for use with the FEGSEM is 100 cm3 at NTP. This is achieved by only having one bespoke 1 cm3 cylinder designed to fit the FEGSEM manifold to be filled to 100 bar (1 x 107 Pa) at any one time. In practice this was achieved by having an external company certify a fill pressure and supply cylinders when they are required. This placed an inherent delay in re-exposing a sample to hydrogen. The gas dosing system on the FEGSEM is shown in Figure 29 and consists of a H2 dosing cylinder, manometer, metering valve, roughing pump and pirani gauge. Sample heating and cooling is also available.
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Io n gun
Figure 29 The Process & Instrumentation Diagram for the FEGSEM gas dosing system
The preparation chamber is nominally 30 L volume, and the gas dosing system is designed to give a maximum dosing pressure of 3.3 mbar (330 Pa). The volume has been calibrated using a known volume of gas and found to be 30.49 L which gives final pressure of 3.28 mbar (328 Pa) of gas in the chamber, i.e. correct within 1 % of the design pressure.
The chamber walls act as the primary containment system preventing spread of radioactive contamination to the surrounding laboratory. The gas dosing cylinders 7 are sized so that when filled to 100 bar (1 x 10 Pa) of H2 the amount of gas present is insufficient to create an explosion or deflagration which might challenge the integrity of the chamber walls. The margin of safety calculated on the gas available from the cylinder volume is in excess of 500 %. This means that should an error in filling pressure of the cylinders occur, a cylinder would need to be filled to a pressure in excess of 500 bar (5 x 107 Pa) to challenge the integrity of the containment by a
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pressure wave arising from explosion should air be inadvertently be allowed into the preparation chamber and a source of ignition be present. A burst disc protects the chamber should the internal pressure increase beyond 300 mbar (3 x 104 Pa) above atmospheric.
The sample can be heated up to 500 °C using resistive heaters built into the sample stage or cooled to -150 °C by passing dry nitrogen gas through a liquid nitrogen cooled heat exchanger and then through the sample stage. No cooling was carried out on either of the samples. The H2 gas dosing operation is carried out by filling the reaction vessel to a maximum pressure of 3.3 mbar (330 Pa) and waiting for a set period or until there is a reduction of pressure on the pirani gauge indicating that gas has been consumed.
The initial work on the mixed phase alloy attempted to grow the precipitates as close to the nucleation time as possible, but following two attempts lasting 60 s and 7200 s where no hydride precipitates were observable using the FEGSEM, it was decided to grow sites which were large enough to study. This was carried out by exposing the sample to 3.3 mbar (330 Pa) of gas until consumption of the gas was evident on the pirani gauge. This process required a further dose of 70320 s and consumed 96 % of the H2 gas available. The chamber was evacuated to base pressure and the sample was analysed in the SEM.
The electropolished Pu sample was exposed to hydrogen for two days with no observable change in pressure on the pirani gauge. The sample was heated to 287 °C at 10 °C min-1 in an attempt to initiate hydride growth as had been shown in the work of Brown, Ockenden and Welch [64], but no change in pressure was evident on the pirani gauge. Following cooling to room temperature, the sample continued to be exposed to the H2 for three days with no apparent change in hydrogen pressure, upon which the chamber was evacuated to the base pressure and the sample was analysed in the SEM. This thermal treatment would have been sufficient to transform the sample to the γ-phase and back via the transformation path αβγβα. The sample remained in hydrogen for sufficient time for a significant proportion to have returned to α-Pu [159].
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Chapter 4 PHILOSOPHY FOR THE
RESEARCH
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4.1 Overall aims
This EngD was devised to investigate plutonium hydride growth phenomena which would contribute new information to the general understanding of corrosion during long term storage. The research on the EngD was guided by the requirement to find three main questions:
4.1.1 Does the overall shape of the hydride evolve isotropically?
Existing reports of plutonium hydride growth behaviour have focussed on the rate measured from consumption of hydrogen, mass gain or in situ microscopy of reaction site growth [88,89]. During this EngD a focussed ion beam (FIB) investigation reported on features of the hydride product formed [107]. The existing work is not exhaustive and a more detailed view of plutonium hydriding is desired because understanding the hydride site morphology will aid the quantification of hydride formed during long term storage.
4.1.2 Does the microstructure provide evidence of growth characteristics?
An investigation into the microstructure of the hydride material within and particularly at the interface of the reaction site with the metal is important to develop an understanding of any mechanisms that may occur when as the hydride reaction site grows. Of particular interest is relationships with the parent metal, as metallurgical features may affect the rate at which hydrogen reaches fresh plutonium metal, leading to continued precipitation of hydride.
4.1.3 Is the hydride morphology directed by strain introduced into material surrounding a dilated hydride?
Plutonium hydride has large volume dilation over that of plutonium metal in both the δ and α-phase, with the PuH2 phase being 1. 7 times larger than δ-Pu and 1.963 times larger than α-Pu. This dilation would necessarily produce a force on the plutonium metal surrounding a hydride reaction site. The nature of the force would be a radial compressive stress and a circumferential tensile stress in the metal. Different metals react in different ways to applied loads. An investigation of the metal
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surrounding the hydride reaction site would be invaluable in determining any growth mechanisms.
4.2 Experimental philosophy
Plutonium is a hazardous, expensive and limited resource, often having surprising results to certain stimuli. It is necessary therefore to develop suitable experimental methods to investigate the morphology of hydride growth using a less hazardous material. Work on plutonium typically requires a full containment system to be used to reduce the hazards to workers. Analogue materials are frequently used to study specific behaviours of hazardous materials to enable more straightforward development of experimental methods. The ideal material analogue is something that reacts to stimuli in the same manner as the material of interest, for example, research into uranium is often carried out using a less hazardous isotope of uranium (238U) [116],which owing to a small (ca 1 % atomic mass difference) has no differences from enriched uranium, and hence retains the microstructure and chemistry of the material of interest. Use of 238U isotope, which is a β emitter with a half life of 4.468 x 109 y, instead of 235U, which is an α emitter with a half life of 7.08 x 108 y, reduces the radioactive dose to operators. In addition 238U is not fissile, which eliminates the threat of criticality. Consequently use of less hazardous materials allows greater flexibility in experimental design.
The use of a lower activity and non fissile isotope is not possible for plutonium and its alloys because the main uses of plutonium utilise the fissile properties of 239Pu therefore very little plutonium has been produced as a less hazardous isotope. 239Pu has a half life of 2.411 x 104 y, non-fissile isotopes with lower radioactivity such as 242Pu (half life 3.75 x 105 y) and 244Pu (half life 8.0 x 107 y) only exist in very small quantities.
Instead, cerium was chosen as a non-radioactive analogous material to δ- stabilised plutonium because it known to form surface hydrides [92]. Cerium is often compared to plutonium because both metals have similarities which arise from the f- band electrons interactions which can be either itinerant or localised depending on the allotrope [160]. Cerium has four allotropes and will undergo a phase change (γ to α) under compression in a manner similar to δ-stabilised plutonium [119]. In the
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presence of oxygen, both Ce and Pu form binary oxides, with a sesquioxide (M2O3) in contact with the metal and a dioxide (MO2) in contact with the oxygen [79,161]. Cerium hydride exhibits comparable thermodynamic properties to plutonium hydride [87,162], suggesting similar hydriding mechanisms. Table 3 highlights the similarities and differences of cerium and plutonium.
The allotropic behaviour of plutonium and cerium are believed to occur because they experience both itinerant and localised electrons. In both cerium and plutonium the valence electrons are the f-series electrons, 4f for cerium and 5f for plutonium. The f orbitals in each system are within the d orbitals and so the atomic bonding for the lanthanides and actinides is heavily influenced by the degree of overlap of the f orbitals between atoms. In the lanthanides, the f electron orbitals tend not to overlap, meaning that the overlapping filled d orbitals take part in the bonding behaviour. In cerium the Wigner-Seitz is similar to the elements with 5d valency for the same number of valence electrons. The isostructural collapse of cerium under compressive stress is described as a system which can transition from one minimum energy state at large volume (f-like) to another at the compressed volume (d-like) [163] as observed from Figure 30 [164]. The transition between itinerant (d-like at low volume) and localised (f-like at large volume) electron behaviour in the actinide occurs in plutonium with the Wigner-Seitz volume between the α-Pu and δ-Pu phases [165].
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Figure 30 The pressure-temperature phase diagram for cerium [164]
Whilst the use of cerium has the potential to help plutonium research significantly it should be remembered that it is not a perfect substitute for δ-stabilised plutonium alloys. Some similarities exist, but so do differences. Care must be taken to consider the work in terms of the similarities and differences, careful selection of when to use cerium as an analogue is a difficult problem. One significant advantage of using analogue materials beyond technique development is that equipment available to use on everyday materials can be used which can help supplement the scarce plutonium data and improve confidence in theories.
Work with cerium is far from straightforward however; the rapid oxidation causes issues with sample surface integrity, with tarnishing occurring in a few minutes. Care was needed during sample preparation to avoid undesirable thermal
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excursions from occurring. All subsequent handling and investigation was planned to minimise exposure to the air.
Parameter Plutonium (δ-Pu) Cerium Unit cell at room temp fcc fcc Valence electrons 5f 4f Reactivity with oxygen Highly reactive Highly reactive Oxide structure PuO2 in contact with air, CeO2 in contact with air, Pu2O3 in contact with metal Ce2O3 in contact with metal Reactivity with hydrogen Yes Yes Temperature induced 6 in unalloyed Pu 4 phases 2, 3 or 6 (dependent on alloy, pressure and temperature) in δ- stabilised Pu.
Pressure induced phase δα at 0.3 GPa (fcc to γα at 0.8 GPa (fcc to fcc) change monoclinic), dependent on alloying content Hazards Radioactive:, α-particle Metal is pyrophoric emitter, half-life 24000 Hydride is pyrophoric years Spontaneous Fission neutron emitter Fissile undergoes fission on capture of a neutron. Risk of criticality Toxic Metal is pyrophoric Hydride is pyrophoric Table 3 Comparison of plutonium and cerium for the parameters important in this work
4.3 Experimentation on cerium
4.3.1 Aims
The aim of the work carried out on cerium was threefold: to develop a suitable methodology for studying plutonium hydride reaction sites; to obtain insight into the growth mechanisms and to measure properties that to help explain the microstructure of the plutonium system. Care must be exercised with each of these aims because the response of plutonium may differ from that of cerium despite the
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apparent similarities of the two systems, because differences may exist in the in the valence bands and microstructural features.
Because using cerium remains difficult and the potential for the use of cerium as a plutonium surrogate has led earlier work [92,132,133,166] to study the reaction of cerium with hydrogen, there was a good starting point to begin investigating hydriding of cerium as a plutonium surrogate. The work on cerium in the literature was not exhaustive and much remained to investigate.
The cerium work was carried out using equipment as similar as possible to that available to plutonium research. The gas rigs were of the same basic design for both systems; with minor adjustments to the plutonium rig to allow for ease of handling in a glove box. This was to give the plutonium research the best chance of success, while allowing the opportunity to find and resolve shortcomings in the experimental design without the restrictions of a glove box should they arise.
4.3.2 Initial work
The first paper included in the thesis (Chapter 5: The morphology and anisotropic growth kinetics of cerium hydride reaction sites) was work performed in collaboration with researchers at AWE during this EngD. In this work a cerium sample was exposed to a pressure of 100 mbar (1 x 104 Pa) of hydrogen for sufficient time for a number of reaction sites to form. The work demonstrated a degree of anisotropy existed in the growth rates of reaction sites leading to oblate hemispherical morphology. The main contribution to the paper was to carry out initial investigation into the microstructural details of the hydride reaction sites using a combination of scanning electron microscopy and ion beam etching. This contribution revealed that in cerium hydride reaction sites, the basic metallic microstructure is retained throughout the hydride and a halo of stressed material was evident surrounding the reaction sites. Macroscopic hardness measurements of a large hydride reaction site and the parent material were also contributed to the paper. The journal article is included in chapter 5 because it defines the starting point of the experimental work carried out for this EngD, and has strongly influenced the direction that the research has taken. This paper confirmed that hydride reaction sites form anisotropically on cerium.
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During this work, the height of deformed material surrounding a number of hydride reaction sites was measured using stereographic projection and found to be in the range of 10-20 µm protrusion from the surface. In most cases, the larger hydride sites had lost the outer layers from the process used to measure the depth of the larger craters. A few smaller hydride reaction sites had intact surfaces and their overall heights were measured to protrude ca 10 µm from the surface.
4.3.3 Microstructure of the cerium hydride reaction sites
The second publication (Chapter 6: The microstructure of cerium hydride growth sites) reports work that specifically investigated the hydride morphology of cerium hydride reaction sites, the extent of hydrogen in the surrounding material and what extent the surrounding material is affected by the dilated hydride.
Investigating the extent of hydrogen in the material surrounding a hydride reaction site requires a microscopic analysis capability that can detect hydrogen. EDS and Wavelength Dispersive Spectroscopy (WDS) which are techniques available using scanning electron microscopy cannot detect hydrogen so SIMS was selected for this investigation. SIMS was used to confirm that small spots of reaction product that were smaller than could be measured using the in situ optical microscope were also hydride reaction sites. Cerium was used as an analogue material to investigate the extent of hydrogen surrounding a hydride reaction site. This work was carried out in collaboration with other researchers at AWE and showed that hydrogen is not evident in any significant concentration except within a hydride reaction site.
Concerns existed about the extent that passivation might alter a hydride reaction site, potentially rendering any subsequent analysis void. The passivation used for these samples was to briefly expose the sample to air in a controlled situation (i.e. away from other flammable materials) by removing it from an argon filled sample bottle and then returning it to the bottle. The previously adherent hydrides were observed visually during this process and they slowly began to protrude further from the sample and the surface of the largest sites spalled from the sample. No further changes were observed after ca 20 min. The reaction was slower than anticipated from the available literature [167], presumably because the hydride was adherent to the surface prior to exposure to air, slowing down the rate at which
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the oxygen could reach fresh hydride material. The extent of oxidation of cerium hydride reaction sites was measured in cross section using EDS is shown in Figure 31. The passivation had transformed approximately 50 % of the depth to oxide, leaving around 10 µm of hydride remaining. The interface between the hydride and the metal was confirmed as unaffected by the passivation process. This allowed confidence in further investigation concerning material close to the interface.
The work from chapter 5 was designed to remove the hydride from the reaction sites to allow a measurement of the aspect ratio between penetration into the metal and surface lateral growth by measurement using extended focus optical microscopy. This was not ideal for investigation of the interface between hydride and metal, therefore improved cross-sectional work was needed to confirm the initial suggestions on that paper. The oil-based polishing method that had been used previously on cerium samples did not give a sufficiently satisfactory surface finish to allow for a detailed metallographic investigation.
For this work, a new metallographic cross sectioning technique was developed, using water based diamond suspensions. All prior information [168,169] suggested that water based grinding and polishing would result in the cerium reacting vigorously with the water, liberating flammable gases, therefore a carefully planned experiment was conducted to establish if this indeed was the case. A 12 mm round coupon, 1 mm thick was hand-ground using a 2500 grit SiC paper with water as a coolant on float glass. The float glass and a beaker of lapping oil were placed nearby on a metal draining board and tongs were available so that if the sample began heating up it could be quenched in the lapping oil. Suitable PPE was used and an appropriate fire extinguisher was available. The cerium sample was polished by a gloved finger in figure of eight motions so that any heating could be felt through the sample. No heating of the sample was detected and a lack of bubbles suggested that no liberation of gases had occurred. Following this trial, a polishing method employing standard water-based diamond suspensions followed by a colloidal silica polish was developed which allowed a satisfactory polished surface to be obtained for metallographic investigation.
The improved polish allowed the interface between the hydride and the metal to be reliably determined and post polishing deformation of the metal surrounding the
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hydride site suggested that the dilation of the hydride introduced a stress into the metal. The microstructure of the reaction sites was imaged and found to adopt that of the parent metal, with intergranular failure occurring in the oxidised portion and transgranular cracking in the remaining hydride.
Figure 31 Top: an optical image of a hydride reaction site cross section the extent of the hydride is the yellow area and would have also been the gray area. Bottom: an EDS oxygen map (green) of the same hydride reaction site overlaid onto the backscattered electron image. The sample was re-polished between the two images.
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4.3.4 Nanoindentation of the cerium/cerium hydride interface
The halo of material around the ion milled cerium hydride reaction sites from chapter 5 and the bulging material from the cross sectional work in chapter 6 (see Figure 31) suggested that there was a compressive stress surrounding the reaction sites. Nanoindentation was used as a direct technique of investigating the stressed region. The results from this work form the paper in chapter 7, Nanoindentation of the cerium/cerium hydride interface. The nanoindentation cantilever was used in tapping mode, prior to indentation, as a complimentary method to indicate that the oxidised hydride and remaining hydride regions (previously identified in chapter 6) had different material properties.
This work indicated that the hydride was harder than the plutonium metal and showed a region under stress underneath a cerium hydride, which supported the suggestion that the cracking which was observed in the surface oxide surrounding the hydride reaction site (made in chapter 6) resulted from the stresses surrounding the dilation of the hydride during growth. The indentations made penetrated as far into the stressed region as they did into unstressed metal, but the relaxation in the stressed region was greater following unloading of indentation tip.
4.4 Experimentation on plutonium
Having successfully carried out hydriding of cerium and analysis of cerium hydride there was confidence that a robust methodology was in place to carry out successful plutonium work.
4.4.1 Active commissioning of the FEGSEM
Prior to commencement of the EngD, the FEGSEM had been commissioned for use with inactive materials to demonstrate safety systems operated as intended. The active commissioning process had started prior to the EngD and was completed to the extent that images could be acquired on plutonium samples.
To commission the FEGSEM for full use on plutonium the ion milling, gas dosing and heating/cooling capabilities of the preparation chamber were
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demonstrated separately and then a combined test was undertaken to show that the capabilities perform as expected.
A homogenised δ-stabilised plutonium sample was sourced, polished then allowed to form an oxide layer (>1 week) from the trace level of oxygen available in the glove box system. The test involved sputtering, imaging, gas dosing with oxygen and heating a Pu sample. A region of the sample was ion milled at normal incidence for 130 min at 3 kV, 15 mA filament current. An etched depth of 3.5 µm was measured from the height of pillars which were formed by shielding offered by dispersed surface particulates that were accrued during handling operations. The sample was then etched for a further 240 min and the grain structure of the sample was revealed Figure 32. Attempts were made to obtain EBSD Kikuchi patterns from the etched crater, but this proved unsuccessful.
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Figure 32 Commissioning results of the ion mill on the FEGSEM used on plutonium. Top: a pillar masked by a particle on the surface allowed the etched depth to be measured after 130 min. Bottom: etched crater with grain detail revealed after 370 min.
The fresh surface was then oxidised and changes in surface morphology 3 7 observed. A small 1 cm gas dosing cylinder filled to 100 bar (1 x 10 Pa) with O2 was connected to the gas dosing panel. The gas dosing system was evacuated
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vacuum using a scroll pump before oxygen was admitted. Once in the 30 litre preparation chamber, the pressure was 3.3 mbarA.
During exposure, the temperature was cycled to 445 °C to demonstrate that gas dosing could operate in conjunction with sample heating. Numerous cracks developed in the surface which was observed most severely at the pillars, indicating that the surface had been considerably altered by the exposure. The cracks in the surface and disintegration of the pillars was likely to have occurred because the ductility of the oxide is lower than the metal and dilation of the oxide would have occurred during the reaction followed by contraction on cooling. This test concluded the active commissioning, which allowed the remainder of the planned EngD work to be carried out.
Figure 33 Cracking and disintegration of etch pillars following a high temperature oxidation at 445 °C.
4.4.2 The morphology and anisotropic growth of plutonium hydride reaction sites
Samples of fully heat treated and stabilised δ-Pu which had already been subjected to hydrogen exposure for another study [85,90] and which had grown a large number of sites (>100) were selected and analysed in the FEGSEM. The work
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consisted of an analysis of the surface, including ion milling of a number of reaction sites. The samples were subsequently removed from the FEGSEM and cross sectioned to reveal the interface between the hydride and the plutonium metal before reanalysis in the FEGSEM. The aim of this work was to optimise the analysis conditions prior to experimentation on the in-situ samples in the FEGSEM and fully characterise the morphology and microstructure of plutonium hydride sites on the commonly used δ-Pu. This work yielded significant results on delta stabilised material, and is the subject of the fourth publication of the EngD, found in chapter 8.
This section is a full metallographic investigation of the hydride reaction sites to a level of detail not previously obtained on plutonium hydride reaction sites, which confirmed some observations for hydride formed on δ-Pu made previously [107] and extended the knowledge surrounding processes which occur in the growth of hydride reaction sites.
Analysis of the surface revealed that the core of the hydride reaction sites was significantly different to that of the parent material and deformation had been caused surrounding the sites. The deformation had cracked the native oxide. Ion milling carried out on the surface of sites aimed to reveal detail about the interface between the hydride in the reaction site and the adjacent metal close to the surface. A mixed phase ahead of the hydride/metal interface was observed, probably caused by thermal redistribution from the passivation process. The use of backscattered electron imaging on electropolished cross sections allowed the extent of hydrided material to be identified, which showed the interface to have a number of discontinuities.
4.4.3 The reaction between hydrogen and electro-refined plutonium observed by in situ electron microscopy
Following the successful analysis of pre-hydrided δ-stabilised Pu, the next stage was to expose plutonium to hydrogen in the FEGSEM and subsequently analyse the reaction product without prior passivation. A sample of electro-refined plutonium was selected for experimentation because this was the purest plutonium sample that was available to the project. The sample was metallographically prepared and analysed prior to exposure to hydrogen. The sample was investigated at room temperature and the room temperature stable alpha-Pu microstructure of
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this sample had torturous grain boundaries and a grain size of between 100-200 µm. The material had numerous micro-cracks arising from the large volume collapse of the β α transition, with the alpha phase being brittle, the stresses cannot be accommodated by plastic deformation of the new phase and internal cracking occurs (see Figure 34). The micro cracks were usually associated with grain boundaries.
Figure 34 An area of the α-Pu matrix with a high density of micro-cracks.
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This sample was slow to react with hydrogen, consuming no quantity of hydrogen measurable using the pirani gauge throughout the test. After two days in hydrogen the sample was heated to 287 °C to ensure hydride formation. This process was evidently successful. Following reduction of the passivating oxide barrier and indeed complete removal in patches, the sample did not subsequently consume all the available hydrogen despite three day’s further exposure at room temperature. The quantity of hydrogen consumed was estimated from the average depth of hydride coverage of hydride (14.8 µm) with the measured void fraction (0.41) to be ca 0.3 % of the total available.
Spallation of the native oxide layer occurred in patches, to accommodate the dilation of the underlying hydride. The hydride remained adhered to the sample until cross sectioning in the glove box could be carried out and the sample re analysed. This allowed the mechanism for the formation of hydride protrusions on this sample to be identified as hydride formed within surface intersecting micro cracks.
This work demonstrated the successful operation of the gas dosing capability of the FEGSEM to carry out hydriding experimentation, which allowed the most detailed view of the morphology of as-formed plutonium hydride to be obtained. This work formed the basis of a letter to the editor suitable for publication which is found in chapter 9.
4.4.4 In situ hydriding of mixed phase α/δ-Pu
As cast plutonium alloys can often be mixed phase due to inhomogeneous distribution of the phase stabilising alloying element. In Pu this is caused by significant differences in diffusion rates between the ε-phase and the δ-phase, leading to a cored microstructure [38] although heat treatment can redistribute the alloying element sufficiently to fully stabilise the δ-phase.
A sample of as-cast 0.3 wt% Ga alloy was available for research, which had been measured using Archimedes’ immersion method to have a density 16.94 g/cm3. The theoretical density of δ-Pu is 15.92 g cm-3, however Ling et al. [170] recently measured fully delta stabilised plutonium produced in the same manner as the sample in this study to have a density of 15.77 g cm-3.The theoretical density of α-Pu is 19.86 g cm-3, but was measured recently between 19.45 and 19.55 g cm-3 by
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Farrow et al. [171]. The rule of mixtures indicates this sample contained 31 % α- phase in good agreement with an estimate of alpha content using image analysis of 36 %.
The sample exhibited a cored, non-equiaxed grain microstructure and the aim was to investigate whether a growth preference existed via α or δ domains (see Figure 35). The inhomogeneous alloy allowed a direct side-by-side comparison of the effect of phase on the propagation of the hydride interface into the surrounding material, principally to investigate whether there was a difference in hydride growth rate as a function of phase. It is possible that the presence of gallium used to stabilise the δ-phase may affect the nucleation and subsequent growth of the hydride reaction sites. As this was an as-cast alloy the gallium content was not homogeneous and no differentiation can be made between the effect of gallium and the effect of the allotrope on the hydriding behaviour. Understanding how the stresses surrounding the hydride reaction site would be accommodated by the matrix of α/δ-Pu is of interest.
Figure 35 The α/δ microstructure of the as cast 0.3 wt% alloy.
The sample was exposed to 3.3 mbar (330 Pa) of hydrogen at room temperature. This sample reacted with 96 % of the available hydrogen in 21.5 hours.
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The majority of the hydriding had occurred on the back face of the sample, forming four large sites which were examined in cross section, exhibiting similar anisotropy to that measured on the fully δ-stabilised samples in chapter 8.
A hydride reaction site had formed on the front face which allowed a detailed view of the interface between the hydride reaction site and the different phases to be investigated. Considerable spallation of hydride reaction product had occurred during the reaction, which had liberally covered the sample surface in hydride reaction product. There was no evidence of secondary nucleation sites arising from the spalled hydride product as has been suggested previously [106]. At the surface, differences were observed between the δ-phase and the α-phase domains with the δ-phase domains distorting plastically and the α-phase domains breaking up. No significant difference was found in the hydriding rate between δ and α-phase plutonium. This work resulted in work suitable for a sixth publication for the EngD, as found in chapter 10.
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Chapter 5 THE MORPHOLOGY AND
ANISOTROPIC GROWTH KINETICS OF CERIUM
HYDRIDE REACTION SITES
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5.1 Cover Page
The author’s role in this publication was limited to post-hydriding analysis using FEGSEM, broad area ion beam etching, Vickers Hardness indentation on a single hydride reaction site and cerium metal and to review and discussion on content of the manuscript.
This journal article has been included because it provided the starting point for all subsequent work.
Word count (paper only): 4383
References [77,85,92,98,133,146,154,162,166,172–186] refer to the references [1–24] from this article.
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Chapter 6 THE MICROSTRUCTURE OF
CERIUM HYDRIDE GROWTH SITES
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6.1 Cover Page
This paper shows the effects of exposure of cerium metal to pure hydrogen at 30 °C.
The author carried out all the work with guidance and discussion with J.P. Knowles. The SIMS work was carried out by a colleague at AWE (N. Montgomery). The paper was written by the author.
Word count: 5984
The references [80,92,98,114,132,133,162,167,177–179,181,187–199] in the main reference section of the thesis refers to the references [1 –26] in this section.
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Chapter 7 PROBING THE CERIUM/CERIUM
HYDRIDE INTERFACE USING
NANOINDENTATION
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7.1 Cover Page
This work was presented at the International Symposium for Metal Hydrogen Systems at Salford in May 2014.
Word count: 2719
This journal article actually investigates the effect on the material surrounding a growing hydride by the hydride. The technique of nanoindentation is used, but on such an air reactive system, this required to take place within a vacuum capable nanoindenter. One such instrument was available in a laboratory adjacent to the sample preparation equipment, minimising the oxide film that could grow on the hydride and metal between completion of polishing and investigation. This was a difficult piece of work, because growing the hydrides and then subsequently cross sectioning through a hydride is non-trivial, but the need for constant vacuum around the sample meant that all imaging needed to be completed with the nanoindenter tip.
An unplanned power outage caused damage to a scroll pump, which exploded back through the turbo and into the chamber following the completion of the second run of indentations. This prevented any further work being carried out on the sample. The final article was published online 7 Aug 2015
The references [98,150–152,154,172–174,182,195,200–202] in the main reference section of the thesis refer to the references [1–13] in this section.
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Chapter 8 THE ANISOTROPIC GROWTH
MORPHOLOGY AND MICROSTRUCTURE OF
PLUTONIUM HYDRIDE SITES
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8.1 Cover Page
This journal article describes the microstructure and morphology of plutonium hydride reaction sites grown in a specific hydriding rig. The samples were then analysed, and cross sectioned.
The hydriding on this paper was carried out by G. McGillivray, I.M. Findlay, M.J. Dawes and J.P. Knowles as part of the study of nucleation rate and incubation times [85]. Sample preparation was carried out by the plutonium research and development Technical Support team with advice from the author. All SEM work and interpretation was carried out by the author. Sample loading was carried out by the author with support from the plutonium research and development Technical Support team.
Word count (paper only): 5130
This manuscript was submitted to the Journal of Nuclear Materials on the 11 Jun 2015, and has been recommended for acceptance subject to modifications. The revised manuscript was submitted on the 4th Sep 2015, and that revised version is included here. Reference number JNM-D-15-00592R1. The manuscript was accepted and published on line and can be found at Journal of Nuclear Materials 469 (2016) 145-152
The references in the text of this chapter are numbered as for the main references of the thesis, but the original list of references for this paper can be found at the end of this chapter
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8.2 Title
The anisotropic growth morphology and microstructure of plutonium hydride reaction sites
8.3 Author names and affiliations
Martin Brierley* a b, John Philip Knowles a, Andrew Sherry b, Michael Preuss b
* Corresponding author [email protected]
a Atomic Weapons Establishment, Aldermaston, Berkshire UK
b University of Manchester, Manchester, UK
© British Crown Owned Copyright 2015/AWE
8.4 Abstract
Plutonium hydride reaction sites grown on δ-stabilised plutonium alloy have been investigated by a combination of scanning electron microscopy, ion beam milling and cross sectional polishing. The reaction sites are oblate and results indicate that this may be the consequence of failure of the native oxide over the metal surrounding the hydride sites. The interior of the hydride reaction sites has a significantly modified microstructure to that of the surrounding metal. Growth of the hydride into the plutonium metal appears to have a strong relationship with the metallurgical features, causing a discontinuous interface. Possible models for anisotropic growth and formation of a discontinuous interface are discussed.
8.5 Keywords
Plutonium hydride, microstructure, corrosion, anisotropic growth, growth mechanism
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8.6 Introduction
Plutonium hydride is pyrophoric on contact with air and the conditions of formation can modify the ignition properties. When plutonium hydride is produced slowly and below 100 °C a finely divided, face centred cubic product forms which can ignite upon contact with air at room temperature. If the hydride is produced at a temperature above 400 °C and several bar of hydrogen pressure then a hexagonal form occurs, which is not pyrophoric in room temperature air [63]. During extended storage of plutonium metal in sealed vessels hydrogen may accumulate as a consequence of packaging material radiolysis and the reaction between the metal and outgassed water. Therefore, a potential to form plutonium hydride exists during long term storage. Concerns exist about the safety of vessels containing plutonium metal stored for long periods and the likelihood of corrosion and corrosion products challenging the integrity of the containment [62]. For continued safe and long term storage of plutonium metal, a comprehensive knowledge of hydriding behaviour is desirable [97].
Following an induction time where no obvious hydride formation occurs, plutonium reacts with hydrogen to form discrete reaction sites at the surface. Early work reported unpredictable initiation behaviour with the induction time varying from a few minutes to several hours [62,64,90,105]. The occurrence of an initiation time and spot-wise growth of plutonium hydride have been attributed to the presence of a protective surface oxide which forms rapidly even in the presence of low concentrations of oxygen. Differences in sample to sample oxide characteristics could account for the variation in initiation time. When a systematic investigation into hydriding induction time initially grew repeatable oxides prior to hydrogen exposure, then reproducible initiation times and nucleation rates were subsequently observed [85]. Variations of the surface oxide thickness account for the spot-wise growth [3]. Plutonium hydride sites have been shown to preferentially nucleate around scribe marks made on the sample surfaces [105,106], which would have locally diminished the protection afforded by the oxide.
Once nucleated, the plutonium hydride reaction sites grow radially at the surface (lateral growth) and also penetrate into the bulk of the material (penetrative growth). When the entire surface has been covered by hydride reaction product
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plutonium hydride continues to form at the bulk penetrative growth rate. A number of reports on plutonium have investigated the kinetics of hydride formation [102,203]. A recent report indicates that while reaction sites tended to be hemispherical, some sites grew as oblate hemispheres [107]. The same report observed a mixed phase in front of the metal/hydride interface. Recent work on cerium hydride reaction sites suggests that anisotropic growth could be reasonably expected [195,200] therefore a detailed microstructural investigation into plutonium hydride reaction sites was conducted.
8.7 Method
8.7.1 Hydriding reaction
Two samples were taken from the set used by McGillivray et al. to determine the pressure dependency of hydride initiation time [85]. Each sample was loaded into a hydriding reaction cell situated within a nitrogen purged glovebox and evacuated to a base pressure of 1 x 10-7 mbar (1 x 10-5 Pa). Heating to 423 K and exposure to 10 mbar (1 x 103 Pa) of oxygen for 15 min achieved a reproducible oxide surface on the sample coupons. The samples were allowed to cool to 298 K naturally.
Hydrogen was generated at >99.9999 % purity and supplied to the reaction cell at pressures between 10 and 1000 mbar (1 x 103 and 1 x 105 Pa). Hydride reaction sites were observed to grow and liberation of particulate material occurred. Following the hydriding step, the reaction cell was evacuated and each sample was passivated using oxygen in the reaction cell to mitigate the risk of unplanned ignition should the sample come into contact with air. The passivation step was carried out by slowly raising the pressure using an MKS flow controller until 100 mbar (1 x 104 Pa) was reached. This limited the vigorous nature of the oxidation reaction; however the reaction liberated further particulate material whilst within the reaction cell. The pressure of oxygen was subsequently raised to 200 mbar (2 x 104 Pa) and held for 1 hour to ensure that the samples would be stable in the same partial pressure of oxygen as in air. Following passivation, each sample was removed from the reaction cell and stored in a dry nitrogen glovebox until it could be analysed.
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8.7.2 Scanning Electron Microscopy
The samples were loaded into a JEOL JSM7000F Field Emission Gun Scanning Electron Microscope (FEGSEM) using a sample transfer system that maintained glovebox atmosphere until the sample was evacuated to a pressure of 10-7 mbar (10-5 Pa). The samples were imaged using secondary electrons. Three- dimensional models were generated from images taken at angles of 0, 5 and 10º using Alicona MeX. A small section of one of the samples was subsequently ion milled to reveal subsurface detail in an adjoining chamber using a VG EX05 ion gun with argon ions at 3 kV, 15 mA for a total of 137 minutes, investigated at intervals of 30 minutes. The sample remained under a vacuum of 10-7 mbar (10-5 Pa) between milling and analysis, to prevent subsequent oxidation.
The samples were removed from the FEGSEM, mounted in epoxy resin and cross sectioned to 1 µm using diamond paste prior to electropolishing in 10 % nitric acid in 2-ethoxyethanol to reveal metallographic grain detail. The samples were coated with evaporated aluminium for conductivity, and re-loaded into the FEGSEM. The cross sections were imaged using backscattered electrons to enable differentiation between the plutonium metal and the lower density hydride which displayed a darker contrast.
8.8 Results
8.8.1 Hydriding experiment
The hydriding experiments were allowed to proceed until a sufficient number of reaction sites existed that allowed in situ measurement of nucleation rate. These results were discussed previously [85,90] and demonstrated the initiation time varied inversely to the hydrogen pressure and the nucleation rate varied linearly with the hydrogen pressure.
8.8.2 Surface Characterisation
The hydride reaction sites formed roughly circular spots on the surface of the sample (Figure 36a), except where overlapping of sites was found. In most cases, the surface over the central region of the hydride reaction site was missing, leaving
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behind a portion of the passivated hydride and a corona of deformed metal. The surface of the samples surrounding the hydride reaction site was liberally covered in a fine particulate with a particle size in the range of 0.5-4 µm.
In each hydride reaction site was a core of material, which was structurally distinct from the surrounding plutonium metal (Figure 36b). Each core was heavily cracked, with the cracks pointing radially away from the centre of the hydride reaction site. The microstructure of the material between the large cracks comprised small platelets also orientated radially with small scale cracking of the hydride product occurring along the inter-platelet boundary (Figure 36c). Metal surrounding the hydride exhibited ductile behaviour, having deformed plastically around the hydride reaction sites, whereas the material within the core was brittle and friable. The cracks in the core are likely to have developed resulting from large stresses imparted on the hydride product during the dilation which occurs when transforming from δ-Pu to PuH2. The brittle hydride product cannot deform plastically in response to the stresses and cracks form along lines of weakness such as inter platelet boundaries.
The metal surrounding the core had been deformed upwards out of the surface plane of the sample from the dilation of the hydride. Measurements taken using stereo pair imaging on sites with a diameter of between 120-300 µm gave a protruding height of 9-14 µm on the remaining corona of material around the edges of the hydride; material in the centre of the sites over the hydride core had spalled away from the sample and could not be measured. The curvature of the surface of this deformed corona has caused cracking to the surface immediately surrounding the hydride reaction site, which implies that the native oxide would also have fractured (Figure 37a). Closer inspection revealed that both radial and concentric cracking had occurred at the hydride edges (Figure 37b). Cracks in the oxide were observed to extend a distance of 1-2 µm away from the deformed material over the surface plane of the plutonium metal. The presence of such cracks have been observed to occur following hydrogen corrosion of cerium [200].
Further investigation of the subsurface structure of hydride reaction sites was carried out by ion milling prior to cross-sectional analysis. Although the debris on the surface from the hydriding and passivation steps caused the formation of pillars
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whilst ion milling, there was sufficient clarity to investigate features within and immediately surrounding the hydride sites. Figure 38a shows the interior of such an ion milled hydride reaction site, the lower density hydride can be identified by the darker region when compared to the expected higher brightness of the surrounding higher density metal matrix. The hydride had a network of cracks in the centre, which could have occurred during growth of the hydride or during the passivation process. Figure 38b shows a halo of material with a mottled appearance immediately surrounding the plutonium hydride reaction site extending into the plutonium metal, suggesting a mixed phase material, most probably δ-Pu + PuH2+x.
8.8.3 Cross-sectional analysis
The material in the hydride reaction sites near to the surface was of platelets aligned towards the centre of the site (Figure 39), consistent with the top down view from Figure 36. This region is the one most likely to have been affected by the passivation process. Closer to the hydride/metal interface, the microstructure of the hydride again consisted of platelets, but the alignment of these platelets appeared to be influenced by the grain from which the hydride was formed.
The shape of the penetration of the hydride into the parent material, as seen in Figure 39, can be described as an oblate hemisphere. The surface radius to depth aspect ratio of the hydride sites (measured from a straight line from the original surface on one side of the hydride to the other) was similar on both samples, with an average of 2.2 and a range from 1.4 to 3.8, based on 12 hydride reaction sites. This aspect ratio measurement confirms that the surface hydride growth on plutonium is anisotropic.
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The samples in this study were examined using backscattered electrons and exhibited a discontinuous interface between the plutonium hydride and the plutonium metal, meaning that the interface cannot be fitted using a simple mathematical function such as a hemisphere or ellipse. In addition, the use of electropolishing has revealed the internal microstructure of the hydride material. The overall discontinuities of the hydride/metal interface occurred at a frequency consistent with the microstructure of the plutonium; in some cases there appeared to be almost no penetration of the hydride into particular grains in contact with the hydride interface, but neighbouring grains at the interface had been almost fully consumed.
Two features appeared to be commonly found protruding into the parent metal from the hydride. The first type of feature is associated with grain boundaries; the second type of feature appears as projections of hydride into the surrounding metal. Enhanced precipitation along grain boundaries was observed from the interface of the hydride reaction sites along grain boundaries into the surrounding plutonium (Figure 40a). Certain grain boundaries appeared to be susceptible to enhanced precipitation while others did not. In many cases, the enhanced precipitation along grain boundaries was frequently observed alongside other metallurgical grain boundary precipitates (e.g. Pu6Fe).
A preferred growth habit was observed from the grain boundaries into the grain interiors. In some cases, the hydride did not penetrate a particular grain on one side but had penetrated the grain from a different direction, suggesting a preferred orientation for hydride growth exists in δ-plutonium. Once growth into the interior of a grain was initiated, high aspect ratio plates or needles grew into the grain interior (Figure 40a, b). These high aspect ratio features were tightly aligned and formed along defined directions within the same crystallite grain. In cases where growth did not appear to occur via high aspect ratio features, differences in grain penetration were found. Figure 40c shows a grain boundary where the hydride has penetrated a distance three to four times as far into one crystallite as into the neighbour on the other side of the grain boundary.
The second type of feature, protruding into the metal, was triangular projections from the hydride into the parent material (Figure 41). The triangular projections consisted of an inner region of platelets aligned with the long axis of the
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projection, surrounded by a tapered sheath of hydride material. The material in the outer sheath was not obviously aligned to the projection or the host crystallite, suggesting a two-stage formation of the projections. The sheaths had formed thickest nearest the junction of the projections with the body of the hydride. The contrast level of the electron backscatter signal indicates that the density of the material within the projections is the same as that in the rest of the hydride at the time of imaging. The projections were measured to be between 8 and 20 µm in length and found within crystallite grains. Unlike in the case of the highly aligned crystals discussed previously, the projections were observed to be unaffected by grain boundaries or triple point junctions.
8.9 Discussion
8.9.1 Accommodation of hydride dilation
A small amount of elastic and plastic strain must exist around the hydride reaction site arising from the harder hydride exerting an outward force upon dilation into the softer surrounding metal. PuH2 has 1.575 times and the face centred cubic form of PuH3 has 1. 27 times the atomic volume of δ-phase metal per plutonium atom, similar to the value of 1.579 for plutonium dioxide [40]. The hydride is ca 2.5 times stiffer [33,204] and ca 7.2 times harder [31,32,94] than the surrounding δ-Pu metal which would cause the metal to accommodate most of the strain associated with the dilation. . The remaining dilation must be accommodated by protrusion of hydride from the surface of the sample since dilation of the hydride is constrained by the material underneath and surrounding the hydride. The degree of protrusion from the surface should therefore be similar to the relative volume per plutonium atom of the hydride compared with that of the parent phase.
Dimensional measurements were taken from hydride site cross sections. The overall depth of the hydride (outer surface to the hydride/metal interface) was compared to the penetrative depth (location of the original surface to the hydride/metal interface) for each site, which yielded an average ratio of 1.43. This is consistent with the dilation considering the strain in the surrounding plutonium, and some loss of surface material from the hydride reaction sites had occurred in most
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cases. Flexure of the native surface oxide to accommodate this protrusion has caused cracking, which extends a short distance over the plutonium metal surrounding the hydride. This negates the protective effect of the oxide over the fresh metal surrounding the hydride.
8.9.2 Growth anisotropy
The ideal shape for a hydride reaction site, based on elastic stresses within the bulk, would be hemispherical [182]. The hydride reaction sites here are demonstrably non hemispherical, so other factors must affect the shape. Cracking of the surface oxide above and immediately around the hydride shown in Figure 37 offers a solution to the anisotropy because it allows hydrogen facile access to plutonium metal immediately surrounding the hydride reaction site as well as through the hydride product [195]. The many inter-platelet boundaries and cracks in the hydride product shown in Figure 36 and Figure 41 may provide a rapid diffusion path for hydrogen through the hydride to the metal interface.
The mechanisms that could cause anisotropic growth of a hydride reaction site at the surface are described in terms of the following four periods of reaction: incubation, early growth, oxide failure and continued growth (Figure 42). The incubation of hydride formation on metals (1) has been discussed in detail elsewhere [3,85,146,189]. Essentially, hydrogen is adsorbed onto the surface of the oxide over- layer before diffusing through the oxide to concentrate in the underlying metal. Eventually super-saturation occurs at the oxide metal interface and a hydride precipitates.
The early growth period (2) begins once the precipitation of hydride has occurred. During this period, growth of the hydride precipitate at the oxide metal interface exerts an outward strain related to the 58 % dilation associated with hydride formation. The dilation is constrained laterally by the surrounding plutonium metal; therefore, a compressive stress is induced in the plutonium. The surface oxide layer provides the only constraint on the dilation of the hydride vertically.
The oxide failure period (3) occurs once dilation of the hydride reaction site distorts the surrounding material sufficiently to fracture the oxide layer above the reaction site.
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Following the initial fracture of the oxide, the continued growth period (4) begins. As the reaction site grows, continued force is exerted into the surrounding metal which distorts plastically. The surface oxide above the distorted metal in the immediate vicinity of the site cracks and no longer provides a barrier between the metal and hydrogen. As hydrogen will readily diffuse through the hydride product and the cracks in the oxide surrounding the sites hydrogen super saturation occurs in the adjacent metal (highlighted in red). The supersaturated regions form further hydride and the overall growth of the hydride sites become a convolution of growth at the hydride/metal interface plus the contribution from growth under the cracked oxide surrounding the reaction site, resulting in anisotropic growth. Hydride product has been observed to spall during the continued growth period [105].
8.9.3 Hydride/metal interface
The discontinuous nature of the hydride/metal interface takes the form of four features: 1) mixed phase region, 2) grain boundaries, 3) crystallite orientation and 4) triangular projections.
8.9.4 Mixed phase hydride
In this study, a mixed phase region ahead of the reaction front was observed following ion etching (Figure 38) similar to that observed by Dinh et al. [107]. It is unlikely that domains of plutonium hydride tens of nanometres in size would exist separated from their neighbours by tens of nanometres over such a large extent (several micrometers) as a normal growth habit when plutonium reacts readily with hydrogen. This would require the solubility limit for hydrogen in plutonium to be exceeded at discrete locations while the surrounding material remained below the terminal solubility limit. It is more probable that the mixed phase region was formed by the thermally driven re-distribution of hydrogen from the reaction site by the inherent heating of the passivation step. Upon oxidation of a hydride site, the local environment is heated, increasing the solubility of hydrogen in the nearby plutonium [87], allowing hydrogen to migrate in solution from the reaction site beyond the reaction front. As the oxidation reaction ends, the heat source is reduced and the local environment cools which results in the precipitation of a mixed phase region ahead of the reaction front. Further experiments are currently underway investigating
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the reaction front interface without utilising a passivation step by exercising an in situ hydrogen dosing, ion beam etching and electron imaging capability.
8.9.5 Grain boundaries
A significant number of grain boundaries were observed to have hydrides extending further into the metal than the surrounding grains (Figure 40). Grain boundaries have a more open structure than the crystal lattice within the grains and offer a lower energy diffusion path than the grains. However, not every grain boundary appeared to have hydride precipitated along them; suggesting either a preference for diffusion along grain boundaries with particular degrees of mismatch across the grains where there is more open volume, or the presence of a grain boundary precipitates that enhances nucleation at some grain boundaries.
Nucleation of hydride at a Pu6Fe precipitate found at a grain boundary has been reported previously [107] but, in the plutonium used here, the grain boundary precipitates were too small for reliable elemental identification.
8.9.6 Crystallite orientation
The orientation of the metal grain to the advancing hydride front appears to influence the growth of the hydride (Figure 39). The logical explanation is the grain orientation relative to the reaction front is critical to transporting the hydrogen into the interior of the grain. Some grains appear to resist hydride growth while others form high aspect ratio hydride features along aligned directions in the crystallite. It is suggested that the dilation of the newly-formed hydride, particularly along grain boundaries, induces local stresses and initiates slip lines within the adjacent crystallite grain along defined planes and hydride precipitation occurs along these slip lines. Preferred hydride precipitation has been observed along slip lines in uranium [205] and in titanium under stress [206].
8.9.7 Triangular Projections
Numerous ca 20 µm length triangular projections of similar morphology were observed extending beyond the reaction front (Figure 41). The projections had a two part structure with an inner core of platelets aligned to the main axis of the projection, surrounded by a sheath of non-aligned hydrided material. The tapering
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thickness of the outer sheath of non-aligned hydride towards the tip suggests that the inner region of the projection allows quicker diffusion than the surrounding metal.
Currently it is not clear how these projections formed, although it is possible they are the result of a pressure-induced δ to α’ phase transformation occurring ahead of the growing hydride interface. Dilation of the hydride imparts a compressive radial stress into the surrounding material; the extent of this stressed zone has been demonstrated underneath a cerium hydride to be between 2-3 µm [207]. If this compressive stress was to exceed a certain level, (ca. 0.3 GPa) an amount of plutonium metal in the region of highest stress could transform from the δ-phase to the denser α’-phase [208]. Calculations suggest that the radial compressive stresses surrounding a growing hydride reaction site could exceed 2.6 GPa [202].
8.9.8 Discontinuous interface mechanism
A discontinuous interface between the hydride and metal was observed on all the reaction sites examined in cross section across the two samples. Since the discontinuities except for grain boundary diffusion appear to be driven by stresses imparted by hydride dilation and constitute ca 10 % of the depth of hydride penetration into the metal at the time the hydride reaction sites were observed, a reasonable inference is that at an early stage the interface has no significant discontinuous features.
A possible mechanism for the occurrence of such an interface is proposed, consisting of four stages (Figure 43): 1) An oblate hydride reaction site of a nominal size surrounded by plutonium crystallite grains (not to scale) that has not yet formed a discontinuous interface is shown in cross section; 2) Grain boundaries offer a more open structure and diffusion preferentially occurs along them, particularly where grain boundary precipitates are present. Hydride precipitates along the grain boundaries where the terminal solubility is exceeded as observed in Figure 40 and the associated dilation of hydride imparts additional stresses locally into the plutonium metal; 3) Plane slip occurs along specific directions in the adjacent plutonium grains to accommodate the stress, causing dislocations along the slip lines which allows additional paths for hydrogen to flow into the bulk of the grain, where further aligned growth of hydride occurs; 4) Continued growth of hydride at the interface exerts an increasing outward radial compressive stress into the surrounding
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metal. Once the stress field surrounding the hydride exceeds a level of compressive stress (ca 0.3 GPa) then projections (shown in red) are induced in the plutonium metal surrounding the hydride possibly by martensitic transformation from δ-Pu to α’- Pu. The projections have a number of inter-platelet boundaries, which act as conduits for further diffusion of hydrogen to the surrounding metal. A tapered sheath of hydride forms surrounding the projections resulting in the triangular features seen in Figure 41. Grain boundary growth and intra-grain aligned growth via slip lines are expected to occur throughout the growth process, periodically raising the stress around the hydride sufficiently to cause triangular projections to occur.
8.10 Conclusions
Samples of plutonium were exposed to hydrogen and formed a number of hydride reaction sites. The surfaces and cross sections of the sites were examined using FEGSEM and ion milling, which showed that the hydride sites were anisotropic with faster lateral growth compared to penetrative growth. Plutonium hydride grows as discrete sites on a plutonium metal surface, which form as circular spots on the surface with a core of material exhibiting different microstructure to the metal. The overall shape of the hydride penetrating into the metal can be described as oblate hemispheres with an average ratio of lateral radius to depth of 2.2.
A mechanism which could cause anisotropic growth is proposed which is consistent with the observations that dilation of the hydride deforms the surrounding material and causes cracking of the surface oxide immediately surrounding the hydride reaction site. Subsequent rapid transport of hydrogen to the underlying metal through cracks in the surface oxide surrounding the reaction site could then accelerate the lateral growth.
A discontinuous interface was observed to exist between hydride reaction sites and the plutonium metal. The crystallographic orientation of grains may be an important factor that influences the rate of advance of the hydride growth front. On the cross-sections of some sites, projections were found to extend from the hydride into fresh metal. A model is proposed that describes the formation of a discontinuous interface arising from the stresses which are imparted by dilation of the hydride.
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8.11 Acknowledgements
AWE plc provided the funding to carry out this research in support of an EngD at the University of Manchester. Ian Findlay and Gordon McGillivray are acknowledged for their experimental expertise and assistance generating the samples. The plutonium R&D operations team are thanked for their technical assistance. Mike Matthews and Nigel Park are thanked for discussions on hydride reaction site geometry.
8.12 References
[1] J.M. Haschke, A.E. Hodges, R.L. Lucas, J. Less-Common Met. 133 (1987) 155. [2] J.L. Stakebake, J. Nucl. Mater. 38 (1971) 241. [3] J.M. Haschke, J.L. Stakebake, in:, L.R. Morss, N.M. Edelstein, J. Fuger (Eds.), Chem. Actin. Transuranium Elem., Springer, Dodrecht, 2006, pp. 3199–3272. [4] F. Brown, H.M. Ockenden, G.A. Welch, J. Chem. Soc. (1955) 3932. [5] G.W. McGillivray, J.P. Knowles, I.M. Findlay, M.J. Dawes, J. Nucl. Mater. 385 (2009) 212. [6] L.N. Dinh, J.M. Haschke, C.K. Saw, P.G. Allen, W. McLean, J. Nucl. Mater. 408 (2011) 171. [7] G.W. McGillivray, J.P. Knowles, I.M. Findlay, M.J. Dawes, J. Nucl. Mater. 412 (2011) 35. [8] J. Glascott, Philos. Mag. 94 (2013) 221. [9] C.K. Saw, J.M. Haschke, P.G. Allen, W. Mclean, L.N. Dinh, J. Nucl. Mater. 429 (2012) 128. [10] J.M. Haschke, T.H. Allen, J. Alloys Compd. 320 (2001) 58. [11] J.L. Stakebake, J. Alloys Compd. 187 (1992) 271. [12] L.N. Dinh, S.K. McCall, C.K. Saw, J.M. Haschke, P.G. Allen, W. McLean, J. Nucl. Mater. 451 (2014) 143. [13] J.P. Knowles, G. Rule, M. Brierley, Corros. Sci. 77 (2013) 31. [14] M. Brierley, J.P. Knowles, N. Montgomery, M. Preuss, J. Vac. Sci. Technol. A 32 (2014) 031402. [15] D.L. Clark, S.S. Hecker, G.D. Jarvinen, M.P. Neu, in:, L.R. Morss, N.M. Edelstein, J. Fuger, J.J. Katz (Eds.), Chem. Actin. Trans. Elem., Third, Springer, Dordrecht, The Netherlands, 2008, pp. 813–1264. [16] Y. Guo, J.-J. Ai, T. Gao, B.-Y. Ao, Chinese Phys. B 22 (2013) 057103. [17] H.M. Ledbetter, R.L. Moment, Acta Metall. 24 (1976) 891. [18] D.W. Wheeler, S.M. Ennaceur, M.B. Matthews, J. Nucl. Mater. 456 (2015) 68. [19] D.W. Wheeler, S.M. Ennaceur, M.B. Matthews, P. Roussel, J. Nucl. Mater. 452 (2014) 509. [20] E.M. Cramer, J.B. Bergin, Plutonium Microstructures Part 2 Binary and Ternary Alloys, Lawrence Livermore National Laboratory, UCRL-53174-2, 1983. [21] Y. Greenbaum, D. Barlam, M.H. Mintz, R.Z. Shneck, J. Alloys Compd. 452 (2008) 325. [22] J. Bloch, M.H. Mintz, J. Alloys Compd. 253-254 (1997) 529. [23] H. Uchida, Int. J. Hydrogen Energy 24 (1999) 861. [24] S. Richmond, J.S. Bridgewater, J.W. Ward, T.H. Allen, IOP Conf. Ser. Mater. Sci. Eng. 9 (2010) 012036.
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[25] J. Bloch, F. Simca, M. Kroup, A. Stern, D. Shmariahu, M.H. Mintz, J. Less-Common Met. 103 (1984) 163. [26] J.D. Boyd, in:, R.I. Jaffee, N.E. Promisel (Eds.), Sci. , Technol. Appl. Titan., Pergamon Press, Edinburgh, 1970, pp. 545–555. [27] M. Brierley, J. Knowles, J. Alloys Compd. 645 (2015) S148. [28] D.R. Harbur, J. Alloys Compd. 444-445 (2007) 249. [29] Y. Greenbaum, D. Barlam, M.H. Mintz, R.Z. Shneck, J. Alloys Compd. 509 (2011) 4025.
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Figure 36 a) A top-down view of a plutonium hydride reaction site. b) The structure of the central core of each site consisted of multiple platelets aligned towards the centre (indicated by cross). c) Radial fractures were found within the core, in all cases this followed the radially oriented platelet boundaries.
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Figure 37 a) The surface oxide layer immediately surrounding the hydride reaction site was observed to have been cracked by the dilation of the hydride site. b) The cracks are oriented in the radial and circumferential directions and extend beyond the edge of the hydride site.
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Figure 38 a) Subsurface structure of a plutonium hydride reaction site revealed by ion beam milling. b) A magnified section (red box in a) of the hydride-metal interface.
Figure 39 Cross section of a typical plutonium hydride reaction site imaged using backscattered electrons. The sites were found to inhabit an oblate shape in
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cross section. Each site had a discontinuous interface with the parent plutonium metal.
Figure 40 a) Growth of hydride into the plutonium bulk following grain boundaries (identified in black), b) infill of the grains occurs by successive growth of aligned linear hydride features, c) preferred growth into one grain compared to its neighbour was observed.
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Figure 41 A number of triangular projections were found along the interface between plutonium and plutonium hydride have a high aspect ratio and consist of an inner region of aligned platelets surrounded by a tapered sheath of non aligned material. Grain boundaries (GB) and a triple point junction (TP) are identified in black.
Figure 42 A suggested schematic of the hydriding process: (1) Incubation, (2) Early growth, (3) Oxide failure, (4) Continued growth.
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Figure 43 A schematic of the growth processes: (1) an oblate hydride reaction site surrounded by plutonium crystallites, (2) grain boundary precipitation, (3) plane slip and aligned hydride growth along slip lines and (4) stress induced triangular projections.
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Chapter 9 THE REACTION BETWEEN
HYDROGEN AND ELECTRO-REFINED
PLUTONIUM OBSERVED BY IN SITU ELECTRON
MICROSCOPY
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9.1 Cover Page
This chapter is included in the thesis because this is the initial work carried out in support of in situ hydriding of plutonium in practical terms virtually all of the plutonium metal that exists in the storage across the world will be in the form of α-Pu or δ-Pu. Other phases cannot be adequately stabilised to room temperature.
This work is the first in situ hydriding work carried out on plutonium, and to ensure chemical purity was performed on unalloyed plutonium. The in situ work allowed the sample to be analysed post exposure to hydrogen without subsequent exposure to oxygen, retaining the hydride in condition in which it had formed. This experiment was allocated instrument time for two weeks against competing work. It was anticipated that the reaction would have occurred within 24 hours of exposure to hydrogen. The sample did not react as expected and after two days with no measureable consumption of gas, a thermal cycle was carried out to initiate the hydriding reaction. Despite this and three further days in hydrogen, no observable change on the pressure gauge was evident. Upon analysis it was evident that a small amount of reaction had occurred.
The author carried out all the SEM and hydriding work on this sample, with sample preparation carried out by the technical support team under the direction of the author.
Word count: 1746
The references in the text of this chapter are numbered as for the main references of the thesis, but the original list of references for this paper can be found at the end of this chapter
This publication was submitted to the Journal of Nuclear Materials as a Short Communication on 10/09/2015. Reference number JNM-D-15-00892. The manuscript was accepted and published on line and can be found at Journal of Nuclear Materials 469 (2016) 39-42
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A letter to the editor of the Journal of Nuclear Materials
9.2 Title
The reaction between hydrogen and electro-refined plutonium observed by in situ electron microscopy
9.3 Author names and affiliations
M. Brierley a b *, J. P. Knowles a, M. Preuss b
* Corresponding author
a Atomic Weapons Establishment, Aldermaston, Berkshire UK
b University of Manchester, Manchester, UK
© British Crown Owned Copyright 2015/AWE
9.4 Abstract
Electro-refined plutonium was reacted with hydrogen within the preparation chamber of a field emission gun scanning electron microscope and in situ images obtained. The plutonium hydride reaction product was observed to have precipitated at the oxide metal interface as angular particulates (ca 2 µm in length) and was also present within micro cracks intersecting the surface.
9.5 Introduction
Safe consignment of plutonium metal requires a thorough understanding of its corrosion in a variety of environments. During extended storage of plutonium metal in sealed vessels hydrogen may accumulate as a result of radiolysis of organic materials such as plastic packaging and by reaction of the metal with outgassed water vapour. Therefore, a potential to form plutonium hydride exists. Concerns regarding hydride formation arise from issues surrounding its pyrophoric nature which has the potential to cause dispersal of plutonium outside of suitable containment [97] and form respirable particulate [209,210]. For the safe and long
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term storage of plutonium, a comprehensive knowledge of plutonium hydriding behaviour is desirable.
Plutonium hydride is extremely reactive and can oxidise even within nominally inert gas atmospheres, therefore limited morphological assessments on the reaction product formed between plutonium and hydrogen have been reported. The present work has assessed the morphology of the hydride formed on pure plutonium using in situ FEGSEM hydriding studies and cross sectional analysis.
9.6 Experimental
All handling, cutting, polishing and loading steps took place within a dry nitrogen glovebox system prior to transfer into a Field Emission Gun Scanning Electron Microscope (FEGSEM). At no time was the sample exposed to the laboratory atmosphere to prevent gross oxidation. A sample transfer system was used for transferring samples between glove box and FEGSEM, which maintains the glove box atmosphere during transit.
A sample of electro refined plutonium was cut to a 2.5 × 10 × 10 mm square from an as-cast billet. The microstructure consisted of a mixture of grain sizes between 100 and 200 µm (Figure 44). The grain boundaries were torturous and a number of sub-grain regions were evident within larger grains, the micro cracks present at the surface are associated with grain boundaries.
Figure 44 Microstructure of the electro-refined plutonium
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Prior to loading into the FEGSEM, the sample was polished to 1 µm then electropolished in 10 % nitric acid in 2-ethoxyethanol to reveal the metallographic grain detail. The analysis and in-SEM hydriding took place within a specially designed JEOL JSM7000F FEGSEM that comprises three vacuum chambers for loading, analysis and preparation (Figure 28). The preparation chamber contains a VG EX05 ion gun, a heating stage and gas doing capabilities allowing samples to be ion milled, heated, cooled and exposed to different gases then analysed without exposing the sample to an external atmosphere.
Figure 45 A three dimensional representation of the field emission gun scanning electron microscope arrangement.
Hydriding took place in the preparation chamber at a pressure of 3.3 mbar
(330 Pa) and the sample was exposed to the H2 for five days. A thermal treatment at 287 °C was used to initiate reaction similar to that previously utilised elsewhere [64]. Upon completion of the exposure, the chamber was evacuated to base pressure (1.1 x 10-4 Pa) and the sample was transferred to the analysis chamber.
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Following hydrogen dosing and in situ analysis, the sample was removed from the FEGSEM and subsequently cross-sectioned, polished to 1 µm, electropolished, coated with aluminium then returned to the FEGSEM for further analysis.
9.7 Results and discussion
9.7.1 In situ surface examination
The in situ analysis of the reacted surface indicated that a degree of reaction had taken place. The native oxide layer had been disturbed (Figure 46a), cracks had formed across the entire surface and regions of the surface had spalled to reveal underlying hydride. It is likely that the surface (i.e. oxide over-layer) spalled owing to the formation of the low density hydride at the oxide-metal interface [3]. High resolution images of the hydride revealed the material to consist of angular particles of irregular shape around 2 µm in length (Figure 46b). To the best of our knowledge, this represents the most resolved image of plutonium hydride reported to date.
The hydride reaction product formed at a free surface formed as a particulate substance, presumably because the interfacial stresses arising from proportional dilation of the hydride are large. The molar volume of cubic PuH2 is 1.963 times that of α-Pu [40]. Since the hydride formed as a continuous layer on the surface, dilation of the hydride during formation would have induced a tensile stress into the α-Pu parallel to the interface plane and an equivalent compressive stress would occur in the hydride at the interface. The tensile stress (p) imparted into the plutonium is estimated to be 106.8 GPa using p = -(V/V0)k, where: V/V0 is the dilation of 1.963 given previously and k is the bulk modulus, taken to be 4.4 GPa for α-Pu [28]. This exceeds the ultimate tensile stress of α-Pu (ca 450 MPa) [211]. However, cracking of the plutonium metal is not likely to occur despite the hydride being stiffer than α-Pu because the hydride forms as a thin coating on a bulk material and the stress is applied over a much smaller area than that of the plutonium metal. The compressive stress in the hydride layer, which is ionic and brittle [212], is accommodated by cracking leading to the formation of the angular particles observed.
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Figure 46 Secondary electron images of the reacted sample showing (a) surface cracks, spalled regions and extensive hydride coverage, (b) fine detail of the hydride product in a spalled region.
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The precipitation and growth of hydride on Ga stabilised δ-Pu characteristically occurs at discrete locations [85,90,107]. On this sample, in addition to coverage across the entire surface, a number of surface protrusions were observed across the surface (Figure 47) which protruded up to 100 µm from the surface and had radial and circumferential cracks extending ca 100 µm into the surrounding material. In general, the surface surrounding the protrusions contained ‘oxide wrinkles’ encircling at distances of ca 100-200 µm.
Figure 47 Secondary electron image of a typical hydride protrusion associated with a micro crack in the plutonium.
9.7.2 Post experimental cross section
Based upon the contrast observed in backscatter electron imaging, the sample in cross section appeared to consist of three materials: an upper layer of oxide (2.2 ± 0.3 µm thick), an inter-layer of hydride (15 ± 5 µm thick) and the bulk plutonium. The hydride layer was present all around the cross section although the upper oxide layer was missing in some places, which is consistent with the surface view from Figure 46.
The layer of hydride covering the entire surface without the reaction running to full consumption of the available hydrogen is an unexpected result. The full surface coverage of hydride is consistent with a change from PuO2 to Pu2O3 in the oxide during the thermal treatment. Pu2O3 is much more permeable to hydrogen than PuO2
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and if this change had occurred in a short timescale (i.e. during heating), then full surface coverage would ensue.
Subsequent cooling of the oxide should not allow it to revert to PuO2 and become a diffusion barrier again. Regardless of the stoichiometry of the oxide; significant cracking and loss of oxide had occurred because of the dilation of the hydride underneath. Once the hydride layer had formed, the oxide could no longer act as a diffusion barrier.
This leaves three possible mechanisms which prevented the continued growth of hydride.
1. The thermal treatment raised the solubility of hydrogen in plutonium for a brief period. Subsequent cooling reduced the solubility and hydride precipitated at the surface. 2. Adherent hydride forms a diffusion barrier 3. The electro refined metal is resistant to hydride formation at the reaction temperatures used
For mechanism 1 to be a reasonable explanation, the surface region would have to absorb hydrogen up to the terminal solubility and then transform to PuH2 once the temperature decreased so that the terminal solubility was exceeded. The terminal solubility of unalloyed Pu is ca 1 at% H in Pu (Figure 48) [87], and reducing the temperature further to the reaction conditions would very likely decrease the terminal solubility further. Stoichiometric PuH2 has 200 at% H; for the plutonium to absorb sufficient hydrogen up to the terminal solubility and then subsequently precipitate a ca 2 µm layer would require that the plutonium was saturated to a depth of ca 400 µm. In this case, it is unlikely that all of the hydrogen would subsequently migrate back to the surface region upon cooling, especially as the diffusion rate would decrease. A more likely microstructure for this scenario would be a mottled effect as observed in Figure 38, where precipitation would occur around local -1 nucleation sites. The reaction Pu = H2 ⇌ PuH2 is exothermic (ΔG = -156 kJ mol ) making the reaction favourable and Pu should form PuH2 on contact with H2. Formation of a layer of hydride on contact with hydrogen is a more likely explanation than an extended zone of plutonium under saturated with hydrogen.
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Figure 48 low concentration PCT curves from unalloyed Pu. The lower four traces are from the δ-Pu temperature range [87].
Adherent hydride was proposed as a diffusion barrier by Bloch and Mintz (mechanism 2) [146], however analysis of the hydride/metal interface on this sample showed that cracks existed in the hydride layer, probably caused by the stresses imparted by the large volume dilation on forming hydride. The cracks were observed to extend close to the interface (Figure 49). Any adherent hydride layer acting as a diffusion barrier must be very thin; however, thin layers can adopt the lattice spacing of the substrate to a greater degree than thick layers. Additional stresses imparted by cooling and changes of phase of the underlying plutonium metal would be very likely to disrupt the hydride film, however. As an explanation for the lack of continuation of hydriding this seems to have a few shortcomings.
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Pu Hydride
Pu Metal
Figure 49 Cracking was visible throughout the hydride product with some cracks in close proximity to the hydride/metal interface.
There remains a possibility that at the lower temperatures of the hydrogen exposure used in this experiment, unalloyed Pu may be more resistant to hydriding than observed during PCT experiments [84,86,87] conducted within the range of temperatures of the δ-Pu phase (mechanism 3).The lower symmetry plutonium phases α, β and γ have a higher density and consequently less open volume.
Reducing the temperature in Figure 48 makes the onset of PuH2 more likely, but at the pressure of the experiment, the PuH2 phase is bound to form. The free energy at the point of saturation taken from Figure 48 is given by ΔG = ΔH-TΔS. A van’t Hoff plot of the above data for the δ-phase suggests that ΔH is -8.31 kJ mol-1 and ΔS is 10.2 J mol-1. Plotting the variation of ΔG for the point of saturation temperature range 400 to 475 °C and extrapolating to the reaction temperature gives the free energy at the point of saturation as ΔG = -11.2 kJ mol-1 at 287 °C and ΔG = -8.5 kJ mol-1 at 23 °C. While the magnitude of the free energy decreases towards lower temperatures it remains negative. These values are within the δ-Pu stability region in unalloyed Pu. There is likely to be a slightly different value in γ-Pu at 287 °C and α- Pu at 23 °C.
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The formation of PuH2 would require energy because the volume increase would impart stresses across the interface. The α-phase in particular is very hard and would require a large amount of energy to distort or break the lattice. While this energy may be much smaller than the exothermic chemical energy produced during the reaction, at lower temperatures where the free energy is far lower, this may form a sufficient activation energy barrier to negate formation of hydride. A combination of adherent hydride and resistance of α-Pu could be possible.
The presence of micro cracks has been reported previously [213] and are inherent in pure Pu because, during cooling from casting, the β-phase to α-phase transition involves a large anisotropic ca 12 % volume contraction between two phases with low ductility. In cross section, the surface protrusions discussed earlier were clearly demonstrated to be features associated with hydride formation within micro cracks intersecting the surface (Figure 50).
Figure 50 Backscatter electron image of a cross section through a surface protrusion reveals hydride had formed within the micro cracks intersecting the surface.
9.8 Conclusions
The methodology required to perform the in situ formation and microscopic analysis of plutonium hydride is described and the initial results of a reaction with electro refined plutonium is reported. Plutonium hydride was observed to precipitate across the entire surface of the sample and was comprised of angular particulates ca 2 µm in length. A number of large protrusions were also observed by in situ analysis,
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which subsequent cross sectional analysis revealed to be surface intersecting micro cracks filled with plutonium hydride.
9.9 Acknowledgements
This work was funded by AWE plc and carried out in conjunction with the University of Manchester. The authors acknowledge the support staff at AWE plc for sample polishing, material transfers and gas dosing support, in particular Lester Derek is acknowledged for the metallographic cross sections reported.
9.10 References (see main references section)
[1] J.M. Haschke, J.L. Stakebake, Handling , Storage , and Disposition of Plutonium and Uranium, in: L.R. Morss, N.M. Edelstein, J. Fuger (Eds.), Chem. Actin. Transuranium Elem., Springer, Dodrecht, 2006: pp. 3199–3272. [2] C.L. Comar, L.A. Sagan, Health Effects of Energy Production and Conversion, Annu. Rev. Energy. 1 (1976) 581–600. doi:10.1146/annurev.eg.01.110176.003053. [3] R.Z. Omar, J.A. Barber, P.G. Smith, Cancer mortality and morbidity among workers at the Sellafield plant of British Nuclear Fuels., Br. J. Cancer. 70 (1994) 1232–1243. doi:10.1038/sj.bjc.6690207. [4] F. Brown, H.M. Ockenden, G.A. Welch, The Prepration and Properties of some Plutonium Compounds, Part I. Plutonium Hydride, J. Chem. Soc. (1955) 3932–3936. doi:10.1039/JR9550003932. [5] J. Glascott, A model for the initiation of reaction sites during the uranium–hydrogen reaction assuming enhanced hydrogen transport through thin areas of surface oxide, Philos. Mag. 94 (2013) 221–241. doi:10.1080/14786435.2013.852286. [6] D.L. Clark, S.S. Hecker, G.D. Jarvinen, M.P. Neu, Plutonium, in: L.R. Morss, N.M. Edelstein, J. Fuger, J.J. Katz (Eds.), Chem. Actin. Trans. Elem., Third, Springer, Dordrecht, The Netherlands, 2008: pp. 813–1264. [7] A. Migliori, I. Mihut, J.B. Betts, M. Ramos, C. Mielke, C. Pantea, et al., Temperature and time- dependence of the elastic moduli of Pu and Pu-Ga alloys, J. Alloys Compd. 444-445 (2007) 133–137. doi:10.1016/j.jallcom.2006.11.157. [8] S.S. Hecker, M.F. Stevens, Mechanical Behavior of Plutonium and Its Alloys, Los Alamos Sci. 26 (2000) 336–355. [9] G.G. Libowitz, Metal Hydrides, in: W.M. Mueller, J.P. Blackledge, G.G. Libowitz (Eds.), Met. Hydrides, Elsevier, 1968: pp. 490–544. doi:10.1016/B978-1-4832-3215-7.50015-0. [10] G.W. McGillivray, J.P. Knowles, I.M. Findlay, M.J. Dawes, The plutonium/hydrogen reaction: The pressure dependence of reaction initiation time, J. Nucl. Mater. 385 (2009) 212–215. doi:10.1016/j.jnucmat.2008.09.046. [11] G.W. McGillivray, J.P. Knowles, I.M. Findlay, M.J. Dawes, The plutonium/hydrogen reaction: The pressure dependence of reaction initiation time and nucleation rate controlled by a plutonium dioxide over-layer, J. Nucl. Mater. 412 (2011) 35–40. doi:10.1016/j.jnucmat.2011.01.123. [12] L.N. Dinh, S.K. McCall, C.K. Saw, J.M. Haschke, P.G. Allen, W. McLean, The plutonium–hydrogen reaction: SEM characterization of product morphology, J. Nucl. Mater. 451 (2014) 143–146. doi:10.1016/j.jnucmat.2014.03.058. [13] R.D. Nelson, Phase transformation, J. Nucl. Mater. 20 (1966) 153–161.
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Chapter 10 IN SITU GROWTH OF HYDRIDE
REACTION SITES ON AN ALPHA/DELTA MIXED
PHASE PLUTONIUM ALLOY
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10.1 Cover Page
This chapter is included in the thesis to demonstrate the effect that a mixed phase alloy has on the formation and microstructure of plutonium hydride reaction sites. This chapter completes the three possible states that plutonium could reasonably be found; the two previous chapters covered the δ-stabilised alloy and α- Pu phases. The information in this chapter was gained using in situ hydriding within a FEGSEM, followed by subsequent analysis to study the role of the phase has on the growth of the hydride. The work in this chapter is ready for submission to a journal as an article.
The author carried out all SEM and hydriding work on this sample, the sample preparation was carried out by the technical support team under the direction of the author.
Word count: 3133
The references in the text of this chapter are numbered as for the main references of the thesis, but the original list of references for this paper can be found at the end of this chapter
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10.2 Title
The reaction between hydrogen and as cast plutonium alloy observed by in situ electron microscopy
10.3 Author names and affiliations
M. Brierley a b *, J. P. Knowles a, M. Preuss b
* Corresponding author
a Atomic Weapons Establishment, Aldermaston, Berkshire UK
b University of Manchester, Manchester, UK
© British Crown Owned Copyright 2015/AWE
10.4 Abstract
A sample of a mixed phase (α/δ) as cast plutonium alloy was exposed to hydrogen at a pressure of 3.3 mbar (330 Pa) at 20 °C and the reaction product was analysed in situ using a field emission gun scanning electron microscope. The exposure resulted in the nucleation and growth of oblate hemispherical reaction sites and high resolution micrographs of unoxidised plutonium hydride product and the reaction interface located in mixed phase regions are reported. The hydride consists of needle like particles (ca 160 nm wide) and the volume expansion associated with hydride formation generated stresses which were accommodated by deforming the ductile δ phase and cracking the brittle α phase.
10.5 Introduction
Safe consignment of plutonium metal requires a thorough understanding of its corrosion in a variety of environments. During extended storage of plutonium metal in sealed vessels hydrogen may accumulate as a result of radiolysis of organic materials such as plastic packaging and by reaction of the metal with outgassed water vapour. Therefore, a potential to form plutonium hydride exists. [64] Concerns
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regarding hydride formation arise from issues surrounding its pyrophoric nature which has the potential to cause dispersal of plutonium outside of suitable containment. [97] For the safe and long term storage of plutonium, a comprehensive knowledge of hydriding kinetics is desirable.
In the unalloyed form plutonium undergoes six allotropic phase transitions from liquid to the room temperature stable α-phase. Considerable changes occur in the plutonium during cooling from the melt with the volume decreasing by ca. 25 % from the lowest density fcc δ-phase to the highest density monoclinic α-phase. The δ-Pu phase can be retained to room temperature by alloying with certain elements, including Al, Ga, Am, and Ce [25,214,215]. Plutonium is usually stabilised in the δ- phase by small amounts of Ga [40] as this makes a plutonium component dimensionally stable with temperature fluctuations [124] and makes it easier to form into components, being soft and ductile compared to the hard, brittle α-Pu [18]. Plutonium alloys undergo a strong coring effect when they cool from the ε to the δ- phase [38], which results in high concentrations of Ga in the grain centres. With continued cooling to room temperature the grain boundary regions, which have insufficient Ga to stabilise them, continue to transform via the γ and β-phases to the α-phase, resulting in a mixed α/δ phase material. Heat treatment is required to redistribute the Ga sufficiently to ensure full stabilisation of the δ-phase. As-cast plutonium metal could therefore be found under storage conditions as a mixed phase alloy therefore, this research assessed the effect of phase on the hydrogen corrosion properties of plutonium. To achieve this an as cast sample was exposed to hydrogen and analysed by in situ scanning electron microscope (SEM) and post experimental cross sections of corrosion features.
10.6 Experimental
All handling, cutting, polishing and loading steps took place within a dry nitrogen glovebox system prior to transfer into a specially designed JEOL JSM7000F field emission gun scanning electron microscope (FEGSEM). The instrument has two adjoining vacuum chambers to accommodate sample loading and sample preparation, including Ar+ ion milling, heating, cooling and exposure to hydrogen. Further details of the instrument and relevant procedures have been described
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elsewhere [216]. Electron Probe Microanalysis (EPMA) was used to generate Pu and Ga maps of the as cast material using a JEOL JXA8200, which utilised the same transfer system as the FEGSEM.
10.6.1 In-FEGSEM hydriding reaction
The exposure of the sample to hydrogen took place in the preparation chamber at a hydrogen pressure of 3.3 mbar (330 Pa), the base pressure of the preparation chamber was 1.1 x 10-4 Pa prior to admittance of hydrogen. The sample was exposed to hydrogen three times successively at progressively longer exposure times of 60 s, 7200 s and 70320 s. Upon completion of each exposure, the gas was evacuated and the base pressure of the chamber was returned to 1.1 x 10-4 Pa. The sample was transferred to the analysis chamber under a vacuum of 10-5 Pa for analysis.
10.6.2 Material
A 3 x 10 x 10 mm sample was cut from a cast billet of a mixed phase 0.32 wt% as-cast Ga alloy prepared following a methodology previously described [170]. The material contained elongated grains typical of as-cast Pu-Ga alloys with δ-Pu in the grain centres of a matrix of α-Pu [170]. The density of the as cast sample was measured to be 16.94 g cm-3 and the rule of phase mixtures indicated the material was comprised of 31 % α-phase with the remainder δ-phase.
10.6.3 Sample preparation
The sample was prepared by mounting in epoxy resin with a 10 x 10 mm face exposed before polishing to 1 µm using diamond paste and lapping oil. The sample was broken out of the mount and electropolished in 10 % nitric acid in 2- ethoxyethanol then loaded into the FEGSEM.
Following the hydrogen dosing procedure and surface analysis, the sample was removed from the FEGSEM and returned to the glove box system where it was mounted in epoxy resin and cross-sectioned by serial polishing to 1 µm and subsequent electropolishing. The cross section was then coated with evaporated aluminium for conductivity and loaded into the FEGSEM for analysis.
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10.7 Results
10.7.1 Sample state before hydrogen exposure
As cast Pu-Ga alloy exists as a two-phase material (Figure 51a). Electron Probe Microanalysis (EPMA) revealed the Ga concentration (Figure 51b) in the as cast materials to be heavily segregated, with negligible Ga existing at the grain boundary regions. For clarity, contiguous regions of α-Pu and grain cores of δ-Pu will be subsequently referred to as domains since each grain comprises two phases. The α-domain crosses the boundary between two grains and there is no reason that the α-phase material either side of a grain boundary should have any related orientation. In cross section, the α-Pu domains were often connected by thinner ligaments as a matrix, but in some cases isolated domains of α-Pu could be found. The electropolish was slightly aggressive and the δ-phase etched more quickly than the α-phase so each phase had differing surface morphologies which aided identification.
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Figure 51 (a) Back scatter electron image showing the microstructure of the mixed phase alloy, the denser α-Pu domains have a brighter contrast. (b) EPMA WDS map of Ga distribution in the sample (different area to a) showing distinct Ga segregation between the grain centres and grain boundaries.
Precipitate inclusions were commonly found throughout the sample and were presumably more resistant to the electropolish than plutonium since they remained proud of the surface. These inclusions took the form of sheets of thickness ca. 100 nm at α/α grain boundaries (Figure 52a). Even though δ/δ grain boundaries were present in the material, no sheet-like precipitates were found at them. Cubic precipitates, around 1 µm size were found within the δ-domains, often near to the boundary with an α-domain (Figure 52b) but these were fewer in number. Both types of precipitate were not reliably identified using EDS although it is believed that the
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sheets were the Pu6Fe eutectic [121], and the cubes were the Pu6Fe phase abutting onto an fcc δ-Pu grain.
Figure 52 SEI micrographs showing (a) a typical thin sheet grain boundary precipitates ca. 100 nm thick at a α/α grain boundaries and (b) a typical cubic precipitate found in δ-domains.
10.7.2 In situ hydriding
The hydrogen dosing was conducted in a staged manner to investigate the possible existence of small early stage nuclei. The first dose provided a 60 s exposure in an attempt to investigate if early stage nucleation occurred on plutonium in a similar manner to that observed on uranium by Owen and Scudamore and later by Arkush et al. [110,217]. No change in the surface was evident at high magnification in the SEM. The sample was returned to the preparation chamber and dose for a subsequent 7200 s exposure before SEM analysis and again, no
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nucleation of hydride reaction sites was evident during analysis at high magnification. A third exposure of the sample to H2 for 70320 s was carried out, taking the total exposure to 77580 s. Reaction of the plutonium with hydrogen was confirmed by a drop in chamber pressure equivalent to 96 % of the available hydrogen and the sample was subsequently analysed in the SEM. This staged dosing, although far from conclusive, may suggest that the reaction sites on plutonium result from a single type of nucleation site unlike the early stage growth found by Owen and Scudamore.
One reaction site had nucleated on the front face of the sample at a region not obscured by the sample holder which had liberated particulate material The site was large and extended under the edge of the sample mount so could not be viewed in its entirety although the visible area was sufficient to enable a detailed investigation.
Significant quantities of plutonium hydride had been ejected from the reaction site during its growth and was scattered across a large proportion of the sample. It has been hypothesised that spalled plutonium hydride may nucleate additional reaction sites underneath it [106], however no additional reaction sites were observed under any of the spalled material. The particulate plutonium hydride was formed from agglomerates of needle-like sub particles with a common orientation (Figure 53), which closely resembles that reported for uranium hydride [218]. The rods were 160 nm wide on average but of variable length in various parts of the sample, determined by the position of fractures transverse to the rods. The spalled particles were between 5-100 µm in size with the greatest frequency occurring around 10-20 µm.
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Figure 53 Secondary electron image of a plutonium hydride particle which shows the material is comprised of agglomerated needles of plutonium hydride ca 160 nm wide with a common orientation.
The formation of plutonium hydride results in an expansion owing to the lower density of the hydride compared to the metallic phases and is known to generate a significant compressive stress [202,207,219] which must be accommodated by the material surrounding the reaction site. The as cast sample was especially chosen to observe how the stress is accommodated in both the α and δ phases.
The compressive stress induced cracks in the surface of the brittle α-domains near to the reaction interface and occurred within 1-3 µm of the interface (Figure 54). The cracks were formed either by stresses applied to the α-domains from the hydride/metal interface directly or by distortion of adjacent δ-domains by the hydride dilation.
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Figure 54 Secondary electron image showing hydride formation within an alpha region and brittle fracture of the alpha phase close to the interface.
The surface adjacent to the hydride/metal interface in the δ-Pu domains was observed to have deformed in response to the dilation of the hydride (Figure 55a). High resolution images show fractures in the oxide over-layer covering the distorted regions (Figure 55b). Distortion of the native oxide has been described as a mechanism by which anisotropic growth can occur on surface hydride reaction sites for cerium [195,220] and plutonium [219].
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Figure 55 Backscatter electron images showing the distortion of δ-domains at the hydride/metal interface
When the stress field imparted by the hydride dilation included both δ and α domains, the distortion was mainly accommodated by the δ-phase, with the α-phase remaining relatively unaffected (Figure 56a). The distortion was exhibited as wrinkles in the surface and where the wrinkles existed, fine scale cracks were seen in the native surface oxide. Figure 56b is a magnified view of the black ringed area of Figure 56a, which shows a narrow α-domain between two δ-domains. The δ- domains are heavily distorted with cracking to the native oxide layer but the α- domain has fractured near to the hydride interface. The stiffer α-domain has also transmitted the stress further into the δ-domains, creating a secondary stress concentration (ringed in white) deeper into the adjacent grains, where further wrinkling in the δ domains and another fracture in the α-domain has occurred.
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Figure 56 Backscatter electron images showing a) a region of mixed phase surface adjacent to the hydride interface, b) a detailed view of the black ringed area in (a) which shows distortion in the δ-domains and fracture in the α-domain.
10.7.3 Post hydriding cross sections
Following unloading of the sample prior to cross sectioning three additional reaction sites were located on the front face of the sample which had been hidden underneath the sample holder. Also, a considerable amount of hydride had formed on the rear side of the sample as 4 reaction sites.
High resolution investigation of the cross-sections through hydride sites was carried using the FEGSEM (Figure 57 and Figure 58). The radii and depths of the sites suggested the sites had grown as oblate hemispheres as previously observed [195,219,220] with aspect ratios between 1.9 and 4.2. A cross section through the
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surface surrounding a hydride reaction site demonstrates the distortion had occurred by the mechanism of plane slip as a number of slip zones are visible in the material to the reaction site (Figure 57). Analysis of the reaction interface did not reveal any preference for growth within α-Pu or δ-Pu domains although the hydride within the site differed in morphology (Figure 58), which may relate to the domains of α and δ. Although it is impossible to say with certainty, it appears that the hydride formed within α domains consisted of large particulates (>10 µm) and had brighter backscattered electron contrast than the hydride formed in δ domains (< 10 µm).
Figure 57 Secondary electron images of the hydride/metal interface at the metal surface.
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Figure 58 Backscatter electron images of cross sections through the same site, (a) highlighting the oblate morphology as the radius (r) is greater than the depth (d) and (b) detailing of the reaction interface.
On occasion where thin sheet grain boundary precipitates had been encountered by the hydride interface; hydride had preferentially precipitated around them in advance of the main hydride interface. The metal region in Figure 59 contains several grain boundary precipitates and hydride has formed around the precipitates, extending 3-6 µm further than the hydride interface. This did not contribute significantly to formation of large discontinuities in the hydride interface because the grain boundary precipitates were relatively small in number and short in length (< 10 µm).
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Figure 59 Advanced precipitation of hydride at a sheet-like α/α grain boundary precipitate, here labelled GBP.
10.8 Conclusions
A sample of a mixed phase (α/δ) as cast plutonium alloy was exposed to hydrogen at a pressure of 3.3 mbar (330 Pa) and resulted in the nucleation and growth of oblate hemispherical reaction sites.
High resolution in situ micrographs of the unoxidised plutonium hydride particulate product revealed plutonium hydride to have structure similar to uranium hydride with a particulate material comprised agglomerated needles of plutonium hydride ca 160 nm wide.
The volume expansion associated with hydride formation generated stresses in the plutonium surrounding the reaction sites sufficient to result in cracks developing within the brittle α phase and the δ phase to undergo ductile deformation.
10.9 Acknowledgements
This work was funded by AWE and carried out in conjunction with the University of Manchester. Thanks go to the technical support staff at AWE for sample preparation and to Mike Matthews and Matt Denness for assisting the gas dosing operations.
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10.10 References
[1] F. Brown, H.M. Ockenden, G.A. Welch, The Prepration and Properties of some Plutonium Compounds, Part I. Plutonium Hydride, J. Chem. Soc. (1955) 3932–3936. doi:10.1039/JR9550003932. [2] J.M. Haschke, J.L. Stakebake, Handling , Storage , and Disposition of Plutonium and Uranium, in: L.R. Morss, N.M. Edelstein, J. Fuger (Eds.), Chem. Actin. Transuranium Elem., Springer, Dodrecht, 2006: pp. 3199–3272. [3] S.S. Hecker, Plutonium and Its Alloys From atoms to microstructure, Los Alamos Sci. 26 (2000) 290 – 335. [4] American Society of Metals, Volume 3, Alloy Phase Diagrams, in: ASM Handbooks Online, 2004. [5] D.C. Miller, J.S. White, The mechanical properties of plutonium-aluminium and plutonium-cerium alloys, J. Nucl. Mater. 10 (1963) 339–345. doi:10.1016/0022-3115(63)90185-5. [6] D.L. Clark, S.S. Hecker, G.D. Jarvinen, M.P. Neu, Plutonium, in: L.R. Morss, N.M. Edelstein, J. Fuger, J.J. Katz (Eds.), Chem. Actin. Trans. Elem., Third, Springer, Dordrecht, The Netherlands, 2008: pp. 813–1264. [7] S.S. Hecker, D.R. Harbur, T.G. Zocco, Phase stability and phase transformations in Pu–Ga alloys, Prog. Mater. Sci. 49 (2004) 429–485. doi:10.1016/S0079-6425(03)00032-X. [8] S.S. Hecker, The Magic of Plutonium : 5f Electrons and Phase Instability, Metall. Mater. Trans. A. 35 (2004) 2207–2222. [9] J.N. Mitchell, M. Stan, D.S. Schwartz, C.J. Boehlert, Phase Stability and Phase Transformations in Plutonium and Plutonium-Gallium Alloys, Metall. Mater. Trans. A. 35A (2004) 2267–2278. [10] M. Brierley, J.P. Knowles, M. Preuss, The reaction product of hydrogen and electro-refined plutonium observed by in situ electron microscopy, 2015. [11] M. Ling, R.F.E. Jenkins, N. Park, Solid state phase transformation study on a series of Pu–Ga alloys containing 0.18–0.63wt.% Ga, J. Alloys Compd. 586 (2014) 709–717. doi:10.1016/j.jallcom.2013.10.112. [12] C.J. Boehlert, R.K. Schulze, J.N. Mitchell, T.G. Zocco, Initial electron backscattered diffraction observations of a plutonium alloy, Scr. Mater. 45 (2001) 1107–1115. [13] L.W. Owen, R.A. Scudamore, A microscope study of the initiation of the hydrogen - uranium reaction, Corros. Sci. 6 (1966) 461–468. doi:10.1016/S0010-938X(66)80053-7. [14] R. Arkush, A. Venkert, M. Aizenshtein, S. Zalkind, D. Moreno, M. Brill, et al., Site related and growth on uranium surfaces I ’, J. Alloys Compd. 244 (1996) 197–205. doi:10.1016/S0925-8388(96)02505-4. [15] C.K. Saw, J.M. Haschke, P.G. Allen, W. Mclean, L.N. Dinh, Hydrogen corrosion of plutonium: Evidence for fast grain-boundary reaction and slower intragrain reaction, J. Nucl. Mater. 429 (2012) 128–135. doi:10.1016/j.jnucmat.2012.05.044. [16] C. Ablitzer, F. Le Guyadec, J. Raynal, X. Génin, a. Duhart-Barone, Influence of superficial oxidation on the pyrophoric behaviour of uranium hydride and uranium powders in air, J. Nucl. Mater. 432 (2013) 135–145. doi:10.1016/j.jnucmat.2012.08.008. [17] Y. Greenbaum, D. Barlam, M.H. Mintz, R.Z. Shneck, Elastic fields generated by a semi-spherical hydride particle on a free surface of a metal and their effect on its growth, J. Alloys Compd. 509 (2011) 4025–4034. doi:10.1016/j.jallcom.2011.01.010. [18] M. Brierley, J. Knowles, Probing the cerium/cerium hydride interface using nanoindentation, J. Alloys Compd. 645 (2015) S148–S151. doi:10.1016/j.jallcom.2015.01.116. [19] M. Brierley, J.P. Knowles, A. Sherry, M. Preuss, The microstructure of plutonium hydride growth sites, AWE plc, Reading, UK, Manuscript submitted for publication, 2015. [20] M. Brierley, J. Knowles, N. Montgomery, M. Preuss, Microstructure of surface cerium hydride growth sites, J. Vac. Sci. Technol. A. 32 (2014) 031402–1–031402–8. doi:10.1116/1.4867475. [21] J.P. Knowles, G. Rule, M. Brierley, The morphology and anisotropic growth kinetics of cerium hydride reaction sites, Corros. Sci. 77 (2013) 31–36. doi:10.1016/j.corsci.2013.07.020.
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Chapter 11 OVERALL DISCUSSION
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The overall discussion will follow the original list of questions from the questions posed in chapter 4, which are repeated here with the original numbering for clarity. Discussion of other aspects of the work which was not guided by these questions is also discussed later.
4.1.1 Does the overall shape of the hydride evolve isotropically?
4.1.2 Do gradients of hydrogen dissolved in the metal direct growth morphology?
4.1.3 Is the hydride morphology directed by strain introduced into material surrounding a dilated hydride?
4.1.4 Does the microstructure provide evidence of growth characteristics?
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11.1 Does the overall shape of the hydride evolve isotropically?
The degree of anisotropy of the hydride sites is described by the ratio of the lateral radius compared to the depth of the hydride reaction sites. For cerium, this ratio was ca 4.0, whilst for fully δ-stabilised plutonium was ca 2.2. The few sites cross sectioned on a mixed δ/α-Pu sample had a similar aspect ratio to that of the fully δ-stabilised Pu, lying in the range of 1.9 to 4.1. On the δ-stabilised plutonium, every hydride reaction site had clearly grown faster laterally than by penetration into the sample. The mixed phase sample had a similar degree of anisotropy despite there being α-Pu domains present.
This work clearly demonstrates that anisotropic growth occurs on both cerium and δ-phase plutonium. Except for the case of one very large hydride reaction site on cerium (Figure 60), which appeared to have nucleated deep into the sample, every site observed had a higher surface radius than depth, giving an overall oblate hemispherical shape. The large site appeared to have nucleated below an existing non isotropic reaction site, where the oxide on the surface could play no part in the continued growth of the hydride. The low test pressure of 10 mbar (1 x 103 Pa) used on this sample may have kept the nucleation and growth rates of cerium hydride low, allowing hydrogen time to percolate into the bulk via a metallurgical defect (not observed) prior to precipitation of the hydride at a favourable nucleation site.
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Figure 60 Top: an unusually large spherical hydride which appears to have formed below a large, flatter surface hydride. Bottom: a typical oblate cerium hydride in cross section showing a two layered oxide/hydride structure of an air-exposed hydride.
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11.2 Does the microstructure provide evidence of growth characteristics?
11.2.1 Cerium
The cerium hydriding reactions appeared to retain the surface over the hydride reaction sites. Grain boundaries and triple points were commonly found in the surface of the reaction sites. These features may be the site of nucleation, however because the sites were grown to a size so that the growth rate could be measured, they exceeded the grain size of the metal. This meant that a number of different grain boundary or triple point features could also have been the site of nucleation, and in a number of cases, these features were not directly in the centre of the hydride reaction site. It is also possible that the grain boundaries and triple points opened up as a result of the dilation of the site placing the outer surface into tensile stress, and were not involved with nucleation.
Cerium hydride retained the grain microstructure of the parent metal which dilated and hardened. Following exposure to oxygen, the hydride reaction sites underwent a degree of spallation, with intergranular fracture occurring between the grains and entire crystallites of oxidised cerium hydride falling out of the reaction sites. The interface between the cerium and the cerium hydride was abrupt, with no evidence of a diffusive zone. The interface was largely continuous, and was observed to cross cerium metal grains with no evidence that any particular grain had accelerated or impeded progress, which suggests that the formation of hydride within cerium metal occurred isotropically with no preference for precipitation along grain boundaries. The CeO2 dendrites found in the parent material did not appear to affect the growth of the hydride.
Needle-like or plate-like laths were observed surrounding the reaction sites in cerium that had been formed on a polished surface, which were postulated to be cerium which had transformed from γ-Ce to α-Ce under the outward radial stress imparted by the dilated hydride. Cracks in the surface, hence the surface oxide, were observed to have occurred directly above the laths, which would allow hydrogen to access underlying metal without having to diffuse through an oxide barrier. This
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would create a region of the surface slightly larger than a reaction site that hydrogen could pass through more easily than intact oxide.
11.2.2 Electro-refined plutonium
The hydride growth on electro-refined plutonium sample was slow. The sample was exposed for two days to hydrogen at 3.3 mbar (330 Pa), during which time it would have been in the form of α-Pu. A thermal treatment to 287 °C was performed to initiate reaction as observed by Brown et al. [64] and allowed to cool to room temperature. This thermal treatment would have been sufficient to transform the sample via the β-phase to the γ-phase. The subsequent transformations on cooling to room temperature would have been via the γβ transformation, described by Mitchell et al. as sluggish [38], however other reports by Nelson [221] about the transformation suggest that transformation to the β-phase should complete in a matter of minutes or hours. Subsequent transformation from the β-phase to the α- phase would then complete rapidly [42,159,221,222], completing in minutes. The sample continued to be exposed to hydrogen for three days following the heat treatment
The growth of hydride on the electro-refined sample favoured a complete surface coverage, with an average thickness of 15 µm, underneath the native oxide. The hydride had also become powdery, with a reasonable proportion of void space, which means that less plutonium was consumed than this estimate. The void space was measured using ImageJ to constitute 41 % of the total volume of the hydride reaction layer. Calculating the depth of plutonium converted to hydride was thus obtained from the following calculation, using the volume of PuH2: 14.8 µm (hydride layer thickness) x 0. 9 % (particle volume) x 1/1.963 (α to PuH2 volume ratio) = 4.5 µm. This was unexpectedly slow. The mixed phase sample had been reacted with hydrogen in the system prior to the alpha-phase sample and had consumed to 96 % of the available hydrogen within 21.55 hours. Subsequent analysis on spalled flakes of surface oxide (Figure 61) revealed the layer to be ca 2.2 µm thick. It is possible that the oxide thickness on this sample had exceeded a critical thickness above which hydride precipitation will not occur [3]. The heat treatment applied to this sample in vacuum would have converted the oxide from mainly PuO2 to mainly
Pu2O3 [69]. The role of Pu2O3 has been discussed in the literature with respect to
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hydriding [105] and, in this case it is possible that the oxide would have stopped acting as a passivating layer and allowed hydride to form over the entire surface.
Figure 61 An edge-on view of a spalled piece of surface oxide showing a change in morphology at around 2.2 µm depth.
In addition to full surface coverage a number of protrusions were observed on the surface of the sample following hydrogen exposure, however these did not have the same oblate morphology as the hydride reaction sites observed on the δ-Pu and Ce samples; instead favouring growth along internal surfaces of the micro cracks which intersected the sample surface. The hydride protrusions at the surface were formed from the metal around the openings which had transformed to hydride. There was more depth to the hydride around the micro cracks and dilation in these regions had pushed the surface upwards preferentially compared to the surrounding surface. The cross sectional work confirmed that bulges were present over the top of micro cracks.
The oxide above a selected region of the sample was locally thinned to observe the microstructure using the ion mill. No preferential formation of hydride was observed in that region. EBSD was attempted in that region; however no Kikuchi patterns were obtained, which suggests that although the oxide had been thinned it was not removed entirely.
A question remains from the slow surface wide coverage as to why, given the reduction and subsequent spallation of the passivating oxide layer and the open
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powdery structure of hydride, did more of the electro-refined-Pu not transform to hydride while exposed to hydrogen? Growth of hydride took place, potentially while the sample was at elevated temperature, which leaves two possible explanations:
(1) Either, the hydride when cooled acted as an extraordinary barrier to diffusion. Given the open structure of the hydride layer formed this appears unlikely, and stresses induced upon cooling and phase change of the plutonium are likely to induce defects in any adherent layer, reducing the effectiveness of a diffusion barrier.
(2) Or, plutonium in the α-phase is relatively unreactive with hydrogen. This explanation also seems unlikely because the reaction of unalloyed plutonium powder has been reported to be identical to that of δ-stabilised Pu 1 wt% Ga alloy [203]. However, no results have been found in the literature for hydriding of bulk unalloyed plutonium at the pressure and temperature regime used in this test.
11.2.3 Delta plutonium
The hydride material that had been formed on δ-stabilised plutonium had a microstructure that was distinct from that of the δ-Pu, consisting of small platelets which appeared to be oriented radially from the centre of the hydride reaction site near the outer surface. The samples that this was observed on had been passivated by careful exposure to an oxygen environment, which could have affected the morphology in the oxidised region. The inter platelet boundaries were observed to be the locus of failure of the hydride material, which is a new finding which does not support previous suggestions that the powder forms as the loss of entire grains from the surface [106]. The inter-platelet boundaries also provide a path of hydrogen to readily diffuse through hydride material, it would seem that earlier discussion of the requirement for superionic conduction through the hydride [97,102] is unnecessary to explain the lack of barrier posed by the hydride reaction product [203].
The δ-Pu samples exhibited spot wise growth, resulting in oblate hemispheres growing near to the interface between the metal and the oxide layer. The aspect ratio of the hydrides had an average value of 2.2. The interface formed with the plutonium metal had three different types of discontinuous features. A discontinuous interface suggests that the microstructure cause differences in the diffusion rate of hydrogen
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and subsequent precipitation of hydride adjacent to the interface. The extent of hydride appeared to be highly dependent on particular directions within the δ-Pu grain structure. These samples were reacted with hydrogen in an external cell at higher pressures between 1 x 104 and 1 x 106 mbar (1 x 106 and 1 x 108 Pa), however there are no currently understood mechanisms that would change the morphology between the pressures used in the external cell and FEGSEM preparation chamber.
11.2.4 Mixed α/δ phase plutonium
The mixed phase sample investigated had approximately 36% area fraction of α-Pu, measured here by image analysis of backscattered electron images of the microstructure. Density measurements carried out previously on this alloy gave a density of 16.94 g/cm3 which corresponds to an alpha volume fraction of 31% (unreported data) which is in reasonable agreement with those obtained by image analysis. The α-phase in this case would contain nominally zero gallium as it has been cooled slowly from the cast, allowing sufficient time at temperature for the gallium to be ejected from the γ and β-phases before the α-phase formed [170]. No micro cracks were present because the transformation stresses would be accommodated by deformation processes within the δ-phase domains during cooling.
There was no significant difference in the position of the interface observed in cross section in either the δ or α domains within the material. Deformation at the surface was almost entirely found within the δ domains as a bulging of the surface with evidence that plane slip had occurred. There were no obvious discontinuities with the interface between the hydride and the metal as found in the fully δ-stabilised alloys. This could be because to fully stabilise plutonium in the δ-phase a thermal treatment is necessary to homogenise the gallium content, thereby reducing the gallium composition of the δ cores in the grain to increase the grain edge composition. This sample was as cast and therefore no thermal treatment had taken place. It is possible that the δ-Pu domains had a higher concentration of gallium than is sufficient to stabilise the δ phase, increasing the stability in these regions against physical stresses caused by the hydride dilation. The α-phase is the most stable
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allotrope under pressure at room temperature and is unlikely to change under compression.
The hydride growth morphology was similar to that of δ-Pu, being of oblate morphology, with similar lateral radius to penetrative aspect ratios. This strongly suggests that morphology of the surface hydrides forms as a function of the weakest link in the system. Once a hydride had precipitated the dilation imparted stress to the surrounding region and the softer δ-Pu deformed to a greater degree than the α-Pu and caused the failure of the oxide layer, which would have reduced the barrier to diffusion there. Loss of the diffusion barrier surrounding a hydride has been proposed as a result of earlier work in this EngD as a mechanism which allows the oblate hemispherical morphology to occur as observed on this sample and both the cerium and δ-stabilised plutonium samples.
Although the hydride interface with the metal had no significant preference for δ-domains over α-domains it may be that the δ-domains induced a secondary effect in the α-domains which kept the rate of growth similar between the two phases. Fracture of the metal in the α-domains was observed, especially at points of high stress adjacent to considerable surface distortion in the δ-domains. The α-phase often became broken up into pieces, which may allow many pathways for hydrogen to enter the metal in the α-domains, maintaining a synchronised overall position of the hydride interface between δ and α domains.
Sheet-like grain boundary precipitates observed in the mixed phase sample gave a small advantage to the growth of the hydride interface with hydride having nucleated at the boundary between the precipitates and the metal. In the metallographic analysis of the sample prior to the hydriding, these precipitates were observed only at α/α grain boundaries, and all were less than 10 µm long. These precipitates did not cause significant advantage to the growth of hydride.
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11.3 Is the hydride morphology directed by strain introduced into material surrounding a dilated hydride?
During surface preparation prior to polishing, a number of grinding and polishing steps were employed to achieve a repeatable surface. Grinding and polishing is well understood by metallurgists to cause deformation to the underlying material in the sample surface. This deformation could influence the oxidation and subsequent hydriding behaviour of the samples. Both Ce and δ-Pu are soft materials and plastic deformation of the surface during polishing is possible, resulting in a Beilby (amorphous) layer at the surface [157,223] with further subsurface strain and deformation able to take place beneath. Care was taken to remove the damage introduced by previous steps on the Ce samples using a final colloidal silica step, which gives a gentle attack polish concurrently etching and polishing the sample with very fine polishing media, leaving a surface with minimal distortion. In the situation where the δ stabilised Pu has sufficient gallium it can be stabilised to room temperature, but where the composition is close to the stability limit compressive mechanical stresses can cause a transformation to the α’ phase (Figure 62). The subsurface microstructure of δ-Pu could potentially comprise plastically and elastically strained material with transformed material present.
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Figure 62 Transformation of δ-Pu to α’-Pu beneath polishing scratches near to grain boundaries with lower gallium concentration in a partly homogenised sample.
The depth of the layers of mechanically induced deformation is not currently well understood in plutonium, however the depth of mechanically induced α’ has been found to be ca 24 µm [118]. The Beilby layer and elastic deformation would have certainly been present in the δ-Pu samples that were ground to 400 grit and then oxidised. The oxidation would have consumed a proportion of the Beilby layer and potentially underlying material. The mixed phase α/δ-Pu sample and the electro refined Pu sample were electropolished to remove distorted material, but the possibility of some remaining cannot be conclusively excluded. Figure 62 shows an example of an insufficiently electropolished sample with α’ existing near scratches. Analysis of the microstructure on the sample used in these tests following electro polishing indicated no α’ phase was present (as seen in Figure 62). Oxidation during transfer in the glove box system would have likely consumed a proportion of any Beilby layer present.
The relative volume of the MH2 hydride phase compared to the parent metal from which it forms is 1.264 for γ-Ce, 1. 7 for δ-Pu and 1.963 for α-Pu. Dilation of the hydride material is constrained by the surrounding parent metal phase, which imparts stresses into the surrounding material. Calculation of the stresses
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surrounding a hydride reaction site has been carried out for hydrides in isotropic bulk and shown to exhibit a compressive radial stress in the material surrounding the hydride but a tensile tangential stress [182]. The radial stresses perpendicular to the surface can be relieved by deformation of the hydride and surrounding material out of the plane of the surface, but radial stresses parallel to the surface cannot be relieved. The extent of the stressed region near the surface surrounding a hydride reaction site is consequently much larger than underneath (Figure 63).
a) b)
Figure 63 a) The radial compressive stresses (depicted in red) and the tangential tensile stresses (green) that surround the hydride at the surface. b) The extent of the stressed zone is anisotropic because of the ability of the forces normal to the surface to push the hydride outwards reduces the resultant force
The size of the stress field has been measured using nanoindentation to be 2- 3 µm underneath a ca. 90 µm cerium hydride reaction site. The extent of lateral stress field at the surface of surrounding cerium hydrides have been estimated to be 10-30 µm from coronas surrounding the sites following removal of the oxide layer using ion beam milling. Transformed laths of material surrounding a ca. 90 µm cerium hydride reaction site was shown to extend ca. 20 µm from the interface with the metal (Figure 64).
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Figure 64 Laths surrounding a cerium hydride reaction site.
Breakup of the α-Pu was observed on the mixed phase alloy without any significant plastic distortion to the metal. There was no evidence of cracking to the surface over any part of the α-Pu without there being breakage of the α-Pu itself. Distortion of the soft metal at the surface surrounding the site was found in the cerium, the δ-stabilised Pu sample and in the δ-Pu domains of the mixed phase alloy. The surface of the distorted metal exhibited fine scale cracking indicating that the native surface oxide had failed above these regions.
The relative volume of hydride to the base metal causes the hydride as a whole to dilate on transformation. When the material transforming is situated within the surrounding matrix, this expansion is constrained unequally by the surrounding material and the material below it, and so the hydride dilation imparts a compressive stress into the surrounding material. Simulations for the magnitude of this stress field
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for a given amount of material transformed to hydride have been calculated by Greenbaum et al. [182,202]. In their work, Greenbaum et al. calculated the minimum energy shape for hydrides to be hemispherical when the hydride is harder than the surrounding matrix. The nominal hardness of cerium and cerium hydride are 0.34
GPa and 2.3 GPa] and the nominal hardness of δ-plutonium, α-plutonium and PuH2 are 0.34 GPa, 0.82 GPa and 1.22 GPa respectively. The hardness values are greater for the hydride in every case which suggests from the calculations of Greenbaum et al. that the hydride reaction sites should be hemispherical; however the reaction sites as measured tended to have anisotropic morphology with a larger lateral radius than depth. Other mechanisms must also be important in the growth of surface hydrides to form a non-hemispherical shape. The different morphologies found in theoretical and experimental work has led to additional mechanisms being suggested. Since early stage precipitates could not be reliably found using cross section the morphology of early precipitates could not be confirmed, the possibility remains that they nucleate spherically or hemispherically as per theory before another process becomes dominant.
11.4 Hydride reaction site growth mechanism
The initiation models discussed earlier describe the initial stages in hydriding that the oxide acts as a diffusion barrier to hydrogen, slowing the rate at which hydrogen reaches the metal interface with the oxide, which causes an initiation time to be observed, defects such as thinner regions, grain boundaries and crystal defects act as preferred pathways for diffusion of hydrogen though the oxide. The current models describing the growth stage suggest that continued growth occurs via failure of the oxide barrier above the hydrides. Hydride is not considered to be a significant barrier to diffusion of hydrogen [203], which is supported by the observation of inter-platelet boundaries found in this investigation. The high diffusion rate of hydrogen in hydride implies that hemispherical sites should form because the hydride would be rapidly saturated with hydride and each point on the interface with the metal would have the same growth characteristics as the others.
This work demonstrates that hydrides grow anisotropically on cerium and δ- plutonium becoming oblate hemispheres at the surface. A mechanism for anisotropic
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growth has been proposed in both cases. The hydride phases in both cerium and plutonium have significantly lower density and greater hardness than the metal phases, which imparts a stress into the surrounding metal. These metals are relatively soft and distort to accommodate the radial outward stress. In some cases, evidence was found to support phase transformation around hydride reaction sites. The distortion of the metal surrounding the hydride has been observed to induce cracking in the surface, which would allow hydrogen easier access to the oxide/metal interface immediately surrounding the hydride site in addition to the hydride/metal interface. This would give an advantage to lateral growth compared to penetrative growth.
The anisotropic growth in cerium and δ-phase plutonium becomes a matter of the rate at which the oxide surrounding the existing hydride reaction site fails to act as a barrier to diffusion compared to the rate at which hydrogen penetrates into the metal via metallurgical diffusion processes.
In the mixed phase alloy, the overall growth mechanism resembled that of δ- Pu, with oblate hemispheres forming at the surface with similar lateral to penetrative aspect ratios. Although the unalloyed plutonium exhibited a thin surface wide coverage, it could be argued that this was the result of the process used to initiate growth. The δ-phase is a ductile phase compared to the α- phase and acts as a weak link, dominating the hydriding behaviour.
11.5 General observations
11.5.1 On passivation
The cerium cross sections that had been exposed to air all formed a two layer hydride structure. Use of light microscopy, backscattered electron imaging and EDS in the SEM, tapping mode AFM and nanoindentation all confirmed differences in the properties of the oxidised region compared to the metal or hydride. There was no change to the hydride reaction sites on cerium visually until they had been exposed to air briefly. EDS of the hydride cross sections revealed that the outer layer was rich in oxygen but there always remained an unoxidised inner layer. The oxidised layer extended roughly 50% of the hydride from the outer surface, leaving the remaining
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half of the hydride untouched. Except on one unusually large hydride reaction site (ca. 0.6 mm diameter), where the oxygen was found to have occupied the majority, but not all, of the hydride reaction site; sufficient cerium hydride remained on that reaction site for macro indentation to be performed.
The oxygen content of the plutonium hydrides could not be measured due to absorption of the oxygen X-rays in the aluminium coating required to ensure conductivity of the sample. Aluminium evaporation is used because it is possible to handle using tweezers in a glovebox, which ensures reliability of process. The plutonium hydrides from the δ-stabilised alloy which were reacted in an external rig had been passivated to ensure safety. Each hydride reaction site in that situation when viewed in cross section appeared to have a crust to it which appeared to be more tightly packed than the hydride underneath. Figure 65 shows examples of three hydride reaction sites with different diameters, each with a crust extending roughly 50% of the hydride reaction site. The proportional depth of these crusts is comparable to the oxygen rich layer in the cerium hydrides. It is reasonable to suggest that this is also the extent of the oxidation of the plutonium hydrides.
The depth of oxidation poses a potential problem for storage. It appears that non-reacted hydride remains within the reaction sites following passivation. There is potential that during handling the oxide crust could be damaged, which could potentially expose fresh hydride to an oxygen rich environment and lead to initiation of a vigorous reaction.
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Figure 65 Three hydride site cross sections of different sizes; site a) ca. 400 µm diameter, site b) ca. 304 µm diameter, site c) ca. 141 µm diameter; each with a crust extending ca. 50% of the depth of the hydride.
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11.5.2 Sites of nucleation
This aim of this study was not to investigate the induction and nucleation phases of hydride precipitates because plutonium hydride forms a powder, which subsequently destroys information about the microstructural features that are involved in the nucleation process. In the course of this work however, a number of features were co-located with hydride reaction sites on the cerium samples and the electro refined plutonium sample suggesting that there may be microstructural features which have influenced the formation of protrusions of hydride on the surface of samples.
The hydride reaction product that was observed to protrude from the surface of the electro-refined sample was associated with the presence of a surface intersecting micro crack in every case observed. The protrusions are not referred to here as hydride reaction sites because the overall reaction of the electro refined sample with hydrogen occurred across the entire surface. The protrusions in this case consisted of reaction product which had formed within micro cracks and had pushed the surface oxide out of the plane of the sample surface more than that which occurred above the surrounding surface hydride coverage.
Cracks were observed in the cerium metal during post reaction analysis and were a nucleation site for a number of cerium hydride reaction sites Figure 66. The majority of reaction sites on the cerium samples however did not have any cracks associated with them.
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Figure 66 Cracks observed in the cerium metal often had hydride reactions sites nucleated along them
In one case having made a cross section of a cerium hydride reaction site, an iron rich inclusion was found using EDS (Figure 67). This was not reported in the paper because it was the only incidence found in a number of hydride reaction sites cross sectioned and appears to have been coincidental; however it is reported here for completeness. It is important to state that in the case of the cerium, the presence of surface hydride reaction sites did not appear dependent on an underlying metallurgical feature because the majority of reaction sites had none apparently connected with them.
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Figure 67. Backscattered electron image of an iron rich inclusion seen in the cross section of an oxidised cerium hydride reaction site. The inclusion was analysed at higher magnification using EDS (red overlay bounded by the black rectangle), the optical image (bottom right) shows the remaining hydride as a yellow tinted area, indicated on the main image as the purple overlay, the green tint is the EDS oxygen map.
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11.6 Future work
Characterise oxide surfaces grown on plutonium samples prepared by
different means and the effect they have on hydriding behaviour
Develop a method of performing EBSD using the FEGSEM on plutonium and
use EBSD to investigate the extent of stress placed into the material
surrounding typical hydride reaction sites in plutonium.
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Chapter 12 CONCLUSIONS
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An ambitious project working on plutonium metal and using hydrogen has been completed. New equipment with the capability to perform in situ hydriding studies of plutonium was commissioned during this study. The first in situ analysis of the hydride product formed has been carried out on electro-refined plutonium, which exhibited an open powder like structure consisting of elongated particles. The expansion of the hydride underneath the native oxide had caused the oxide to spall. In situ analysis of hydride formed on a mixed phase alloy revealed a slightly different structure where platelets were co-aligned; large cracks had occurred in the hydride, splitting the hydride product into fingers. The ability to carry out hydriding in situ and prevent subsequent oxidation of the hydride reaction product is unique to this instrument.
The use of cerium was beneficial in developing a working methodology on a more simple system and then applying it to the more complex system of plutonium, Comparisons could then be drawn between the two metals which allowed some of the more complex behaviour of the plutonium system to be understood. Cerium exhibited no preference to crystal orientation for the growth of the hydride into the metal, retaining the crystal structure of the parent metal. Spallation of the hydride occurred by intergranular cracking after exposure to air.
A detailed surface analysis and microstructural investigation of plutonium hydride which had formed on δ-stabilised plutonium showed a strong dependence of the hydride morphology with crystal orientation. The hydride had formed as platelets which were aligned along particular directions in the grains of the parent metal.
The overall morphology of the hydride reaction sites was not hemispherical in either the cerium or the δ-plutonium cases. The material properties for γ-Ce and δ- Pu are sufficiently similar to expect similar mechanisms to take place on both and while a degree of variability between the two systems existed; the hydride reaction site morphology found on each is similar. Cracks in the oxide layer were observed on cerium above features found surrounding the sites at the surface, which exhibited preferred orientations and were likely to be from a compressive stress induced phase transformation. Cracks were also evident above a corona of deformed material in the plutonium.
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The anisotropic growth occurred because lateral growth rate was more rapid than the penetrative growth rate. Two significant factors appear to control the morphology and structure of the hydride product formed on both metals: the oxide provides a barrier to diffusion, which in addition to controlling the induction and nucleation stages, continues to prevent surface wide coverage from occurring favouring continued growth at existing hydride reaction sites; the stress field caused by the volume dilation when the metal becomes hydride affects the continued growth morphology of the hydride, disrupting the oxide surrounding the hydride reaction sites and providing a facile path for hydrogen to reach the metal, favouring lateral growth.
The situation for electro-refined plutonium is somewhat different than found for both γ-Ce and δ-Pu because the electro-refined plutonium did not form hydrides in the same manner as either of the other material types. The α-Pu material was more resistant to reaction with hydrogen requiring a heating cycle to 287 °C to initiate hydriding. Auto-reduction of the oxide from PuO2 to Pu2O3 would have occurred during the heat treatment and Pu2O3 does not act as a diffusion barrier. Consequently the growth of hydride was across the entire surface and within surface intersecting micro cracks.
Mechanisms for anisotropic hydride growth have been proposed as a result of this EngD which apply observations of cracks in the surface over distorted and phase-transformed material surrounding hydride reaction sites to explain the advantage of lateral growth over penetrative growth. This model for the anisotropic growth of hydrides on systems where the native oxide acts as a diffusion barrier requires the parent material to be ductile, hydride to have a significant decrease in density and a greater hardness to deform the surrounding parent metal. This deformation leads to a failure of the oxide diffusion barrier surrounding a hydride reaction site, which advantages the lateral growth rate.
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