The Pennsylvania State University

The Graduate School

College of Earth and Mineral Sciences

DEPOSITION, CHARACTERIZATION, AND THERMOMECHANICAL FATIGUE OF

NICKEL ALUMINIDE AND RUTHENIUM ALUMINIDE THIN FILMS

A Dissertation in

Materials Science and Engineering

by

Jane A. Howell

© 2010 Jane A. Howell

Submitted in Partial Fulfillment of the Requirements for the Degree of

Doctor of Philosophy

December 2010

The dissertation of Jane A. Howell was reviewed and approved* by the following:

Suzanne E. Mohney Professor of Materials Science and Engineering Dissertation Co-Advisor Co-Chair of Committee

Christopher L. Muhlstein Associate Professor of Materials Science and Engineering Dissertation Co-Advisor Co-Chair of Committee

Joseph R. Flemish Senior Scientist and Professor of Materials Science and Engineering

Clifford J. Lissenden Professor of Engineering Science and Mechanics

Joan M. Redwing Professor of Materials Science and Engineering Chair, Intercollege Graduate Degree Program in Materials Science and Engineering

*Signatures are on file in the Graduate School

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ABSTRACT

Intermetallic thin films have properties that make them attractive for applications such as metallizations, high temperature coatings, microelectromechanical systems, and diffusion barrier layers. In this study B2 aluminide films (NiAl and RuAl) have been deposited and characterized. Both intermetallics could be deposited at temperatures near room temperature using co- sputtering with an as-deposited resistivity of 45.5 ± 1.5 μΩcm for NiAl and 157 ± 4 μΩcm for RuAl. Ni/Al multilayers with a wavelength of 30 nm and below were fully reacted to form NiAl after annealing for 2 h at 400°C. These films had a resistivity of 15.5-26.7 μΩcm (wavelengths from 15.4-30 nm) after a 4 h anneal at 400°C, and the lower values of resistivity correspond to films with larger wavelengths. In order for Ru/Al multilayers to be fully reacted at 400°C, wavelengths of less than 10 nm were required as well as longer annealing times (more than 11 h). RuAl from the 400°C reaction of Ru/Al multilayers had a resistivity of 71.6-123 μΩcm. The lowest resistivity obtained for the B2 thin films was 11 ± 0.1 μΩcm, which was obtained for NiAl after a 20 min anneal at 800°C.

Both NiAl and RuAl have excellent bulk oxidation resistance, and in this study it was found that the oxidation resistance of the intermetallic thin films was superior to Al, Ni, and Ru metal films. The intermetallic films showed no observable surface changes (light microscopy) up to 500ºC in flowing oxygen and were conductive to higher oxidation temperatures (800°C for RuAl, 850ºC for NiAl) than Ni (500ºC), Al (600ºC), and Ru (800ºC, although vaporization may have begun at ~700ºC).

The intermetallics NiAl and RuAl along with Ru and Au have been patterned into thin line structures and then tested using an alternating current (100 Hz) to induce thermomechanical fatigue (TMF) with 200 thermal cycles per second. RuAl samples were able to withstand higher cyclic values of ΔT than NiAl for comparable times to failure. Both NiAl and RuAl were able to withstand higher values of ΔT than gold. The ΔT for tests on NiAl ranged from ~300-520°C (Tmax: 400-600°C), with a time to failure of 100’s of hours when ΔT was near 300°C (Tmax: ~400°C). Samples of NiAl that had a lower resistivity were able to withstand higher current densities due to reduced Joule heating. RuAl had a longer time to failure than NiAl at high testing temperatures (ΔT > ~350°C), but trends in the data indicate that the time to failure at lower temperature may be higher for NiAl. Curves of the resistance as a function of time were plotted for all of the samples during the course of the AC tests. These curves showed three distinct regions, which in the case of the intermetallic films can be ascribed to heating, the production of defects (dislocations and vacancies), and crack growth (and possibly voids). It was determined that by measuring the slope of the R(t) curve at the beginning of the test, the time to failure and the length of time spent in each of the three regions could be estimated. This will enable future tests to be paused at specific locations along the R(t) curves so that damage formation and crack growth may be studied.

The use of Ni- and Ru-aluminide films combined with 2 different etchants for the fabrication of MEMS has also been investigated. Using a conventional wet etch for SiO2 sacrificial layers (HF) led to cracking and/or buckling depending on the stress state in the film (tensile/compressive). However, using a gas phase etchant for silicon sacrificial layers, XeF2, free-standing regions could be formed that were crack free. Resonators were fabricated from co-sputtered NiAl and

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RuAl, and annealed multilayers of Ni/Al and Ru/Al, and first out-of-plane bending mode resonance was observed by using XeF2 etching. The best results were obtained for as-deposited NiAl that was co-sputtered at 1.5 mTorr and was under a compressive stress of ~0.83 GPa. While the devices were not completely flat, they were free-standing, and improvements are expected by decreasing the stress in the co-sputtered films by increasing the sputtering pressure. The results for RuAl resonators are also promising as the films can withstand high compressive stresses (~1.5 GPa, calculated from edge buckling), and improved performance from co- sputtered RuAl is expected by increasing the sputtering pressure in order to decrease the film stress. Thermal actuators fabricated from as-deposited co-sputtered NiAl showed ~18 m of motion with an applied current of 70 mA, and very little curvature/buckling was noted. The effect of the high compressive stress is less pronounced in the doubly supported actuators than in the singly supported resonators. Improvements in processing should lead to improvements in the device yield and curvature, but the results presented plus the high strength, high electrical conductivity, good oxidation/corrosion resistance, and moderate toughness make NiAl and RuAl suitable materials for MEMS.

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TABLE OF CONTENTS

LIST OF FIGURES ...... ix LIST OF TABLES ...... xxvi ACKNOWLEDGEMENTS ...... xxix

CHAPTER 1 ...... 1 INTRODUCTION AND BACKGROUND ...... 1 I. INTRODUCTION ...... 2 II. INTERMETALLIC COMPOUNDS ...... 3 A. Selection of Intermetallics ...... 4 III. OVERVIEW OF NiAl ...... 6 A. Mechanical Properties of NiAl ...... 7 B. Electrical Properties of NiAl ...... 11 IV. OVERVIEW OF RuAl ...... 13 A. Mechanical Properties of RuAl ...... 14 B. Electrical Properties of RuAl ...... 15 V. DISCUSSION ...... 16 REFERENCES ...... 17

CHAPTER 2 ...... 27 NICKEL ALUMINIDE FILM FABRICATION ...... 27 I. INTRODUCTION ...... 28 A. Ni-Al Films from the Reaction of Ni/Al Multilayers ...... 28 B. Ni-Al Films from Co-Deposition and Alloyed Sources/Targets ...... 31 II. EXPERIMENTAL PROCEDURE ...... 35 III. RESULTS AND DISCUSSION ...... 37 A. Ni-Al Films from the Reaction of Ni/Al Multilayers ...... 39 B. Co-Sputtered Ni-Al Films ...... 49 IV. CONCLUSIONS...... 57 REFERENCES ...... 59

CHAPTER 3 ...... 62 RUTHENIUM ALUMINIDE FILM FABRICATION ...... 62 I. INTRODUCTION ...... 63

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II. EXPERIMENTAL PROCEDURE ...... 63 III. RESULTS AND DISCUSSION ...... 64 A. Co-Sputtered RuAl...... 64 B. RuAl Formed from Annealed Ru/Al Multilayers ...... 68 IV. CONCLUSIONS...... 80 REFERENCES ...... 81

CHAPTER 4 ...... 82 OXIDATION OF RUAL AND NIAL THIN FILMS...... 82 I. INTRODUCTION ...... 83 II. EXPERIMENTAL PROCEDURE ...... 86 III. RESULTS AND DISCUSSION ...... 87 A. Oxide Surface Morphology and Depth Profiles...... 88 B. Electrical Characterization ...... 92 IV. CONCLUSIONS...... 97 REFERENCES ...... 99

CHAPTER 5 ...... 103 ALTERNATING CURRENT INDUCED THERMOMECHANICAL FATIGUE OF NIAL AND RUAL THIN LINES ON OXIDIZED SILICON ...... 103 I. INTRODUCTION ...... 104 A. Background: Alternating Current (AC) Thermomechanical Fatigue (TMF) Testing ...... 104 B. Temperature During AC TMF ...... 105 C. Literature Results: AC TMF Testing ...... 108 D. Comments on AC TMF Literature ...... 112 E. Comparison of AC and DC Electrical Tests ...... 115 F. Limitations to AC TMF Testing ...... 117 II. EXPERIMENTAL PROCEDURE ...... 118 A. Sample Preparation ...... 118 B. Sample Testing...... 120 C. Temperature Measurement ...... 122 1. IR camera measurements ...... 122 2. Peak resistance and thermocouple measurements ...... 126 3. Whole waveform test ...... 127

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4. Comparison of the three techniques ...... 128 D. Data Analysis ...... 130 III. RESULTS AND DISCUSSION ...... 131 A. Ruthenium ...... 131 B. NiAl...... 136 1. NiAl from annealed e-beam deposited 20 nm wavelength multilayers ...... 136 2. NiAl from annealed sputter deposited 30 nm wavelength multilayers: deposition 1 ...... 138 3. NiAl from annealed sputter deposited 30 nm wavelength multilayers: deposition 2 ...... 140 4. NiAl from annealed sputter deposited 30 nm wavelength multilayers: doped with 0.5% Ag ...... 144 5. Co-sputtered NiAl ...... 148 6. Summary of the NiAl data ...... 150 C. RuAl ...... 152 1. Annealed 25 nm wavelength Ru/Al multilayers ...... 152 2. Co-sputtered RuAl ...... 156 3. Summary of the RuAl TMF data ...... 160 D. Gold...... 161 E. Comparison of the TMF Results for the Different Materials Tested ...... 164 F. Analysis of the Resistance as a Function of Time Curves ...... 165 G. Comparison between AC and DC Tests ...... 175 IV. CONCLUSIONS...... 178 APPENDIX I: RESISTANCE AS A FUNCTION OF TEMPERATURE ...... 180 APPENDIX II: CURVE FIT RESULTS ...... 184 REFERENCES ...... 189

CHAPTER 6 ...... 192 NICKEL ALUMINIDE AND RUTHENIUM ALUMINIDE FILMS FOR MICROELECTROMECHANICAL SYSTEMS ...... 192 I. INTRODUCTION ...... 193 II. EXPERIMENTAL PROCEDURE ...... 194 III. DEPOSITION AND ANNEALING OF INTERMETALLIC STRUCTURAL FILMS ... 196 IV. PATTERNING AND RELEASE OF INTERMETALLIC STRUCTURAL FILMS ...... 198 A. Resistance of Blanket Films to HF ...... 199

B. Resistance of Patterned Films to XeF2...... 203 vii

V. ELECTROSTATICALLY AND THERMALLY ACTUATED MEMS ...... 206 VI. CONCLUSIONS...... 211 REFERENCES ...... 213

CHAPTER 7 ...... 217 CONCLUSIONS AND SUGGESTIONS FOR FUTURE WORK ...... 217 I. SUMMARY ...... 218

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LIST OF FIGURES

CHAPTER 1

Figure 1 – Effect of alloying additions on the resistivity of gold. Adapted from [9]...... 3

Figure 2 – (a) Ni-Al phase diagram. Adapted from [148]. (b) The B2 crystal structure of NiAl [147]...... 7

Figure 3 – The effect of temperature on the strain to failure for NiAl tested in tension. (a) Results from [149] for NiAl with a grain size of 84 μm. (b) Results from [151] for NiAl with 50.3% Al...... 8

Figure 4 – The effect of grain size on the tensile (a) elongation and (b) yield and fracture strength for NiAl with 49% Al tested at 400°C. Data from [18]...... 9

Figure 5 – The effect of temperature on the yield stress and tensile strength of NiAl tested in tension. (a) Data on NiAl with 50.3% Al [151]. (b) NiAl with a grain size of 84 μm [149]...... 9

Figure 6 – The effect of grain size on the properties of NiAl with 48.9% Al when tested in compression. Data from [17]...... 10

Figure 7 – The effect of stoichiometry on the yield stress for NiAl tested in compression. Data from [17]...... 10

Figure 8 – The effect of film thickness on (a) the fracture strength for Al-rich thin films and (b) the yield stress for Ni-rich thin films. Arrows indicate films that did not yield during testing. Films were in tension. Data from [19, 152]...... 11

Figure 9 – Summary of the literature data on the room temperature resistivity of bulk NiAl as a function of the Ni content. Data compiled from [43, 48, 153-154]...... 12

Figure 10 – Resistivity of NiAl thin films as a function of the film thickness. Data from [157]. 12

Figure 11 – Ruthenium-aluminum phase diagram adapted from [158]...... 13

Figure 12 – The yield stress as a function of temperature for RuAl samples with a grain size of ~150-200 μm and tested in compression. All samples contained a second phase, which was mostly located at the grain boundaries. Data from [161]...... 15

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Chapter 2

Figure 1 – Schematic showing a Ni/Al multilayer as-deposited. The first and last layers represent half of an aluminum layer. The diagram is representative of the multilayers deposited in this work, with the thickness of the individual layers varying depending on the composition and wavelength of the sample. The number of layers deposited was adjusted so that the total thickness was ~200 nm...... 29

Figure 2 – Figure plotting some of the literature data on the reaction of Ni/Al multilayers with compositions within the single phase B2 NiAl phase field. Curves are shown for the beginning of intermetallic formation, the beginning of the formation of NiAl, and the point at which the film is completely reacted to form NiAl. The data are curve fit with the equation y = a(1-bx). Data compiled from [3-4, 7, 9-12]...... 30

Figure 3 – Thin film phase diagram for Ni-Al films deposited by co-evaporation as a function of substrate temperature. Adapted from [20]. The dashed lines represent extrapolations of the experimental data...... 32

Figure 4 – Phases present as a function of substrate temperature in Ni-Al films evaporated simultaneously from Ni and Al sources by resistance heating. Data reproduced from [22]...... 33

Figure 5 – Variation in the as-deposited film resistivity as a function of the wavelength of the Ni/Al multilayer for sputtered and e-beam deposited samples. For clarity the error bars have been omitted, but in all cases lie within the symbols...... 39

Figure 6 – (a) Low and (b) high magnification TEM images of a 30 nm wavelength as-deposited Ni/Al multilayer film deposited by sputtering. In the images the light layers are aluminum, and the darker layers are nickel. The as-deposited film thickness was 179 ± 2 nm as determined from profilometry. The inset in (a) is the electron diffraction pattern of the film, and the scale bar in the diffraction pattern represents 2 nm-1...... 40

Figure 7 – (a) Bright and (b) dark field images of a 30 nm wavelength Ni/Al multilayer after annealing at 200ºC for 4 hours. The inset in (a) shows the electron diffraction pattern for the film and the scale bar represents 5 nm-1...... 41

Figure 8 – Grazing angle XRD scans of annealed Ni/Al multilayer films deposited by e-beam evaporation. The results shown are for a 25 nm wavelength Ni/Al annealed at (a) 400ºC for 1 min, (b) 400ºC for 5 min, and (c) 425ºC for 1 min, and (d) for a 33 nm wavelength Ni/Al annealed at 425ºC for 1 min. The peaks marked with a square could be either NiAl or Ni2Al3. 42

Figure 9 – Cross-sectional TEM images of a 30 nm wavelength sputtered Ni/Al film annealed at 400ºC for 4 hours. (a) A low magnification image of the film with the diffraction pattern shown in the inset (scale bar 5 nm-1), and (b) a higher magnification image of the film-substrate interface. The images are from the same deposition as those shown in Figure 6 for the as-

x deposited film. The thickness as determined from profilometry was 164 ± 1 nm after annealing...... 43

Figure 10 – Film thickness of a 25 nm wavelength multilayer as a function of annealing time and temperature...... 44

Figure 11 – The resistivity of a 25 nm wavelength Ni/Al multilayer film as a function of annealing time and temperature. For clarity the error bars have been omitted, but in all cases lie within the symbols...... 44

Figure 12 – Tapping mode AFM images of the surface of e-beam evaporated Ni/Al multilayer films. (a) An as-deposited 25 nm wavelength Ni/Al film, (b) an as-deposited 33 nm wavelength Ni/Al film, and (c) a 33 nm wavelength Ni/Al film annealed at 400ºC for 10 minutes...... 46

Figure 13 – Resistivity of the Ni/Al multilayer samples after annealing at 400°C. For clarity the error bars have been omitted, but in all cases lie within the symbols. (a) Resistivity as a function of the annealing time at 400°C for samples deposited by e-beam (solid symbols) and sputtering (open symbols). Numbers on the graph refer to the wavelength of the sample. (b) Resistivity as a function of the multilayer wavelength for Ni/Al films annealed at 400°C for 12 min (sputtered samples) or 10 min (e-beam samples)...... 47

Figure 14 – Results of high temperature anneals on two Ni/Al multilayer sputtered films. The wavelengths of the two films are (a) 15 nm and (b) 30 nm. For clarity the error bars have been omitted, but in all cases lie within the symbols...... 48

Figure 15 – Cross-sectional TEM images of a 30 nm wavelength sputtered Ni/Al film annealed for 1 minute at 600ºC followed by 30 sec at 1000ºC. This is a section of film deposited in the same run as the images shown in Figure 6 for the as-deposited case and in Figure 9 for a film annealed at 400ºC. This film corresponds to the blue diamond data point shown in Figure 14b.48

Figure 16 – Resistivity of the annealed Ni/Al multilayer films after annealing at high temperature. The annealing temperatures are indicated in the figures. (a) The lowest resistivity obtained after high temperature annealing as a function of the initial multilayer wavelength and (b) as a function of the resistivity after annealing at 400°C. For clarity the error bars have been omitted, but in all cases lie within the symbols...... 49

Figure 17 – Grazing angle XRD scans of Ni-Al films co-sputtered at 7 mTorr Ar pressure. (a) An as-deposited Ni-Al film with ~39.7% Ni, (b) an Ni-Al film with ~39.7% Ni annealed at 600ºC for 10 minutes, (c) a film with ~48.5% Ni annealed for 30 min at 400ºC, and (d) an ~27.0% Ni Ni-Al film annealed at 600ºC for 10 minutes...... 50

Figure 18 – Resistivity as a function of nickel composition for Ni-Al films co-sputtered at 7 mTorr Ar pressure. Also shown are the phases expected to be present according to the

xi equilibrium phase diagram. For clarity the error bars have been omitted, but in all cases lie within the symbols...... 51

Figure 19 – Tapping mode AFM images of a Ni-Al film co-sputtered at 7 mTorr Ar pressure. The film is ~39.7% Ni and is shown (a) as-deposited, and (b) after 90 minutes annealing at 400ºC...... 52

Figure 20 – (a) TEM image and (b) diffraction pattern of an as-deposited NiAl film co-sputtered at 1.5 mTorr Ar pressure...... 53

Figure 21 – Grazing angle XRD scans of 2 Ni-Al films co-sputtered at 1.5 mTorr Ar pressure with approximate Ni contents of (a) 32.7 at% and (b) 78.4 at%. Both films were annealed at 400ºC for 90 minutes. The solid and dashed vertical lines show the expected peak locations for Ni and Ni3Al respectively...... 54

Figure 22 – Resistivity data as a function of nickel content for Ni-Al films co-sputtered at 1.5 mtorr Ar pressure. Shown in the darker gray shading are the single phase fields according to the equilibrium phase diagram, and the lighter gray areas show regions that are expected to be 2 phase. For clarity the error bars have been omitted, but in all cases lie within the symbols...... 54

Figure 23 – Comparison of the resistivity after annealing at 600ºC for 1 min of Ni-Al films deposited at 7.0 and 1.5 mTorr. For clarity the error bars have been omitted, but in all cases lie within the symbols...... 55

Figure 24 – Two Ni-Al films deposited at 7.0 mtorr that cracked during annealing at 400ºC. The film in (a) is ~27% Ni and was annealed for 60 min, and the film in (b) is ~31.4% Ni and was annealed for 1 min. The darker regions are sections of the substrate, which is visible in areas where the film has peeled away...... 56

Figure 25 – Optical microscope images with Nomarski filtering of Ni-Al films co-sputtered at 1.5 mTorr Ar pressure and then annealed at 600ºC. Approximate nickel compositions for the films in atomic percent are: (a) 40.1%, (b) 38.2%, (c) 36.4%, (d) 34.6%, (e) 32.8%, (f) 30.0%, (g) 29.1%, (h) 27.3%, and (i) 23.6%...... 57

Chapter 3

Figure 1 – XRD results from two co-sputtered RuAl films deposited at (a) 1.5 mTorr and (b) the first two diffraction peaks from films deposited at 1.5 and 7.0 mTorr argon pressures. The film deposited at 1.5 mTorr is as-deposited and the 7.0 mTorr film is after annealing at 400ºC for 60 min...... 65

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Figure 2 – Cross-sectional TEM images (a, c, d) of an as-deposited co-sputtered RuAl film deposited at 1.5 mTorr Ar pressure. Also shown for comparison are (b) a cross-section of a Ru/Al multilayer and (e) an NiAl film co-sputtered at 1.5 mTorr. The inset in (b) shows the diffraction pattern of the co-sputtered film with the scale bar representing 5 nm-1...... 66

Figure 3 – Resistivity of RuAl films co-sputtered at 1.5 mTorr Ar pressure. Error bars for the annealed samples are smaller than the symbol size...... 67

Figure 4 – Resistivity of (a) co-sputtered RuAl films in comparison to (b) NiAl co-sputtered films. The NiAl films were annealed for 30 min at 400°C. Error bars are within the symbol size...... 68

Figure 5 - Cross-sectional TEM images of (a) an as-deposited Ru/Al multilayer with a wavelength of 25 nm, (b) the same Ru/Al film after annealing at 200°C for 4 h, (c) an as- deposited Ni/Al multilayer with a wavelength of 30 nm, and (d) the same Ni/Al film after annealing for 4 h at 200°C. The lighter layers are aluminum and the darker layers are ruthenium/nickel (except in the dark field image in d). The insets show the diffraction patterns of the films and the scale bar length is 5 nm-1...... 70

Figure 6 – Grazing angle XRD of multilayer Ru/Al films before and after annealing at 400ºC. The 25.0 and 15.4 nm films were annealed for 11 hours and the 9.6 nm film for 5 hours. Each data set is offset from the previous data set by 100 CPS...... 71

Figure 7 – Resistivity as a function of annealing time at 400ºC for a 25 nm wavelength Ru/Al multilayer film. The inset shows data from a 30 nm wavelength Ni/Al multilayer film also annealed at 400ºC (axes and units are the same in the main graph and the inset). Error bars are within the symbol size...... 72

Figure 8 – TEM image and diffraction pattern from a 25 nm wavelength Ru/Al multilayer film annealed at 400ºC for 193 hours. The inset in (a) shows a 30 nm wavelength Ni/Al multilayer after annealing for 4 h at 400ºC...... 73

Figure 9 – Arrangement of (a) the Ru/Al layers prior to annealing (273 ± 2 nm total thickness) and (b) the Ru(Al)/RuAl layers after annealing for 4 h at 400°C (261 ± 2 nm total thickness). .. 73

Figure 10 – Resistivity as a function of annealing time at 400ºC for Ru/Al multilayer films with wavelengths between 6.7 and 25.0 nm. Error bars are within the size of the symbol...... 75

Figure 11 – Auger depth profiles of a 15.4 nm Ru/Al multilayer (a) before and (b) after annealing for 11 h at 400ºC...... 76

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Figure 12 – Resistivity as a function of annealing temperature for a 15.4 nm wavelength Ru/Al multilayer. Each anneal was 30 min. Error bars are not shown, but in all cases lie within the size of the symbol...... 77

Figure 13 – Resistivity of the Ru/Al multilayers after annealing at temperatures between 800°C and 1100°C. In all cases, the standard deviation in the measurement of the resistivity was less than 0.2 μΩcm for 9 measurements...... 77

Figure 14 – (a) Cross-sectional TEM image of a 25 nm wavelength Ru/Al film annealed at 1000ºC for 2 min. (b) An electron diffraction pattern of the reaction layer shown in (a). The inset in (a) shows a 30 nm wavelength Ni/Al multilayer that was also annealed at 1000ºC...... 78

Figure 15 – A cross-sectional TEM image of a 25 nm wavelength Ru/Al film annealed at 1000ºC for 2 minutes, and the EDS spectra taken from one of the grains (left) and from the reaction layer (right)...... 79

Figure 16 – Comparison of the resistivity of multilayer and co-sputtered RuAl films after annealing at 400, 800, and 1100ºC. The Ru/Al multilayer had a wavelength of 25 nm. Error bars are not shown, but are smaller than the symbol size...... 80

Chapter 4

Figure 1 – Two thermal actuators fabricated from an as-deposited co-sputtered NiAl thin film. The Si substrate was patterned with conventional photolithography, and after deposition free- standing structures were created by etching the Si with XeF2 gas...... 84

Figure 2 – TEM micrographs of an (a) NiAl film fabricated from an annealed (400ºC for 4 h) Ni/Al multilayer with a 30 nm wavelength deposited by sputtering and (b) an as-deposited RuAl film deposited by co-sputtering...... 87

Figure 3 – Light microscopy images of the typical evolution of the surface morphology seen in NiAl films during annealing in oxygen. Figures (a-d) represent cumulative 1 h anneals from 100ºC to (a) 600ºC, (b) 700ºC, (c) 750ºC, and (d) 800ºC. The film imaged was fabricated from a sputter deposited Ni/Al multilayer with a wavelength of 30 nm...... 89

Figure 4 – Light micrographs of the surface of a NiAl film after oxidation. The film was annealed cumulatively for 1 h in oxygen from 100ºC to (a) 800ºC and (b) 850ºC. The film was fabricated from an e-beam evaporated Ni/Al multilayer with a wavelength of 25 nm...... 89

Figure 5 – Auger depth profiles of a NiAl film formed from a multilayer with a 33 nm wavelength Ni/Al (e-beam deposited) after oxidation to 850ºC. The depth profiles shown in (a)

xiv are for an area in between oxide particles and in (b) are for an area centered on an oxide particle. Locations of the 2 scan areas are shown in (c) where box 1 corresponds to (a) and box 2 is the scan area represented in (b)...... 90

Figure 6 – Comparison of the surface morphology of co-sputtered (a) RuAl and (b) NiAl thin films. Both were annealed in oxygen from 100ºC to 800ºC...... 91

Figure 7 – Auger depth profiles of a co-sputtered RuAl film after oxidation to (a) 800ºC and (b) 850ºC...... 92

Figure 8 – 4-pt probe measurements as a function of annealing temperature in oxygen for co- sputtered RuAl and NiAl films in comparison to Al, Ni, and Ru films. The data are presented as both the value of the normalized resistance after a given anneal, and as a percent resistance change compared to the sample prior to oxidation. In (a) and (b) the nearly vertical line for Ni represents a loss of conductivity (the film was conductive at 450°C, but not at 500°C). For clarity, the last data point for Al is not shown in (b) (Al at 600°C exhibits a 2×106 % increase in resistance)...... 93

Figure 9 – 4-pt probe measurements as a function of oxidation temperature for Ni-Al films. The data are presented as both the value of the normalized sheet resistance and as a percent resistance change compared to the sample prior to oxidation. The data for oxidation temperatures less than 400ºC are not shown, as there was no change in this region. Error bars are within the size of the symbols...... 94

Figure 10 – The resistance increase after oxidation to 850ºC as a function of the initial film thickness. The data for the multilayer samples (open squares) are plotted according to Eq. 3 (see text for details). The fit to the data yields an R2 value of 0.919...... 96

Figure 11 – Comparison of the film surface morphology after oxidation to 850ºC. The films were imaged using the same illumination conditions. The images show NiAl films formed from annealed Ni/Al multilayers with a wavelength of (a) 20 nm, (b) 25 nm, (c) 30 nm, and (d) a co- sputtered NiAl film. The upper right of each image shows the initial film thickness as well as the percent increase in resistance after oxidation to 850ºC...... 97

Chapter 5

Figure 1 – Schematics showing two possible test structures for AC testing. Most test structures used in the literature are a variation of one of the above lines. On the bottom figure, the current (or voltage) is applied at the two end contact pads and the voltage drop (or current) is measured across the two side contact pads...... 104

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Figure 2 – Schematic showing the method used to set-up an oscillating temperature and stress in a metal thin film sample. The exact position of the stress curve will depend on the initial stress state of the film and the difference in CTE between the film and substrate...... 104

Figure 3 – Schematics of the temperature oscillations of a metal line on a substrate heated by an AC signal with a low frequency (bottom) and a high frequency (top)...... 106

Figure 4 – The effects of the applied power and AC frequency on the temperature during AC thermal cycling for Cu lines 800 μm long, 8 μm wide, and 300 nm thick tested on oxidized silicon. (a) The minimum and maximum temperatures as a function of the peak applied power for a testing frequency of 100 Hz. (b) ΔT as a function of the AC frequency for a peak applied power of 4.5 W. Data from [3, 9]...... 107

Figure 5 – Time to failure at an AC of 100 Hz for Al-1Si and Cu lines on oxidized silicon substrates. Samples were 800 μm long, 3.3 (Al) or 2 (Cu) μm wide, and 0.5 (Al) or 0.7 (Cu) μm thick. Data from [1]...... 109

Figure 6 – Lifetime data for a set of 20 copper interconnects tested using an AC of 100 Hz and an RMS current density of 17 MA/cm2. The lines were fabricated using the dual Damascene approach and were 60 μm long, 250 nm thick, and 4.5 μm wide. Data from [14]...... 110

Figure 7 – Fatigue data obtained using an AC of 100 Hz for Cu lines of (a) varying thickness and grain size [7] and (b) of different line width [8]...... 111

Figure 8 – Thermal fatigue data of gold lines on quartz substrates tested with an AC of 50 Hz. The samples were 50 μm long, 2 μm wide, and 200 nm thick. The inset shows the same data plotted with a linear scale on the x-axis rather than a log scale (axes and units are the same in the main graph and the inset). Data from [19]...... 112

Figure 9 – Plot of the change in resistance as a function of the number of cycles for a sample of notched tool steel cycled with alternating deflection (bending of cantilever up and down). The arrow indicates the onset of cracking as determined from the point at which the resistance begins to more rapidly increase. Adapted from [26]...... 113

Figure 10 – (a) Graph of Johnson’s equation (Eq. 5) using the following parameters: a0 = 0 μm, W = 8 μm, and y = 150 μm. (b) Schematic diagram of the sample modeled in Eq. 5. For clarity the initial crack is shown as an ellipse rather than as a slit with infinitesimal height...... 114

Figure 11 – Comparison of AC and DC electrical tests on Al-2Si (a-d) and Cu (a). (a) MTTF as a function of frequency for Al-2Si and Cu at a peak current density of 1.5×107 A/cm2 and an ambient temperature of 275°C, (b) MTTF as a function of frequency for Al-2Si at a peak current density of 4.5×107 A/cm2 and an ambient temperature of 250°C, (c) MTTF for AC and DC tests on Al-2Si as a function of the peak current density, and (d) MTTF for AC and DC tests as a function of the ambient temperature for a constant applied current. Data from [29-30]...... 117

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Figure 12 – Geometry of the metal lines used to characterize the thermomechanical fatigue behavior of the thin films...... 118

Figure 13 – Layout of a group of six metal lines on a 6 mm chip. The 6 mm chip is affixed to a 40-pin package for wire bonding and then AC thermomechanical fatigue testing...... 119

Figure 14 – Section of a TMF sample attached to a 40 pin package by Al-1Si wire bonds. The square contact pads are 200 μm × 200 μm...... 120

Figure 15 – Front panel for the LabView program that recorded voltage data during TMF testing...... 121

Figure 16 – Schematic of the TMF setup shown with the 21x voltage divider circuit...... 122

Figure 17 – IR image of a NiAl sample at a hotplate temperature of 100°C prior to calibration of the sample emissivity. The square contact pads are 200 μm × 200 μm...... 123

Figure 18 – (a) Line profiles of the average temperature along the length of a ruthenium line taken after 5 minutes at each current. The profiles run the length of the line from the top current pad (0 μm) to the bottom current pad (800 μm), which is shown by the white arrow in (b). The images in (b), (c), and (d) correspond to the IR camera images prior to testing, and after 5 minutes at 32 mA and 44 mA, respectively. The temperature scale in the IR images refers to the temperature of the silicon substrate...... 124

Figure 19 – (a) Line profiles across the width of a NiAl line as a function of the applied current after 5 minutes at each current. The location of the line profile is shown as the white line in (b). The peak values were obtained using the emissivity of the metal line, whereas the rest of the data were obtained using the emissivity of the silicon substrate. The images correspond to IR images taken a) prior to testing, (b) after 5 minutes at 34 mA, and (c) after 5 minutes at 61 mA. The temperature scales in (b-d) are for the silicon substrate...... 125

Figure 20 – Minimum, maximum, and average line temperature as a function of the applied current for a NiAl sample. Data obtained from the line profiles shown in Figure 19, where Tmin is estimated from the substrate temperature, Taverage is the peak in the line profiles, and Tmax is estimated by: - . The lines are curve fits using the function ...... 126

Figure 21 – Temperature as a function of time (left) and current (right) obtained during a step test on a sample of ruthenium sputtered at 1.5 mTorr. The inset to the figure on the left shows the current as a function of time during the step test (x-axis is the same in the main graph and the inset). The curve fits to the temperature as a function of current data use the equation: ...... 127

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Figure 22 – Voltage and current as a function of time acquired during a whole waveform test on a sample of co-sputtered RuAl...... 127

Figure 23 – (a) 5 thermal cycles acquired during the whole waveform test and the curve fit to the data, and (b) the maximum and minimum temperature as a function of time obtained from curve fits similar to those shown in (a). The solid and dotted lines in (b) are curve fits to the data using the equation T=atb, where a and b are fit parameters...... 128

Figure 24 – Temperature as a function of the applied current for a NiAl sample, where the temperature was measured using resistance, thermocouple, and IR measurements. The inset shows an enlarged view of the Tmin measurements (axes and units are the same in the main graph and the inset)...... 129

Figure 25 – The ΔT (Tmax-Tmin) as a function of the peak temperature for NiAl and Ru determined using peak resistance and thermocouple measurements, and for RuAl using the whole waveform test...... 129

Figure 26 – (a) Resistivity as a function of temperature calibrations performed prior to testing and after TMF testing to a 15% increase in resistance. (b) The error in temperature obtained after a 15% increase in resistance has been observed if the initial resistivity as a function of temperature curve is used to calculate the temperature. The above tests were performed on a Ni/Al multilayer sample annealed at 800°C. Error bars are within the size of the symbols. .... 130

Figure 27 – Resistance as a function of time for ruthenium sputtered at (a) 1.5 mTorr and (b) 7 mTorr...... 133

Figure 28 – Comparison of the lifetime data for the two sets of ruthenium samples as a function of (a) current density and (b) the initial peak temperature during AC thermal fatigue testing... 133

Figure 29 – SEM images of a ruthenium line deposited at 1.5 mTorr and tested at 66.4 mA (10.64 MA/cm2) until failure, which occurred after 1,728,000 thermal cycles. (a) Low magnification image of the region where failure occurred. (b) High magnification image of the bottom of the failure region...... 134

Figure 30 – SEM images of a ruthenium line deposited at 7 mTorr and tested at 56.9 mA (7.08 MA/cm2) until failure, which occurred after 21,960,000 thermal cycles. (a) and (c) show high magnification images of regions near the bottom and top of the failure, respectively. (b) Low magnification image of the region where failure occurred...... 135

Figure 31 – SEM images of a ruthenium line deposited at 7 mTorr and tested at 58.7 mA (7.30 MA/cm2) until failure, which occurred after 9,792,000 thermal cycles. (a) Low magnification image of the region where failure occurred. (b) A high magnification image of the bottom side of the failure region...... 135

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Figure 32 – (a) Relative resistance as a function of time for the e-beam multilayer NiAl. The inset shows the same data, but zoomed in to highlight tests with shorter lifetimes (axes and units are the same in the main graph and the inset). (b) Lifetime plot as a function of current density. Temperature data are included in the inset...... 137

Figure 33 – SEM images of a failed NiAl TMF sample tested at 76.0 mA (8.60 MA/cm2). Failure occurred after 756,000 thermal cycles. The sample was annealed at 400°C followed by a 700°C anneal. (a) Low magnification image of the entire sample, and (b) a higher magnification image of the failure location...... 138

Figure 34 – (a) Relative resistance as a function of AC testing time for annealed 30 nm wavelength Ni/Al multilayers with the low temperature and high temperature anneal. The inset shows the same data, but zoomed in to highlight the tests with shorter times to failure (axes and units are the same in the main graph and the inset). (b) SEM image of a region near the failure site for the sample annealed at low temperature. (c) and (d) SEM images of 2 of the samples annealed at high temperature after TMF testing...... 140

Figure 35 – Lifetime data plotted as a function of (a) current density and (b) temperature for the samples from anneals 1-4...... 142

Figure 36 – (a) Relative resistance as a function of testing time (the inset shows a zoomed in view of the data at short times) and (b) the time necessary to reach a given amount of resistance change as a function of the peak temperature at the start of the test. Data are for samples from anneal 1...... 142

Figure 37 – SEM image of the failure region in a sample from anneal 1...... 143

Figure 38 – SEM images of failed samples from (a) anneal 4, (b) anneal 2, and (c) anneal 4, where the top image is an enlarged region to the right of the failure and the bottom image shows the failure location...... 144

Figure 39 – Nomarski light microscope images of the current pads of annealed Ni/Al multilayer TMF samples (a) un-doped and (b) doped with 0.5% Ag. The samples were annealed at 1000°C for 30 s...... 145

Figure 40 – Results from the NiAl-0.5Ag TMF tests. The numbers within the circled regions refer to the line width of the samples in μm. (a) The time to failure as a function of the current density. The data plotted as a function of temperature is shown in the inset. The data from a 3.3 μm wide un-doped NiAl sample is also included. (b) The temperature during TMF testing as a function of the applied current density. The circled data point in the 3.8 μm region corresponds to one of the 4.0 μm samples...... 147

Figure 41 – (a) Relative resistance as a function of testing time for 3 samples of NiAl-0.5Ag. (b- d) SEM images of samples tested until failure. The insets show the resistance (in Ω) as a

xix function of time (in h) for the samples in the images. Ti in (a-d) refers to the peak temperature measured at the start of the test...... 148

Figure 42 – (a) Relative resistance as a function of time curves for co-sputtered NiAl of varying line width, and (b) the lifetime data plotted against the applied current density with the inset showing the lifetime as a function of temperature (x-axis is the same in the main graph and the inset)...... 150

Figure 43 – SEM images of samples of co-sputtered NiAl with line widths of (a) 3.3 μm and (b) 5.2 μm that were cycled until failure...... 150

Figure 44 – Summary of the NiAl TMF data. The numbers in the plots refer to individual sample sets and are ranked according to the average resistivity of the sample set. Samples 1-6 have an average resistivity in the range of 20-24 μΩcm and samples 7-12 have an average resistivity of 38-49 μΩcm. The open symbols (7-9) are for the co-sputtered samples, and the closed symbols are for the annealed multilayer samples. Arrows indicate sample that did not fail. (a) Time to failure plotted against the initial peak temperature (left) and ΔT (right). The solid line is a linear fit to the entire data set plotted as log(t) vs. T. (b) Time to failure plotted as a function of the applied current density...... 151

Figure 45 – The current density resulting in failure in 10 h as a function of the average resistivity. The numbers in the plot refer to the same numbers as are seen in Figure 44...... 152

Figure 46 – Time to failure for the annealed Ru/Al multilayers as a function of (a) current density and (b) temperature. The solid line is a fit to anneals 1, 2, and 4, and the dashed line is a fit to the data from anneal 3...... 154

Figure 47 – Results from anneal 1. (a) Resistance as a function of time (the inset shows a zoomed in view of the data at short times) and (b) the time to a given increase in resistance as a function of the peak temperature at the beginning of the test...... 155

Figure 48 – Resistance as a function of time curves for samples from anneal 3...... 155

Figure 49 – SEM images of the annealed Ru/Al multilayers from (a) anneal 1 (both sides of the failure), (b) anneal 2, (c) anneal 3 (2 different magnifications), and (d) anneal 4...... 156

Figure 50 – Results from the repeatability tests on RuAl co-sputtered at 1.5 mTorr and annealed at 1000°C for 1 min. (a) Plot of the failure probability for the four different currents tested and (b) the time to failure as a function of the testing current for failure probabilities of 25%, 50%, and 75%...... 157

Figure 51 – Select images from the failure region (a, b, c, and e) and from a region far from the failure (d) for co-sputtered RuAl annealed at 1000°C for 1 min and then thermally cycled until

xx failure. The inset in (a) shows an enlarged view of the region close to the failure. The boxed region in (d) highlights the damaged patch...... 158

Figure 52 – Close-up views of the damaged patches observed in RuAl co-sputtered samples tested at (a) 89.4 mA (0.2 h to failure) and (b) 89.5 mA (0.117 h to failure)...... 159

Figure 53 – Relative resistance as a function of testing time for the 6 RuAl co-sputtered samples tested at 82.0-82.1 mA that failed between the voltage pads. For clarity, the curves in (a) are re- plotted in (b) with an offset of 0.016 in the y-direction between each curve...... 160

Figure 54 – Results from all of the co-sputtered RuAl samples. (a) Time to failure as a function of the current density and (b) the time to failure as a function of the peak temperature...... 160

Figure 55 – Summary of the results from TMF tests on RuAl. (a) Time to failure as a function of the peak temperature (left) and the ΔT (right) experienced during testing and (b) time to failure as a function of the AC current density. The lines in (a) are linear fits to the data obtained from plots of log(t) vs. T...... 161

Figure 56 – Resistance as a function of time curves for gold sputtered at (a) 4 mTorr and (b) 7 mTorr Ar pressure. The sharp changes in resistance as indicated by the regions between the arrows shown in (b) are due to a slightly fluctuating current...... 162

Figure 57 – ESEM images of gold TMF samples cycled until failure. (a) Sample deposited at 4 mTorr and tested at 137.2 mA (11.53 MA/cm2). Failure occurred after ~37,008,000 thermal cycles. (b) Sample deposited at 7 mTorr Ar pressure and tested at 135.9 mA (15.74 MA/cm2). Failure occurred after ~12,600,000 thermal cycles. (c) The same sample shown in (b) showing a larger region of the sample...... 163

Figure 58 – AFM image of a blanket gold film sputtered at 1.5 mTorr Ar pressure...... 163

Figure 59 – Summary of the results from TMF testing in terms of the peak temperature and ΔT experienced during thermal cycling. (a) Plot of Tmax (and ΔT) versus the time to failure and (b) a schematic of the same data showing approximate boundaries to the various data sets. RuAl* refers to the annealed multilayer from anneal 3...... 164

Figure 60 – Summary of the current density results from TMF tests of the different materials tested. The table shows details of the samples represented in the plot. The data points circled in the plot are for tests that were stopped prior to failure (indicated by the arrow)...... 165

Figure 61 – Schematics showing the 3 different types of curves obtained during TMF testing, the equation used to fit the data, and an example of a curve fit to an experimental data set...... 166

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Figure 62 – An example of the procedure used to determine the transitions between regions I, II, and III. (a) The entire data set and (b) the region encompassed by the lines in (a). The thick red line shows the location of the center of curvature for this dataset. The dashed black lines show the last point where the r2 value is greater than 0.999 when the size of the curve fit is increased by 10 data points in each direction. The solid black lines show the last point where r2 is greater than 0.999 when the curve fits are increased by 5 data points in the negative x-direction and 14 data points in the positive x-direction...... 167

Figure 63 – Example data sets and curve fits showing the three different trends noted in the R(t) curves. Also shown is the equation used to obtain the curve fits as well as lines showing the obtained transition points between the three regions...... 168

Figure 64 – Results from the analysis of curve fits to the NiAl and RuAl data showing three regions of increasing resistivity. (a) The time in region I as a function of the maximum test temperature, (b) the time in region I as a function of the time to failure, (c) the slope in region I as a function of the maximum test temperature, and (d) the slope in region I as a function of the time to failure. The insets in (a) and (c) show the time in region I and the slope in region I as a function of the current density, respectively (the y-axes have the same units in the insets and the main graphs)...... 170

Figure 65 – Results from the analysis of curve fits to the NiAl and RuAl data showing three regions of increasing resistivity. (a) The time in region II as a function of the maximum test temperature, (b) the time in region II as a function of the time to failure, (c) the slope in region II as a function of the maximum test temperature, and (d) the slope in region II as a function of the time to failure. The insets in (a) and (c) show the time in region II and the slope in region II as a function of the current density, respectively (the y-axes have the same units in the insets and the main graphs)...... 171

Figure 66 – Results from the analysis of curve fits to the NiAl and RuAl data showing three regions of increasing resistivity. (a) The time in region III as a function of the maximum test temperature, (b) the time in region III as a function of the time to failure, (c) the slope in region III as a function of the maximum test temperature, and (d) the slope in region III as a function of the time to failure. The insets in (a) and (c) show the time in region III and the slope in region III as a function of the current density, respectively (the y-axes have the same units in the insets and the main graphs)...... 172

Figure 67 – Figure showing the two different values of the resistance due to crack growth ( and - ), which were used to calculate the crack length using Johnson’s equation...... 173

Figure 68 – Analysis of the curve fits to all of the NiAl and RuAl data showing 3 regions of increasing resistance. (a) The relationship between the slope in regions I and II, and (b) the relationship between the slopes in region I and III. The solid lines are guides to the eye and represent the case where the slopes are equal...... 174

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Figure 69 – The relationships between the slope in region I and (a) the time in region I and (b) between the slope in region I and the time in region II...... 174

Figure 70 – Comparison between AC and DC tests conducted on RuAl co-sputtered at 1.5 mTorr and then annealed at 1000°C for 1 min. (a) An example of a DC R(t) curve compared to an AC curve with a similar time to failure, and (b) the increase in resistance observed during electrical testing as a function of the time to failure...... 175

Figure 71 – SEM images of 4 co-sputtered RuAl samples tested under DC conditions until sample failure...... 177

Figure 72 – SEM images of RuAl TMF samples in regions away from the failure location for samples tested until failure using (a) AC and (b) DC stressing conditions...... 178

Figure 73 – Comparison of the time to failure under AC and DC stressing conditions for co- sputtered RuAl...... 178

Figure A1 – Comparison of two different methods of determining the TCR for a set of samples. The graph on the left shows the resistance as a function of temperature for 5 different TMF lines labeled lines 2-5. The line number in the graph corresponds to the number of rectangles located either directly above or below the TMF line on the 6 mm chip (schematic shown on the right). The curve fit to the data from all 5 TMF lines is shown as the thick black line in the plot, with the corresponding equation given in the upper right of the graph. Equations obtained for linear fits to each TMF line are shown adjacent to the corresponding line in the schematic diagram to the right. Taking an average of the slopes from the 5 linear fits results in a value of 0.1931, which is the same as the value obtained for the single curve fit. The data in this figure are the same as anneal 4 shown in the left hand graph of Figure A4...... 180

Figure A2 – Resistance as a function of temperature for (a) ruthenium samples deposited at 7 mTorr and (b) ruthenium samples deposited at 1.5 mTorr...... 181

Figure A3 – Resistance as a function of temperature for (a) annealed Ni/Al multilayers with a wavelength of 20 nm deposited by e-beam and (b) annealed Ni/Al multilayers with a wavelength of 30 nm (deposition 1) deposited by sputtering...... 181

Figure A4 – Resistance as a function of temperature for (a) annealed Ni/Al multilayers with a wavelength of 30 nm (deposition 2) deposited by sputtering and (b) annealed Ni/Al multilayers with a wavelength of 30 nm, doped with 0.5% Ag, and deposited by sputtering...... 182

Figure A5 – Resistance as a function of temperature for (a) co-sputtered NiAl and (b) annealed Ru/Al multilayers with a wavelength of 25 nm deposited by sputtering...... 182

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Figure A6 – Resistance as a function of temperature for the co-sputtered RuAl samples tested with (a-b) AC and (b) DC conditions. The table presents the results from the curve fits...... 183

Figure A7 – Resistance as a function of temperature for gold deposited at (a) 7 mTorr and (b) 4 mTorr...... 183

Chapter 6

Figure 1 – Cross-sectional TEM images of films formed from annealed multilayers and co- sputtering. (a) Ni/Al 30 nm wavelength multilayer deposited by sputtering and annealed at 400ºC for 4 h. (b) Ru/Al 25 nm wavelength multilayer deposited by sputtering and annealed at 400ºC for 193 h. (c) As-deposited NiAl co-sputtered at 1.5 mTorr. Indexed diffraction patterns are shown in the insets of the figures (blue-RuAl, yellow-Ru, white-NiAl), and the scale bar in each is 5 nm-1. The diffraction patterns in (a) and (c) only have diffraction rings corresponding to NiAl (white), whereas many of the rings in (b) can be indexed as either RuAl (first number- blue) or Ru (second number-yellow). Rings that can be clearly identified as RuAl are the (100), (111), (200), and (310) planes, whereas there is only 1 ring that can be clearly identified as belonging solely to Ru, i.e. the (105) plane...... 198

Figure 2 – Light microscope image of a NiAl film immersed in HF for 10 min. The film was deposited as a Ni/Al multilayer with a 25 nm wavelength using e-beam evaporation. After deposition the film was annealed in an RTA for 10 min at 400ºC, cooled to 350ºC and held for 2 min, cooled to 300ºC and held for 2 min, and then cooled to 40ºC...... 200

Figure 3 – NiAl film co-sputtered at 1.5 mTorr Ar pressure and then exposed to HF for 10 min. The images represent different sections of the same sample: (a) as-deposited material, (b) higher magnification of the as-deposited NiAl, (c) annealed at 400ºC for 10 min, and (d) annealed at 600ºC for 1 min. In (c) and (d), the lighter gray represents areas of exposed substrate...... 201

Figure 4 – (a) Ru/Al multilayer film annealed at 400ºC and then immersed in HF for 10 min. The film is a 6.7 nm wavelength multilayer annealed for 29 h. RuAl films co-sputtered at 1.5 mTorr Ar pressure and then exposed to HF for 10 min: (b) as-deposited film, (c) film annealed at 400ºC for 30 min, and (d) high magnification image of the film annealed at 600ºC for 1min. . 202

Figure 5 – Patterned co-sputtered NiAl and RuAl films on silicon after exposure to XeF2. The SEM images are for (a) 1.5 mTorr NiAl as-deposited, (b) 7.0 mTorr NiAl as-deposited, (c) as- deposited RuAl co-sputtered at 1.5 mTorr Ar pressure, and (d) RuAl deposited at 1.5 mTorr and annealed for 4 h at 400ºC...... 204

Figure 6 – Annealed Ru/Al multilayer films deposited by sputtering and then exposed to XeF2. Both samples were annealed for 11 h at 400ºC. The images shown are for (a) a 9.6 nm wavelength film and (b) a 25 nm wavelength film...... 205

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Figure 7 – Surface profile obtained from optical interferometry of an as-deposited RuAl film co- sputtered at 1.5 mTorr and then etched with XeF2. The stress in the film is calculated as ~1.5 GPa compressive...... 206

Figure 8 – Representative light microscope images of co-sputtered NiAl (a-d) and Ni/Al multilayer (e-g) resonators in the as-deposited condition (a) and after annealing for 4 h at 100ºC (b and e), 200ºC (c and f), and 400ºC (d and g). The co-sputtered resonators annealed at 400ºC were not etched as the resonators cracked and peeled after annealing. The scale bar in (a) applies to all of the images...... 207

Figure 9 – Nomarski optical images of NiAl resonators from 2 different depositions conducted at 1.5 mTorr. Both samples are as-deposited, but the film in (a) was deposited from 2 sequential depositions totaling 34 min and (b) was deposited from 3 sequential depositions for a total of 33.5 min. The samples were kept in vacuum for ~30 min in between each sequential deposition...... 208

Figure 10 – Representative images of resonators fabricated from co-sputtered RuAl (a-c) and Ru/Al multilayers (d-f). The resonators are shown in the as-deposited condition (a), after annealing for 4 h at 200ºC (d) and 400ºC (b and e), and after annealing for 1 min at 600ºC (c and f). The scale bar in (a) applies to all of the images...... 210

Figure 11 – Light microscope images showing the range of motion of a thermal actuator as the current is ramped down from 70 mA to 0 mA. The actuator was fabricated from an as-deposited NiAl film that was co-sputtered with an Ar pressure of 1.5 mTorr...... 211

xxv

LIST OF TABLES

CHAPTER 1

Table 1 – Comparison of the strength of some alloyed intermetallics to titanium and nickel alloys. Data from [20]...... 4

Table 2 – Compilation of property data for intermetallics having a resistivity lower than ~35 μΩcm. Data for several pure metals are included as a comparison. Hardness data for the pure metals are for the metals after annealing...... 5

Table 3 – Stress -strain data for RuAl alloys in compression [50, 164]...... 15

Chapter 2

Table 1 – Summary of some of the data on Ni/Al multilayer films with compositions in the B2 NiAl phase field. The temperatures should be considered an approximation as some were taken at the beginning of DSC exotherms, while some were taken at the end of the DSC exotherm, and others are from XRD data...... 29

Table 2 – Data on some of the Ni-Al films deposited and studied. The “Best Annealing Conditions” are those that resulted in the lowest resistivity...... 38

Table 3 – Indexed rings from the diffraction pattern in Figure 20...... 53

Chapter 3

Table 1 – Measured d-spacings from the electron diffraction pattern shown in Figure 2c compared to the expected values for stoichiometric RuAl...... 67

Chapter 5

Table 1 – Details of the ruthenium samples tested and the testing conditions...... 132

Table 2 – Details on the multilayers deposited and tested. NF refers to samples that did not fail...... 136

xxvi

Table 3 – Sample details, testing conditions, and results for 30 nm wavelength annealed Ni/Al multilayers...... 138

Table 4 – Sample details for the second set of 30 nm wavelength Ni/Al multilayers, the testing conditions, and TMF test results. All heating ramp rates were 50°C/s...... 141

Table 5 – Details of the NiAl-0.5Ag samples tested and the testing results...... 145

Table 6 – Overview of the co-sputtered NiAl samples and the TMF results...... 149

Table 7 – Sample details and results for annealed Ru/Al 25 nm wavelength multilayers tested using AC TMF. NF refers to a sample that did not fail in the time indicated...... 153

Table 8 – Details of the gold samples, testing conditions, and results of thermal fatigue testing...... 161

Table A1 – Results of the curve fits for NiAl samples fabricated from annealed Ni/Al multilayers with a wavelength of 20 nm deposited by electron-beam evaporation...... 184

Table A2 – Results of the curve fits for NiAl samples fabricated from annealed Ni/Al multilayers with a wavelength of 30 nm deposited by sputtering – deposition 1...... 184

Table A3 – Results of the curve fits for NiAl samples fabricated from annealed Ni/Al multilayers with a wavelength of 30 nm deposited by sputtering – deposition 2...... 184

Table A4 – Results of the curve fits for NiAl samples doped with 0.5% Ag fabricated from annealed Ni/Al multilayers with a wavelength of 30 nm deposited by sputtering. Samples were annealed at 400°C for 4 h...... 185

Table A5 – Results of the curve fits for co-sputtered NiAl films annealed at 400°C for 4 h. .... 185

Table A6 – Results of the curve fits for RuAl fabricated from annealed Ru/Al multilayers with a wavelength of 25 nm...... 186

Table A7 – Results of the curve fits for co-sputtered RuAl...... 187

Table A8 – Results of the curve fits to ruthenium samples deposited by sputtering and annealed at 400°C for 4 h...... 188

xxvii

Chapter 6

Table 1 – Resistivity of Ni-Al and Ru-Al Films after Annealing at 400ºC for up to 193 h...... 197

Table 2 - Summary of the results from the ruthenium and nickel aluminide resonator fabrication. All of the anneals were for 4 h except the 600ºC anneal, which was for 1 min...... 211

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ACKNOWLEDGEMENTS

Completion of this project would not have been possible without the help from members of the Penn State Materials Science and Engineering Department and from Northrop Grumman. First, I would like to thank my advisors, Suzanne Mohney and Chris Muhlstein, for the opportunity to work on this project and for their guidance and support. I would like to thank the faculty, staff, and students in the department of Materials Science and Engineering for all their help and support. I am especially grateful for assistance from R. Kirkpatrick for etching, E. Lysczek, B. Dormaier, and B. Downey for photolithography, Q. Zhang for TEM, B. Liu for Auger depth profiles, and A. Romasco for optical interferometry. I also need to thank all the past and present members of the Mohney and Muhlstein research groups for all of their help and support in the completion of this project. I would also like to thank the Penn State Electro-Optics Center for use of their IR camera.

I would especially like to thank several members of Northrop Grumman for their, assistance and guidance throughout the course of the project: Jack Hawkins, Rob Young, Silai "Krish" Krishnaswamy, and Marc Sherwin. Parts of this research were also facilitated by the Pennsylvania State University Materials Research Institute Nanofabrication Lab and the National Science Foundation Cooperative Agreement No. 0335765, National Nanotechnology Infrastructure Network, with Cornell University.

I would also like to thank all of my friends, especially those in the Materials Science Department, for making my time at Penn State enjoyable. Finally, I would like to thank my parents, Paul and Wendy Howell, for their love, support, understanding, and patience throughout my time at Penn State.

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CHAPTER 1

INTRODUCTION AND BACKGROUND

1

I. INTRODUCTION

Thin films are increasingly important due to their vital roles in electronics, micro- electromechanical systems (MEMS), optics, and as hard coating materials. The properties required for these applications are wide-ranging: high strength, oxidation resistance, wear resistance, stability at elevated temperatures, fatigue resistance, and high electrical conductivity. A class of materials that can offer this array of properties is intermetallics, which are compounds composed of two or more metals, or occasionally metals plus a metalloid or the nonmetals C and N. Intermetallic compounds typically have long range ordered crystal structures that are different from the constituent elements, and have properties that place them between metals and ceramics. Many intermetallics are already used in a variety of applications and have the potential for use in many more. For example, transition metal silicides are currently used in microelectronics [1-3]. Nitride intermetallic films such as ZrN and TiN are used for a number of different applications such as hard coatings on machine parts and cutting tools, diffusion barriers, electrical contacts, optical coatings for heat mirrors, and decorative coatings [4-6]. Intermetallic diboride thin films are of interest for diffusion barriers in electronics [7], and TiB2 has been used as a coating for cutting tools and is of interest as an electrode material for ultra-high temperatures [8]. While there are already a number of applications utilizing thin film intermetallics, a larger knowledge base is necessary so that their full potential can be realized.

For a thin film application requiring high electrical conductivity, a natural choice of material would be one of several pure metals (for example, gold, aluminum, or copper). When the application involves a film on a substrate exposed to conditions of fluctuating temperature, the material selection process can be complex. In such a case, the difference in thermal expansion between the metal and substrate will lead to stresses that may plastically deform the metal. In the event that these stresses vary with time, damage can progressively accumulate leading to failure of the film by processes generally termed “fatigue”. When there are both stresses and temperature that vary with time, the process is termed thermomechanical fatigue.

As a general rule, materials with a higher yield (and ultimate) strength are more resistant to fatigue degradation. Materials that have a higher melting point maintain their strength more effectively at elevated temperatures and are more resistant to the coupled damage accumulation mechanisms found in thermomechanical fatigue. To increase the strength of metal films, alloying elements can be added, which impede the motion of dislocations, but this often leads to an undesirable increase in the resistivity of the metal due to scattering of the electrons by the impurity (see Figure 1). An alternative to alloying the metal is to use an intermetallic compound, many of which have a low resistivity due to the long range ordered structure, in place of the pure metal. When selecting an intermetallic for such an application, two criteria can be used: closely matching the coefficient of thermal expansion to the substrate to minimize stresses during thermal cycling, or using a high strength material that can withstand the thermal stresses without failure. Both of these strategies are to be employed in the current research, but it is expected that even with large differences in thermal expansion, intermetallic films should have better thermomechanical fatigue (TMF) resistance than pure metals.

2

25 Gallium Tin Vanadium 20

cm) 15

10 Resistivity ( Resistivity 5

0 0 2 4 6 8 10 Atomic Percent Alloying Addition Figure 1 – Effect of alloying additions on the resistivity of gold. Adapted from [9].

The first phase of this research involved an extensive literature review, and based on a number of criteria, intermetallics believed to be suitable for conditions of fluctuating temperature, such as those experienced in many electronic devices and components, were chosen. Once materials of interest have been identified, ideal low temperature processing conditions that minimize the thin film resistivity will be determined, as well as the oxidation resistance of the films and the thermomechanical fatigue resistance. The TMF resistance of the thin film intermetallics will be evaluated by passing an alternating current through a thin film line patterned on a substrate, a relatively new technique, which has been used mainly on copper and aluminum [10-15]. The final phase of this research will be an evaluation of the applicability of the intermetallics in MEMS devices.

II. INTERMETALLIC COMPOUNDS

Intermetallic compounds are materials composed of two or more metals that are bonded together with a certain stoichiometry. They typically have long-range ordered crystal structures and often exist within relatively narrow composition ranges. Intermetallics are considered to be a class of materials between metals and ceramics as they have properties belonging to both classes of materials: electrical and thermal conductivity similar to metals and the high strength and low fracture toughness generally associated with ceramics.

There are more than 25,000 known intermetallic compounds [16] with widely varying properties, but there are some characteristics shared by the majority of intermetallic compounds. Intermetallics tend to have high strength at room and elevated temperatures and high wear resistance. They have high oxidation and corrosion resistance, are stable at elevated temperatures, and have low densities. The above combination of properties has made 3 intermetallics an attractive candidate for high temperature structural materials, but their use has been limited by their extreme brittleness. Of the tens of thousands of known intermetallics, only a few dozen have been shown to possess any room temperature ductility. The brittleness of intermetallics at room temperature is due to the small number of slip systems available, which is related to the ordered structure and maintaining this ordered structure during deformation. Thin film intermetallic applications may be able to circumvent the problem of brittleness as the ductility of intermetallics often increases with decreasing grain size [17-18] and thin films tend to have a very small grain size. Another benefit to thin film intermetallics is an increased yield and fracture strength with decreasing film thickness [19]. Due to this phenomena intermetallics in thin film form may be able to withstand larger stresses than their bulk counterparts prior to failure.

In some cases, alloying has led to improvements in ductility and fracture toughness of bulk intermetallics. Table 1 compares the strengths of intermetallic alloys with titanium alloys and nickel-base superalloys. The strength of the intermetallic alloys can exceed those of the superalloys, but even with alloying the fracture toughness and ductility of the intermetallics are significantly lower.

Table 1 – Comparison of the strength of some alloyed intermetallics to titanium and nickel alloys. Data from [20]. Ni-base Ti-alloys Ti Al-alloys TiAl-alloys 3 superalloys Density (g/cm3) 4.5 4.1-4.7 3.7-3.9 4.5 Yield Strength (MPa) 380-1150 700-990 400-650 250-1310 Tensile Strength (MPa) 480-1200 800-1140 450-800 620-1620 Strain to Fracture (%) 10-25 2-26 1-4 3-50 Fracture Toughness 25-100+ 13-42 10-20 25 (MPa m1/2)

Aside from their mechanical properties, which have motivated the study of intermetallics for structural applications, intermetallics tend to be good electrical conductors, motivating the study of intermetallics for various electrical applications. In some cases the resistivity of the intermetallic is lower than the resistivity of the pure metals that comprise the intermetallic. When the intermetallic compound is stable over a range of compositions, the resistivity will generally be a minimum at the stoichiometric composition.

A. Selection of Intermetallics

Several criteria were considered when choosing a selection of intermetallics to test. The first property investigated was resistivity, which yielded literature values for more than 150 intermetallics. To obtain a smaller subset of materials, only materials with a resistivity of less than ~35 μΩcm were further investigated for their potential usefulness to this study. A list of the materials that were initially under consideration is shown in Table 2.

4

Table 2 – Compilation of property data for intermetallics having a resistivity lower than ~35 μΩcm. Data for several pure metals are included as a comparison. Hardness data for the pure metals are for the metals after annealing. Melting Resistivity (μΩcm) Thermal Fracture Youngs Temp Hardness Material Thin Expansion Toughness Modulus References (C) Bulk* (GPa) Film (10-6/C) (MPa m1/2) (GPa) [21-22] Si 7.6 [21] Al 660 2.65 2.9-3.1 23.1 0.2 71 [21, 23] Au 1064 2.35 2.6 14.2 0.3 79 [21, 24] Cu 1080 1.67 2.0 16.5 0.4 117 [21, 25] Ni 1455 6.84 12-19 13.4 0.6 200 [21, 25] Pt 1768 10.6 26-28 8.8 0.4 170 [21, 26] Ru 2334 7.6 11-40 6.4 2.0-3.4 432 [21, 27-28] Ti 1668 42 55-60 8.6 0.7 120 [21, 29] W 3422 5.65 7.7-8.7 4.5 1-2.4 3.4 411 [21, 23, 30] AuAl2 1060 8 (sc) - 11.3-12 - - - [31-33] CuAl 563 8 12 - - - - [34-35] CuAl2 591 7-8 7-11 20 - - 106 [34-39] Cu4Al3 570 13 - - - - - [35] [34-35, 38, 40- Cu9Al4 873 13-26 13-40 18.2 - - - 41] NiAl 1638 10-11 - 11.9-15.1 6-14 5 194-237 [42-48] RuAl 2000 13-15 - 5.5 - 4 280 [49-52] TiAl3 1350 16 - 10.7 - 8-12 216 [46, 53-56] [46, 54-55, 57- ZrAl3 1580 17 - 12.1 - 6 202-205 58] CrB2 2200 56-91 21 6-8 - 13-22 77 [59-63] HfB2 3380 8-15 125 6.8-7.7 - 11 440-530 [61-62, 64-67] MgB2 1550 12 200 - 1.4 12 243 [68-71] Mo2B5 2140 17-18 - 8.6 - 23 672 [46, 62, 72] [46, 60, 62-63, NbB2 2950 12-13 32-137 8 - 18-20 637 68, 73-74] Ni3B 1156 26 - 10-17 1.4-2 12 151-182 [46, 75-76] Pd3B 890 10 - - - 5 - [46, 77] ScB2 2250 7-15 - 6.8-7.6 - 13-17 - [46, 78-79] [46, 60, 62, 64, TaB2 3037 14 68 7.9 - 22 257 73] TiB2 3225 9-20 16-150 8.3 5.5-7 18 526-579 [62, 64, 80-85] VB2 2747 38 16 7.6 - 26-29 268 [46, 60, 62, 86] W2B5 2365 33-39 21 7.8 - 26 775 [46, 60, 62, 72] [46, 62, 66, 68, ZrB2 3250 7-10 25-162 5.9 3.5 23 489-552 83-84, 87-89] NbC 3613 20-75 - 7.2 21 338 [46, 90-91] TaC 3985 15-30 - 7.1 17 285 [46, 90, 92-93] AuGa2 491 13 (sc) - 16 [31, 33] PtGa2 922 - 18 - [94] AuIn2 541 - 8-13 17.5-18 [33, 95-97] CuIn2 - - 14 - [98] DyIn3 1150 15 (sc) - 18.8 [99-100] GdIn3 1175 14 (sc) - 17.4-18.6 [100-102] HoIn3 1130 13 (sc) - - [102] LuIn3 960 13 (sc) - - [99, 102]

5

PdIn 1285 - 18 - [103-104] Pd5In3 946 - 17 - [105] TbIn3 1140 16 (sc) - - [99] ZrIn2 >900 20 - - [106] TiN 3290 - 24-90 9.4 - 20 350 [6, 46, 107-111] ZrN 2960 - 11.4 6.1-7 5.9-6.6 13-14 380-382 [21, 107, 112] CoSi2 1326 32 14-21 9.5-10.1 - 5 183 [46, 113-118] [46, 117, 119- MoSi2 2020 22 59-100 8.3 3-4.5 9-11 438-442 125] [46, 121, 126- NbSi2 1940 6.3 38-150 8.4 1.9 10 362-364 130] NiSi 992 - 12-18 - - 10 132 [131-135] [46, 131-132, Ni2Si 1255 - 24 16.5 5.7 13 - 136] Pd2Si 1394 - 30-35 13.2-23 - - 86 [1-2, 121, 137] [46, 117, 123, TaSi2 2040 38 34-50 9.5 - 15-16 357-361 127, 138-139] [116-117, 123, TiSi2 1500 17 14-24 9.9 1.9 6 256-278 138, 140-142] [46, 57, 121- VSi2 1677 13 50-64 11.2 - 8 331-343 123, 143-144] [46, 117, 121- WSi2 2160 38 23-33 6.0-6.3 3.2 8 465-471 123, 138, 145- 146] *sc indicates material that was a single crystal

The materials listed in Table 2 were next analyzed with respect to thermal expansion and fracture toughness where data was available. Originally borides were considered as potential materials due to their low thermal expansion coefficients and low resistivity. Unfortunately, due to their high melting points, obtaining thin films of these materials with low resistivity is only possible when deposition is followed by high temperature anneals (temperatures 800-1100°C) [7]. Some of the silicides have lower melting temperatures than the borides, thus making it possible to obtain thin films with lower resistivity with low temperature anneals, but in general the resistivity values of the silicides were not as encouraging as those of the aluminides. The aluminides also have lower melting temperatures than the borides, and with one exception have high coefficients of thermal expansion when compared to common substrates (oxidized silicon will be used for this study). The only aluminide that was identified that had a low thermal expansion coefficient (and low resistivity) was RuAl, and this intermetallic along with the similar compound NiAl will be the intermetallics under investigation in this research.

III. OVERVIEW OF NiAl

NiAl is one of several intermetallic compounds in the Ni-Al system (see Figure 2a). It has the B2 crystal structure (Figure 2b) and the stoichiometric alloy melts at 1638°C. There is a large composition range over which NiAl is the stable intermetallic: at room temperature this ranges from 41-55% Al. Because of its low density, high thermal and electrical conductivity, high melting temperature, good oxidation resistance, and high elastic modulus, NiAl is considered a promising material for many applications such as electronic metallizations, high-temperature coatings, surface catalysts, high-current vacuum circuit breakers, turbine blades, and high temperature aerospace applications and propulsion systems [147]. NiAl exhibits some limited room temperature ductility although it is considered to be brittle at room temperature and undergoes a ductile-to-brittle transition in the range of 300-600°C [147]. For an intermetallic

6 compound NiAl has a high fracture toughness, which ranges from 4-14 MPa at room temperature and 10-50 MPa as the temperature increases to 350-400ºC [42, 147].

1700 a)

L

1400 b)

C)

NiAl Ni

1100

Al

3

Ni Temperature (

800 3

Al

2

3

Ni

3

NiAl

Al 5

500 Ni 0 20 40 60 80 100

Composition (at% Ni) Figure 2 – (a) Ni-Al phase diagram. Adapted from [148]. (b) The B2 crystal structure of NiAl [147].

A. Mechanical Properties of NiAl

At room temperature when tested in tension, stoichiometric NiAl shows limited ductility (due to only 3 independent slip systems) with up to 2% strain to failure [149-151]. This value increases as the grain size decreases and as the temperature increases. In the range of 300-600°C NiAl undergoes a ductile-to-brittle transition (DBT) and considerable elongation to failure can be observed above the DBT (>30%) [149, 151]. The increasing ductility in NiAl has been proposed to be due to the activation of dislocation climb [147]. The effect of temperature on the strain to failure is shown in Figure 3.

7

75 50 a) b)

60 40

45 30 range of values for the same alloy in

30 20 compression

Strain to Failure (%) Failure to Strain Strain to Failure (%) Failure to Strain 15 10

0 0 0 100 200 300 400 500 0 100 200 300 400 500 600 700 Temperature (°C) Temperature (°C) Figure 3 – The effect of temperature on the strain to failure for NiAl tested in tension. (a) Results from [149] for NiAl with a grain size of 84 μm. (b) Results from [151] for NiAl with 50.3% Al.

A similar transition is observed as the grain size decreases [18, 151]. For NiAl with 49% Al tested in tension at 400°C, ~2-3% strain to failure was observed when the grain size was large (45-125 μm), but increased to more than 40% when the grain size was decreased to 8 μm [18]. It was determined that the transition for brittle to ductile failure occurred at a grain size of ~20 μm for a testing temperature of 400°C. No ductility is observed for off-stoichiometry NiAl tested in tension at room temperature due to the generated point defects constraining dislocation motion [150-151]. Another effect of decreasing grain size is an increasing yield and fracture strength, which occurs due to the grain boundaries hindering the motion of dislocations [18]. When tested at 400°C the fracture strength increased from less than 400 MPa when the grain size was large (45-125 μm) to more than 800 MPa as the grain size was decreased to 8 μm. The effect of grain size on the percent elongation and the yield and fracture strength is shown in Figure 4. The values of yield and fracture strength in Figure 4 are for alloys tested at 400°C and are expected to be larger at room temperature as both decrease with increasing temperature (see Figure 5).

8

50 900 a) Yield Stress Fracture Strength 40 750

30 600

20 450 Stress (MPa) Stress

Percent Elongation (%) Elongation Percent 10 300

b) 0 150 0 30 60 90 120 150 0 30 60 90 120 150 Grain Size ( m) Grain Size ( m) Figure 4 – The effect of grain size on the tensile (a) elongation and (b) yield and fracture strength for NiAl with 49% Al tested at 400°C. Data from [18].

400 250 a) b) Yield Stress Tensile Strength 200 300

150

200

100

Stress (MPa) Stress Stress (MPa) Stress

100 50

Yield Stress Ultimate Tensile Stress 0 0 0 100 200 300 400 500 600 700 0 100 200 300 400 500 Temperature (°C) Temperature (°C) Figure 5 – The effect of temperature on the yield stress and tensile strength of NiAl tested in tension. (a) Data on NiAl with 50.3% Al [151]. (b) NiAl with a grain size of 84 μm [149].

Similar trends are seen for NiAl tested in compression although the alloys are more ductile in compression. Ni-rich samples, which show no strain to failure in tension at room temperature, show more than 6% strain to failure for alloys with a large grain size (350 μm) when tested in compression [17]. The effect of grain size on the strain to failure and the yield and fracture strength is shown in Figure 6 for Ni-rich NiAl. The values for the yield and fracture strength are expected to be higher than stoichiometric NiAl as strength is noted to increase with deviations from stoichiometry (see Figure 7).

9

1000 20 a) Yield Stress b) Fracture Strength 800 16

600

12

400 Stress (MPa) Stress

200 (%) Failure to Strain 8

0 0 100 200 300 400 0 100 200 300 400 Grain Size ( m) Grain Size ( m) Figure 6 – The effect of grain size on the properties of NiAl with 48.9% Al when tested in compression. Data from [17].

1200 45% Al 48% Al 49% Al 50% Al 900

600

Yield Stress (MPa) Stress Yield 300

0 0 100 200 300 400 500 Grain Size (m) Figure 7 – The effect of stoichiometry on the yield stress for NiAl tested in compression. Data from [17].

Significant improvements in strength have been observed for thin film NiAl. Al-rich thin films have a fracture strength of over 2 GPa if the film thickness is decreased to 0.4 μm [19]. Ni-rich thin films have yield stresses that are higher than 1.98 GPa for a 0.2 μm thick film [152]. The effect of film thickness on the yield and fracture strength of NiAl thin films is shown in Figure 8.

10

2400 2400 a) NiAl Cu 2000 Al 2000

1600 1600

1200

1200

800 Yield Stress (MPa) Stress Yield

Fracture Strength (MPa) Strength Fracture 800 400 b) 400 0 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 0.0 0.8 1.6 2.4 3.2 Film Thickness ( m) Film Thickness (m) Figure 8 – The effect of film thickness on (a) the fracture strength for Al-rich thin films and (b) the yield stress for Ni-rich thin films. Arrows indicate films that did not yield during testing. Films were in tension. Data from [19, 152].

B. Electrical Properties of NiAl

NiAl has a low bulk resistivity at room temperature of 8-11 μΩcm [43, 48, 153-154]. A very low value of ~5.2 μΩcm has been obtained for a stoichiometric single crystal of NiAl [153]. A minimum in the room temperature resistivity of bulk NiAl is obtained at the stoichiometric compound [48, 154], or at slightly Al-rich compositions (~49.6% Ni) [43, 153]. The resistivity increases with increasing deviations from stoichiometry, with a more rapid increase in resistivity observed for Ni-rich compositions [153]. A summary of the literature data on the room temperature bulk resistivity of NiAl is shown in Figure 9.

The resistivity of NiAl thin films has been studied by several groups, but there are no literature results on the resistivity of RuAl thin films. Values for the resistivity of NiAl range from 24-45 μΩcm for ordered NiAl thin films [155-157] and 53.9 μΩcm for partially ordered films [156]. As NiAl remains ordered up to the melting temperature, disordered films could only be produced by deposition onto substrates cooled by liquid nitrogen to -123°C [156]. Even with the decreased substrate temperature the film was still mostly ordered, but a reduction in grain size from 50 nm to 20 nm was obtained when the substrate temperature was decreased from 507°C to -123°C. Thus, the increased resistivity for the “disordered” film (from 31.9 μΩcm to 53.9 μΩcm) is a combination of decreased order and decreased grain size. As-deposited films sputtered from a compound target onto room temperature substrates had a resistivity of 45 μΩcm [155], while the lowest value of 24 μΩcm was obtained for a 100 nm thick film co-deposited in a molecular beam epitaxy (MBE) system (Al effusion cell, and Ni electron beam evaporation) heated to 600°C [157]. As the film thickness of the MBE film decreased the resistivity increased significantly. This is shown in Figure 10 for films with a thickness ranging from 1.5 nm to 100 nm [157].

11

50 Caskey (1973) Butler (1969) Jacobi (1971)

40 Yamaguchi (1968) cm) cm)

30 

20 Resistivity ( Resistivity

10

0 44 48 52 56 60 64 Composition (at% Ni)

Figure 9 – Summary of the literature data on the room temperature resistivity of bulk NiAl as a function of the Ni content. Data compiled from [43, 48, 153-154].

400

300

cm) cm) 

200 Resistivity ( Resistivity 100

0 0 25 50 75 100 Film Thickness (nm) Figure 10 – Resistivity of NiAl thin films as a function of the film thickness. Data from [157].

12

IV. OVERVIEW OF RuAl

RuAl is one of several intermetallic compounds that form in the Ru-Al system (see Figure 11). RuAl, like NiAl, has the B2 crystal structure, and is the equilibrium compound for ruthenium contents ranging from ~46.2-50.2 at% at room temperature [158]. The stoichiometric compound melts congruently at 2060ºC, which is more than 400°C higher than NiAl [158]. Interest in this alloy stems from its high electrical and thermal conductivity, good toughness (compared to other intermetallics), high strength, good oxidation resistance, high melting temperature, and excellent corrosion resistance. RuAl has been found to be resistant to attack from aqueous media such as nitric acid, aqua regia, HF, sulfuric acid, FeCl3, NaOH, and HCl [16]. The only solution found to date that attacks RuAl is NaOCl [16].

2400

Ru(Al)

C)

1600

RuAl

5

Al

3

2

13

Al

2

Al

Ru

4

Ru 2

800 Ru

Temperature (

RuAl

6 RuAl 0 0 20 40 60 80 100

Composition (at% Ru) Figure 11 – Ruthenium-aluminum phase diagram adapted from [158].

More is known about NiAl, but in a few studies, RuAl has exhibited more room temperature ductility than NiAl and is expected to have a higher fracture toughness [159-160]. Although RuAl has shown higher ductility than NiAl, many of the studies on RuAl are not on single phase RuAl, but on a two phase mixture of RuAl and Ru(5%Al) [49, 52, 159, 161]. The 2-phase microstructure noted in many of the RuAl studies is mainly due to complications with melt processing due to the high melting temperature of RuAl, attack on the refractory lining, and volatilization of aluminum [162]. These complications will be avoided in the solid-state reactions used in this work to deposit thin films of RuAl. Many attribute the high ductility of RuAl to a layer of ruthenium at the grain boundaries, which inhibits grain boundary cracking, a predominant mode of failure in intermetallic compounds. In support of this statement is the fact

13 that aluminum-rich RuAl has no ruthenium at the grain boundaries, and is more brittle than ruthenium-rich and stoichiometric RuAl [49]. The mechanical and electrical properties will be discussed in more detail below, and the oxidation of RuAl will be reviewed in the Chapter 5.

A. Mechanical Properties of RuAl

Significant research on bulk intermetallics for high temperature use has been conducted by Fleischer et al. who identified close to 300 intermetallics that melt at or above 1500ºC, and have performed screening tests on 90 different binary compounds for possible use at elevated temperatures [163]. RuAl was 1 of only 5 compounds that were identified that warranted further research. The intermetallics proposed for further research were those that showed high 1 toughness [164] at room temperature: RuAl, RuSc, IrNb, and off-stoichiometry Ru11Ta9 and Re24Ti5 [58]. In a later publication it was concluded that of these 5 alloys RuAl showed the most promise for high temperature applications [163]. Further research on RuAl showed that deviations from stoichiometry to the Al-rich side resulted in a brittle material, whereas Ru-rich samples were more ductile [58], which has been proposed to be due to the second phase -Ru being present at the grain boundaries in Ru-rich RuAl. It has also been observed that adding 0.5% boron increases the strength and ductility of RuAl, but additions of boron of 1% or more resulted in a brittle alloy [58]. Table 3 shows compression stress-strain data for RuAl alloys as the composition is modified by changing the ratio of Ru/Al and by adding boron. Significant plasticity is noted in the Ru-rich alloys as well as the boron-doped alloys with at least 50 at% Ru. Compression tests on near-stoichiometric single phase RuAl result in values of ~200-400 MPa for the yield stress, ~800-1000 MPa for the ultimate strength, and ~10% strain to failure [50, 161-162, 164-165]. Increasing the ruthenium content and the presence of a second phase at the grain boundaries increases the ultimate strength to ~1500-2500 MPa and the strain to failure to 20-25%. The yield stress as a function of temperature for RuAl with various compositions is shown in Figure 12. All alloys contained a second phase at the grain boundaries. The data show that increasing the ruthenium content and adding boron to RuAl both increase the yield strength. The yield strength in compression is similar for NiAl and RuAl; however RuAl appears to have a higher ultimate strength than NiAl.

1 Toughness measured as “chisel toughness” on a relative scale from 0-3, but was found to correlate well with measured mechanical properties. For 11 materials with toughness and compression data, the 3 samples with a toughness of 0 or 1 experienced less than 0.5% strain, whereas 6 of 8 samples with a toughness of 2 or 3 exhibited greater than 0.5% strain. A toughness value of 3 was observed for RuAl, RuSc, and IrNb, and values of 2 were obtained for off-stoichiometry Ru11Ta9 and Re24Ti5. 14

Table 3 – Stress -strain data for RuAl alloys in compression [50, 164]. Un-doped RuAl RuAl with 0.5 at% boron § § § § Composition (at% Al) max (%) max (MPa) max (%) max (MPa) 40† 29 2150 - - 42† 31 2100 - - 47† 17 1500 33.6 3000 50 9 1000 22.7 2100 53 0 170 6.6 800 †Composition is in the two-phase field: Ru-RuAl § max and max refer to the strain and stress at which cracks began to propagate

1000 RuAl Ru Al 53 47 Ru Al B 49.5 50 0.5 Ru Al B 800 52.5 47 0.5

600 Yield Stress (MPa) Stress Yield 400

200 0 200 400 600 800 1000 Temperature (°C) Figure 12 – The yield stress as a function of temperature for RuAl samples with a grain size of ~150-200 μm and tested in compression. All samples contained a second phase, which was mostly located at the grain boundaries. Data from [161].

B. Electrical Properties of RuAl

RuAl has a low bulk resistivity of 13-15 μΩcm [49, 51] at room temperature, which is slightly higher than the bulk resistivity of NiAl. It is interesting to note that changing the stoichiometry in RuAl towards the ruthenium-rich side of the phase diagram has very little effect on the resistivity [49]. The resistivity of RuAl alloys with ruthenium contents from 50-53% ranged only from 13.3 μΩcm for the Ru-rich alloy to 14.3 μΩcm for the stoichiometric alloy. As both of these alloys contained some porosity and second phase -Ru, a slight decrease in resistivity may be possible for single-phase, fully dense samples. The effect of a second phase on the resistivity can be seen in the data of Hermann et al. [166] where a single crystal with 48.75 at% Al had a resistivity of 22 cm, and a 2-phase polycrystalline sample had a resistivity of 30- cm (48.40-48.94 at% Al). Changes in resistivity towards the aluminum-rich side were hard to 15 ascertain as the aluminum-rich RuAl tested was very brittle and contained cracks and porosity. Smith et al. [51] determined that the resistivity as a function of temperature for RuAl was linear from 20-1000ºC with a room temperature resistivity of 15.3-19.4 cm (nominal composition: 48-52 at% Al) and a resistivity at 1000ºC of 51.5-60.9 cm (temperature coefficient of resistance: 21.6-25.2 x 10-4/ºC).

V. DISCUSSION

It is believed that this research will increase the understanding of the properties of intermetallics, especially in the area of thin films. The majority of intermetallic research has in the past centered on bulk samples or single crystals, and there are very few published results on the formation of RuAl thin films [167-168]. NiAl has been more widely studied. The majority of intermetallics are brittle, this being a key reason that their use is not more widespread. Thin film applications may be an area that could find many uses for intermetallics as studies on NiAl indicate that the fracture stress increases as the film thickness decreases [19]. For example, NiAl alloys containing 52.2% Al showed an increase in fracture stress from approximately 580 MPa to 1590 MPa when the film thickness was decreased from 3.0 to 0.6 μm. Suggested uses for RuAl include corrosion protection coatings, spark-plug electrodes, gas turbine components, bond coats in oxidation protective coatings, and electrical contacts [16, 158]. Several of these proposed uses are thin film applications and therefore studies on the processing and properties of RuAl are of extreme importance.

Another important aspect of this research is in the area of the thermomechanical fatigue behavior of thin films. The method of AC TMF testing is relatively new, and the results presented herein will increase understanding of the technique as well as extend the range of materials tested using this method. This testing technique makes it possible to compile a large amount of data in a relatively short amount of time on the thermal fatigue behavior of NiAl and RuAl intermetallic films. To our knowledge, this study will be the first on the TMF behavior of thin film intermetallics.

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[142] R. Rosenkranz, G. Frommeyer, and W. Smarsly, Microstructures and properties of high melting point intermetallic Ti5Si3 and TiSi2 compounds. Materials Science & Engineering A: Structural Materials: Properties, Microstructure and Processing, 1992. A152(1-2): p. 288-294. [143] R.L. Fleischer. Mechanical properties of diverse high-temperature compounds-thermal variation of microhardness and crack formation. in High-Temperature Ordered Intermetallic Alloys III. Symposium, 29 Nov.-1 Dec. 1988. 1989. Boston, MA, USA: Mater. Res. Soc. [144] F. Nava, O. Bisi, and K.N. Tu, Electrical transport properties of V3Si, V5Si3, and VSi2 films. Physical Review B, 1986. 34(9): p. 6143-6150. [145] O. Knotek, R. Elsing, and H.-R. Heintz, Measurement of the coefficients of thermal expasion of plasma-sprayed coatings, in Thermal Spray: Advances in Coatings Technology, D.L. Houck, Editor. 1988, ASM International. p. 181-183. [146] I.J. Shon, D.H. Rho, H.C. Kim, and Z.A. Munir, Dense WSi2 and WSi2-20 vol.% ZrO2 composite synthesized by pressure-assisted field-activated combustion. Journal of Alloys and Compounds, 2001. 322(1-2): p. 120-126. [147] D.B. Miracle, Physical and mechanical properties of NiAl. Acta Metallurgica et Materialia, 1993. 41(3): p. 649-684. [148] F. Lechermann and M. Fahnle, Ab-initio statistical mechanics for the phase diagram of NiAl including the effect of vacancies. Physica Status Solidi B, 2001. 224(2): p. R4-R6. [149] F. Ebrahimi and T.G. Hoyle, Brittle-to-ductile transition in polycrystalline NiAl. Acta Materialia, 1997. 45(10): p. 4193-204. [150] P. Nagpal and I. Baker, Room temperature fracture of FeAl and NiAl. Materials Characterization, 1991. 27(3): p. 167-173. [151] K. Vedula and P.S. Khadkikar, Effect of stoichiometry on low temperature mechanical properties on B2 NiAl alloys. High Temperature Aluminides and Intermetallics, 1990: p. 197-216. [152] P. Wellner, G. Dehm, O. Kraft, and E. Arzt, Size effects in the plastic deformation of NiAl thin films. Zeitschrift fuer Metallkunde, 2004. 95(9): p. 769-778. [153] S.R. Butler, J.E. Hanlon, and R.J. Wasilewski, Electric and magnetic properties of B2 structure compounds: NiAl, CoAl. Journal of the Physics and Chemistry of Solids, 1969. 30: p. 1929-1934. [154] H. Jacobi and H.-J. Engell, Defect structure in non-stoichiometric -(Ni, Cu)Al. Acta Metallurgica, 1971. 19: p. 701-711. [155] T.J.S. Anand, H.P. Ng, A.H.W. Ngan, and X.K. Meng, Temperature-coefficient-of- resistance characteristics of sputter-deposited NixAl1-x thin films for 0.5

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[159] D.C. Lu and T.M. Pollock, Low temperature deformation kinetics of ruthenium aluminide alloys. Materials Research Society Symposium - Proceedings, 1999. 552: p. 7-11. [160] T.M. Pollock, D.C. Lu, X. Shi, and K. Eow, A comparative analysis of low temperature deformation in B2 aluminides. Materials Science and Engineering A, 2001. 317(1-2): p. 241-248. [161] T.K. Nandy, Q. Feng, and T.M. Pollock, Elevated temperature deformation and dynamic strain aging in polycrystalline RuAl alloys. Intermetallics, 2003. 11(10): p. 1029-1038. [162] I.M. Wolff and G. Sauthoff, Role of an intergranular phase in RuAl with substitutional additions. Acta Materialia, 1997. 45(7): p. 2949-2969. [163] R.L. Fleischer, C.L. Briant, and R.D. Field, Tough, ductile high-temperature intermetallic compounds: results of a four-year survey. Materials Research Society Symposium - Proceedings, 1991. 213: p. 463-474. [164] R.L. Fleischer, R.D. Field, and C.L. Briant, Mechanical properties of high-temperature alloys of AlRu. Metallurgical Transactions A (Physical Metallurgy and Materials Science), 1991. 22A(2): p. 403-414. [165] F. Mücklich, N. Ilic, and K. Woll, RuAl and its alloys, Part II: Mechanical properties, environmental resistance and applications. Intermetallics, 2008. 16(5): p. 593-608. [166] R. Hermann, H. Vinzelberg, G. Behr, K. Woll, F. Mücklich, and B. Büchner, Magnetic field controlled floating-zone crystal growth and properties of RuAl single crystal. Journal of Crystal Growth, 2008. 310(18): p. 4286-4289. [167] K.W. Liu and F. Mucklich, Formation and characterization of periodically structured B2-aluminides thin films by laser interference irradiation. Materials Research Society Symposium - Proceedings, 2003. 753: p. 6.5.1-6.5.6. [168] W.K. Shen, J.H. Judy, and J.-P. Wang. In situ epitaxial growth of ordered FePt (001) films with ultra small and uniform grain size using a RuAl underlayer. in 49th Annual Conference on Magnetism and Magnetic Materials. 2005. Jacksonville, Florida (USA): AIP.

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CHAPTER 2

NICKEL ALUMINIDE FILM FABRICATION

27

I. INTRODUCTION

Nickel aluminide thin films have been fabricated by several groups using pre-alloyed sources/targets, co-deposition of elemental materials (co-sputtering and co-evaporation), and by the reaction of multilayers of nickel and aluminum. In the sections that follow, the literature on the fabrication of nickel aluminide films will be summarized.

A. Ni-Al Films from the Reaction of Ni/Al Multilayers

Many research groups have studied the solid-state formation of Ni-Al intermetallics from Ni/Al multilayers [1-13]. Although numerous studies have been published, there are discrepancies between the various research groups as to the early stages of intermetallic formation from multilayer thin films. This could be due to differences in microstructure arising from different deposition conditions, or to the resolution of the techniques used to study the films.

As-deposited multilayers of nickel and aluminum have been described as: unmixed [10], having some nickel diffusion into the aluminum layers, but no aluminum diffusion into the nickel layers [13], having a 2.5 nm reaction layer at the interfaces [7], a 1 nm solid solution at the interface [6], and an amorphous layer at the interfaces [9]. During annealing and irrespective of composition, the intermetallics typically form sequentially beginning with the more aluminum rich phases, and progressing towards more nickel rich phases as the annealing temperature increases (NiAl3→Ni2Al3→NiAl…). The final phase that forms is always the one corresponding to the composition of the multilayer. The remaining discussion on Ni/Al multilayers will only cover those multilayers with compositions leading to films that are single phase NiAl. Although the more Al-rich phases tend to form first, NiAl3 is not always noted as the first phase to form, and if the Al/Ni bi-layer thickness (hereafter referred to as the wavelength – see Figure 1) is 10 nm or less NiAl is the first and only phase to form [4, 7, 9]. The first phase to form in larger wavelength multilayers was determined to be either NiAl3 [3, 9-12], a metastable phase (Ni2Al9) [4] that was first noted in Ni-Al alloys rapidly quenched from the melt [14], or a non-equilibrium 63% Al NiAl phase [7]. Blobaum et al. [2] studied the formation of Ni2Al9 in a series of multilayer films, and found that Ni2Al9 was the first intermetallic to form for wavelengths of 25 nm to 200 nm, and NiAl3 formed first when the wavelength was decreased to 12.5 nm. It could be that the microstructure of the starting multilayer films is causing the difference in reaction paths. Jeske et al. [5-6] studied two different Ni/Al bi-layers, one deposited by sputtering and the other by electron-beam evaporation. At 200ºC the e-beam sample formed a uniform layer ~3 nm thick with a composition of ~60% Al, whereas the sputter deposited sample did not show this layer and instead exhibited Ni wetting of the aluminum grain boundaries, and the eventual formation of NiAl3 at 250ºC (which is the same temperature at which NiAl3 formed in the e- beam sample). They concluded that the difference in the early reaction was due to the different microstructures of the samples (larger grains in the case of the e-beam sample).

28

nickel aluminum

bi-layer thickness/wavelength

Figure 1 – Schematic showing a Ni/Al multilayer as-deposited. The first and last layers represent half of an aluminum layer. The diagram is representative of the multilayers deposited in this work, with the thickness of the individual layers varying depending on the composition and wavelength of the sample. The number of layers deposited was adjusted so that the total thickness was ~200 nm.

The transformation temperatures for intermetallic formation have been shown to depend on the wavelength of the film, and shift to higher temperatures as the wavelength increases. For example, in a 60Ni/40Al film with a 10 nm wavelength, NiAl formed at 175ºC, but was not formed at temperatures of up to 370ºC in a sample with a wavelength of 80 nm [4]. A summary of some of the literature data on intermetallic formation from Ni/Al multilayers is shown in Table 1.

Table 1 – Summary of some of the data on Ni/Al multilayer films with compositions in the B2 NiAl phase field. The temperatures should be considered an approximation as some were taken at the beginning of DSC exotherms, while some were taken at the end of the DSC exotherm, and others are from XRD data.

Wavelength First Intermetallic Forms NiAl Forms NiAl Reaction Annealing Time † Reference (nm) (ºC) (ºC) Complete (ºC) (min, °C/min) 5 150 NiAl 150 275 10 min [9] 10 175 NiAl 20°C/min [4] 10 80-90 63% Al NiAl 40°C/min [7] 20 ~100 63% Al NiAl 307 40°C/min [7] 27.6 230 NiAl3 (Ni2Al9?) 275-362 362 20°C/min [11] 30 125 NiAl3 175 225 10 min [9] 50.3 160 NiAl3 250 330 45 min [10] 80 250 Ni2Al9 > 370 20°C/min [4] 140 150 NiAl3 325-425 500 10 min [9] 500 270 NiAl3 500 600 50°C/min [12] 520 350 NiAl3 425 30 min [3] †Those listed with a heating rate refer to samples heated in a DSC/DTA

While there is considerable scatter in the data (due to different measurement techniques, the number of annealing temperatures studied, etc.), in general the larger the wavelength is, the higher the annealing temperature for the formation of the first intermetallic, for the formation of NiAl, and for the complete reaction to NiAl. This is shown schematically in Figure 2, where

29 curves are shown for the temperatures at which intermetallics begin to form, NiAl begins to form, and the reaction to NiAl is complete as a function of the wavelength of the multilayer. Curve fits to the data are also included and are in good agreement with the data obtained in the present study. For the wavelengths analyzed in this chapter (20-33 nm), the linear fits give values of ~150°C, 250°C, and 320°C for the beginning of intermetallic formation, the beginning of NiAl formation, and the complete reaction to NiAl.

700 First intermetallic forms 600 NiAl forms Reaction to NiAl complete

500

400

300

Temperature (ºC) Temperature 200

100

0 10 100 Bi-Layer Thickness (nm) Figure 2 – Figure plotting some of the literature data on the reaction of Ni/Al multilayers with compositions within the single phase B2 NiAl phase field. Curves are shown for the beginning of intermetallic formation, the beginning of the formation of NiAl, and the point at which the film is completely reacted to form NiAl. The data are curve fit with the equation y = a(1-bx). Data compiled from [3-4, 7, 9-12].

The grain size of the reacted films was shown by Rothhaar et al. to be ~45 nm after annealing at 330ºC for 45 minutes [10]. They also found that some nickel had diffused into the silicon substrate after annealing at 330ºC. This was ascribed to the fact that Ni and Si react to form silicides at temperatures of over 300ºC. In fact, the majority of studies have been carried out on multilayers deposited onto Si substrates [2, 4, 8-11, 13] even though Ni2Si can be formed at temperatures as low as 270ºC [15]. It is possible to suppress the reaction of Ni with Si if an oxide layer is present between the two. Liehr et al. [16] found that no reaction occurred between Ni and SiO2 up to ~800ºC if the oxide was greater than or equal to 20 Å thick. At temperatures of greater than 800ºC, Ni diffused through the SiO2 layer, possibly through pinholes in the oxide, and reacted with the Si to form NiSi2. As the thickness of the oxide layer decreased, so did the reaction temperature.

30

In a study by Zalar et al. [13], Auger electron spectroscopy (AES) was used to study the reaction of Ni/Al films (5 layers of: 25 nm Ni, 38 nm Al) that were heated in a DSC at 40ºC/min. Their films had nickel as the top layer, and after heating to 550ºC, the top two layers appeared to be oxidized, and largely un-reacted. The top layer was Ni (~65%) and O (~35%), while the second layer contained Al (~60%), Ni (~10%), and O (~30%). This two-layer oxidized region does not seem to occur when Al is the top layer. In this case there is a small single surface oxide layer containing mostly Al (~46%) and O (~41%), but with a small concentration of Ni (~13%) [1], which is similar in composition to the second oxide layer observed by Zalar et al. [13].

Besides conventional and DSC annealing, laser mixing [1, 17], hydrogen plasma annealing [8], and ion beam mixing [8] have also been used to study the reaction of Ni/Al multilayers. Laser mixing appears to be an attractive method for creating composite, or graded coatings, but not for the formation of homogeneous films. In the laser mixed Ni/Al studied by Bhattacharya et al. [1], B2 NiAl was formed away from the center of the laser beam, and the grain size could be tailored (from 115 nm to 400 nm) by changing the laser power. Near the center of the beam the face centered tetragonal (FCT) phase was formed. In a separate study by Liu at al. [17], they found a mixture of fine-grained Ni and Al (20 nm), and larger grained B2 NiAl (40 nm) in laser irradiated bi-layers of Ni and Al.

Mozetic et al. [8] used hydrogen plasma annealing to prevent oxidation of the Ni/Al multilayers (5 layer: 25 nm Ni, 38 nm Al) and found that although the periodicity of the film disappeared at ~527ºC, no NiAl had been formed in the fully mixed film. In fact, analysis of the films by extended x-ray absorption fine structure (EXAFS) showed a very similar local environment around the Ni atoms in the annealed and pure Ni films. In ion beam mixed samples, the periodicity disappeared at lower temperatures (~277ºC) than the hydrogen plasma annealed samples, but homogeneous films were not formed for temperatures up to 327ºC (Ni-rich at the surface, and Al-rich at the substrate). The films were found to be ~64% NiAl and ~36% Ni at 227ºC, and ~79-86% NiAl and 14-21% un-reacted Ni at temperatures between 277°C and 327ºC. There does not appear to be any benefit to incorporating ion beam mixing during annealing of Ni/Al multilayers as the results at 327ºC with ion beam mixing do not appear to be any better than annealing without ion beam mixing, where a 50.3 nm wavelength film was completely reacted and homogeneous by 330ºC [10].

B. Ni-Al Films from Co-Deposition and Alloyed Sources/Targets

As an alternative to the reaction of multilayers of Ni and Al to form intermetallic films, co- evaporation, evaporation from alloyed sources, co-sputtering, and sputtering from alloyed targets can also be used to obtain intermetallic films. Depending on the composition and deposition conditions, intermetallic films can be formed without the need for post-deposition annealing. It is also possible to form non-equilibrium phases from these techniques.

Several papers by Andersson et al. [18-21] have studied the formation of Ni-rich (0-55 at% Al) Ni-Al films deposited by the co-evaporation of elemental nickel and aluminum in an electron- beam system at various substrate temperatures. They found that the solubility of Al in Ni could be increased during co-evaporation to a maximum of ~30% at 87ºC in comparison to the equilibrium value of ~7% at room temperature. For films deposited at 297ºC, the films

31 containing intermetallics were either fully disordered (BCC NiAl rather than B2 NiAl) or partially disordered [18]. This trend likely holds for lower temperatures as well. From their studies, they formed a phase diagram (see Figure 3) describing the phases present in Ni-Al films as a function of composition and substrate temperature.

900

800 NiAl Ni(Al) Ni2Al3 700 Ni3Al NiAl Ni(Al) Ni3Al 600 Ni3Al NiAl 500 NiAl Ni Al Ni(Al) 2 3 NiAl 400 Ni Al 3 Ni(Al) 3 NiAl 300 Ni2Al3

0 10 20 30 40 50 60 70 Concentration Al (at.%) Figure 3 – Thin film phase diagram for Ni-Al films deposited by co-evaporation as a function of substrate temperature. Adapted from [20]. The dashed lines represent extrapolations of the experimental data.

Comparing the phase diagram in Figure 3 to the equilibrium phase diagram shows that at low temperatures there is a large deviation from the expected phases, and for compositions of greater than 30 at% aluminum there is still a large deviation between the two phase diagrams for deposition temperatures of 517ºC. For example, the single-phase field for the B2 NiAl is narrower in the case of the co-evaporated thin films (45-52% Al) than in the equilibrium phase diagram (41-55% Al). The grain size [19] for the films deposited at 87ºC was less than 50 nm, with the exception of films with less than 5% Al, which had grain sizes ranging from 50 to 70 nm. A similar study was carried out by Kizuka et al. [22] on thin films deposited simultaneously from Ni and Al by resistance heating. A summary of their results is shown in Figure 4.

32

NiAl 3 Ni Al 400ºC Ni Al 2 3 2 3 NiAl NiAl Ni Al 3 Al 300ºC NiAl3 NiAl Ni2Al3

Al NiAl 200ºC 3 NiAl NiAl3 NiAl

Ni 100- Al NiAl Al 25ºC Ni Ni

20 30 40 50 60 70 80 Composition of Ni (at %) Figure 4 – Phases present as a function of substrate temperature in Ni-Al films evaporated simultaneously from Ni and Al sources by resistance heating. Data reproduced from [22].

The results from Kizuka et al. are similar to the results of Andersson et al., although a direct comparison is difficult as the exact boundary of phase fields was not given by Kizuka, and some of the depositions were carried out at lower temperatures. The two studies also used very different deposition rates (2 nm/min in the case of Kizuka, and 600-1000 nm/min in the case of Andersson). Even with the above-mentioned differences, both studies show that there is a large supersaturation of Al in Ni that decreases with increasing temperature, and both also show that single-phase intermetallic films are generally not formed at low temperatures when deposited by co-evaporation, and that the deposited intermetallics are not fully ordered. Ordering of the films deposited by Kizuka occurred when the films were heated to just above 400ºC. They also demonstrated an increased solubility of Ni in Al as compared to the equilibrium phase diagram (~0.02% at room temperature).

In the studies by Andersson et al. on co-evaporated films, it was noted that a unique surface feature began to appear at 34% Al [21], increased as the aluminum content increased, and then disappeared at ~50% Al. In this composition range, the film surface was covered in “outgrowths” that Andersson et al. described as “small and curved plates that were standing up from the plane of the substrate” [18]. These surface outgrowths all had the same orientation and from SEM images appear to be ~2-3 m long. A similar surface morphology was found for films deposited onto various substrates (cast iron, copper, steel, aluminum), but was not noted for films deposited from alloyed sources [21].

Johansson et al. [21] compared the surface morphology of co-evaporated films and films evaporated from a single alloyed source at substrate temperatures of ~150ºC. For aluminum contents of less than 10% the films were smooth, and the 2 sets of films were similar in

33 appearance. As the aluminum content increased to ~15% for co-evaporation, and to 10-20% for the alloyed sources, cracks appeared in both films (due to tensile stresses), with the crack density increasing as the aluminum content increased until a maximum was reached at ~30% for co- evaporation, and 30-40% for alloyed-sources. Cracks remained in films deposited from the alloyed source as the aluminum content was increased to 50%, whereas the co-evaporated films were covered in surface outgrowths. Cracking and surface outgrowths are not problems that have been encountered in Ni-Al films deposited by sputtering.

As with films deposited by co-evaporation, co-sputtered films frequently contain non- equilibrium phases. Cantor and Cahn [23] studied co-sputtered Ni-Al films deposited onto liquid nitrogen cooled substrates (~77-80K). As with the co-evaporated films, the content of Ni in Al (21%), and Al in Ni (16%) were increased as compared to the equilibrium phase diagram. The single phase NiAl region (21.2-57.5% Al) was also increased in comparison to equilibrium (41- 55% Al), although the alloy was in its disordered state. The only intermetallic formed at these deposition temperatures was disordered NiAl. The estimated grain size for the films was 5-10 nm.

Similar results were obtained by Sumiyama et al. [24] on the formation of Ni-Al films co- sputtered at ~77ºC. Significant solid-solutions were observed for both Ni in Al (15%), and Al in Ni (30%), and a large single phase NiAl region was observed (35-65% Al). As with Cantor et al., the only intermetallic formed was NiAl, but unlike Cantor et al. the NiAl was ordered, which was likely due to the slight increase in deposition temperature (ordered from 35-60% Al, but disordered from 60-65% Al). The key difference between the 2 studies was the occurrence of an amorphous film from 65-80% Al in the films co-sputtered by Sumiyama. These films crystallized when heated to ~327ºC. Similar results were found by Michaelsen et al. [7] on Ni- Al films co-sputtered onto water and liquid nitrogen cooled sapphire substrates. In the region between a supersaturated solid solution of Ni in Al and the single phase B2 NiAl, they found that films consisted of a mixture of amorphous and crystalline material. The only intermetallic formed between 48 and 88% Al was NiAl.

Co-sputtered films of B2 NiAl tend to be columnar with a small as-deposited grain size (~10 nm without substrate heating) [25], that increases upon annealing (0.5-0.9 μm after 1 hr at 600ºC) [26], and with deviations from stoichiometry. Good adhesion was obtained for NiAl films deposited onto glass, aluminum, nickel, and stainless steel substrates [25]. Wellner et al. [26] measured the stress in NiAl films co-sputtered onto Si substrates after annealing at 600ºC for 1 hour as a function of aluminum concentration. All of the films had large tensile stresses and films with a high aluminum content cracked due to the film stresses, indicating the brittle nature of Al-rich NiAl. Films with less than 50.5% Al had a tensile stress of ~1 GPa, irrespective of composition. The stress was seen to decrease as the aluminum content increased to 50.5% and higher, which was likely due to the cracking observed in these films.

NiAl films deposited from compound targets [27] have similar grain sizes (13-17 nm) as co- sputtered NiAl films (10 nm), and the grain size depends slightly on sputtering power and pressure. Higher pressure and power both lead to increased grain size, but the high-pressure deposited films have voids at the grain boundaries, whereas the low pressure deposited films are more uniform and dense. The transition from a voided film to a dense film as the sputtering

34 pressure decreases is a common occurrence in sputter deposited films [28-29]. The density of sputter deposited platinum films varied from 20.8 g/cm3 to 9 g/cm3 as the sputtering pressure was increased from 10 mTorr to 200 mTorr [29]. The decreased density was due to increased voiding at the grain boundaries. This was also seen in palladium films where the columnar grains were densely packed with no evidence of voids from 5-15 mTorr sputtering pressure, with increased voiding as the pressure was increased from 20-30 mTorr [28]. Not only did the number of voided boundaries increase as the pressure increased, but the width of the voided boundaries also increased. Also associated with increasing sputtering pressure is the transition from compressive to tensile film stresses. This has been demonstrated for several metals including Ti, Ni, Mo, Ta, Al, W, Cr, V, Zr, and Nb [30-32]. The above effects are believed to be due to a “shot-peening” effect that occurs as the sputtering pressure is decreased, which increases the film density while at the same time creating compressive stresses in the film. As the pressure is decreased the mean-free path between collisions increases and the energy of the metal atoms arriving at the substrate increases. The transition from tensile to compressive stresses as the sputtering pressure is decreased also results in a lowering of the resistivity of the deposited films [30-32]. This is a consequence of the decreased scattering due to the reduction of voided grain boundaries.

The above discussion shows that NiAl films may be deposited by several different methods and the morphology of the film depends on the type of deposition as well as the deposition conditions. Two of the above techniques (co-sputtering and annealed multilayers) have been studied to determine the best conditions for fabricating low-resistivity samples. The techniques used in this chapter will be used in future chapters to examine the effects of processing on the thermal fatigue of thin metal lines and the applicability of NiAl films to microelectromechanical systems (MEMS).

II. EXPERIMENTAL PROCEDURE

The materials used for electron beam evaporation of the thin films were all high purity materials available from thin film and vacuum supply companies. Two different sources of evaporation material were used for the aluminum. The first used was a 99.9995% purity aluminum starter source fabricated to fit the crucible liner from Kurt J. Lesker. The starter source had a diameter at the top of 0.802 inches, a 15-degree angle and was 0.375 inches thick. After the useful lifetime of the crucible liner, new source material was obtained in pellet form from Alfa Aesar and all subsequent depositions were done using this material. The aluminum pellets from Alfa Aesar were 99.9999% purity, with a 3.175 mm diameter and 3.175 mm length (stock number 42333-18). The nickel source material was also pellets from Alfa Aesar with the same dimensions as the aluminum and with a purity of 99.995% (stock number 44326-18).

Thin film samples were deposited onto sections of 4-inch diameter oxidized silicon wafers obtained from University Wafer, Inc. Boston, MA. Wafers had a (100) orientation, were boron doped, and had a thickness of 450 μm with a 1 μm oxide layer. Crucible liners were used in the copper hearth for the lower melting point/more reactive metals. Nickel was deposited using a graphite crucible, and aluminum was deposited using a TiB2/BN crucible (Kurt J. Lesker EVCEB-4INT). After loading the source material and substrates into the e-beam system, the

35 chamber was allowed to pump down for 6-8 hours. The chamber pressure prior to deposition was typically ~9.8×10-8 Torr.

Multilayer samples of Ni-Al were fabricated by the sequential deposition of aluminum and nickel from the elemental source materials contained in crucible liners located in a water-cooled copper hearth. The first and last layer deposited was always aluminum, and the thickness of the individual layers was calculated so that the resulting film was 49.5, 50.0, or 50.5 atomic percent aluminum. Films of 50% aluminum were deposited at wavelengths of 20 nm, 25 nm, and 33.3 nm, and the off-stoichiometry films were deposited at a wavelength of 25 nm. Deposition rates were kept between 2.5 Å/sec and 3.5 Å/sec for each layer, and the total thickness of the multilayer samples was approximately 200 nm prior to annealing. After removing the samples from the deposition chamber, the thickness of the film was checked with a profilometer and the samples were sectioned into smaller pieces for annealing and resistivity studies.

Samples for sputtering were loaded into the chamber and then the chamber was allowed to pump for at least 10 hours prior to sample deposition. Thin films were deposited onto oxidized silicon substrates similar to those used for electron beam evaporation at a chamber pressure prior to deposition of less than 10-7 mbar. Prior to sample deposition, the cold trap was filled with liquid nitrogen and the targets were pre-sputtered for at least 5 minutes. Some samples had an additional 10 minute pre-sputter and 30 minute chamber purge prior to the 5 minute pre-sputter to further reduce oxygen contamination in the samples. Multilayer samples of Ni-Al were sputtered from nickel (99.99%) and aluminum (99.9995%) elemental targets at 7 mTorr argon pressure onto rotating substrates. The first and last layers deposited were aluminum and the thickness of the aluminum and nickel layers were varied so as to alter the wavelength or the composition of the sample. Ni-Al multilayers were deposited at wavelengths between 15.4 nm and 30 nm. The total sample thickness prior to annealing was between 150 and 200 nm and was verified by profilometry (Tencor P-10), as was the thickness of the samples after annealing. During sample deposition, the target voltage was adjusted so that the deposition rate was between 1.5 Å/sec and 2 Å/sec for each layer. This corresponded to voltage/current settings of approximately 440 V/480 mA for aluminum and 410 V/250 mA for nickel. These values varied depending on the thickness of the target.

Thin films of NiAl were also deposited by the simultaneous sputtering from two elemental targets rather than the sequential deposition used for the multilayer samples. Films were co- sputtered at argon pressures of 7 mTorr and 1.5 mTorr, which was the lowest pressure at which a plasma could be sustained in the sputtering system. Films were deposited at varying power ratios of Al/Ni and the phases present were determined using grazing angle x-ray diffraction (XRD) with Cu Kα radiation. The power settings that resulted in the correct phase and minimized the resistivity within that phase were those used to deposit the stoichiometric NiAl films. Once the correct power settings were determined, annealing and resistivity studies were carried out on the films.

After deposition, the resistivity of the samples was measured using a 4-point probe set-up (Keithley 2400 SourceMeter) with a current of 10 mA. Nine measurements were recorded and then the samples were sectioned into ~29 mm × ~14 mm pieces for annealing. Short time anneals (t < 10 minutes) were performed in a rapid thermal annealer (RTA) and longer anneals

36 were performed in a tube furnace. In either case, the annealing gas was ultra high purity argon, which was gettered for oxygen prior to entering the annealing chamber. In the RTA, samples were heated at a rate of 50ºC/sec to the annealing temperature, whereas samples annealed in the tube furnace took about 5 minutes to reach the annealing temperature. Samples were annealed cumulatively at temperatures between 400ºC and 1000ºC for sufficient time so that further annealing caused no reduction in the resistivity. In between anneals, nine resistivity measurements were taken using the 4-point probe. Some samples were examined with XRD (Philips MPD, 0.5-2° grazing angle, 1.25 s/step, 0.02° step size, 30 kV, 50 mA, Cu Kα), atomic force microscopy (AFM) in tapping mode, and transmission electron microscopy (TEM: JEOL 2010F). The data from XRD were smoothed using five point adjacent averaging in Origin v6.1052. All values given for the resistivity are the average of nine measurements, with the error bars representing the standard deviation of the nine measurements. In the case where the error bars are smaller than the symbol size they are omitted from the graph.

III. RESULTS AND DISCUSSION

Ni-Al films of varying composition have been deposited using electron beam evaporation (e- beam) and sputtering. A brief overview of the NiAl films and one Ni2Al3 film are given in Table 2, and will be discussed in more detail below. For multilayer films, the majority of the results are for 25 nm wavelength films as early results on e-beam deposited samples showed this was the largest wavelength at which NiAl could be obtained by annealing at 400°C. Further results on sputter deposited NiAl showed that increasing the wavelength slightly to 30 nm still resulted in a completely reacted NiAl film after annealing at 400°C. NiAl fabricated from sputter deposited and e-beam deposited multilayers yielded similar values of resistivity whereas co- sputtered NiAl resulted in higher values, especially when depositing at higher argon pressures. Deviations from stoichiometry also significantly affect the resistivity. Reference to Table 2 shows that an excess of 0.5 at.% Ni results in an increase in resistivity of 105%, whereas an increase of 28% was noted for an excess of Al of 0.5 at.% (for 25 nm wavelength Ni/Al after 4 h at 400°C).

37

Table 2 – Data on some of the Ni-Al films deposited and studied. The “Best Annealing Conditions” are those that resulted in the lowest resistivity. Resistivity Best Annealing Thickness (nm) Material Wavelength (cm) Conditions Description As- % (nm)* 400ºC Temperature Time Annealed§ Lowest† Deposited§ change 4 hrs† (ºC) (min) 214 (4) 204 (2) -5.0 20.0 (21.4) 16.2 1000 0.5 (0.1) E-beam 192 (2) 176 (2) -8.5 25.0 (24.0) 18.3 13.3 700 10 multilayer (0.1) (0.1) 196 (2) 179 (1) -8.4 33.3 (32.6) 24.2 14.2 1000 0.5 (0.1) (0.0) E-beam multilayer 194 (2) 181 (2) -6.7 25.0 (24.5) 37.6 (0.1) 50.5% Ni E-beam multilayer 204 (3) 191 (2) -6.7 25.0 (25.5) 23.5 (0.0) 50.5% Al E-beam multilayer 190 (2) 185 (3) -2.2 25.0 (23.7) 50.5 (0.0) Ni2Al3 Co-sputter 221 (3) 221 (3) 0.0 - 148 34.3 1000 2.5 7 mtorr Ar (1) (0.2) Co-sputter 204 (1) 204 (1) 0.0 - 24.1 18.7 1000 0.5 1 mtorr Ar (0.1) (0.0) 197 (3) 180 (2) -8.5 15.4 (15.1) 16.3 1000 1 (0.0) 193 (5) 173 (2) -10.1 20.0 (19.3) 14.6 1000 0.5 Sputtered (0.0) multilayer 193 (2) 173 (2) -10.2 25.0 (24.1) 15.3 1000 0.5 (0.1) 179 (2) 164 (1) -8.6 30.0 (29.8) 15.5 11.0 800 20 (0.0) (0.1) Sputtered multilayer 289 (3) 263 (2) -9.2 30.0 (28.7) 15.7 12.8 800 0.5 (0.0) (0.1) 0.5% Ag § Numbers in parentheses are the standard deviation of at least 7 measurements. * Values are the expected value for a given deposition. Numbers in parentheses are the actual value calculated from the measured film thickness and number of layers. † Numbers in parentheses represent the standard deviation of nine measurements.

38

A. Ni-Al Films from the Reaction of Ni/Al Multilayers

The formation of NiAl from multilayers of Ni and Al was studied using XRD, 4-point probe resistivity measurements, and cross-sectional imaging by TEM. Both sputtered and e-beam deposited multilayers were studied, and the e-beam samples had a lower resistivity as-deposited and at low annealing temperatures (350-425°C). After elevated temperature annealing (700- 1000°C), no difference in resistivity between the two deposition methods was noted. Figure 5 shows the variation in the as-deposited resistivity as a function of wavelength for samples deposited by sputtering and by e-beam. As can be seen from the data, the sputtered multilayers have a higher resistivity than the e-beam multilayers, which could indicate that some mixing has occurred at the interfaces during sputtering. This could explain some of the differences noted in the literature for reaction temperature and phase formation, as the deposition conditions affect the state of the un-reacted multilayer. Nonetheless, after annealing to 700-1000°C both sets of films achieve similar values of resistivity.

30 sputtered Ni/Al e-beam deposited Ni/Al

25 cm)

20 

15 Resistivity ( Resistivity 10

5 10 15 20 25 30 35 40 Wavelength (nm) Figure 5 – Variation in the as-deposited film resistivity as a function of the wavelength of the Ni/Al multilayer for sputtered and e-beam deposited samples. For clarity the error bars have been omitted, but in all cases lie within the symbols.

TEM images of an as-deposited sputtered multilayer are shown in Figure 6. The images show clearly defined layers of nickel and aluminum, with some waviness within the layers. Results from XRD and electron diffraction indicate that the only phases present are nickel and aluminum.

39

a) b)

25 nm 10 nm

Figure 6 – (a) Low and (b) high magnification TEM images of a 30 nm wavelength as-deposited Ni/Al multilayer film deposited by sputtering. In the images the light layers are aluminum, and the darker layers are nickel. The as- deposited film thickness was 179 ± 2 nm as determined from profilometry. The inset in (a) is the electron diffraction pattern of the film, and the scale bar in the diffraction pattern represents 2 nm-1.

XRD was used to study the formation of Ni-Al intermetallics during the low temperature annealing of the Ni/Al multilayers. Determination of the phases present was complicated by the fact that the majority of peaks for NiAl and Ni2Al3 lie very close to one another. This is why a multilayer with the composition of Ni2Al3 was deposited, so that the peak positions from a Ni2Al3 film could be compared to Ni/Al films with a composition of NiAl. Although the peak positions of NiAl and Ni2Al3 are similar, Ni2Al3 has a considerably higher resistivity than NiAl (Table 2). By combining the results obtained from XRD, TEM, profilometry, and resistivity measurements, the intermetallics present in each film were determined. TEM images from a 30 nm wavelength Ni/Al multilayer after annealing for 4 h at 200ºC are shown in Figure 7. At this temperature the film is partially reacted, and electron diffraction from the film indicated the presence of NiAl3 possibly with some un-reacted Ni and Al. As the images appear to show 3 different phases, the film is believed to be alternating layers of Ni, Al, and NiAl3. The appearance of NiAl3 prior to Ni2Al3 or NiAl is consistent with the literature, where the most aluminum-rich phases form first, followed by the more nickel-rich phases until the equilibrium phase(s) are formed. The TEM images obtained on the multilayer annealed at 200ºC showed a sharp interface between the film and the substrate, and no reaction with the SiO2 was detected.

40

a) b)

10 nm 20 nm

Figure 7 – (a) Bright and (b) dark field images of a 30 nm wavelength Ni/Al multilayer after annealing at 200ºC for 4 hours. The inset in (a) shows the electron diffraction pattern for the film and the scale bar represents 5 nm-1.

Multilayers annealed between 350°C and 425ºC showed only the presence of Ni2Al3 or NiAl. No oxides or silicides from a reaction with the substrate were observed by XRD at annealing temperatures up to 425ºC. Some of the XRD results are shown in Figure 8 for e-beam deposited multilayers with wavelengths of 25 nm (Figure 8a-c) and 33 nm (Figure 8d). From the results in Figure 8, it can be seen that annealed multilayers contain both NiAl and Ni2Al3, and that the amount of Ni2Al3 decreases with increasing annealing time (Figure 8a and b), increasing annealing temperature (Figure 8a and c), and decreasing multilayer wavelength (Figure 8c and d). For the 25 nm wavelength Ni/Al film, no Ni2Al3 was detected after annealing at 400ºC for 2 hours. It appears that the 33 nm wavelength Ni/Al sample is not completely reacted and needs higher temperature anneals to convert the film entirely to NiAl. This observation would explain why the 33 nm wavelength NiAl has a resistivity 32.3% higher than the 25 nm wavelength NiAl after annealing at 400ºC, but only 6.7% higher after annealing at higher temperatures. As the annealing temperature increases, the amount of Ni2Al3 in the 33 nm wavelength film decreases and the resistivity decreases, becoming closer to the resistivity of the annealed 25 nm wavelength film.

41

350 600 a) NiAl b) NiAl 300 Ni2Al3 500 Ni2Al3

250 NiAl/Ni2Al3 NiAl/Ni2Al3 400 200 300 150

Intensity (CPS) Intensity 200 100 (CPS) Intensity

50 100

0 0 40 60 80 100 120 40 60 80 100 120 Angle (2) Angle (2)

500 c) NiAl 500 d) NiAl

Ni2Al3 Ni2Al3

400 NiAl/Ni2Al3 400 NiAl/Ni2Al3

300 300

Intensity (CPS) Intensity 200 200 Intensity (CPS) Intensity

100 100

0 0 40 60 80 100 120 40 60 80 100 120 Angle (2) Angle (2) Figure 8 – Grazing angle XRD scans of annealed Ni/Al multilayer films deposited by e-beam evaporation. The results shown are for a 25 nm wavelength Ni/Al annealed at (a) 400ºC for 1 min, (b) 400ºC for 5 min, and (c) 425ºC for 1 min, and (d) for a 33 nm wavelength Ni/Al annealed at 425ºC for 1 min. The peaks marked with a square could be either NiAl or Ni2Al3.

TEM images of a 30 nm wavelength sputtered Ni/Al film are shown in Figure 9 after annealing for 4 h at 400ºC. The images show that there is complete mixing between the Ni and Al layers with no trace of the layer structure remaining, and electron diffraction confirmed that the film is completely reacted to form NiAl. The interface between the film and the substrate is still sharp with no indication of a reaction between the two. From TEM images it was estimated that the grain size of the 30 nm wavelength Ni/Al films after annealing for 4 h at 400°C was in the range of 20-60 nm. Comparing the data in Figure 8 and Figure 9 indicates that films with a wavelength of 30 nm and less are fully reacted after annealing at 400ºC, whereas films with a wavelength of 33 nm and higher will not be fully reacted after annealing at 400ºC.

42

a) b)

25 nm 20 nm

Figure 9 – Cross-sectional TEM images of a 30 nm wavelength sputtered Ni/Al film annealed at 400ºC for 4 hours. (a) A low magnification image of the film with the diffraction pattern shown in the inset (scale bar 5 nm-1), and (b) a higher magnification image of the film-substrate interface. The images are from the same deposition as those shown in Figure 6 for the as-deposited film. The thickness as determined from profilometry was 164 ± 1 nm after annealing.

Another indication of a higher degree of reaction with increasing time and temperature is the change in film thickness from the as-deposited multilayer. As the nickel and aluminum react to form NiAl (Eq. 1) or Ni2Al3 (Eq. 2) there is a reduction in volume that leads to a reduced film thickness.

Ni  Al  NiAl Eq. 1

2Ni  3Al  Ni2Al3 Eq. 2

By using the molecular weight and density of the materials, Eq. 1 leads to an 11.7% reduction in volume, and Eq. 2 results in a 3.6% reduction in volume assuming all reactions go to completion. This observation is in agreement with the data from Table 2 where the Ni2Al3 multilayer film has a much lower reduction in film thickness (2.2%) after annealing than all of the NiAl multilayer films (5% to 10.2%). Figure 10 shows the thickness of a 25 nm wavelength Ni/Al multilayer film after annealing as a function of annealing time. The data in Figure 10 show that the film thickness decreases as the annealing time and temperature increase. This trend is in agreement with the decreasing amount of Ni2Al3 with increasing time and temperature as noted with XRD, and is what would be expected as any Ni2Al3 that is present in the film is converted to NiAl. The decreasing film thickness and decreasing content of Ni2Al3 with increased annealing time and temperature mirrors the decrease in resistivity with increased annealing time and temperature. The resistivity as a function of annealing time and temperature is shown in Figure 11 for a 25 nm wavelength Ni/Al film.

43

350ºC 195 400ºC 425ºC

190

185

180 Film Thickness Film (nm)

175

0 1 2 3 4 5 Annealing Time (min) Figure 10 – Film thickness of a 25 nm wavelength multilayer as a function of annealing time and temperature.

45 350ºC 400ºC 40 425ºC 4 hours at 400ºC

35

cm) 

30

25 Resistivity ( Resistivity

20

15 0 2 4 6 8 10 12 Annealing Time (min)

Figure 11 – The resistivity of a 25 nm wavelength Ni/Al multilayer film as a function of annealing time and temperature. For clarity the error bars have been omitted, but in all cases lie within the symbols.

44

The decreasing resistivity as a function of increasing annealing time and temperature is likely due to the decreased amount of Ni2Al3 present in the film, and not to any appreciable increase in the grain size with annealing, as is shown in the AFM images of the surface of Ni/Al multilayer films in Figure 12. The images show that there is likely only a slight increase in the grain size during annealing at 400°C. There actually appears to be a larger increase in grain size that occurs when increasing the wavelength from 25 nm to 33 nm than there is when annealing an as- deposited multilayer at 400ºC. The approximate grain sizes for the images in Figure 12 are 15- 40 nm for the as-deposited 25 nm wavelength multilayer, 35-50 nm for the as-deposited 33 nm wavelength multilayer, and 35-60 nm for the 33 nm wavelength multilayer annealed at 400ºC for 10 minutes. Even though the grain size is larger for the larger wavelength sample, the resistivity is higher, which is likely due to the increased amount of Ni2Al3, as mentioned previously. The images and data in Figure 12 also show that the roughness increases slightly when increasing the wavelength from 25 nm to 33 nm and also after annealing. Both increase the roughness by approximately 10%.

45

a) b)

500 nm 500 nm

c)

rms ra Image (nm) (nm) a) 1.183 0.924

b) 1.295 1.027

c) 1.430 1.123

500 nm

Figure 12 – Tapping mode AFM images of the surface of e-beam evaporated Ni/Al multilayer films. (a) An as- deposited 25 nm wavelength Ni/Al film, (b) an as-deposited 33 nm wavelength Ni/Al film, and (c) a 33 nm wavelength Ni/Al film annealed at 400ºC for 10 minutes.

An increasing grain size with increasing wavelength for low temperature annealing could explain the differences noted in the resistivity of the multilayer samples. The resistivity of the multilayer samples after annealing at 400°C is shown in Figure 13 as a function of the annealing time and film wavelength. The data show that as the wavelength decreases the resistivity increases, so there is no benefit to decreasing the wavelength of the multilayer, unless the resistivity of the film is not as important as keeping a low annealing temperature, since smaller wavelengths favor lower reaction temperatures. For annealing at 400°C the resistivity decreases as the wavelength increases until a point is reached at which the film is no longer completely reacted. At this point

46 the resistivity increases as the film is a combination of NiAl and Ni2Al3 (Figure 13b). One benefit that does arise due to a decreasing wavelength is the decreased surface roughness seen in the AFM images in Figure 12. Another observation from the data in Figure 13 is the lower resistivity for the e-beam deposited samples.

a) b) Sputtered 50 50 E-beam deposited

15.4 nm

NiAl NiAl cm)

40 cm) 40 +

  20 nm Ni2Al3 25 nm 30 30

20 nm

Resistivity ( Resistivity

Resistivity ( Resistivity

25 nm 20 20 30 nm

0 5 10 15 20 25 16 20 24 28 32 Anneal Time (min) Multilayer Wavelength (nm) Figure 13 – Resistivity of the Ni/Al multilayer samples after annealing at 400°C. For clarity the error bars have been omitted, but in all cases lie within the symbols. (a) Resistivity as a function of the annealing time at 400°C for samples deposited by e-beam (solid symbols) and sputtering (open symbols). Numbers on the graph refer to the wavelength of the sample. (b) Resistivity as a function of the multilayer wavelength for Ni/Al films annealed at 400°C for 12 min (sputtered samples) or 10 min (e-beam samples).

For all of the multilayer films deposited, the resistivity decreases as the annealing time and temperature increases until a minimum in resistivity is found, and then the resistivity begins to increase, as time and/or temperature are further increased. This transition occurs at annealing temperatures of 600ºC and above, but has not been observed for anneals of up to 4 hours at 400ºC. This trend is shown in Figure 14 for two sputtered films with wavelengths of 15 nm and 30 nm. It is believed that this increase in resistivity is due to reactions between the film and the substrate when the film is annealed for too long or at too high of a temperature. This is confirmed in the TEM images shown in Figure 15 for a film annealed at 600ºC for 1 min and then 1000ºC for 30 s. The images indicate a reaction occurring between the NiAl film and the SiO2 of the substrate. Also evident from the images is an increased grain size in comparison to the film annealed at 400ºC, with most of the grains running from the substrate to the surface for the higher temperature anneal. From diffraction patterns, the products from the reaction of the NiAl film with the SiO2 substrate have tentatively been identified as Al2O3 and Si, with no conversion of the NiAl to other Ni-Al intermetallics. EDS has also indicated that some Ni may be present in the reaction layer.

47

60 35 a) b) 400ºC 600ºC 800ºC 50

30 1000ºC

cm) cm)

40 25





30 20

Resistivity ( Resistivity Resistivity ( Resistivity 20 400ºC 15 600ºC 800ºC 1000ºC 10 10 0 5 10 15 20 0 5 10 15 20 Annealing Time (min) Annealing Time (min) Figure 14 – Results of high temperature anneals on two Ni/Al multilayer sputtered films. The wavelengths of the two films are (a) 15 nm and (b) 30 nm. For clarity the error bars have been omitted, but in all cases lie within the symbols.

25 nm

25 nm

Figure 15 – Cross-sectional TEM images of a 30 nm wavelength sputtered Ni/Al film annealed for 1 minute at 600ºC followed by 30 sec at 1000ºC. This is a section of film deposited in the same run as the images shown in Figure 6 for the as-deposited case and in Figure 9 for a film annealed at 400ºC. This film corresponds to the blue diamond data point shown in Figure 14b.

48

The resistivity for all of the Ni/Al multilayer films after annealing at high temperature (700- 1000°C) is shown in Figure 16. The results show that the lowest resistivity obtained for a given film depends on the initial wavelength of the multilayer: longer wavelength films tend to result in a film with a lower resistivity. There is also a correlation to the resistivity obtained after low temperature annealing (400°C). Films with a low resistivity after annealing at 400°C also obtained lower values of resistivity after high temperature annealing and at lower annealing temperatures. This lower resistivity could be due to grain size or due to contamination of the smaller wavelength films. Since the multilayers required switching of the sources/targets between each layer, the length of time to deposit the smaller wavelength samples was longer due to the increased number of layers. It is possible that more oxygen was incorporated in the smaller wavelength films during growth and in the time between the deposition of one layer and the subsequent layer. For those films where the resistivity was still relatively high after annealing at 700-800°C, anneals were carried out at 1000°C in order to further decrease the resistivity. This increased temperature anneal did decrease the resistivity of the higher resistivity films; however, due to reactions with the substrate, resistivity values for films requiring an anneal at 1000°C were not as low as the resistivity of films that obtained a low resistivity at lower temperature.

18 18 a) b)

1000 C cm)

cm) 16 16

 

14 1000 C 14

12 12

Lowest Resistivity ( LowestResistivity Lowest Resistivity ( LowestResistivity 700-800 C 700-800 C 10 10 12 16 20 24 28 32 36 15 20 25 30 35 40 45 50 Multilayer Wavelength (nm) Resistivity after 400ºC Anneal (cm) Figure 16 – Resistivity of the annealed Ni/Al multilayer films after annealing at high temperature. The annealing temperatures are indicated in the figures. (a) The lowest resistivity obtained after high temperature annealing as a function of the initial multilayer wavelength and (b) as a function of the resistivity after annealing at 400°C. For clarity the error bars have been omitted, but in all cases lie within the symbols.

B. Co-Sputtered Ni-Al Films

Ni-Al films of varying composition were fabricated via co-sputtering by varying the ratio of the power applied to the aluminum and nickel targets. The compositions of the films were determined by a combination of XRD and resistivity measurements. XRD scans for several of the films co-sputtered at 7 mTorr Ar pressure are shown in Figure 17. As can be seen from the data in Figure 17, the as-deposited films were crystalline, which is often not the case for co- sputtered intermetallic films. The resistivity data for the films shown in Figure 17, along with data from all films sputtered at 7 mTorr is shown in Figure 18.

49

120 400

a) Ni2Al3 b) Ni2Al3

100 NiAl/Ni2Al3 Substrate Substrate 300 80

60 200

40

Intensity (CPS) Intensity Intensity (CPS) Intensity 100 20

0 0 15 20 25 30 35 40 45 50 55 15 20 25 30 35 40 45 50 55 Angle (2) Angle (2)

400 200

c) NiAl d) Ni2Al3 NiAl3

Substrate NiAl3/Ni2Al3 Substrate 300 150

200 100

Intensity (CPS) Intensity Intensity (CPS) Intensity 100 50

0 0 30 40 60 80 100 120 20 25 30 35 40 45 50 Angle (2) Angle (2) Figure 17 – Grazing angle XRD scans of Ni-Al films co-sputtered at 7 mTorr Ar pressure. (a) An as-deposited Ni-Al film with ~39.7% Ni, (b) an Ni-Al film with ~39.7% Ni annealed at 600ºC for 10 minutes, (c) a film with ~48.5% Ni annealed for 30 min at 400ºC, and (d) an ~27.0% Ni Ni-Al film annealed at 600ºC for 10 minutes.

50

500 NiAl + Ni Al Ni Al Ni Al 3 2 3 2 3 2 3 NiAl + NiAl

400 cm)

300 

200 Resistivity ( Resistivity 100

As-deposited 600ºC - 1 minute 0 25 30 35 40 45 50 55 60 Composition (at% Ni) Figure 18 – Resistivity as a function of nickel composition for Ni-Al films co-sputtered at 7 mTorr Ar pressure. Also shown are the phases expected to be present according to the equilibrium phase diagram. For clarity the error bars have been omitted, but in all cases lie within the symbols.

The resistivity of the films co-sputtered at 7 mTorr is high even after annealing, and is considerably higher than the annealed multilayer films. As-deposited at 7 mTorr a NiAl film with 48.5% Ni had a resistivity of 318 ± 1 μΩcm, which decreased to 148 ± 1 μΩcm after 4 h at 400°C and 34.3 ± 0.2 μΩcm after 2.5 min at 1000°C. It is believed that the high resistivity is due to the films not being completely dense, and there being some voids or porosity at the grain boundaries. AFM images of one of the co-sputtered films are shown in Figure 19. The images indicate that there may be some areas around the grains that are not completely dense, and annealing at 400°C has little to no effect on this. A film with voids at the grain boundaries may indicate that the films deposited at 7 mTorr contain tensile stresses. This was confirmed in etching studies (see Chapter 6 on MEMS). As with the multilayer samples, the grain size is small: in the images shown the grain size is ~40-65 nm, which is similar to the low temperature annealed Ni/Al multilayers. The surface roughness of the co-sputtered film is more than twice that of the multilayer films, and shows a slight increase with annealing (~3%).

51

a) b)

rms Image (nm)

a) 3.237

125 nm 125 nm b) 3.349

Figure 19 – Tapping mode AFM images of a Ni-Al film co-sputtered at 7 mTorr Ar pressure. The film is ~39.7% Ni and is shown (a) as-deposited, and (b) after 90 minutes annealing at 400ºC.

To increase the density of the co-sputtered films and potentially decrease their resistivity, the Ar sputtering pressure was decreased to ~1.5 mTorr, which was the lowest pressure at which a plasma could be sustained. As was the case for the 7 mTorr co-sputtered films, at 1.5 mTorr it was necessary to vary the ratio of power applied to the two targets to obtain films with the desired stoichiometry. Also similar to the 7 mTorr Ni-Al films, the as-deposited 1.5 mTorr Ni- Al films were crystalline as can be seen from the TEM data for a stoichiometric NiAl film shown in Figure 20. The images show that the co-sputtered films have columnar grains with no evidence of voiding at the grain boundaries, with an estimated grain size of 10-25 nm for the co- sputtered NiAl as-deposited at 1.5 mTorr Ar pressure. This grain size is smaller than that of the films as-deposited at 7 mTorr, but is similar to data from the literature for co-sputtered NiAl and NiAl deposited from compound targets (10-17 nm grain size) [25-26]. As with the samples deposited at 7 mTorr, annealing at 400°C had little effect on the grain size. After annealing at 400°C the grain size was in the range of 15-25 nm, which is smaller than the grain size of the multilayers after annealing at 400°C. This explains the slightly higher resistivity noted for the co-sputtered samples. The indexed rings from the diffraction pattern are shown in Table 3, and are compared to the expected lattice spacings for the crystallographic planes with Miller indices (hkl) for stoichiometric NiAl. As can be seen from the data, the as-deposited film is most likely NiAl.

52

a) b)

100 nm 5 1/nm

Figure 20 – (a) TEM image and (b) diffraction pattern of an as-deposited NiAl film co-sputtered at 1.5 mTorr Ar pressure.

Table 3 – Indexed rings from the diffraction pattern in Figure 20. d Spacing Measured (Å) NiAl Expected (Å, hkl) 2.830 2.8869 (100) 2.021 2.0413 (110) 1.669 1.6671 (111) 1.447 1.4437 (200) 1.292 1.2913 (210) 1.179 1.1788 (211) 1.025 1.0209 (220)

XRD data from the films showed the phases expected according to the phase diagram except for nickel compositions beyond the NiAl single phase field (Ni > 59%). An example is seen in Figure 21, which shows the XRD results for a 78.4% Ni Ni-Al film annealed at 400ºC. At this composition, the expected phases are Ni and Ni3Al, and the expected locations for these phases are also indicated in the figure. As the majority of the peaks associated with Ni3Al are not seen in the XRD scan, the film could be composed of a non-equilibrium solid solution of aluminum in nickel. This is consistent with prior studies on co-deposited films where an increased solubility of Al in Ni was observed compared with the equilibrium phase diagram [20, 22]. The resistivity of this Ni-Al thin films shows an increase in resistivity with annealing rather than a decrease as was seen for all other compositions. The resistivity as a function of composition and annealing is shown for the 1.5 mTorr Ni-Al films in Figure 22.

53

200 125

a) Ni2Al3 NiAl3 b) Substrate NiAl /Ni Al Substrate Unidentified 3 2 3 100 x 150

75

100 x

50

Intensity (CPS) Intensity Intensity (CPS) Intensity 50 x 25

0 0 20 25 30 35 40 45 50 30 40 60 80 100 120 Angle (2) Angle (2) Figure 21 – Grazing angle XRD scans of 2 Ni-Al films co-sputtered at 1.5 mTorr Ar pressure with approximate Ni contents of (a) 32.7 at% and (b) 78.4 at%. Both films were annealed at 400ºC for 90 minutes. The solid and dashed vertical lines show the expected peak locations for Ni and Ni3Al respectively.

200 NiAl3 Ni2Al3 NiAl Ni5Al3 Ni3Al

160 cm)

120



80 Resistivity ( Resistivity 40

As-deposited 600ºC - 1 min 0 20 30 40 50 60 70 80 Composition (at% Ni) Figure 22 – Resistivity data as a function of nickel content for Ni-Al films co-sputtered at 1.5 mtorr Ar pressure. Shown in the darker gray shading are the single phase fields according to the equilibrium phase diagram, and the lighter gray areas show regions that are expected to be 2 phase. For clarity the error bars have been omitted, but in all cases lie within the symbols.

The resistivity for the 1.5 mTorr Ni-Al films is significantly lower than the 7.0 mTorr Ni-Al films for all compositions deposited. As-deposited NiAl films had a resistivity of 318 ± 1 μΩcm when deposited at 7 mTorr, which decreased to 45.5 ± 1.5 μΩcm when the pressure was reduced to 1.5 mTorr. After annealing at 600°C the 7.0 mTorr film has a resistivity that is almost 5 times higher than the 1.5 mTorr film. It is likely that this reduction in resistivity is primarily due to an

54 increase in the density of the films induced by increased “shot-peening” at the lower argon pressure. This observation would suggest that compressive stresses are present in the as- deposited films deposited at 1.5 mTorr. The presence of compressive stress was confirmed in etching studies conducted to establish the viability of NiAl for MEMS (see Chapter 6 on MEMS for details). A comparison of the resistivity of the two sets of films after annealing for 1 min at 600ºC is shown in Figure 23.

500 1.5 mtorr NiAl Ni5Al3 Ni3Al 7.0 mtorr

400 NiAl3 Ni2Al3 cm)

300



200 Resistivity ( Resistivity 100

0 20 30 40 50 60 70 80 Composition (at% Ni) Figure 23 – Comparison of the resistivity after annealing at 600ºC for 1 min of Ni-Al films deposited at 7.0 and 1.5 mTorr. For clarity the error bars have been omitted, but in all cases lie within the symbols.

Another difference between the films deposited at 7 mTorr and 1.5 mTorr is that several of the NiAl3-Ni2Al3 films deposited at 7 mTorr cracked during annealing (see Figure 24). No cracking was observed for any of the films deposited at 1.5 mTorr. Cracking may be due to a combination of effects such as differences in thermal expansion between the two phases and the substrate, the tensile stresses already present due to the higher pressure deposition process, and the increased porosity observed at the grain boundaries of the films deposited at 7 mTorr.

55

a) b)

25 μm 50 μm

Figure 24 – Two Ni-Al films deposited at 7.0 mtorr that cracked during annealing at 400ºC. The film in (a) is ~27% Ni and was annealed for 60 min, and the film in (b) is ~31.4% Ni and was annealed for 1 min. The darker regions are sections of the substrate, which is visible in areas where the film has peeled away.

Rather than the cracking seen in Al-rich films deposited at 7 mTorr, Al-rich films deposited at 1.5 mTorr exhibited significant grain growth after annealing at 400°C and 600ºC. The films are seen to have large grains that can be observed with optical microscopy (see Figure 25); in some cases the grain sizes approach 50 μm. A large increase in grain size explains the significant reduction in resistivity noted for these films upon annealing. The phases observed in this composition region (Ni2Al3, NiAl3) have much lower melting points (1133ºC, 854ºC) than NiAl (1638ºC). The lower melting points may explain the extensive grain growth of these films and the presumably reduced grain growth of NiAl as the grains were too small to observe optically. The annealing temperatures of 400°C and 600°C represent much higher homologous temperatures (annealing temperature/melting temperature) for NiAl3 (0.60, 0.78) and Ni2Al3 (0.48, 0.62) than NiAl (0.35, 0.46). Even with these large grain sizes, the resistivity of these films is still somewhat larger than the NiAl films, which have the lowest as-deposited and annealed resistivity. One exception is the Al-NiAl3 film (23.6% Ni), which had a resistivity of 11.2 ± 0.1 μΩcm after annealing for 30 min at 400°C and 11.0 ± 0.1 μΩcm after 1 min at 600°C. This value is lower than the bulk value listed for NiAl3 of 15 μΩcm [33], which is most likely due to the lower resistivity of aluminum.

56

a) b) c)

10 μm 10 μm 10 μm

d) e) f)

10 μm 50 μm 10 μm

g) h) i)

50 μm 10 μm 5 μm

Figure 25 – Optical microscope images with Nomarski filtering of Ni-Al films co-sputtered at 1.5 mTorr Ar pressure and then annealed at 600ºC. Approximate nickel compositions for the films in atomic percent are: (a) 40.1%, (b) 38.2%, (c) 36.4%, (d) 34.6%, (e) 32.8%, (f) 30.0%, (g) 29.1%, (h) 27.3%, and (i) 23.6%.

IV. CONCLUSIONS

NiAl films have been successfully fabricated from annealed Ni/Al multilayers and from the co- sputtering of nickel and aluminum targets. The method used to deposit the multilayers (electron beam evaporation or sputtering) affects the as-deposited resistivity and the resistivity after annealing at 400°C (e-beam samples have a lower resistivity), but not the resistivity after annealing at higher temperatures (700-1000°C). Multilayers with a wavelength of 30 nm and below were fully reacted to form NiAl after annealing for 2 h at 400°C with a resistivity of 15.5-

57

26.7 μΩcm obtained after a 4 h anneal at 400°C. Lower values of resistivity were obtained for the higher wavelength films, and this reduction may be due to a slight increase in grain size with increasing wavelength. Increasing the annealing temperature to 800°C resulted in a decreased resistivity (lowest: 11.0 ± 0.1 μΩcm), but reactions with the substrate were observed by TEM after annealing at 1000°C. Reaction with the substrate likely occurred at annealing temperatures as low as 600°C and was noted as an increase in resistivity after sufficiently long annealing times.

Co-sputtered NiAl films did not achieve as low of a resistivity (lowest: 18.7 ± 0.0 μΩcm) as the annealed Ni/Al multilayers, but co-sputtering is an attractive deposition method as the films are crystalline as-deposited and do not require additional annealing to form NiAl. Films deposited at 7 mTorr contained tensile stresses and had a high as-deposited resistivity: 228-459 μΩcm for Ni- Al films with nickel contents ranging from 27.0-56.7 at%. Voids present along some of the grain boundaries were likely the reason for the high resistivity seen in the NiAl films deposited at 7 mTorr. Annealing at 400°C caused a decrease in resistivity for most of the films, but the values were still significantly higher than annealed Ni/Al multilayers with values ranging from 85.7-313 μΩcm for nickel contents of 27-48.5%. The higher resistivity is likely due to the fact that annealing did not cause any noticeable decrease in the voiding noticed at grain boundaries. The lowest resistivity obtained for NiAl films deposited at 7 mTorr was 34.3 ± 0.2 μΩcm for a film annealed at 1000°C for 2.5 min. Decreasing the sputtering pressure resulted in films with compressive stresses and a much reduced resistivity. As-deposited the 1.5 mTorr films had a resistivity in the range of 43.2-179 μΩcm for Ni-Al films with Ni contents ranging from 23.6- 78.4 at%. The as-deposited NiAl film had a resistivity (45.45 μΩcm) not much higher than the 7 mTorr deposited film after it was annealed at 1000°C. After annealing for 30 min at 400°C, the resistivity of the 1.5 mTorr films had decreased to 11.2-75.3 μΩcm for Ni contents in the range of 23.6-67.5 at%. The lowest resistivity NiAl film deposited at 1.5 mTorr was 18.7 ± 0.0 μΩcm after annealing for 30 s at 1000°C. The lower resistivity of the films deposited at 1.5 mTorr is due to a more dense microstructure compared to the 7 mTorr films, with no evidence of voids at grain boundaries. In terms of resistivity the annealed multilayer films were superior to the co- sputtered films, but co-sputtering produced NiAl films at significantly lower temperatures: near room temperature for co-sputtering versus 400°C for a 30 nm wavelength Ni/Al multilayer.

58

REFERENCES

[1] P. Bhattacharya, K. Chattopadhyay, P. Bellon, and K.N. Ishihara, Phase evolution and microtwins in the Ni-Al multilayers upon annealing and laser mixing. Materials Science & Engineering A: Structural Materials: Properties, Microstructure and Processing, 2004. 375-377: p. 1277-1284. [2] K.J. Blobaum, D. Van Heerden, A.J. Gavens, and T.P. Weihs, Al/Ni formation reactions: characterization of the metastable Al9Ni2 phase and analysis of its formation. Acta Materialia, 2003. 51: p. 3871-3884. [3] E.G. Colgan and J.W. Mayer, Sequence of phase formation in Ni/Al contrasted with Ni/Si. Materials Research Society Symposium - Proceedings, 1986. 54: p. 121-126. [4] A.S. Edelstein, R.K. Everett, G.Y. Richardson, S.B. Qadri, E.I. Altman, J.C. Foley, and J.H. Perepezko, Intermetallic phase formation during annealing of Al/Ni multilayers. Journal of Applied Physics, 1994. 76(12): p. 7850-7859. [5] T. Jeske and G. Schmitz, Influence of the microstructure on the interreaction of Al/Ni investigated by tomographic atom probe. Materials Science and Engineering A, 2002. 327(1): p. 101-108. [6] T. Jeske, M. Seibt, and G. Schmitz, Microstructural influence on the early stages of interreaction of Al/Ni-investigated by TAP and HREM. Materials Science and Engineering A, 2003. 353(1-2): p. 105-111. [7] C. Michaelsen, G. Lucadamo, and K. Barmak, The early stages of solid-state reactions in Ni/Al multilayer films. Journal of Applied Physics, 1996. 80(12): p. 6689-6698. [8] M. Mozetic, A. Zalar, J. Jagielski, I. Arcon, and P. Panjan, Characterization of NiAl thin layers by AES and EXAFS. Surface and Interface Analysis, 2002. 34: p. 365-368. [9] J. Noro, A.S. Ramos, and M.T. Vieira, Intermetallic phase formation in nanometric Ni/Al multilayer thin films. Intermetallics, 2008. 16(9): p. 1061-1065. [10] U. Rothhaar, H. Oechsner, M. Scheib, and R. Muller, Compositional and structural characterization of temperature-induced solid-state reactions in Al/Ni multilayers. Physical Review B, 2000. 61(2): p. 974-979. [11] F. Takahashi and A.L. Greer, Interfacial reactions in Al/Ni multilayers. Materials Science Forum, 1998. 269-272(pt 2): p. 601-606. [12] A. Ustinov, L. Olikhovska, T. Melnichenko, and A. Shyshkin, Effect of overall composition on thermally induced solid-state transformations in thick EB PVD Al/Ni multilayers. Surface and Coatings Technology, 2008. 202(16): p. 3832-3838. [13] A. Zalar, S. Hofmann, D. Kohl, and P. Panjan, Characterization of intermetallic phases and oxides formed in annealed Ni/Al multilayer structures. Thin Solid Films, 1995. 270(1-2): p. 341-345. [14] A. Tonejc, D. Rocák, and A. Bonefacic, Mechanical and structural properties of Al-Ni alloys rapidly quenched from the melt. Acta Metallurgica, 1971. 19(4): p. 311-316. [15] E.G. Colgan, M. Maenpaa, M. Finetti, and M.-A. Nicolet, Electrical characteristics of thin Ni2Si, NiSi, and NiSi2 layers grown on silicon. Journal of Electronic Materials, 1983. 12(2): p. 413-422. [16] M. Liehr, H. Lefakis, F.K. LeGoues, and G.W. Rubloff, Influence of thin SiO2 interlayers on chemical reaction and microstructure at the Ni/Si(111) interface. Physical Review B (Condensed Matter), 1986. 33(8): p. 5517-25.

59

[17] K.W. Liu and F. Mucklich, Formation and characterization of periodically structured B2-aluminides thin films by laser interference irradiation. Materials Research Society Symposium - Proceedings, 2003. 753: p. 6.5.1-6.5.6. [18] K.A.B. Andersson, H.T.G. Nilsson, and S.E. Karlsson, Composition-dependent structural and morphological changes in nickel-aluminium films prepared by co-evaporation of the pure metals. Thin Solid Films, 1979. 58(2): p. 345-57. [19] H.T.G. Hentzell, B. Andersson, and S.-E. Karlsson, Grain size and growth of Ni-rich Ni- Al alloy films. Acta Metallurgica, 1983. 31(12): p. 2103-2111. [20] H.T.G. Hentzell, B. Andersson, and S.-E. Karlsson, Crystallographic structure of Ni-rich Ni-Al films prepared by co-evaporation. Acta Metallurgica, 1983. 31(7): p. 1131-1140. [21] B.O. Johansson, J.-E. Sundgren, H.T.G. Hentzell, B. Andersson, and S.-E. Karlsson, Differences in nickel-aluminium films prepared by different evaporation methods. Vacuum, 1981. 31(6): p. 247-52. [22] T. Kizuka, N. Mitarai, and N. Tanaka, High-resolution electron microscopy of nanocrystalline Ni-Al alloys: instability of ordered structure and dynamic behaviour of grain boundaries. Journal of Materials Science, 1994. 29(21): p. 5599-606. [23] B. Cantor and R.W. Cahn, Metastable alloy phases by co-sputtering. Acta Metallurgica, 1976. 24: p. 845-852. [24] K. Sumiyama, Y. Hirose, and Y. Nakamura, Magnetic and electrical properties of nonequilibrium Ni-Al alloys produced by vapor quenching. Physica Status Solidi A, 1989. 114: p. 693-704. [25] Y. Ding, Y. Zhang, D.O. Northwood, and A.T. Alpas, PVD NiAl intermetallic coatings: microstructure and mechanical properties. Surface and Coatings Technology, 1997. 94- 95: p. 483-489. [26] P. Wellner, O. Kraft, and E. Arzt, The role of chemical composition for the ductility and microstructure of thin NiAl films. Materials Research Society Symposium - Proceedings, 2002. 695: p. 411-416. [27] Y.-N. Hsu, D.E. Laughlin, and D.N. Lambeth, Effects of sputtering conditions on texture, microstructure and magnetic properties of the CoCrPt/NiAl thin films. Materials Research Society Symposium - Proceedings, 1998. 517: p. 199-204. [28] B.T. Sullivan and R.R. Parsons, A spectroellipsometric investigation of the effect of argon partial pressure on sputtered palladium films. Journal of Vacuum Science & Technology A: Vacuum, Surfaces, and Films, 1987. 5(6): p. 3399-3407. [29] J.H. Thomas, Effect of pressure on dc planar magnetron sputtering of platinum. Journal of Vacuum Science & Technology A: Vacuum, Surfaces, and Films, 2003. 21(3): p. 572- 576. [30] D.W. Hoffmann and J.A. Thornton, The compressive stress transition in Al, V, Zr, Nb and W metal films sputtered at low working pressures. Thin Solid Films, 1977. 45: p. 387-396. [31] D.W. Hoffmann and J.A. Thornton, Internal stresses in sputtered chromium. Thin Solid Films, 1977. 40: p. 355-363. [32] J.A. Thornton and D.W. Hoffmann, Internal stresses in titanium, nickel, molybdenum, and tantalum films deposited by cylindrical magnetron sputtering. Journal of Vacuum Science and Technology, 1977. 14(1): p. 164-168.

60

[33] J.B. Dunlop, G. Gruner, and A.D. Caplin, Dilute intermetallic compounds II: Properties of aluminum rich aluminum-transition metal phases. Journal of Physics F: Metal Physics, 1974. 4: p. 2203-2217.

61

CHAPTER 3

RUTHENIUM ALUMINIDE FILM FABRICATION

62

I. INTRODUCTION

Due to a combination of properties including high strength, good oxidation and corrosion resistance, and metallic electrical conductivity, ordered intermetallic thin films show promise for a wide range of applications. Specifically, the B2 intermetallics NiAl and RuAl are thought to be useful for thin film applications such as MEMS, microelectronics, high-temperature coatings, diffusion barrier layers, and protective layers for mechanical, thermal, and chemical protection [1-3]. One feature that makes NiAl and RuAl particularly attractive materials is that at room temperature both materials exhibit some ductility [3-4], unlike the majority of ordered intermetallic compounds, which tend to fail prior to yielding. NiAl has been well studied, but RuAl has only recently attracted attention due to its high electrical and thermal conductivity, good toughness, high strength, good oxidation resistance, high melting temperature, and excellent corrosion resistance. The bulk resistivity of RuAl is 13-65 cm [5-6], and it has demonstrated resistance to attack from many chemicals such as aqua regia, FeCl3, NaOH, and mixtures of HNO3 and HF [7]. The fracture toughness is expected to be higher than NiAl, with an estimated value of ~10 MPa m [8].

The processing and properties of thin film NiAl have been studied in detail, but there are few studies on thin film RuAl. RuAl has been formed in regions of an Ru/Al bi-layer using laser irradiation [1], and RuAl films have been co-sputtered [9] for use as an underlayer for growing epitaxial FePt. A more recent study [10] fabricated near single phase RuAl from 2 m thick Ru/Al multilayers with an 88 nm wavelength. The authors determined that Al was the dominant diffusing species (as is the case in Ni/Al multilayers [11]), and found that annealing temperatures below 350ºC resulted in the formation of a solid solution. The formation of RuAl6 occurred at 350ºC, and no further phase formation was noted until ~550ºC when RuAl started to form. After annealing to 630ºC the film was mostly RuAl, but some un-reacted Ru was also present.

The aim of this chapter is to fabricate RuAl thin films at low temperatures and to determine the electrical properties. Additional chapters will explore the thermomechanical fatigue and oxidation resistance, and demonstrate the use of NiAl and RuAl thin films in MEMS. In this chapter, RuAl films have been fabricated by co-sputtering and by annealing multilayers of sputtered ruthenium and aluminum. The films were co-sputtered using two different argon pressures, and multilayer films were deposited with varying wavelengths. Intermetallic formation was studied as a function of annealing time, temperature, and deposition method. The films have been characterized by 4-point probe resistivity measurements, x-ray diffraction (XRD), Auger electron spectroscopy (AES), and transmission electron microscopy (TEM), and the results compared to NiAl films prepared using the same equipment and methods. The results will show that RuAl can be fabricated at near room temperature and films with a resistivity as low as 19.3 cm can be obtained after high temperature anneals (1100°C).

II. EXPERIMENTAL PROCEDURE

All samples studied were deposited in a vacuum chamber with a base pressure in the range of 6-9 × 10-8 Torr. During deposition the sample was rotated to increase film uniformity. No substrate heating was applied, although the temperature of the substrate increased due to self-heating and is in the range of 40-70ºC by the end of the deposition. Ru-Al films were deposited using

63 sputtered Ru/Al multilayers and co-sputtering onto (100) oxidized silicon wafers. Ru-Al films formed from the reaction of Ru/Al multilayers had wavelengths from 6.7 to 25 nm, and films were co-sputtered at 1.5 mTorr and 7.0 mTorr argon pressure. Targets for sputtering were aluminum (99.9995%) and ruthenium (99.95%) elemental targets from Kurt J. Lesker (Al) and Williams Advanced Materials (Ru). After deposition, some of the samples were annealed in ultra-high purity (UHP) Ar in a rapid thermal annealer (RTA) for short anneals (1-10 min), or in a tube furnace for longer anneals (1+ h). After annealing, the resistivity and film thickness were measured using the 4-point probe method (Keithley 2400 SourceMeter, 9 measurements at 10 mA current) and mechanical profilometry (Tencor P-10, 5-10 measurements). Additionally, some of the films were studied with grazing angle XRD (Philips MPD, Cu Kα radiation), AES (Physical Electronics 670), and TEM (JEOL 2010F). Cross-sectional samples for TEM were prepared using a Fischione Ion Mill. The magnitude of stress for films under a compressive stress was measured using the buckling method (for details see Chapter 6). All data was plotted and where appropriate curve fit using Origin v6.1052. XRD data were smoothed using 5 point adjacent averaging. All values given for the resistivity are the average of nine measurements, with the error bars representing the standard deviation of the nine measurements. In the case where the error bars are smaller than the symbol size they are omitted from the graph. NiAl films were fabricated using the same equipment and methods, and more details can be found in Chapter 2.

III. RESULTS AND DISCUSSION

In this section we describe the results of microstructural and electrical studies on ruthenium aluminide films deposited by co-sputtering and the annealing of Ru/Al multilayers. These results are compared to those obtained on nickel aluminide films (Chapter 2) that have been deposited and analyzed using the same equipment and procedures. RuAl films can be fabricated using the same methods as those used to produce NiAl, but due to the higher melting point of RuAl the films have a smaller grain size (co-sputtered films) or require higher annealing temperatures/shorter wavelengths (multilayers) to become fully reacted.

A. Co-Sputtered RuAl

RuAl films were prepared via co-sputtering from ruthenium and aluminum targets at 1.5 and 7.0 mTorr Ar pressure. As was done with Ni-Al co-sputtered films, the ratio of the sputtering power was varied until the stoichiometric RuAl intermetallic was formed. The conditions that yielded a RuAl intermetallic were ascertained by XRD (Figure 1), and the sputtering condition that resulted in a stoichiometric RuAl film was established when a minimum in resistivity was obtained (Figure 3).

XRD data for two co-sputtered RuAl films deposited at 1.5 mTorr and 7.0 mTorr are shown in Figure 1, for as-deposited RuAl (1.5 mTorr) and for an annealed RuAl film (7.0 mTorr). Figure 1a shows the data for a 1.5 mTorr sputtered film and corresponds to the resistivity data in Figure 3 for a sputtering power ratio (Al/Ru) of 1.39. The XRD results show that the as-deposited and annealed co-sputtered films are crystalline, as was also the case for co-sputtered NiAl films. No peaks corresponding to oxides, other Ru-Al intermetallics, elemental ruthenium or aluminum, or Ru-Al-Si ternary phases were observed. When comparing the data from the films co-sputtered at

64 low and high pressures it can be seen that the film sputtered at the higher pressure has broader peaks (Figure 1b), even after annealing, which may indicate that the film deposited at 7.0 mTorr has a smaller grain size. Also notable from Figure 1b is the decreased intensity of the first diffraction peak in comparison to the second diffraction peak for the film deposited at 7.0 mTorr. This could indicate that the film is partially disordered (BCC) as the first diffraction peak corresponds to the (100) reflection, which would have a decreased intensity if the film were partially disordered.

1000 1000 100 a) RuAl b) 1.5 mTorr - as-deposited Substrate 7.0 mTorr - annealed 800 800 80

600 600 60

400 400 40 Intensity (CPS) Intensity

200

200 20 (CPS) Intensity Annealed As-Deposited Intensity (CPS) Intensity As-Deposited

0 0 0 20 30 40 50 60 80 100 120 20 24 28 41 42 43 44 45 46 Angle, 2 (º) Angle, 2 (º) Figure 1 – XRD results from two co-sputtered RuAl films deposited at (a) 1.5 mTorr and (b) the first two diffraction peaks from films deposited at 1.5 and 7.0 mTorr argon pressures. The film deposited at 1.5 mTorr is as-deposited and the 7.0 mTorr film is after annealing at 400ºC for 60 min.

The crystallinity of the as-deposited films was also confirmed from TEM analysis, which is shown in the data in Figure 2 and Table 1 for a film deposited at 1.5 mTorr Ar pressure. There was no indication of a second phase or of any amorphous material in the as-deposited films. The TEM images and diffraction data in Figure 2 and Table 1 show that the film as-deposited at 1.5 mTorr is crystalline with a columnar microstructure with grains of ~5-20 nm diameter. All of the diffraction rings can be indexed as RuAl, and the d-spacings for the crystallographic planes with Miller indices (hkl) are all within 0.1% to 0.4% of the values for bulk stoichiometric RuAl [12]. The interface between the film and substrate is sharp, and the film appears to have less surface roughness than the as-deposited and annealed Ru/Al multilayers (see Figure 2b). The morphology of the as-deposited RuAl is similar to the as-deposited NiAl film co-sputtered at 1.5 mTorr (Figure 2e).

65

a) c)

50 nm

b)

100 nm 50 nm

d) e)

10 nm 100 nm

Figure 2 – Cross-sectional TEM images (a, c, d) of an as-deposited co-sputtered RuAl film deposited at 1.5 mTorr Ar pressure. Also shown for comparison are (b) a cross-section of a Ru/Al multilayer and (e) an NiAl film co- sputtered at 1.5 mTorr. The inset in (b) shows the diffraction pattern of the co-sputtered film with the scale bar representing 5 nm-1.

66

Table 1 – Measured d-spacings from the electron diffraction pattern shown in Figure 2c compared to the expected values for stoichiometric RuAl. d Spacing Measured (Å) RuAl Expected (Å, hkl) [12] 2.938 2.950 (100) 2.093 2.086 (110) 1.477 1.475 (200) 1.329 1.319 (210) 1.207 1.204 (211) 1.041 1.043 (220) 0.936 0.933 (310)

The resistivity as a function of annealing time and temperature is shown in Figure 3 for films deposited at 1.5 mTorr, and the data show a clear minimum at the stoichiometric compound with a decrease in resistivity as the annealing temperature increases. For the stoichiometric RuAl, the resistivity drops from 157 ± 4 cm for the as-deposited film to 19.6 ± 0.1 cm after annealing at 1100ºC for 1 min. Even after annealing to 1100ºC, the resistivity of the RuAl film is still ~50% higher than the bulk value, although the value for the resistivity is similar to the lowest achieved in NiAl co-sputtered films (18.7 ± 0.0 cm). The elevated resistivity in both of the co-sputtered films is likely a combination of small grain size and reactions with the substrate as the incorporation of silicon in the film during annealing has been seen in multilayer films annealed at 1000ºC (see below). The high temperatures that were used to increase the grain size and decrease the resistivity led to an increase in the silicon concentration, which in turn increased the resistivity.

300 As-deposited 400°C - 30 min 600°C - 1 min

250 cm)

200



150 Resistivity ( Resistivity

100

Ru-Rich Al-Rich 50 1.0 1.2 1.4 1.6 1.8 2.0 2.2 Sputter Power Ratio (Al/Ru) Figure 3 – Resistivity of RuAl films co-sputtered at 1.5 mTorr Ar pressure. Error bars for the annealed samples are smaller than the symbol size.

67

A comparison of the resistivity of RuAl films deposited at 1.5 and 7.0 mTorr as a function of annealing time at 400ºC is shown in Figure 4a. The resistivity of the RuAl films is considerably higher than NiAl films co-sputtered using the same techniques (Figure 4b); however, both show a decreased resistivity when sputtering at 1.5 mTorr in comparison to 7.0 mTorr. The as- deposited resistivity of the RuAl films deposited at 1.5 mTorr and 7.0 mTorr is 157 ± 4 cm and 1827 ± 9 cm, respectively, in comparison to 45.5 ± 1.5 cm and ~318 cm for similar NiAl films. After annealing for 30 min at 400ºC, the resistivity of the RuAl film deposited at 7.0 mTorr is approximately 13 times higher than the resistivity of the film deposited at 1.5 mTorr (1436 ± 2 cm compared to 107 ± 0 cm). This is most likely due to an increase of film density due to compressive stresses that generally occur as the sputtering pressure is lowered as well as the smaller grain size in the 7.0 mTorr films. It is possible that a further decrease in resistivity could be achieved if a plasma could be sustained at pressures below 1.5 mTorr. The presence of compressive stresses in the 1.5 mTorr films has been confirmed due to the buckling of fabricated free-standing structures (see Chapter 6: “Resistance of Patterned Film to XeF2”). Using optical interferometry to measure the period and amplitude of the buckling, the stress was calculated to be 1.53 ± 0.03 GPa in the 1.5 mTorr as-deposited co-sputtered RuAl films, which is almost double the stress found in as-deposited NiAl films (0.83 ± 0.1 GPa). The stress in the 7.0 mTorr films was not determined, but is expected to be tensile as was seen in similar co-sputtered NiAl films.

1600 500 1.5 mtorr a) NiAl Ni5Al3 Ni3Al 7.0 mtorr 1500 400 NiAl3 Ni2Al3

1400

cm) cm)

300

 

1300

200

150

Resistivity ( Resistivity Resistivity ( Resistivity

100 100 1.5 mtorr Ar pressure 7.0 mtorr Ar pressure b) 50 0 0 10 20 30 40 50 60 20 30 40 50 60 70 80 400°C Anneal Time (min) Composition (at% Ni) Figure 4 – Resistivity of (a) co-sputtered RuAl films in comparison to (b) NiAl co-sputtered films. The NiAl films were annealed for 30 min at 400°C. Error bars are within the symbol size.

B. RuAl Formed from Annealed Ru/Al Multilayers

RuAl films were also produced by annealing alternating layers of ruthenium and aluminum using procedures identical to those used to produce NiAl films from Ni/Al multilayers. A TEM image of an as-deposited 25 nm wavelength multilayer film is shown in Figure 5a. As with the Ni/Al multilayer shown in Figure 5c, the Ru/Al multilayer shows distinct layers of ruthenium and aluminum with some waviness clearly visible within the layers. The diffraction pattern of the as- deposited multilayers (inset to Figure 5a) showed no reaction of the Ru and Al layers prior to annealing. The low temperature formation of RuAl from multilayers like the one shown in

68

Figure 5a was studied using XRD and TEM. Figure 5b shows a TEM image of the 25 nm wavelength film after annealing at 200ºC for 4 hours. The diffraction pattern (inset to Figure 5b) showed that the film was still composed of Ru and Al, but from the images it appears that the interfaces may not be as sharp and the aluminum layers appear to be slightly thinner than before annealing with no noticeable change in the thickness of the Ru layers. The film after annealing at 200ºC for 4 h is in a state of compression (818 MPa +/- 298), which was determined from sections of free-standing film that were buckled (see Chapter 6: “Resistance of Patterned Film to XeF2”). The lack of a reaction at 200ºC is in contrast to a 30 nm wavelength Ni/Al sample (Figure 5d), where a reaction between the layers had already occurred at 200ºC, producing NiAl3 and un-reacted Ni. This is likely due to the fact that RuAl has a higher melting point than NiAl, so it will likely require higher annealing temperatures and/or smaller wavelengths to produce single phase RuAl films.

69

a) b)

5 nm 5 nm

c) d)

10 nm 20 nm

Figure 5 - Cross-sectional TEM images of (a) an as-deposited Ru/Al multilayer with a wavelength of 25 nm, (b) the same Ru/Al film after annealing at 200°C for 4 h, (c) an as-deposited Ni/Al multilayer with a wavelength of 30 nm, and (d) the same Ni/Al film after annealing for 4 h at 200°C. The lighter layers are aluminum and the darker layers are ruthenium/nickel (except in the dark field image in d). The insets show the diffraction patterns of the films and the scale bar length is 5 nm-1.

XRD results are shown in Figure 6 for wavelengths of 25.0, 15.4 and 9.6 nm. The as-deposited films are composed of crystalline aluminum and ruthenium as was confirmed from TEM analyses. After annealing at 400ºC for 5 to 11 hours, the films had partially transformed to RuAl, but some un-reacted aluminum and/or ruthenium is still observed in the XRD data. This could be due to the low annealing temperature, or to an insufficient annealing time. From the difference in intensity of the Al(100)/Ru(100) peak and the RuAl(100) peak it is believed that a

70 larger fraction of RuAl is present in the smaller wavelength films, with the 9.6 nm wavelength sample being almost completely reacted after 5 h at 400ºC. This agreesRuAl with the literature data (110) for Ni/Al multilayers, which shows that it is possible to obtain single phase NiAl at lower temperatures as the wavelength of the film is reduced [11, 13].

600 RuAl (110) Ru Al (200) (002) Ru (101) SiO2 substrate RuAl Al (100) 400 (100) Ru (100)

25.0 nm – no anneal

25.0 nm – annealed

200 Intensity (CPS) Intensity

15.4 nm – annealed

9.6 nm – annealed 0 20 30 40 50 Angle (2) Figure 6 – Grazing angle XRD of multilayer Ru/Al films before and after annealing at 400ºC. The 25.0 and 15.4 nm films were annealed for 11 hours and the 9.6 nm film for 5 hours. Each data set is offset from the previous data set by 100 CPS.

Resistivity data from a 25.0 nm wavelength film shows that reactions may still be occurring in the Ru/Al multilayers even after close to 200 hours of annealing at 400ºC (see Figure 7). Fitting the data in Figure 7 to a second order exponential decay (chosen for the good fit to the data, R2=0.999) suggests that it may take close to 1000 h for the resistivity to reach a constant value. This is significantly longer than for a Ni/Al multilayer fabricated in the same way, for which the reaction is complete in less than 4 hours when annealing at 400ºC.

71

150

30 Ni/Al

27

125 24

cm) cm) 21

 18

100 15 0 1 2 3 4

Resistivity ( Resistivity 75

Ru/Al 50 0 50 100 150 200 Time at 400ºC (h) Figure 7 – Resistivity as a function of annealing time at 400ºC for a 25 nm wavelength Ru/Al multilayer film. The inset shows data from a 30 nm wavelength Ni/Al multilayer film also annealed at 400ºC (axes and units are the same in the main graph and the inset). Error bars are within the symbol size.

A TEM image and diffraction pattern (Figure 8) obtained from the Ru/Al multilayer with a wavelength of 25 nm after annealing for 193 hours at 400ºC confirm that the film is not fully reacted. TEM images show that a layered structure is still present after annealing at 400ºC, and the diffraction pattern shows that the film is composed of Ru and RuAl, with no evidence of un- reacted Al. Prior to annealing, there were 11 layers of Ru and 12 layers of Al (10 full-size and 2 half-size, see Figure 9). After annealing there are 10 layers of RuAl and 11 layers of Ru, which is believed to contain a significant amount of substitutional Al. From the TEM images, it appears that the thickness of the RuAl layers is ~5-8 nm, and the thickness of the Ru(Al) layers is ~14-16 nm. By comparing the thickness of the un-reacted (273 ± 2 nm), partially reacted (261 ± 2 nm), and fully reacted (247 ± 2 nm) films, it is estimated that ~45% of the film has been reacted to form the B2 RuAl phase after annealing for 193 hours at 400ºC. As a decrease in film thickness of ~9.5% occurs for the complete reaction of Ru/Al to RuAl, a 45% reaction would lead to a total of 114 nm of RuAl and 147 nm of Ru(Al). By dividing the values for the total thickness of the layers by the number of layers, an estimated thickness of 11.4 nm and 13.3 nm is obtained for the individual layers of RuAl and Ru(Al), respectively. These values are similar to what was estimated from the TEM images. The value of 11.4 nm of RuAl would indicate that films with wavelengths of less than ~12.5 nm should be completely reacted to form RuAl after annealing at 400°C. As the Ru/Al film with a wavelength of 25 nm is not completely reacted after annealing at 400°C and no un-reacted aluminum was detected, it is likely that the compositions of RuAl and Ru(Al) deviate from those shown in the binary equilibrium phase diagram. This has been seen for Ni/Al multilayers where the first phase to form has been proposed to be an aluminum rich non-equilibrium NiAl containing ~63% Al [14].

72

a) b)

20 nm 20 nm 2 nm-1

Figure 8 – TEM image and diffraction pattern from a 25 nm wavelength Ru/Al multilayer film annealed at 400ºC for 193 hours. The inset in (a) shows a 30 nm wavelength Ni/Al multilayer after annealing for 4 h at 400ºC.

a) b)

Aluminum Ruthenium RuAl Ru(Al)

Figure 9 – Arrangement of (a) the Ru/Al layers prior to annealing (273 ± 2 nm total thickness) and (b) the Ru(Al)/RuAl layers after annealing for 4 h at 400°C (261 ± 2 nm total thickness).

73

Another difference in the two annealed metallic multilayers is that the resistivity of NiAl formed from Ni/Al multilayers is considerably closer to the bulk NiAl value than the Ru-RuAl film is to the RuAl bulk value. The resistivity of the NiAl film is 1.8 times the bulk value when annealed at 400ºC, compared to the Ru/RuAl film, which has a resistivity of 4.8 times the bulk value when annealed at 400ºC. This is likely due to the incomplete reaction of Ru with Al at 400ºC leading to the layered structure of Ru and RuAl.

Resistivity data for Ru/Al multilayers of various wavelengths after annealing at 400ºC are shown in Figure 10. The resistivity of the as-deposited films (not shown) decreased from 315 ± 3 cm, to 144 ± 1 cm, and then to 53.9 ± 0.4 cm and 17.3 ± 0.1 cm as the wavelength was increased from 6.7 nm to 9.6, 15.4, and finally to 25.0 nm. As would be expected, the increased number of layers (decreased layer thickness) in the shorter wavelength films results in a higher resistivity due to the increased number of interfaces. During annealing of the multilayers there may be several processes occurring simultaneously that can affect the resistivity: (1) grain growth and the elimination of defects, which would lead to a decrease in resistivity, (2) mixing of Ru and Al without a phase transformation, which would lead to an increase in resistivity, (3) the creation of additional interfaces as Ru/Al reacts to form RuAl, which would lead to an increase in resistivity, (4) the elimination of interfaces as the reaction nears completion, which would lead to a decrease in resistivity, and (5) the creation of a higher resistivity phase, RuAl, from two lower resistivity phases, Ru and Al, which would also lead to an increase in resistivity. Depending on which processes are occurring, the resistivity may increase or decrease. The effect of (1) can be demonstrated by annealing a single layer of Al or Ru. This leads to a decrease in resistivity of ~7% for Al, and ~37% for Ru after annealing at 400ºC for 5 min and 4 h, respectively. During the low temperature annealing of Ru/Al multilayers, the 6.7 nm wavelength film showed a decrease in resistivity of ~62% after a 40 h anneal at 400ºC (when compared to the as-deposited film). It is suspected that this film is completely reacted to form RuAl as the resistivity has decreased significantly and is similar to the co-sputtered film annealed at 400ºC for 4 h, which is known to be composed entirely of RuAl (121 ± 0 cm: 6.7 nm wavelength, 103 ± 0 cm: co-sputtered). Also, the wavelength of the film is approximately the same as the thickness of the layers of RuAl formed in the reaction of the 25 nm wavelength film. The resistivity of the reacted film is considerably higher than the bulk value, which is likely due to a small grain size in the film. The diffusion distances necessary to create RuAl in the 6.7 nm wavelength multilayer are relatively short, and as the annealing temperature is low in comparison to the melting point (Tm) of RuAl (T/Tm = 0.29), it is unlikely that there is significant grain growth after RuAl is formed. The 9.6 nm wavelength film is also believed to be fully reacted after annealing at 400ºC for 40 h. The resistivity after 40 h at 400°C is lower than the as-deposited resistivity (by 51%), which would indicate that either (1) grain growth and the elimination of defects, or (4) the elimination of interfaces as the reaction nears completion are occurring. A small amount of un-reacted Al/Ru was detected after a 5 h anneal, but after annealing for 40 h at 400ºC the resistivity is significantly lower than at 5 h (33%). The resistivity of the film with a wavelength of 9.6 nm was 70.5 ± 0.1 μΩcm after annealing for 40 h at 400°C. This value is significantly higher than the bulk value, but is lower than the value for the film with a wavelength of 6.7 nm. This is consistent with the results obtained on Ni/Al films: for films that were completely reacted, a lower resistivity was obtained for films with a larger wavelength.

74

210 25.0 nm wavelength 15.4 nm wavelength 180 9.6 nm wavelength 6.7 nm wavelength

150

cm)  120

90 Resistivity ( Resistivity

60

30 0 5 10 15 20 25 30 400°C Anneal Time (h) Figure 10 – Resistivity as a function of annealing time at 400ºC for Ru/Al multilayer films with wavelengths between 6.7 and 25.0 nm. Error bars are within the size of the symbol.

The 15.4 nm and 25 nm films are not fully reacted after annealing at 400ºC, and both have a higher resistivity than the as-deposited films. The extent of the reaction at 400ºC for the 15.4 nm wavelength film can be seen in the Auger depth profiles of the film before and after annealing (Figure 11). Figure 11 shows that there is a clear modulation in composition prior to annealing and a slight modulation after 11 h at 400ºC. Due to the incomplete reaction and interface scattering, the resistivity of the films with wavelengths of 15.4 nm and 25 nm are higher than the bulk value for RuAl. Although the Auger data indicates that mixing may have occurred in the as-deposited film, this could be an artifact of the resolution of the instrument and the depth profiling.

75

80 80 a) Ru b) Ru Al Al O O Si Ru Si 60 Ru 60 O O

40 40

Al Si Si Al

20 20

Concentration (a.u.) Concentration Concentration (a.u.) Concentration

0 0 0 150 300 450 600 750 900 0 200 400 600 800 1000 Sputter Time (s) Sputter Time (s) Figure 11 – Auger depth profiles of a 15.4 nm Ru/Al multilayer (a) before and (b) after annealing for 11 h at 400ºC.

From the data in Figure 12 it can be seen that the resistivity of the 15.4 nm wavelength film decreases rapidly with annealing temperature until ~500-550ºC. At this point, the resistivity drops much less rapidly with increasing annealing temperature. It is believed that in the range of 500-550ºC, the 15.4 nm wavelength film has been completely transformed to RuAl, and the slow decrease in resistivity beyond this temperature is due to grain growth. The 25 nm wavelength film is likely completely reacted at a temperature between 900ºC and 1000°C. This is the point at which the resistivity of the 25 nm wavelength film becomes less than the 15.4 nm wavelength film. The sample is definitely single phase RuAl after 2 min at 1000ºC as confirmed by cross- sectional TEM (Figure 14), which shows that the 25 nm wavelength film annealed at 1000ºC for 2 min is completely reacted to form a single phase RuAl intermetallic film.

76

250 30 min anneals

200

cm) 150



100 Resistivity ( Resistivity 50

0 400 500 600 700 800 Temperature (ºC) Figure 12 – Resistivity as a function of annealing temperature for a 15.4 nm wavelength Ru/Al multilayer. Each anneal was 30 min. Error bars are not shown, but in all cases lie within the size of the symbol.

60 6.7 nm 9.6 nm 50 15.4 nm 25 nm

40

cm)  30

20 Resistivity ( Resistivity

10

0 800ºC-7 min 900ºC-7 min 1000ºC-2 min 1100ºC-2 min Annealing Conditions Figure 13 – Resistivity of the Ru/Al multilayers after annealing at temperatures between 800°C and 1100°C. In all cases, the standard deviation in the measurement of the resistivity was less than 0.2 μΩcm for 9 measurements.

77 a) b)

(111)Si

RuAl

Reaction layer (220)Si 50 nm 50 nm SiO2 Figure 14 – (a) Cross-sectional TEM image of a 25 nm wavelength Ru/Al film annealed at 1000ºC for 2 min. (b) An electron diffraction pattern of the reaction layer shown in (a). The inset in (a) shows a 30 nm wavelength Ni/Al multilayer that was also annealed at 1000ºC.

The lowest resistivity obtained from reacted Ru/Al multilayers was 19.3 ± 0.1 cm, which is almost identical to the resistivity of the co-sputtered RuAl film (19.6 ± 0.1 cm). The reason for the higher resistivity of the RuAl film in comparison to the bulk resistivity (~50-60% higher than bulk RuAl) may be due to the RuAl film containing some silicon (see Figure 15). The lowest resistivity obtained from annealed Ni/Al multilayers was 11.0 ± 0.1 cm, which is only 10% higher than the bulk value of NiAl. The lower thin film resistivity for NiAl is due to the fact that the Ni/Al multilayers reached a minimum in resistivity at a lower annealing temperature (800ºC) due to the lower melting temperature (NiAl-1638ºC, RuAl-2060ºC), so reactions with the substrate were minimized. The images in Figure 14 show that the film appears completely reacted after annealing at 1000ºC and none of the initial layer structure depicted in Figure 5 is observed. Electron diffraction from one of the grains (not shown) indicates that the film is RuAl. Some of the grains reach from the substrate to the surface, but not as many as for a Ni/Al film annealed at 1000ºC. In the NiAl film the majority of the grains extend through the thickness and extend laterally for 100-500 m. As with the Ni/Al film annealed at 1000ºC, there is a reaction layer that forms between the RuAl film and the substrate, although the morphology in the case of RuAl is more planar than in the case of NiAl. There is also no increase in resistivity noted with the occurrence of a reaction layer, in contrast to what was seen for NiAl, where for annealing temperatures of 600ºC and above the resistivity initially decreased with annealing time and then began to increase with further annealing. This did not happen for RuAl films although it is possible that the anneals were not carried out for sufficient time for this to occur. A fast-Fourier transform of a TEM image of the reaction layer shows the presence of crystalline silicon and possibly Al, although a small amount of ruthenium was also indicated by EDS (Figure 15). The RuAl film also contains some silicon as an impurity, but the film is still single phase RuAl. Auger depth profiles showed no indication of a reaction with the substrate for annealing at 400ºC and 800ºC.

78

Ru 2000 Ru Ru Si Al Si 1500 1000 Si

Counts 1000 Al Counts 500

500 O Ru Cu Ru Ru Ru Ru Ru Cu Ru Cu Ru Ru Ru Cu Ru 0 50 nm 0 0 Energy (keV) 5 10 20 0 Energy (keV) 305 10 40 15 20 Energy (keV) Energy (keV) Figure 15 – A cross-sectional TEM image of a 25 nm wavelength Ru/Al film annealed at 1000ºC for 2 minutes, and the EDS spectra taken from one of the grains (left) and from the reaction layer (right).

A comparison of the resistivity of a 25 nm wavelength annealed Ru/Al multilayer film and the 1.5 mTorr co-sputtered RuAl is shown in Figure 16. At low annealing temperatures (400ºC), the multilayer film anneals more slowly than the co-sputtered film, but does achieve a lower resistivity, although at this point the film is a combination of Ru and RuAl. When the annealing temperature is increased to 800ºC, the resistivity as a function of annealing time for the two films shows very similar trends, with a slightly higher resistivity for the multilayer film. At 1100ºC, the data for the two films are almost identical.

79

150

multilayer after 1 hr at 400 C

120

400 C

cm) 90 

60 multilayer after 193 hr at 400 C Resistivity ( Resistivity 30 800 C

1100 C Co-sputter - 1.5 mtorr Ar Ru/Al multilayer 0 0 5 10 15 20 25 30 35 Annealing Time (min)

Figure 16 – Comparison of the resistivity of multilayer and co-sputtered RuAl films after annealing at 400, 800, and 1100ºC. The Ru/Al multilayer had a wavelength of 25 nm. Error bars are not shown, but are smaller than the symbol size.

IV. CONCLUSIONS

RuAl films have been fabricated using co-sputtering, and by annealing multilayers of Ru and Al. Films co-sputtered at 1.5 mTorr Ar pressure are crystalline as-deposited with small columnar grains having a diameter of ~5-20 nm. The resistivity of the as-deposited film was 157 ± 4 cm, which decreased to a minimum of 19.6 ± 0.1 cm after annealing at 1100ºC for 1 min. For the complete reaction of Ru/Al multilayer films it was necessary to go to higher temperatures, or shorter wavelengths compared to Ni/Al multilayers. Shorter wavelength films (6.7 and 9.6 nm) are likely fully reacted after annealing at 400ºC for 40 h with a resistivity of 121 ± 0 cm and 71.6 ± 0.1 cm, respectively. Longer wavelength films (15.4 and 25.0 nm) are not fully reacted at 400ºC, although the Ru-RuAl films have a resistivity lower than the single- phase co-sputtered RuAl after annealing at 400ºC. These films are likely single phase after annealing at about 500ºC and 900-1000ºC for the 15.4 nm and 25 nm wavelength samples, respectively. The lowest resistivity in RuAl films is somewhat higher than in similar NiAl films (19.3 ± 0.1 cm in comparison to 11.0 ± 0.1 cm). In order to fabricate RuAl films, co- sputtering is the best method as no further anneals are necessary to create single phase crystalline films. However, both methods of film fabrication require high temperature anneals in order to decrease the resistivity.

80

REFERENCES

[1] K.W. Liu and F. Mucklich, Formation and characterization of periodically structured B2-aluminides thin films by laser interference irradiation. Materials Research Society Symposium - Proceedings, 2003. 753: p. 6.5.1-6.5.6. [2] D.B. Miracle, Physical and mechanical properties of NiAl. Acta Metallurgica et Materialia, 1993. 41(3): p. 649-684. [3] R.D. Noebe, R.R. Bowman, and M.V. Nathal, Physical and mechanical properties of the B2 compound NiAl. International Materials Reviews, 1993. 38(4): p. 193-232. [4] R.L. Fleischer, C.L. Briant, and R.D. Field, Tough, ductile high-temperature intermetallic compounds: results of a four-year survey. Materials Research Society Symposium - Proceedings, 1991. 213: p. 463-474. [5] S.A. Anderson and C.I. Lang, Thermal Conductivity of Ruthenium Aluminide (RuAl). Scripta Materialia, 1998. 38(3): p. 493-497. [6] E.G. Smith and C.I. Lang, High temperature resistivity and thermo-EMF of RuAl. Scripta Metallurgica et Materialia, 1995. 33(8): p. 1225-1229. [7] I.M. Wolff, Toward a better understanding of ruthenium aluminide. JOM, 1997. 49(1): p. 34-39. [8] T. Reynolds and D. Johnson. Solidification processing and fracture behavior of RuAl- based alloys. in 2004 MRS Fall Meeting, November 29, 2004 - December 1, 2004. 2005. Boston, MA, United states: Materials Research Society. [9] W.K. Shen, J.H. Judy, and J.-P. Wang. In situ epitaxial growth of ordered FePt (001) films with ultra small and uniform grain size using a RuAl underlayer. in 49th Annual Conference on Magnetism and Magnetic Materials. 2005. Jacksonville, Florida (USA): AIP. [10] K. Woll, R. Chinnam, and F. Mucklich, Thin-film synthesis and cyclic oxidation behavior of B2-RuAl. Materials Research Society Symposium - Proceedings, 2009. 1128. [11] A.S. Edelstein, R.K. Everett, G.Y. Richardson, S.B. Qadri, E.I. Altman, J.C. Foley, and J.H. Perepezko, Intermetallic phase formation during annealing of Al/Ni multilayers. Journal of Applied Physics, 1994. 76(12): p. 7850-7859. [12] Powder Diffraction File. JCPDS. card no. 29-1404. [13] J. Noro, A.S. Ramos, and M.T. Vieira, Intermetallic phase formation in nanometric Ni/Al multilayer thin films. Intermetallics, 2008. 16(9): p. 1061-1065. [14] C. Michaelsen, G. Lucadamo, and K. Barmak, The early stages of solid-state reactions in Ni/Al multilayer films. Journal of Applied Physics, 1996. 80(12): p. 6689-6698.

81

CHAPTER 4

OXIDATION OF RUAL AND NIAL THIN FILMS

82

I. INTRODUCTION

Ordered intermetallic thin films show promise as materials for microelectromechanical systems (MEMS) due to a combination of properties including high strength, good oxidation and corrosion resistance, and metallic electrical conductivity. Texas Instruments has successfully demonstrated the use of intermetallics for micromirrors [1] and amorphous TiAl3 for their digital micromirror devices [2], and we have fabricated thermal actuators and resonators from NiAl and RuAl thin films (see Chapter 6). For the use of materials in optical applications (such as the aforementioned micromirrors) good surface quality (low roughness) is necessary, and if the application requires elevated temperatures or harsh environments, the evolution of the surface with increasing temperature and the oxidation and corrosion resistance become important. More general requirements for MEMS include a low electrical resistivity, low stress level/stress gradient, a fracture toughness greater than 1 MPa m , ease of micromachining, thermal and mechanical stability, and low processing temperatures (T < 450ºC) to be compatible with complementary metal oxide semiconductor (CMOS) processing [3]. It is the last criteria of low temperature processing that current materials systems (Si/Ge/SiGe, and Si/SiO2/Si) are unable to easily overcome and still be low in stress or have a low stress gradient. Other MEMS materials include metals such as gold, aluminum, and nickel, but pure metals are often prone to fatigue, plastic deformation, and creep [3]. Nickel also tends to be mechanically and thermally unstable and is not oxidation resistant.

NiAl and RuAl were chosen as suitable MEMS materials due to a combination of low resistivity, high strength, moderate toughness, good oxidation and corrosion resistance [4-5], and in the case of NiAl, demonstrated low temperature thin film processing. Additionally, both exhibit some ductility [6-7], unlike the majority of intermetallic compounds, and both the strength and ductility of NiAl can be increased without a loss of conductivity by small (up to 1 at%) alloying additions of silver [8]. The room temperature fracture toughness of NiAl is 4-14 MPa m , which increases to 10-50 as the temperature increases to 350-400ºC [4, 9]. These values are higher than for silicon, which has a room temperature fracture toughness of 0.8-1.1 MPa m [10-13] that shows a moderate increase to 3.3 at 925ºC. There are no values cited in the literature for SiGe, but the fracture toughness should fall between the values of Si and Ge (Ge fracture toughness: 0.60-0.64 [14-15]). The processing of NiAl thin films has been well studied, and the literature shows that NiAl thin films can be fabricated easily at temperatures as low as 77ºC [16-23]. We have shown that both NiAl and RuAl can be deposited near room temperature (see Chapters 2 and 3). Additionally, the film stresses are easily controlled in co-sputtered films by adjusting the sputtering pressure. NiAl films (thin film: 24- 54 cm) [22, 24-26] and bulk RuAl (bulk: 13-65 cm) [27-28] are more conductive and more oxidation resistant than Si (bulk resistivity: 100 mcm) [29] and SiGe (resistivity: 0.5-75 mcm) [30-31]. NiAl and RuAl are also resistant to various forms of chemical attack, making them useful in severe environments. RuAl is resistant to aqua regia, FeCl3, NaOH, and mixtures of HNO3 and HF [5]. The results of a study on the use of RuAl and NiAl in MEMS will be discussed in Chapter 6, but an example of two thermal actuators fabricated from NiAl are shown in Figure 1.

83

100 μm

Figure 1 – Two thermal actuators fabricated from an as-deposited co-sputtered NiAl thin film. The Si substrate was patterned with conventional photolithography, and after deposition free-standing structures were created by etching the Si with XeF2 gas.

Even though the B2 aluminides, such as NiAl and RuAl, are being considered for many thin film applications (MEMS; microelectronics; high-temperature coatings; diffusion barrier; and layers for mechanical, thermal, and chemical protection [4-5, 32]), relatively few studies have been carried out on characterizing the elevated temperature performance of B2 intermetallic thin films. Studying the oxidation resistance of thin films is important due to the fact that bulk and thin film samples often show different oxidation rates, as was seen for the case of a Ni-8Cr-3.5Al alloy that showed a decreased oxidation rate in thin film form in comparison to the bulk [33]. In some cases reducing the grain size has also been shown to increase the oxidation resistance [34], which could be beneficial for thin film applications where smaller grain sizes are typically encountered.

Although this study is primarily concerned with the surface morphology and resistance as a function of oxidation temperature, a brief review of the oxidation of NiAl and RuAl is first presented. The oxides present after oxidation can vary with processing as well as oxidizing conditions. Changes in the oxide phases present can lead to changes in the film surface morphology as well as the electrical properties, both of which are of interest in this study.

In general, the majority of aluminides with high aluminum content are considered to be oxidation resistant as they form a slow-growing, protective Al2O3 layer during oxidation. The oxidation resistance of bulk NiAl has been studied in depth from room temperature up to 1400ºC, but little has been published on the oxidation of single-phase NiAl thin films, and all of the studies have been on relatively thick films (tens of microns) [35-36]. The high temperature oxidation of NiAl has been studied in detail. Bulk polycrystalline NiAl forms metastable -Al2O3 (a smooth and dense oxide) at oxidation temperatures between 700ºC and 850ºC, which transforms to another metastable phase, -Al2O3, at higher temperatures (an oxide with whisker and blade-like morphologies), and then to the stable phase, -Al2O3 (an oxide with ridges at the grain boundaries; forming a cell-like appearance) [37]. Below the external oxide scale is a region of

84

Ni-rich NiAl. The data at lower temperatures are not as complete, but Ohtsu et al. [38] determined that at oxidation temperatures of 200ºC and 500ºC, three different oxides were present, a top layer of NiAl2O4 containing a small amount of NiO, and a bottom layer of Al2O3. The morphology of the oxide layers was not reported.

The oxidation of single crystal NiAl showed some differences compared to polycrystalline materials [39]. NiAl single crystals showed the formation of NiAl2O4, -Al2O3, and -Al2O3 in the early stages of oxidation at 800ºC, whereas polycrystalline samples produced only a -Al2O3 scale between 700ºC and 850ºC. Growth rates and oxide morphology (mostly a platelet-like oxide morphology; differences in orientation and aspect ratio) were also seen to vary with crystal orientation. The oxidation of NiAl thick films and nanoparticles also differs slightly compared to polycrystalline bulk NiAl. A NiAl cladding layer (~15 m thick) oxidized at 1000ºC [35] showed a two-layer oxide (an outer layer of NiO and an inner Al2O3 layer; no description of the oxide morphology), and NiAl nanoparticles oxidized at 600ºC and 800ºC showed the presence of -alumina as well as NiO (no description of the oxide morphology) [40]. The exact scale composition obtained in a given study appears to be a function of processing as well as the time and temperature of oxidation.

The literature on the oxidation of RuAl is not as comprehensive as that on NiAl, and only the temperature range from 1000ºC to 1400ºC has been studied [41-44]. McKee et al. [43] determined that the oxidation resistance of RuAl in the range from 1000ºC to 1400ºC was not as good as for NiAl, as the rate of formation of alumina on RuAl was higher than on NiAl. As with NiAl, RuAl forms an external alumina scale during high temperature oxidation. The external alumina scale has been described as having a needlelike or bladelike morphology [41, 44]. Unlike NiAl, a second phase is formed below the alumina scale, -Ru, which occurs due to the depletion of aluminum beneath the oxide scale, and the limited solubility range of single phase RuAl. There are few published studies on the formation of RuAl thin films [32, 45], and no studies have been performed on the oxidation of RuAl thin films.

In this paper we establish the oxidation resistance of NiAl and RuAl thin films in flowing oxygen in terms of the surface morphology and electrical resistance, and compare the results to those from aluminum, nickel, and ruthenium films. The results will show that NiAl and RuAl can be used in an intermediate temperature range for at least 1 h (up to ~500-550ºC) when the surface condition is of paramount importance, and to higher temperatures for at least 1 h (up to ~800- 850ºC) when the surface smoothness is not as important, but when electrical conductivity needs to be maintained. In comparison, maximum service temperatures for durations of 1 h for Ni and Ru are approximately 500ºC (loss of conductivity) and 700ºC (possible oxide volatilization), respectively. Although this study is primarily concerned with the surface quality (qualitatively, the smoothness and reflectivity) and electrical performance, some discussion on the oxide composition will be tied to observations from the literature combined with Auger electron depth profiles of the constituent elements.

85

II. EXPERIMENTAL PROCEDURE

All samples studied were deposited on oxidized (100) silicon wafers in a vacuum chamber with a base pressure of 6-9×10-8 Torr. Prior to oxidation, all samples were annealed in a tube furnace in ultra-high purity (UHP) Ar for 4 h at 400ºC, and then the resistivity and film thickness were measured using the 4-point probe method (Keithley 2400 SourceMeter, 9 measurements, 10 mA current) and mechanical profilometry (Tencor P-10, at least 7 measurements). Additionally, some samples were studied with grazing angle x-ray diffraction (XRD; Philips MPD), tapping mode atomic force microscopy (AFM; Digital Instruments Dimension 3100), and transmission electron microscopy (TEM; JEOL 2010F). Cross-sectional samples for TEM were prepared using a Fischione Ion Mill. The oxidation of the thin films was examined by annealing in 100 sccm UHP oxygen from 100ºC to 850ºC in a tube furnace. The films were annealed cumulatively in 50ºC increments from 100ºC until the films were no longer conductive (resistivity greater than ~20 cm), or until the temperature reached 850ºC. Annealing at each temperature was performed for 1 h. After each 1 h anneal the samples were cooled to room temperature in oxygen. The resistance was then measured and the surface of the film was examined with light microscopy. Depth profiles after oxidation were obtained on RuAl (800ºC and 850ºC) and NiAl (850ºC) using Auger electron spectroscopy (AES; Physical Electronics 670). All values given for the resistivity (or normalized sheet resistance) are the average of nine measurements, with the error bars representing the standard deviation of the nine measurements. In the case where the error bars are smaller than the symbol size they are omitted from the graph.

Ni-Al and RuAl intermetallic thin films, as well as Ni, Ru, and Al films were studied. The Ni-Al intermetallic films were fabricated using annealed multilayers of Ni/Al that were deposited by electron beam evaporation (e-beam) or sputtering, as well as co-sputtering of elemental targets. Source materials for e-beam deposition were aluminum (99.9999% purity) and nickel pellets (99.995% purity) from Alfa Aesar. Targets for sputtering were nickel (99.99%), aluminum (99.9995%) and ruthenium (99.95%) elemental targets from Kurt J. Lesker Co. (Al, Ni) and Williams Advanced Materials (Ru). During deposition the sample was rotated to improve film uniformity. No substrate heating was applied, although the temperature of the substrate increased due to self-heating and is in the range of 40-70ºC by the end of the deposition. The thickness of the deposited films was 164-293 ± 2 nm, and the resistivity prior to oxidation was 15.7-28.6 ± 0.1 cm (NiAl), 50.5 ± 0.0 cm (Ni2Al3), 103 ± 0 cm (RuAl), 4.11 ± 0.00 cm (Al), 10.5 ± 0.1 cm (Ni), and 17.2 ± 0.0 cm (Ru). More details on the NiAl and RuAl films can be found in Chapters 2 and 3, respectively.

As the thickness of the samples prior to oxidation varied from 164 nm to 293 nm, electrical data from the oxidized samples are not presented as sheet resistance, but as a normalized sheet resistance (sheet resistance × initial film thickness), in order to account for the varying film thickness. The units used are the same as resistivity (cm), but this is not a true resistivity since the thickness of the conductive portion of the film changes as the film oxidizes. NiAl films formed from the reaction of multilayers had a nominal bi-layer thickness (wavelength) prior to annealing of 20, 25, 30, and 33 nm. Additionally, a Ni2Al3 film was fabricated from an annealed Ni/Al multilayer (25 nm wavelength), and a 0.5 at% silver-doped NiAl film was fabricated from a Ni/Al/Ag multilayer (30 nm wavelength). After annealing for 4 h, XRD indicated that samples with a 25 nm wavelength were single phase (within the detection limits), and TEM on the 30 nm

86 wavelength film confirmed that this sample was single phase NiAl with grains ranging in size from ~20-60 nm (Figure 2a). a) b)

50 nm 50 nm

Figure 2 – TEM micrographs of an (a) NiAl film fabricated from an annealed (400ºC for 4 h) Ni/Al multilayer with a 30 nm wavelength deposited by sputtering and (b) an as-deposited RuAl film deposited by co-sputtering.

The co-sputtered NiAl film was deposited using an Ar pressure of 1.5 mtorr, and was determined from XRD and TEM analysis to be a single phase of B2 NiAl in the as-deposited and annealed conditions. The as-deposited co-sputtered films exhibited columnar grains similar to those shown in Figure 2b for the RuAl co-sputtered film. The grain size of the annealed co-sputtered NiAl was not measured, but is expected to be similar to the annealed multilayer films as the resistivity of the co-sputtered film (24.1 ± 0.1 cm) was in the same range as that of the multilayer NiAl films (15.6-28.6 ± 0.1 cm). The RuAl co-sputtered film was also deposited at 1.5 mtorr Ar pressure, and was also single phase and crystalline as-deposited with a grain size of ~5-10 nm (Figure 2b). The grain size was not measured after annealing, but grain growth is expected to be limited due to the relatively low homologous temperature during annealing (T/Tm = 0.3, where T and Tm are the annealing and melting temperatures in K, respectively). Also, the resistivity of the film is significantly higher than the bulk value [27], ~8.6 times higher, compared to 2.4 times higher for the NiAl thin film compared to bulk NiAl [46-47]. While we believe that the elevated resistivity in the RuAl film is most likely due to the nanoscale grain size, a deviation from the expected stoichiometry can also cause an increase in resistivity due to constitutional point defects.

Nickel, ruthenium, and aluminum films were also oxidized as a comparison to the intermetallic films. The aluminum and nickel films were deposited by electron-beam evaporation, whereas the ruthenium film was deposited by sputtering with an Ar pressure of 1.5 mtorr. The nickel and ruthenium films both contained a small amount of aluminum (~0.5 at%), which was used as an adhesion layer.

III. RESULTS AND DISCUSSION

Prior to oxidation, the NiAl films were optically smooth with no visible surface features, and a root mean square (RMS) surface roughness of less than 1.5 nm (2 μm × 2 μm area) for the multilayer films, as determined from tapping mode AFM images. The roughness of the co- sputtered films was not measured, but from cross-sectional TEM images is expected to be

87 slightly higher than the annealed multilayer films. The film deposition was not optimized in order to decrease the roughness further, but even without optimization these values are less than what is seen in the literature for SiGe films (1.8-5.4 nm RMS) [48-49]. Annealing of the multilayer samples prior to oxidation caused no significant change in the RMS roughness for a multilayer with a 33 nm wavelength. The grain size after multilayer annealing was ~20-60 nm. In comparison, the nickel surfaces were visibly rough after annealing and showed a large increase in grain size during annealing from less than 50 nm to ~1 m after 1 min at 400ºC.

A. Oxide Surface Morphology and Depth Profiles

After oxidation, all of the NiAl films studied showed a similar progression of surface morphology, an example of which is shown in Figure 3 from the onset of noticeable oxidation to 800ºC. For the NiAl films, no change in the surface morphology from the pre-oxidized state was observed with light microscopy for temperatures up to and including 500ºC. For temperatures between 550ºC and 600ºC, small particles, presumably an oxide, formed on the surface of the film. These particles grew slowly in size and number up to ~700ºC, at which point the rate of growth increased rapidly up to 850ºC. This is similar to what was observed by Wang et al. [40] for the oxidation of NiAl nanoparticles, which showed almost no weight gain due to oxidation up to ~400ºC, a slow increase in weight gain from 400ºC to 700ºC, followed by rapid weight gain from 700ºC to 900ºC. This suggests that an upper limit for the use of NiAl films as optical components would be ~500ºC, but higher temperatures can be used for short times in applications where surface quality is not as critical. Figure 4 shows the change in the surface when the annealing temperature is increased from 800ºC to 850ºC.

88

a) b)

50 μm 50 μm

c) d)

50 μm 50 μm

Figure 3 – Light microscopy images of the typical evolution of the surface morphology seen in NiAl films during annealing in oxygen. Figures (a-d) represent cumulative 1 h anneals from 100ºC to (a) 600ºC, (b) 700ºC, (c) 750ºC, and (d) 800ºC. The film imaged was fabricated from a sputter deposited Ni/Al multilayer with a wavelength of 30 nm.

a) b)

5 μm 5 μm

Figure 4 – Light micrographs of the surface of a NiAl film after oxidation. The film was annealed cumulatively for 1 h in oxygen from 100ºC to (a) 800ºC and (b) 850ºC. The film was fabricated from an e-beam evaporated Ni/Al multilayer with a wavelength of 25 nm.

89

The images in Figure 4 show that the oxide particles have doubled or tripled in size with increasing oxidation temperature from 800ºC to 850ºC, and will likely connect to form a continuous film if the temperature is further increased to 900-950ºC. Also noted with increasing the oxidation temperature from 800ºC to 850ºC is the appearance of a lighter region on the surface of the film. The difference in shading is believed to be due to areas of a thicker (lighter gray) and thinner (darker gray) oxide scale. Auger depth profiles of NiAl after oxidation to 850ºC are shown in Figure 5.

80 80 a) Ni b) Ni c) Al Ni Al O O Ni Si 60 60 Si O O 1 2

40 Al Si 40

Al

Al Al Composition(a.u.) 20 (a.u.) Composition 20 1 Si 2

1 m 0 0 0 500 1000 1500 2000 0 500 1000 1500 2000 Sputter Time (s) Sputter Time (s) Figure 5 – Auger depth profiles of a NiAl film formed from a multilayer with a 33 nm wavelength Ni/Al (e-beam deposited) after oxidation to 850ºC. The depth profiles shown in (a) are for an area in between oxide particles and in (b) are for an area centered on an oxide particle. Locations of the 2 scan areas are shown in (c) where box 1 corresponds to (a) and box 2 is the scan area represented in (b).

The depth profiles in Figure 5 show that reaction with the substrate has occurred, and that the distribution of Ni, Al, Si, and O are very different in the regions with (Figure 5b – region 2) and without (Figure 5a – region 1) the oxide particles. Although the surface of both regions contains predominately aluminum and oxygen (presumably a form of Al2O3), a signal from silicon becomes evident much closer to the surface, after ~100 s of sputtering, in the area under the particles compared to ~1000 s for regions away from the particles. These results show that a more significant reaction has occurred with the substrate in regions directly below the particles (Si is present almost throughout the film), whereas the reaction with the substrate elsewhere is restricted to a narrower layer between the film and substrate. Both regions contain oxygen throughout the film, although there appears to be more oxygen in the region below the particles in comparison to the areas between the particles. Both regions have a layer that is still rich in Ni and Al, which contains a little oxygen in the case of region 1, and oxygen and silicon in the case of region 2. In region 2 there are also 2 layers between the Al2O3 and NiAl(O, Si): one containing Al, O, and Si and the second containing Al, O, Si, and Ni. The exact stoichiometry of these layers has not been determined. Although the films below the oxide are still conductive (see discussion on electrical characterization), extensive reaction with the substrate has occurred, which may limit the service temperature of NiAl on SiO2 when used in an oxidizing environment. If alternate substrate materials are identified that do not react with NiAl, it may be possible to obtain a similar oxide structure to that shown in Figure 5a (without the reaction layer with the substrate) across the surface of the entire film.

90

The AES profiles showed that after oxidation to 850ºC the surface of the film is likely composed of Al2O3 (quantitative analysis with standards was not performed). The composition of the oxide as it begins to form the discrete particles (~550ºC) is unclear, and the morphology does not correspond to the oxide morphologies described in the literature for , , or -Al2O3, which tend to form at higher temperatures (700ºC and above). To the best of our knowledge, the morphology of the oxide phases reported to form on NiAl at lower temperatures (NiO and NiAl2O4) have not been reported. At the temperatures illustrated in the light micrographs of Figure 4 (800-850ºC), the expected oxide for bulk polycrystalline NiAl would be -Al2O3, the morphology of which has been described as smooth and dense [37]. This oxide may correspond to the region between the particles.

As with the NiAl films, the RuAl films prior to oxidation were very smooth and reflective with a grain size of less than 50 nm, but the evolution of the surface with oxidation differs significantly. The RuAl film after oxidation had a more uniform and fine-grained surface than the NiAl film, but light microscopy did show a change in the surface morphology at a slightly lower temperature in comparison to NiAl (500ºC in comparison to 550ºC). At 500ºC, the RuAl surface shows some roughening, which increases slightly as the temperature increases, but no discrete oxide particles are observed. The surface of the film remains practically unchanged for oxidation temperatures between 600ºC and 800ºC. Figure 6 shows the comparison of the surface morphology of RuAl and NiAl co-sputtered films after annealing in oxygen up to 800ºC. For certain applications, RuAl may be preferable to NiAl since it has a more homogeneous and reflective surface after short-term oxidation up to 800ºC.

a) b)

5 μm 5 μm

Figure 6 – Comparison of the surface morphology of co-sputtered (a) RuAl and (b) NiAl thin films. Both were annealed in oxygen from 100ºC to 800ºC.

AES depth profiles after the oxidation of RuAl to 800ºC and 850ºC are shown in Figure 7. Unlike NiAl, no reaction with the substrate was detected for oxidation temperatures up to 850ºC, and a layered oxide was produced containing alternating layers of Al2O3 and a Ru-rich layer. The Ru-rich layer appears to contain both aluminum and oxygen near the surface, but is nearly pure ruthenium near the substrate, especially after oxidation at 850ºC. A similar banded oxide has been seen in the high temperature oxidation of RuAl [41-42, 50], although this is usually

91 associated with two-phase Ru-RuAl alloys, whereas the film studied here has been confirmed through TEM to be single phase RuAl prior to oxidation.

100 100 a) Ru b) Ru Al Ru Al 80 Ru O 80 O Si Si

60 O 60 O

40 Si 40 Si

Al Al

Composition(a.u.) Composition (a.u.) Composition 20 Al 20 Ru Ru

0 0 0 500 1000 1500 2000 2500 0 500 1000 1500 2000 2500 Sputter Time (s) Sputter Time (s) Figure 7 – Auger depth profiles of a co-sputtered RuAl film after oxidation to (a) 800ºC and (b) 850ºC.

B. Electrical Characterization

The normalized sheet resistance does not include contributions from the oxide scale as the probe tips likely poke through the oxide when the oxide is thin (in the early stages of oxidation). Hence, the values shown in the following figures refer to the conductive film below the oxide. Once the oxide has thickened to a point where the probe tips can no longer fully penetrate the oxide scale (noted with Al and Ni), then the resistance becomes very high.

As the samples were annealed at only 400ºC prior to oxidation, there are a number of competing factors that can influence the resistance of the sample. When the samples are oxidized at temperatures above the annealing temperature, the resistance may increase or decrease depending on which factor(s) are dominant. A decreasing resistance is due to annealing effects such as 1) grain growth, 2) the elimination of defects, and 3) phase changes. Conversely, an increasing resistance is due to the effects of oxidation: 1) a decreased thickness of the conductive layer, 2) diffusion of oxygen into the film from the surface, 3) diffusion of silicon into the film from the substrate, 4) changing film stoichiometry as the oxide layer may be enriched in one of the elements (creating point defects), and 5) phase changes. Phase changes can occur below the oxide layer for a multi-component film if the region below the oxide becomes depleted in one element. The change in resistance depends on the difference in resistivity of the two phases, and thus is included as a factor that could either increase or decrease the resistance. A phase change may also occur if the films are not completely reacted prior to oxidation.

92

105 5000 Al Al Ni Ni Ru 4 Ru

cm) 4000 10 NiAl-cosputter NiAl-cosputter

 RuAl-cosputter RuAl-cosputter

103 3000

2 2000

10 Percent Change (%) Change Percent 1000

101 Normalized Resistance ( Resistance Normalized

0 100 0 150 300 450 600 750 900 0 150 300 450 600 750 900 Temperature (°C) Temperature (°C) Figure 8 – 4-pt probe measurements as a function of annealing temperature in oxygen for co-sputtered RuAl and NiAl films in comparison to Al, Ni, and Ru films. The data are presented as both the value of the normalized resistance after a given anneal, and as a percent resistance change compared to the sample prior to oxidation. In (a) and (b) the nearly vertical line for Ni represents a loss of conductivity (the film was conductive at 450°C, but not at 500°C). For clarity, the last data point for Al is not shown in (b) (Al at 600°C exhibits a 2×106 % increase in resistance).

Figure 8 shows the 4-pt probe data as a function of the oxidation temperature for the co-sputtered NiAl and RuAl films, along with Al, Ni, and Ru films for comparison. The data in Figure 8 show that Ni becomes non-conductive after cumulative oxidation (1 h each) to 500ºC, whereas Al is conductive up to 600ºC, Ru and RuAl to 800ºC, and NiAl to 850ºC. Although Ru and RuAl show similar electrical properties at 800ºC, optical images of the Ru (not shown) indicate the possibility of voids at the film surface between 700 and 800ºC. This could be due to the volatilization of RuO3 and RuO4, which was suggested to occur between 700 and 800ºC in Ru films by Lisker et al. [51], who showed that after Ru films were oxidized completely to the RuO2 phase, further oxidation resulted in resistivity values higher than the bulk value of RuO2. This finding was ascribed to the loss of material from the film, resulting in a decrease in the film thickness, and an apparent increase in the film resistivity. The above data show that the intermetallics NiAl and RuAl perform better than Al, Ni, and Ru in an oxidizing atmosphere. In terms of electrical properties, NiAl performs better than RuAl, as it remains conductive to higher temperatures.

93

1000 NiAl-multi-20 nm NiAl-multi-20 nm NiAl-multi-25 nm NiAl-multi-25 nm

NiAl-multi-33 nm NiAl-multi-33 nm cm) Ni Al -multi-25 nm Ni Al -multi-25 nm 100 2 3 2 3  NiAl-cosputter NiAl-cosputter NiAl-multi-30 nm NiAl-multi-30 nm

NiAl-0.5Ag-multi-30 nm NiAl-0.5Ag-multi-30 nm

Percent Change (%) Change Percent

100 NormalizedResistance (

10 400 500 600 700 800 900 775 800 825 850 875 Temperature (°C) Temperature (°C) Figure 9 – 4-pt probe measurements as a function of oxidation temperature for Ni-Al films. The data are presented as both the value of the normalized sheet resistance and as a percent resistance change compared to the sample prior to oxidation. The data for oxidation temperatures less than 400ºC are not shown, as there was no change in this region. Error bars are within the size of the symbols.

Figure 9 shows the resistivity data as a function of oxidation for NiAl films fabricated using different methods. The co-sputtered NiAl data is included for comparison as is a Ni2Al3 film. In general, the behavior of the resistance during cumulative oxidation can be described as two regions: one of decreasing resistance with increasing oxidation temperature, followed by an increasing resistance with increasing oxidation temperature. For the majority of the films, the changeover from decreasing to increasing resistance occurs between 550ºC and 600ºC, which is just slightly higher than when changes in the surface morphology were detected (~550ºC). The co-sputtered film shows only a very minor decrease (~3%) in resistance, which may indicate that limited grain growth occurs in this sample. Oxygen diffusion along the columnar grain boundaries is likely rapid (in comparison to bulk diffusion) and may effectively pin the grain boundaries and inhibit grain growth. Two of the films show a more significant region of decreasing resistance; the annealed 20 nm wavelength and 33 nm wavelength Ni/Al multilayers. XRD indicated that the multilayer with a wavelength of 33 nm might have had a small amount of Ni2Al3 present prior to oxidation. The conversion of Ni2Al3 to NiAl would result in a decrease in resistance. Regardless of the when the changeover from increasing to decreasing resistance occurs, by 800ºC all of the films show an increasing resistance with increasing oxidation temperature. By 850ºC, all of the films have a resistance higher than the resistance of the film prior to oxidation. There was one exception to the generalized two-region behavior: Ni2Al3. This sample showed an initial increase in resistance, then a decrease in resistance, followed by another increase in resistance. The initial and final increase in resistance is due to oxidation, which according to light microscopy occurs at a slightly lower temperature for Ni2Al3 than NiAl (500ºC in comparison to 550ºC). It is possible that the oxidation of Ni2Al3 results in a phase change, as the composition range where Ni2Al3 is stable as a single phase is narrower than for NiAl. If the oxide forming on Ni2Al3 is alumina, then oxidation may cause a decrease in resistance as the higher resistivity phase (Ni2Al3) transforms to the lower resistivity phase (NiAl). The resistance as a function of oxidation temperature for Ni2Al3 is likely a complex mixture of opposing factors.

94

The surface morphology of the oxidized Ni2Al3 was similar to that of NiAl, but the oxide particles were larger in the case of Ni2Al3.

While there was some variation in the resistance as a function of oxidation temperature for the NiAl samples, there was a clear difference in the percent increase after oxidation to 850ºC (ranging from 90-770% for the NiAl films as shown in Figure 9b). This observation is believed to be due to differences in the starting film thickness, causing (1) a faster increase in resistance in the thinner films, possibly combined with (2) different rates of oxidation for films of a different thickness. The effect of the initial film thickness on the measured resistance can be demonstrated as follows. Eq. 1 shows how the resistance, R, is calculated from the 4-point probe measurements [52],

V  R   Eq. 1 I 4.532t where V is the measured voltage, I is the applied current (typically 10 mA),  is the resistivity of the metal, and t is the thickness of the metal. During oxidation, the resistivity of the metal is changing as well as the thickness of the metal, both of which affect the value of the measured resistance. This is shown in Eq. 2

 R 2  R1   2 t1  % Increase in R     100   1  100 Eq. 2  R1   1t 2  where the subscripts 1 and 2 refer to the film prior to oxidation and the film after oxidation to a certain temperature, respectively. Assuming that all of the NiAl films show a similar change in the resistivity due to oxidation, then the ratio of 2 to 1 can be considered a constant. Similarly, if the conductive region of the films (not the oxide) all decrease in thickness by a similar amount (x) during oxidation, then the thickness t2 can be written as t1-x, and Eq. 2 can be simplified to the form shown in Eq. 3. (The assumption of an equal decrease in film thickness due to oxidation will be discussed below).

 t1  % Increase in R  C 1  100 Eq. 3  t1  x 

Eq. 3 gives an approximation of the expected increase in the measured film resistance at a certain temperature for films with varying initial film thickness. The experimental data for annealed Ni/Al multilayers are plotted and curve-fit using this equation in Figure 10. Also shown in the figure but not included in the fit are data for 2 other samples (to be discussed later).

95

1000 Annealed Ni/Al Co-sputtered NiAl 800 Annealed Ni/Al + Ag Eq. 3

600

400

200 Resistance Increase (%) Increase Resistance

0 150 175 200 225 250 275 300 Film Thickness (nm) Figure 10 – The resistance increase after oxidation to 850ºC as a function of the initial film thickness. The data for the multilayer samples (open squares) are plotted according to Eq. 3 (see text for details). The fit to the data yields an R2 value of 0.919.

Figure 10 shows that Eq. 3 fits the data from the annealed Ni/Al multilayers fairly well even though several assumptions are incorporated in the equation, and a sharply increasing resistance should be expected as the film thickness drops below ~200 nm. Plotting the data in this way also allows for a comparison to be made between the annealed multilayers, co-sputtered NiAl, and the silver-doped annealed multilayer. From the data in Figure 9, it would initially appear that the co- sputtered NiAl and the silver-doped NiAl have similar oxidation behavior to the annealed multilayer samples, as the normalized resistance and the resistance increase after oxidation to 850ºC are bounded by the data from the multilayer samples. However, when the data are plotted as a function of the film thickness prior to oxidation, it becomes clear that after 1 h at 850ºC the silver-doped and co-sputtered films both show a greater percentage increase in resistance due to oxidation when compared with un-doped multilayers of equivalent film thickness. This observation is consistent with the literature data, which showed an increase in the initial oxidation kinetics for a co-sputtered thick film with columnar grains as compared to a bulk sample (both had the same thickness oxide after long term oxidation) [36], and increased oxidation for silver-doped NiAl in the temperature range from 900ºC-1100ºC [53-54].

96

a) 293 nm b) 176 nm +90% +518%

5 μm 5 μm

c) 164 nm d) 204 nm +769% +446%

5 μm 5 μm

Figure 11 – Comparison of the film surface morphology after oxidation to 850ºC. The films were imaged using the same illumination conditions. The images show NiAl films formed from annealed Ni/Al multilayers with a wavelength of (a) 20 nm, (b) 25 nm, (c) 30 nm, and (d) a co-sputtered NiAl film. The upper right of each image shows the initial film thickness as well as the percent increase in resistance after oxidation to 850ºC.

Light microscopy also indicates that there may be a thicker oxide on the co-sputtered film in comparison to the annealed multilayers after oxidation to 850ºC (see Figure 11d) as there is more of the lighter region on the surface. In fact, the amount of the lighter surface region also appears to increase as the film thickness decreases (see Figure 11a-c). As we previously linked this to a thickening of the oxide, this may indicate a thicker oxide on the thinner NiAl films.

IV. CONCLUSIONS

We have described the oxidation resistance of several thin film metallic materials. We observed the surface morphology and electrical resistance for NiAl and RuAl B2 intermetallic thin films that are attractive for a variety of thin film applications, and more specifically for MEMS. The intermetallic thin films were compared to Al, Ni, and Ru thin films, and were found to be superior to the pure metal films in terms of maintaining a smooth surface before and after oxidation and the electrical resistance after oxidation. The intermetallic films showed no observable surface changes (light microscopy) up to 500ºC in flowing oxygen and were conductive to higher temperatures (800-850ºC) than Ni (500ºC), Al (600ºC), and Ru (800ºC),

97 although vaporization may have begun at ~700ºC in the case of the Ru film. After oxidation, the RuAl film still had a very reflective surface, but a distinct haziness was noted on all of the Ni-Al films. Both materials after being oxidized to 850ºC had a surface composed solely of aluminum and oxygen (likely Al2O3), but the composition below the surface was significantly different. The RuAl had an oxide scale composed of alternating Ru-rich and alumina layers, whereas the oxide on NiAl had a more complex structure due to reactions with the substrate. Although some surface roughening was observed in the RuAl film, the surface was still relatively smooth, reflective, and homogeneous in comparison to NiAl, and showed little change in surface morphology from the onset of oxidation (500ºC) up to 800ºC. In contrast, NiAl showed the appearance of discrete surface oxides at ~550ºC, with a more uniform oxide forming between these particles, which began to noticeably thicken at 850ºC. Determining conditions that delay the formation of this thicker oxide, such as doping, and the elimination of columnar grain boundaries, may increase the service temperature of NiAl thin films.

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[36] S. Yang, F. Wang, Z. Sun, and S. Zhu, Influence of columnar microstructure of a sputtered NiAl coating on its oxidation behavior at 1000C. Intermetallics, 2002. 10: p. 467-471. [37] M. Brumm and H.J. Grabke, The oxidation behaviour of NiAl-I. Phase transformations in the alumina scale during oxidation of NiAl and NiAl-Cr alloys. Corrosion Science, 1992. 33(11): p. 1677-1690. [38] N. Ohtsu, M. Oku, K. Obara, S. Ito, T. Shisido, and K. Wagatsuma, Oxidation behavior of NiAl alloy at low temperatures. Surface and Interface Analysis, 2007. 39: p. 528-532. [39] J. Doychak, J.L. Smialek, and T.E. Mitchell, Transient oxidation of single-crystal -NiAl. Metallurgical Transactions A (Physical Metallurgy and Materials Science), 1989. 20(3): p. 499-518. [40] Z. Wang, W. Tian, and X. Li, Oxidation behavior of NiAl nanoparticles prepared by hydrogen plasma-metal reaction. Materials Chemistry and Physics, 2008. 107(2-3): p. 381-384. [41] P.J. Bellina, A. Catanoiu, F.M. Morales, and M. Ruhle, Formation of discontinuous Al2O3 layers during high-temperature oxidation of RuAl alloys. Journal of Materials Research, 2006. 21(1): p. 276-286. [42] F. Cao, T.K. Nandy, D. Stobbe, and T.M. Pollock, Oxidation of ruthenium aluminide- based alloys: The role of microstructure and platinum additions. Intermetallics, 2007. 15: p. 34-43. [43] D.W. McKee and R.L. Fleischer, Oxidation behavior of advanced intermetallic coatings. Materials Research Society Symposium - Proceedings, 1991. 213: p. 969-974. [44] F. Soldera, N. Ilic, S. Brannstrom, I. Barrientos, H. Gobran, and F. Mucklich, Formation of Al2O3 scales on single-phase RuAl produced by reactive sintering. Oxidation of Metals, 2003. 59(5/6): p. 529-542. [45] W.K. Shen, J.H. Judy, and J.-P. Wang. In situ epitaxial growth of ordered FePt (001) films with ultra small and uniform grain size using a RuAl underlayer. in 49th Annual Conference on Magnetism and Magnetic Materials. 2005. Jacksonville, Florida (USA): AIP. [46] G.R. Caskey, J.M. Franz, and D.J. Sellmyer, Electronic and magnetic states in metallic compounds - II. Electron transport and magnetic susceptibility in NiAl and FeAl. Journal of the Physics and Chemistry of Solids, 1973. 34: p. 1179-1198. [47] Y. Yamaguchi, D.A. Kiewit, T. Aoki, and J.O. Brittain, Electrical Resistivity of NiAl, CoAl, NiGa, and CoGa. Journal of Applied Physics, 1968. 39(1): p. 231-232. [48] T.-H. Ahn, I.-S. Yeo, T.-K. Kim, M.-S. Joo, H.-S. Kim, J.-J. Kim, J.-H. Joung, and J.W. Park, Effects of Ge Content on the Oxidation Behavior of Poly-Si1-xGex Layers for Gate Electrode Application. Journal of the Electrochemical Society, 2001. 148(2): p. G50- G54. [49] P.S. Chen, S.W. Lee, and K.F. Liao, Growth of high-quality relaxed SiGe films with an intermediate Si1-yCy layer for strained Si n-MOSFETs. Materials Science and Engineering: B, 2006. 130(1-3): p. 194-199. [50] F. Soldera, N. Ilic, N.M. Conesa, I. Barrientos, and F. Mucklich, Influence of the microstructure on the formation of alumina scales on near stoichiometric RuAl produced by arc melting. Intermetallics, 2005. 13: p. 101-107. [51] M. Lisker, T. Hur'yeva, Y. Ritterhaus, and E.P. Burte, Effect of annealing in oxygen atmosphere on morphological and electrical properties of iridium and ruthenium thin

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films prepared by liquid delivery MOCVD. Surface and Coatings Technology, 2007. 201: p. 9294-9298. [52] D.K. Schroder, Semiconductor Material and Device Characterization. 2 ed. 1998, New York: John Wiley & Sons. [53] X.J. Zhang and Y. Niu, Oxidation of four NiAl-Ag alloys at 900C in 1 atm O2. Materials Science Forum, 2005. 475-479: p. 775-778. [54] X.J. Zhang, Z.G. Zhang, and Y. Niu, Effect of Ag on the oxidation of -NiAl at 1173 to 1373K in 105Pa oxygen atmosphere. High Temperature Materials and Processes, 2005. 24(6): p. 359-368.

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CHAPTER 5

ALTERNATING CURRENT INDUCED THERMOMECHANICAL FATIGUE OF NIAL AND RUAL THIN LINES ON OXIDIZED SILICON

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I. INTRODUCTION

A. Background: Alternating Current (AC) Thermomechanical Fatigue (TMF) Testing

A recent method for studying the thermomechanical fatigue (TMF) of thin films involves passing an alternating current (or voltage) through a thin metallic line [1-3] (see Figure 1 and Figure 2). Due to Joule heating the metal line heats up and because of the alternating current, experiences temperature oscillations at twice the frequency of the applied current. Any mismatch in the coefficient of thermal expansion between the substrate and the line will cause stresses in the line, which will also oscillate at twice the frequency of the current.

Figure 1 – Schematics showing two possible test structures for AC testing. Most test structures used in the literature are a variation of one of the above lines. On the bottom figure, the current (or voltage) is applied at the two end contact pads and the voltage drop (or current) is measured across the two side contact pads.

1. Pass an AC through a thin metal line Joule Heating 2. Oscillating temperature at twice the AC Current frequency (ΔT)

Different Thermal Expansion Coefficients

3. Oscillating stresses (Δσ) at twice the AC Temperature frequency

• Changes in the film stress can be approximated Stress by Δσ  E  ΔT 0 20 40 60 80 substrate film Time (ms) Figure 2 – Schematic showing the method used to set-up an oscillating temperature and stress in a metal thin film sample. The exact position of the stress curve will depend on the initial stress state of the film and the difference in CTE between the film and substrate.

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This method is attractive for studying the TMF behavior of materials as it allows for tests to be conducted in a matter of hours or days rather than weeks or months. For example, accelerated temperature cycling (ATC) and rapid temperature cycling (RTC), which are used to study the TMF behavior of solders, have heating/cooling cycles on the order of 1000 and 100 seconds, respectively [4]. In the time that it takes to heat up a sample in an ATC test, more than 100,000 cycles (10,000 for RTC) could be completed using AC testing. This rapid cycling is partly due the method of testing, where the heating rate is set by the frequency of the current, and also due to the small volume that needs to be heated (thin film).

While the test method is mainly used to study damage formation under conditions of fluctuating temperature and stress, Keller et al. [5] have used this method to “approximate” the ultimate tensile strength (UTS) of aluminum. They used a modified Basquin equation (Eq. 1), where is the stress amplitude, is the fatigue strength coefficient (~UTS), is the mean stress, Nf is the number of stress reversals to failure, and b is the fatigue strength exponent.

Eq. 1

They used the above equation to fit their AC TMF results for Al-1Si and obtained a value of = 250 ± 40 MPa, which agreed well with results from microtensile tests of specimens taken from the same wafer (UTS = 239 ± 4 MPa). This finding is consistent with the general trend for metals that the fatigue life is correlated with the ultimate tensile strength [6].

B. Temperature During AC TMF

During AC TMF testing the ΔT experienced by the metal line is typically in the range of 100- 300°C and varies depending on testing conditions and sample geometry, with the minimum temperature in the range of 70-130°C and the maximum temperature between 200-400°C [3, 7]. However, a peak temperature of ~800°C was observed in 90 nm wide copper interconnects fabricated by the dual Damascene approach [8].

When the AC frequency is low (less than ~15 Hz with Si as the substrate) there is sufficient time for heat to be conducted away from the line by the substrate prior to the next current cycle, and the difference in temperature between the metal line and the substrate will be proportional to the power. At higher frequencies there is insufficient time for the substrate to conduct all of the heat away from the line prior to the next current cycle, so the metal line will not return to the ambient temperature between current cycles. In this case ΔT will no longer be proportional to the power, but will also depend on the frequency. The temperature of the sample will still oscillate at twice the frequency, but will be superimposed on a slow transient caused by heating of the substrate [3]. After a sufficient length of time (1-2 h) the transient will saturate to a value T0 (70-100°C) upon which ΔT will be superimposed [3]. Schematics for low and high frequency testing are shown in Figure 3.

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ΔT

T0

Temperature (au) Temperature Temperature (au) Temperature

ΔT

TimeTime (au)(au)

Figure 3 – Schematics of the temperature oscillations of a metal line on a substrate heated by an AC signal with a low frequency (bottom) and a high frequency (top).

No analytical solutions exist for the temperature as a function of time for the case of a metal line on a substrate being heated by an AC signal, but experimental results indicate that ΔT is proportional to the peak applied power when testing at constant frequency. When testing is conducted at a constant applied power ΔT is proportional to (see Figure 4), where f is the frequency of the AC signal. This proportionality holds for frequencies up to ~6 kHz, where the metal acts as a line heat source: i.e., the depth of penetration of the heat conducted by the substrate in each heating cycle is much larger than the line width. Above ~6 kHz the metal acts more as a planar heat source: i.e., the depth of penetration of the conducted heat in each heating cycle is less than the width of the line. Experimental results from the literature showing the relationships between temperature and power, and temperature and frequency are shown in Figure 4 for Cu lines tested on oxidized silicon.

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140 300 T min 1 T T  49.4W  2.6 max max max T   ln f  120 2 250

200 100 T (°C) T

150  T 80

Temperature (°C) Temperature 100 60 Tmin 17.3Wmax  23.1 50 a) b) 40 2 3 4 5 6 10 100 1000 10000 100000 Power (W) Frequency (Hz) Figure 4 – The effects of the applied power and AC frequency on the temperature during AC thermal cycling for Cu lines 800 μm long, 8 μm wide, and 300 nm thick tested on oxidized silicon. (a) The minimum and maximum temperatures as a function of the peak applied power for a testing frequency of 100 Hz. (b) ΔT as a function of the AC frequency for a peak applied power of 4.5 W. Data from [3, 9].

Due to the large contact pads present at the ends of the line (similar to those seen in Figure 1), temperature gradients are present along the length of the line. Finite element analysis conducted by Monig et al. [3] indicated that the temperature reached 90% of the peak value (located at the center) at 25 μm from the contact pads and 99% of the peak value at a distance of 140 μm from the contact pads. Their analysis also showed that a temperature increase of only a few degrees occurred when a crack was present that encompassed one third of the line width.

As the temperature excursions are too rapid for most conventional temperature measurement methods, the actual temperature of the metal line is determined during the test by measuring the resistivity of the metal. By using Ohm’s law, the resistance of the metal line can be determined (Eq. 2), and from this the resistivity (Eq. 3).

V  IR Eq. 2 RA   Eq. 3 l

In Eq. 2 and Eq. 3, V and I are the voltage and current measured at the side contact pads, A is the cross-sectional area (wt), and l is the distance between the contact pads. As temperature changes lead to changes in resistivity, the temperature can be determined as long as the relationship between the resistivity and temperature for the material is known. In many metals, the relationship between temperature and resistivity often follows the relationship given in Eq. 4 [10].

  0 10 T T0  Eq. 4

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In Eq. 4 0 is the resistivity of the metal at a reference temperature T0 (often 273K or 293K),  and T are the measured resistivity and applied temperature, and 0 is the temperature coefficient of resistance (TCR) at the reference temperature. The TCR is the fractional change in the resistivity of the metal with a change in temperature at T0, and Eq. 4 is only valid for temperature ranges in which the TCR is constant.

C. Literature Results: AC TMF Testing

Several studies have been published on AC TMF of Al-1Si [1-2, 5, 11-12], Cu [1, 3, 5, 7-8, 13- 16], and Au [17-20] lines. As would be expected, the severity of damage increased and the time to failure decreased with increasing current (or voltage). It was also noted that narrower lines could sustain a higher current density than wider lines for equivalent times to failure [2, 8]. The increased current density for the narrower lines is likely a combination of geometry (narrower lines have a higher current density, but a lower current yielding decreased Joule heating) and increasing strength as the dimensions are reduced. In all cases localized melting was observed at the failure sites due to current crowding and Joule heating [2-3, 8].

Aluminum fatigue samples showed regions of no damage, light damage, and severe damage along the length of the line [2]. Periodic surface wrinkling was observed [1-2, 5, 11-12] with some extrusions along the side of the line in severely damaged regions [2]. The wrinkles became more regularly spaced with a larger amplitude as the line width increased and the damage was confined to individual grains, with more damage noted in the larger grains [2, 5, 12]. TEM cross-sectional analysis of an intrusion/extrusion in a failed line showed a thickness of 0.73 μm through the extrusion and a thickness of 0.28 μm through the valley (initial thickness ~0.5 μm), and in some instances the thickness could be less than 100 nm in the trough regions [2, 11]. Also occurring during testing was grain growth [1, 5, 11] and reorientation [1, 11-12]. During 320 s of testing at a ΔT of ~200°C the average grain size increased from 1.4-2.4 μm. Local grain growth was noted in this sample after only 10 s of testing, and after failure (695 s) some grains had increased in diameter by more than 6 times. Keller et al. [5] attributed the grain growth to strain induced boundary migration, where grains with significant “residual plasticity” are consumed by grains with lower “residual plasticity” thereby minimizing the strain energy. This was supported by the fact that the more damaged grains grew and had few remaining dislocation segments (most left the surface forming the topography – intrusions and extrusions). Relatively flat grains had a higher density of dislocations. Prior to testing the Al lines had a <111> fiber texture, which degraded during the AC testing as some of the deformed grains rotated by more than 40° [11-12].

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24 Cu

) Al-1Si 2

21

18 Peak Current Density (MA/cm Density Current Peak

15 100 101 102 103 104 105 106 107 Time to Failure (s) Figure 5 – Time to failure at an AC of 100 Hz for Al-1Si and Cu lines on oxidized silicon substrates. Samples were 800 μm long, 3.3 (Al) or 2 (Cu) μm wide, and 0.5 (Al) or 0.7 (Cu) μm thick. Data from [1].

Copper had an increased lifetime compared to Al-1Si (see Figure 5), but showed similar surface modifications: surface wrinkling and occasional grain growth [1, 3, 7-8, 13, 15-16]. The time to failure was seen to strongly depend on temperature (both Tmax and ΔT) with no failure after 5×107 cycles at a ΔT=120°C, and failure in as few as 1×105 cycles for a ΔT=190°C (at an AC frequency of 100 Hz) [3]. Similarly, increasing the ΔT from 110-200°C at an AC of 10 kHz decreased the lifetime from 12 days to 10 seconds [15]. The determination of the effect of AC frequency on fatigue life and damage accumulation is complicated by the fact that higher currents are required to reach an equivalent value of ΔT at a higher frequency (refer to Figure 4). Thus, tests conducted with a constant ΔT will not be oscillating between the same 2 temperatures if the tests are carried out at different frequencies. The higher frequency test will have a higher current density leading to a higher Tmin and Tmax. For example, at a ΔT of 190°C the average sample temperature was 15°C higher when testing at 10 kHz rather than 100 Hz [7]. Data from 20 copper interconnects tested at the same current density were fit to several common lifetime distribution models and it was found that the most appropriate model was a lognormal distribution for both 90 nm and 4.5 μm wide interconnects [14]. One set of data is shown in Figure 6, where for a constant RMS current density of 17 MA/cm2 failure times ranged from ~270-5600 s.

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100

80

60

40

Failure Probability (%) Probability Failure 20

0 100 1000 10000 Time to Failure (s) Figure 6 – Lifetime data for a set of 20 copper interconnects tested using an AC of 100 Hz and an RMS current density of 17 MA/cm2. The lines were fabricated using the dual Damascene approach and were 60 μm long, 250 nm thick, and 4.5 μm wide. Data from [14].

The damage from AC TMF tests on copper lines was localized within individual grains, and depended on grain orientation, grain size, and film thickness [7, 15]. The damage began as wrinkles near grain and twin boundaries and then grew to encompass the entire grain [7, 13, 15]. <100> oriented grains grew and consumed neighboring grains, whereas <111> grains showed no grain growth and had wrinkles with a larger amplitude and a more regular spacing [7, 15-16]. Due to the larger amplitude wrinkles, failure typically occurred in <111> grains because of more severe current crowding (a higher current density). In thicker films with a larger grain size, parallel surface wrinkles formed with the intrusions extending almost to the substrate [3]. In thinner films with a smaller grain size wrinkled grains were only occasionally observed and only in the largest grains. Fatigue damage in these samples occurred mostly as thinned grains, hillocks, and grain boundary grooving [3]. The number of cycles to failure was up to 2 orders of magnitude higher for 100 nm thick copper compared to 300 nm thick copper samples (see Figure 7a). A similar effect is seen with decreasing grain size (see Figure 7a) and line width (see Figure 7b) [7-8]. When the line width was decreased from 4.5 μm to 90 nm, wrinkle formation was completely eliminated and the current density was more than 1.5 times higher for equivalent lifetimes. Monig et al. [3] determined that as the grain size and/or film thickness decreased, the fatigue damage formation became dominated more by diffusive mechanisms (surface and boundaries) and less by dislocation glide.

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350 0.49 60 a) 100 nm thick - 0.4 m grains b) 90 nm wide lines

300 nm thick - 0.5 m grains ) 4.5 m wide lines 300 0.42 2 300 nm thick - 1.5 m grains 50

250 0.35 Strain (%) 40

200 0.28

T (°C) T  30 150 0.21

20

100 0.14 RMS Current Density (MA/cm Density Current RMS

50 0.07 10 4 5 6 7 8 10 10 10 10 10 102 103 104 105 106 Cycles to Failure Time to Failure (s) Figure 7 – Fatigue data obtained using an AC of 100 Hz for Cu lines of (a) varying thickness and grain size [7] and (b) of different line width [8].

AC TMF tests have also been conducted on Au lines (see Figure 8) [17-20], although comparisons to the tests on copper and aluminum are difficult as the tests were conducted at a different frequency, on material with a much smaller grain size, using a different substrate material, and the temperatures were not measured (they were calculated using finite element analysis and the authors’ developed analytical solution). The grain size of the gold lines studied by Zhang et al. [17-20] is in the range of 57-121 nm, which is much smaller than the grain size of the copper and aluminum lines discussed above, and this is believed to cause the difference in damage formation between the Au and Al/Cu samples. No surface wrinkling was noted in the gold fatigue samples, but voids were seen to form and grow perpendicular to the line direction. Other damage noted was microcracks at triple junctions, wedge-shaped pits, grain growth, and grain thinning. The difference in damage morphology between the Au and Al/Cu samples was proposed to be due to the small grain size inhibiting dislocation motion and promoting enhanced diffusion along the grain boundaries.

111

18 18

16 ) 2 16 14

12

14 10 8

6

12 4 0 1x107 2x107 3x107 4x107

10

8

6 Peak Current Density (MA/cm Density Current Peak

4 105 106 107 108 Cycles to Failure Figure 8 – Thermal fatigue data of gold lines on quartz substrates tested with an AC of 50 Hz. The samples were 50 μm long, 2 μm wide, and 200 nm thick. The inset shows the same data plotted with a linear scale on the x-axis rather than a log scale (axes and units are the same in the main graph and the inset). Data from [19].

While the majority of tests performed using this method monitor the current and voltage as a function of time during the test, few report on the evolution of the voltage (or current) signal as the test proceeds. Tests conducted by Keller et al. on Al-1Si [2] showed an increase in resistance at the start of the test due to gradual heating, followed by a region of constant resistance that lasted for the majority of the test until just prior to failure when the resistance showed a rapid increase. Conversely, dual damascene copper interconnects tested by Biesemans et al. [13] showed an initial slight decrease in resistance followed by an increasing resistance until failure, and those tested by Moreau et al. [8] were reported to show a small linear increase in resistance (<5%) with time until just prior to failure, at which point an exponential increase in resistance was obtained.

D. Comments on AC TMF Literature

While numerous experiments have been conducted using the method of AC TMF, the tests all lack a simple method for monitoring the progression of damage formation during the course of a test. Damage formation has been tracked by conducting tests in an SEM [3, 7, 15-16] or a focused ion beam (FIB) system [17-20]. However, these tests typically do not truly track the progress of damage formation as only “snapshots” are obtained by stopping the tests periodically to image the sample. Another downside to this type of test is the need for an SEM or FIB dedicated to running these experiments, which may not be practical as tests can span a number of days or even longer. Experiments tracking the formation of damage in bulk specimens were conducted in the 1960’s and 70’s and are particularly well suited to AC TMF tests as they rely on tracking changes in the samples electrical resistivity.

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Very large changes in resistivity were observed for molybdenum single crystals deformed uniaxially in tension at room temperature [21]. The resistivity more than doubled for a tensile strain of 0.5%, and increased by more than an order of magnitude after testing to a strain of 4%. By annealing at a temperature low enough to remove only the vacancies it was determined that in the range of strain from 0.5-10%, dislocations accounted for 42-58% of the increase in resistivity. Calculations by Vladimirov et al. [22] resulted in similar results for the increase in resistivity due to dislocations and vacancies in aluminum (similar increases in resistivity due to vacancies and dislocations), but determined that microcracks increased the resistivity much more significantly than dislocations and vacancies (almost an order of magnitude larger increase). Experiments conducted at room temperature under cyclic loading conditions yielded results ranging from 80% of the resistivity increase resulting from dislocations (bulk Ni, push-pull fatigue, [23]) to 20% due to dislocations (Cu wire, tension-relaxation, [24]). Increases of resistivity at failure ranged from ~1.3% (Cu foil, bending, [25]) to 21% (bulk tool steel, alternating deflection, [26]). Several of the experiments show similar behaviors of the resistivity as a function of the number of cycles: an initial region of gradually increasing resistivity, followed by a second region with a more rapidly increasing resistivity [25-27]. These 2 regions have been ascribed to the accumulation of defects (dislocations and point defects) and to crack growth, respectively. Zaharia et al. determined from the point of rapidly increasing resistance that the onset of cracking occurred in tool steel after a 0.2-1.6% increase in resistance, which was slightly before cracking could be detected optically. For the plot shown in Figure 9, which shows an example of the data obtained by Zaharia et al., a 0.8% increase in resistance was observed prior to the onset of cracking and a more rapidly increasing resistance.

25

20

15

10 Resistance Increase (%) Increase Resistance

5

0 0 10 20 30 40 50 Cycles (1000's) Figure 9 – Plot of the change in resistance as a function of the number of cycles for a sample of notched tool steel cycled with alternating deflection (bending of cantilever up and down). The arrow indicates the onset of cracking as determined from the point at which the resistance begins to more rapidly increase. Adapted from [26].

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The rapidly increasing resistance seen in the aforementioned articles has been calculated by Johnson [28] as the potential drop across a crack as the crack length increases. For a sheet of width W with a center slit of length 2a0 and infinitesimal height, the relation obtained by Johnson is:

Eq. 5

In Eq. 5 V(a) and V(a0) are the voltage drops across a crack of length a and the initial crack of length 2a0, and y is the distance from the crack to the voltage probes. This equation results in the curve shown in Figure 10 for the change in voltage as a function of crack length. The shape of the curve is reminiscent of the last region of the resistance data obtained by Zaharia et al. [26].

5 a) b)

4 ) 0 y a potential 3

leads V(a)/V(a a0

2

1 W 0 1 2 3 4 Crack Length, a (m)

Figure 10 – (a) Graph of Johnson’s equation (Eq. 5) using the following parameters: a0 = 0 μm, W = 8 μm, and y = 150 μm. (b) Schematic diagram of the sample modeled in Eq. 5. For clarity the initial crack is shown as an ellipse rather than as a slit with infinitesimal height.

The above discussion has shown that considerable increases in resistance can be obtained during fatigue at room temperature due to the formation of dislocations, point defects, and cracks. Also, the onset of cracking during a test may be detected as the point at which a rapidly increasing resistance is observed. As mentioned previously, the majority of the AC TFM tests conducted in the literature record the voltage and current as a function of time during the tests in order to obtain the temperature of the line, but very few actually report on any trends observed in the voltage drop (or resistance) across the line during the test. Those that report on changes in

114 resistance [8, 13] only briefly discuss the observations and do not explain the trends noted in the data.

While it may be possible to distinguish between the formation of dislocations/vacancies and cracks during an AC TMF test by monitoring the sample resistance, it may not be possible to ascertain what mechanism(s) are responsible for the damage without the use of additional testing techniques combined with microstructural characterization. Due to the elevated temperature, cyclic loading conditions, and electrical current present during the tests, there are many processes that may be contributing to the damage formation (or removal) such as fatigue, creep, electromigration, thermal migration, and annealing. All of the above processes, except annealing, would result in an increase in the resistance. Separating out the contribution to the resistance from each of these processes would be challenging, but determining regions where annealing, or dislocation/point defect production, or cracking dominate should be possible by monitoring changes in the sample resistance. Tracking the resistance during a test may enable the monitoring of damage accumulation and lifetime estimations.

E. Comparison of AC and DC Electrical Tests

While using an alternating current to study the fatigue of small structures is a relatively new technique, the effect of AC on metallic lines has been studied as it relates to electromigration, and comparisons between AC and DC lifetimes have been made. The tests have shown that for a constant current density, the AC lifetime is significantly longer than the DC lifetime [13, 18, 29- 30], and the AC lifetime increases with increasing frequency [30]. Tests based on AC TMF assume that electromigration is limited due to the relatively short times for atom transport prior to each current reversal, and on the basis that the damage does not look like typical electromigration damage [3, 15] such as void formation near the cathode and extrusions near the anode [31].

A special test structure designed by Monig et al. [3] was used to thermally cycle a copper line on a substrate without current flow. Damage was formed in this sample without current flow, although it was noted to be different than the samples with current flow. This was explained due to a smaller grain size in the special test structure. Since there was no current flow in the copper line during this test the damage was due to the cyclic temperatures/stresses and not to electromigration.

Tao et al. [29-30] have studied the behavior of Al-2Si and Cu interconnects exposed to AC signals with varying frequency and DC signals. DC tests on Al-2Si showed a continuously increasing resistance with increasing test time, with a value of R/R0 of ~2.66 at the end of the test. For very low frequencies (period=2 h) the Al-2Si showed a very large increase in resistance during the first half cycle (+ current). During the following half cycle (- current) a drop in resistance was noted. As the test progressed the resistance was seen to rise and fall on top of an increasing background, which was ascribed to incomplete damage healing during the current reversals. At the end of the test R/R0 was ~1.07. When the period was increased to 0.5 h similar behavior was observed, except that more complete healing was observed with each current reversal (background increased more slowly). At a very high frequency (10 MHz) the resistance was essentially constant throughout the test. Similar tests were also conducted on TiN [32], and

115 the results compared to Al-2Si when the samples were stressed with an AC with a period of 2 h. In contrast to Al-2Si, no damage healing was evident in TiN: the resistance showed a continuous increase with test time with no decrease in resistance during the current reversals. The damage in TiN was believed to be due not to electromigration, but possibly to thermal migration from temperature gradients across the length of the line. Under DC conditions TiN was able to withstand significantly higher current densities than Al-2Si due to the absence of significant electromigration.

In AC tests, when the AC frequency is less than 1/2MTTFDC (MTTF=median time to failure) the samples fail in the first half cycle and the AC lifetime is equal to the DC lifetime [30]. As the frequency is increased past 1/2MTTFDC the 2 lifetimes are no longer equal. At low frequency incomplete damage healing occurs due to finite physical changes that occur in each half cycle (change in grain size/shape, change in void size, etc.). Thus mass transport in opposite half cycles will not be mirror images of one another and damage will begin to accumulate, albeit more slowly than for a pure DC signal. As more complete damage healing occurs for higher frequencies, the AC lifetime will increase with frequency until very high frequency where the lifetime saturates (see Figure 11b). The high frequency lifetimes of Al-2Si and Cu were more than 1000 and 500 times higher than the DC lifetimes, respectively. The effect of increasing frequency on the time to failure is shown in Figure 11a-b for Al-2Si and Cu samples. For Al-2Si, an increase in lifetime of ~40 times was observed at 100 Hz for conditions of 250°C and 4.5×107 A/cm2. When the frequency was further increased to 1 MHz, the increase in lifetime was ~1400 times the DC lifetime. Increases in lifetime of ~1400 and >21000 times the DC lifetime were observed for frequencies of 100 Hz and 1 MHz, respectively, when the testing conditions were 275°C and 1.5×107 A/cm2. As can be seen from the data in Figure 11a-d, the increase in the AC lifetime in comparison to the DC lifetime is dependent not only on the frequency, but also on the applied current and the ambient temperature.

116

104 102 a) b)

101

102

100

-1 MTTF (h) MTTF 10 100 (min) MTTF

10-2 Al-2Si Cu 10-2 10-3 10-6 10-3 100 103 106 109 10-1 102 105 108 Frequency (Hz) Frequency (Hz) 7 10 4 2.0x10 c) AC-10 Hz d) AC-10 Hz 6 DC 10 DC

4 105 1.5x10

104

1.0x104

103

MTTF (s) MTTF MTTF (s) MTTF

2 10 3 5.0x10

101

0 0.0 10 0 1x107 2x107 3x107 4x107 5x107 100 150 200 250 300 350 Current Density (A/cm2) Temperature (°C) Figure 11 – Comparison of AC and DC electrical tests on Al-2Si (a-d) and Cu (a). (a) MTTF as a function of frequency for Al-2Si and Cu at a peak current density of 1.5×107 A/cm2 and an ambient temperature of 275°C, (b) MTTF as a function of frequency for Al-2Si at a peak current density of 4.5×107 A/cm2 and an ambient temperature of 250°C, (c) MTTF for AC and DC tests on Al-2Si as a function of the peak current density, and (d) MTTF for AC and DC tests as a function of the ambient temperature for a constant applied current. Data from [29-30].

F. Limitations to AC TMF Testing

While the benefits of the technique include the ease of sample preparation and equipment setup, similar service conditions to electrical components, and the ability to acquire a lot of data in a short amount of time, there are some limitations to the technique, which have been discussed by Monig et al. [3]. Temperature gradients are present along the length of the line due to the large contact pads. This can cause errors in determining the temperature by voltage measurements and may lead to thermal migration in addition to thermal fatigue. The higher temperature near the center of the lines causes damage to be more severe in this region than near the contact pads [16- 19]. Another problem is associated with the noise in the current and voltage signals as they pass through zero. This makes it more difficult to determine the minimum temperature. A thermocouple attached to the substrate near the sample may be used to determine the minimum temperature, but has been shown to read ~10-20°C lower than the actual metal minimum

117 temperature. Using this testing method the temperature and strain cannot be independently varied, as the strain is determined by the ΔT and the difference in thermal expansion coefficients between the metal line and the substrate. Varying the substrate material may lead to different combinations of ΔT and strain, but the Tmin and Tmax temperatures will also be affected, leading to different average sample temperatures, due to the different thermal conductivities and heat capacities of different substrates. Stress cannot be measured during testing, so if the stress in the line changes during testing, calibrations based on wafer curvature measurements made prior to testing will not be accurate and cannot be used to calculate the stress range experienced during testing. As the tests use a combination of high temperature and high current density, several factors may be contributing to damage formation such as electromigration, thermal migration, and other diffusive processes. The damage due to electromigration is expected to be minimal due to the relatively high frequency of the applied current (see above discussion), but as damage begins to accumulate temperature gradients will form and thermal migration may become important, especially towards the end of a test.

II. EXPERIMENTAL PROCEDURE

A. Sample Preparation

For testing the thermomechanical fatigue behavior of the materials, metal lines of varying widths were deposited on oxidized (100) silicon substrates with 1 μm of thermally grown oxide. The geometry of the metal lines is shown in Figure 12. The pads at the end of the line are for the application of the AC signal and the two pads to the side of the line are for measuring the voltage. The current was measured using an external resistor. For testing, six metal lines (all of the same line width) were grouped together on a chip of ~6 mm × ~6 mm. The chip layout is shown in Figure 13.

3-6 μm

200 μm 300 μm

300 μm

800 μm

Figure 12 – Geometry of the metal lines used to characterize the thermomechanical fatigue behavior of the thin films.

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Figure 13 – Layout of a group of six metal lines on a 6 mm chip. The 6 mm chip is affixed to a 40-pin package for wire bonding and then AC thermomechanical fatigue testing.

In the diagram of Figure 13, the small rectangular pads above and below the six lines serve two purposes: to identify the metal lines when viewing the samples in the SEM, and to check the thickness of the deposited samples using a profilometer. The large number in the center of the chip is the row-column designation from the 4-inch wafer. This number corresponds to a certain line width (1.5, 3.0, 4.5, or 6.0 μm). The numbered rectangular pads around the periphery are the gold pads of the package used for wire bonding.

The metal lines were fabricated using conventional photolithography methods, followed by film deposition, and then liftoff. Some samples were patterned with a single thick (~1.5 μm) layer of photoresist that could be removed with acetone. However, the majority of the samples were patterned with a two-layer resist that consisted of LOR5A on the bottom and SPR3012 on the top. Use of the LOR5A resist required liftoff at 60°C using Remover PG. After liftoff, samples were rinsed with acetone, methanol, and DI water, and dried with N2. The TMF samples were subsequently annealed at temperatures ranging from 400-1100°C in either a tube furnace for long anneals (4 h) or in a rapid thermal annealing furnace (RTA) for short anneals (< 10 min). All anneals (except where noted) were carried out in ultra-high purity (UHP) Ar with titanium or zirconium used to getter oxygen from the system.

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After annealing, the metal thickness was measured using mechanical profilometry (error in measurement of 2-3 nm), and the lines were examined using a light microscope to ensure that they were clean, completely lifted off, and free from debris or scratches. Then the resistivity of the samples as a function of temperature was measured using a probe station equipped with a hot stage (the standard deviation of the resistivity measurements was typically less than 0.1 μΩcm). Six measurements of the resistivity were obtained for each sample at room temperature, 50°C, 100°C, and 200°C. The average at each temperature was calculated and then the resistivity as a function of temperature for each 6 mm chip (6 TMF lines) was plotted and fit with a linear curve fit to obtain the temperature coefficient of resistance (TCR). For results from the resistance as a function of temperature measurements for the samples tested see Appendix I. These measurements were used to monitor the peak temperature during TMF testing. Next, the 6 mm chips were loaded and affixed with silver paint into a 40-pin side-brazed dual in-line ceramic package for wire bonding. All wire bonds were performed using Al-1Si wire with a 50 μm diameter (Figure 14).

Figure 14 – Section of a TMF sample attached to a 40 pin package by Al-1Si wire bonds. The square contact pads are 200 μm × 200 μm.

B. Sample Testing

The 40-pin package containing the wire bonded ~6 mm × 6 mm chip was inserted into a 40 pin universal zero insertion force dual in-line package test socket, which was attached to a printed circuit board that was also attached to 40 terminal blocks. The terminal blocks that were attached to the current pads were connected to an amplifier that was used to convert the voltage signal from a waveform generator to a current signal. The waveform generator was operated with a sinusoidal voltage with no DC offset. The terminal blocks for the voltage pads were connected to an SCXI voltage reader (SCXI 1000) so that the voltage drop across the sample could be monitored throughout the test. The voltage drop across the external resistor was also monitored using the SCXI 1000. The data from the voltage reader was collected in a text file

120 using a LabView program that recorded 5 values of the peak voltage every 10 or 30 s (10 s for the first 30 min, and every 30 s thereafter). The program stopped recording data if the detected frequency was not in the specified range, which occurred due to sample failure. The front panel for the LabView program is shown in Figure 15.

Figure 15 – Front panel for the LabView program that recorded voltage data during TMF testing.

For some samples the voltage drop was too high for the SCXI 1000 so a voltage divider circuit was included in the setup for those samples. Depending on the sample the voltage was divided by a factor of 6 or 21 (or 1 if no voltage divider was inserted) before being recorded by the voltage reader. A schematic of the test setup is shown in Figure 16.

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Voltage Divider Amplifier Waveform Generator – 100Hz SCXI Voltage Kepco BOP 36-6M 100 k Reader HP 33120A Kepco BOP 50-2M 5 k SCXI 1000

100 m

16 

Figure 16 – Schematic of the TMF setup shown with the 21x voltage divider circuit.

Once the correct sample was attached to the terminal blocks, the waveform generator was set to the testing frequency (100 Hz) and the voltage was set to the lowest value (50 mV). The amplifier, in current mode, was then turned on and the LabView program was started with a data collection rate of every 10 s. At this point the current output to the sample was ~32 mA and this was slowly increased by ~0.5 mA/s to the desired test settings by increasing the voltage output from the waveform generator. Once the desired current was reached the test was left to run until sample failure (open circuit), stopping the program only to switch the data collection interval from 10 s to 30 s after ~30 min of testing. After testing, the metal lines were examined in an environmental scanning electron microscope (FEI Quanta 200), where the failure location was imaged and measurements of the line width were obtained. Errors in the current and line width were less than 0.3 mA and 0.1 μm, respectively.

C. Temperature Measurement

Temperature during the TMF tests was obtained from 3 different techniques. The maximum temperature (Tmax) was obtained using measurements from an IR camera as well as from the recorded values of the peak voltage and current. The minimum temperature (Tmin) of the TMF lines was obtained from IR camera measurements, data from full waveform V(t) and I(t) data, and thermocouple measurements. The difference between Tmax and Tmin is the cyclic temperature range of the test, ΔT, and will determine the stress range during testing. Details from the three techniques are given below.

1. IR camera measurements – An IR thermal imaging camera (FLIR ThermaCAM S40) was used to image the TMF lines during testing, and the images were used to determine the temperatures obtained during testing. Samples were tested by incrementally increasing the current every 5 minutes and recording an image every 60 s during the 5 min hold starting at 0 s. Prior to imaging the samples during thermal testing, the emissivity of the wafers and samples was determined by heating the samples to various temperatures using a hotplate. Figure 17

122 shows an image obtained prior to calibrating the image for emissivity. Using the software from the IR camera allowed for calculation of the emissivity by selecting a region of known temperature and inputting values for the known sample and environmental temperatures. This was done for images similar to that shown in Figure 17 by choosing an area of the substrate and calculating the emissivity for the substrate (Si/SiO2=1) and then selecting one of the contact pads to calculate the emissivity of the metal (Au=0.32, Ru=0.12, NiAl=0.2). Once this was done the temperature of the substrate and the metal line could be determined separately by using the appropriate value of emissivity. The substrate temperature was used to estimate the minimum temperature of the metal line during testing, and the temperature of the metal line was used to determine the average temperature of the line (Taverage). The average value of the line temperature is determined rather than the peak temperature because the frame rate of the camera is not high enough to capture the temperature fluctuations due to the alternating current.

Figure 17 – IR image of a NiAl sample at a hotplate temperature of 100°C prior to calibration of the sample emissivity. The square contact pads are 200 μm × 200 μm.

Two types of line profiles were obtained from the IR images: one along the length of the line and the second perpendicular to the line. Line profiles along the line could be used to view the evolution of temperature with increasing current (see Figure 18), or the evolution of temperature with increasing time at a constant current. Both could be used to track “hot spots” that may be forming due to localized damage within the line.

123

200 a)

150 b)

100

Temperature (°C) Temperature 50

18 mA 24 mA 27 mA 32 mA 36 mA 40 mA 44 mA 0 0 200 400 600 800 Distance (m) c) d)

Figure 18 – (a) Line profiles of the average temperature along the length of a ruthenium line taken after 5 minutes at each current. The profiles run the length of the line from the top current pad (0 μm) to the bottom current pad (800 μm), which is shown by the white arrow in (b). The images in (b), (c), and (d) correspond to the IR camera images prior to testing, and after 5 minutes at 32 mA and 44 mA, respectively. The temperature scale in the IR images refers to the temperature of the silicon substrate.

The line profiles perpendicular to the line (Figure 19) were used to determine the temperature of the substrate and the difference in temperature between the substrate and the average line temperature. The temperature of the substrate immediately adjacent to the line was used as a measure of the minimum line temperature, and the average and minimum line temperatures were used to estimate the maximum line temperature. An example of the minimum, average, and maximum temperatures as a function of current is shown in Figure 20 for a sample of NiAl.

124

140

a) 34 mA 39 mA 120 44 mA 48 mA b) 51 mA 100 54 mA 61 mA

80

60 Temperature (°C) Temperature

40

20 -150 -100 -50 0 50 100 150 Distance from Sample (m) c) d)

Figure 19 – (a) Line profiles across the width of a NiAl line as a function of the applied current after 5 minutes at each current. The location of the line profile is shown as the white line in (b). The peak values were obtained using the emissivity of the metal line, whereas the rest of the data were obtained using the emissivity of the silicon substrate. The images correspond to IR images taken a) prior to testing, (b) after 5 minutes at 34 mA, and (c) after 5 minutes at 61 mA. The temperature scales in (b-d) are for the silicon substrate.

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250 T min T average T 200 max

150

100 Temperature (°C) Temperature

50

0 30 35 40 45 50 55 60 65 Current (mA) Figure 20 – Minimum, maximum, and average line temperature as a function of the applied current for a NiAl sample. Data obtained from the line profiles shown in Figure 19, where Tmin is estimated from the substrate temperature, Taverage is the peak in the line profiles, and Tmax is estimated by: - . The lines are curve fits using the function .

2. Peak resistance and thermocouple measurements – Measurement of the peak values of current and voltage were used to determine the maximum line temperature during TMF testing. This method was applied to every sample tested as the peak values of voltage and current were recorded for every test. The peak values of voltage and current were converted to peak resistance using Ohm’s law (V=IR), and since the resistance as a function of temperature was measured prior to testing, the temperature could be calculated from the resistance measurements. Prior to the testing of each sample set, one test was run where the current was incrementally increased every 5 or 10 minutes until sample failure. These tests were used to plot the peak temperature as a function of current, and also served as a current limit for each sample. For some samples, a thermocouple was attached to the silicon chip. The thermocouple was placed ~0.5-1 mm from the center of the line being tested, and values of the substrate temperature were recorded every minute during testing, and 30 s after an increase in current. Results from one of these tests are shown in Figure 21.

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500 60 500 T 55 max T 400 min 50 400

45 Current (mA) Current 300 40 40 60 80 100 300

200

50 ΔT 200

45 ΔT

Temperature (°C) Temperature Temperature (°C) Temperature 40 100

T - sample 35 max T substrate 30 0 40 50 60 70 80 90 100 30 35 40 45 50 55 60 65 Time (min) Current (mA) Figure 21 – Temperature as a function of time (left) and current (right) obtained during a step test on a sample of ruthenium sputtered at 1.5 mTorr. The inset to the figure on the left shows the current as a function of time during the step test (x-axis is the same in the main graph and the inset). The curve fits to the temperature as a function of current data use the equation: .

3. Whole waveform test – One sample was tested where instead of acquiring only the peak values of current and voltage, the entire waveform was obtained. Data were recorded for 0.025 s at a sampling rate of 10000/s every 10 s, which resulted in 2.5 AC cycles, or 5 thermal cycles, being recorded every 10 s. An example of the voltage and current as a function of time obtained during this test are shown in Figure 22.

15 90

10 60

5 30

0 0

Voltage (V) Voltage

-5 (mA) Current -30

-10 -60

-15 -90 0.000 0.005 0.010 0.015 0.000 0.005 0.010 0.015 Time (s) Time (s) Figure 22 – Voltage and current as a function of time acquired during a whole waveform test on a sample of co- sputtered RuAl.

The V(t) and I(t) data were converted to R(t) and T(t) using the same procedures as described in the previous section. As there was a slight offset (phase lag) between the voltage and current data (1-2 data points, 0.0001-0.0002 s) it was necessary to shift the current data so that the peaks in the two data sets were more closely aligned. The data for the temperature as a function of

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time were fit to a sine wave using , where A1, B2, C, and W1 are fit parameters referring to the amplitude, x-offset, y-offset, and peak width, respectively. The y- offset was necessary to account for a small DC component in the AC signal (~0.5 mA). An example of the T(t) data is shown in Figure 23a, which shows that there is some noise in the data in the region of the minimum temperature. This is due partly to the fact that this is where the current and voltage are zero, and also to a small amount of remaining offset in the voltage and current data. The curve fits were used to extract the minimum and maximum line temperature as a function of test time, which are shown in Figure 23b for the first 30 min of testing.

500 470 T min T max 400 460

450 300 440

200 430

70

Temperature (°C) Temperature Temperature (°C) Temperature

100 60 a) b) 0 50 1206.400 1206.405 1206.410 1206.415 1206.420 1206.425 0 5 10 15 20 25 30 Time (s) Time (min) Figure 23 – (a) 5 thermal cycles acquired during the whole waveform test and the curve fit to the data, and (b) the maximum and minimum temperature as a function of time obtained from curve fits similar to those shown in (a). The solid and dotted lines in (b) are curve fits to the data using the equation T=atb, where a and b are fit parameters.

4. Comparison of the three techniques – For samples where more than one of the above techniques were used, it was possible to compare the results obtained from them. This is shown in Figure 24 for a sample of NiAl, which was analyzed with resistance measurements, thermocouple measurements, and IR images. The minimum temperature results agree well with one another and there are only a few degrees difference between the thermocouple and IR data. There is more discrepancy between the maximum temperature data and this is mostly due to the 18 micron pixel size of the IR camera compared to the ~3 micron line width. Thus the IR data for the maximum temperature would be expected to be a combination of the line and substrate temperature, and will not accurately predict the maximum temperature of the line. A comparison of the temperature data obtained from tests on NiAl, Ru, and RuAl is shown in Figure 25. The data in the plot shows that the ΔT obtained during a given test is closely correlated with Tmax, and can be estimated just by measuring the peak temperature. As the peak temperature was measured in every test, this allows the ΔT to be calculated even for tests where the minimum temperature was not measured.

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500 T - IR camera max T - IR camera substrate T - R measurement 400 max T - thermocouple substrate

60

50 300

40

30

20 200 30 40 50 60

Temperature (°C) Temperature

100

0 25 30 35 40 45 50 55 60 65 Current (mA) Figure 24 – Temperature as a function of the applied current for a NiAl sample, where the temperature was measured using resistance, thermocouple, and IR measurements. The inset shows an enlarged view of the Tmin measurements (axes and units are the same in the main graph and the inset).

500 NiAl - resistance + thermocouple Ru - resistance + thermocouple 400 RuAl - whole waveform

300 T (°C) T  200

100 From a curve fit to all 3 data sets: ΔT = 0.90459*Tmax – 15.10665 0 0 100 200 300 400 500 600 T (°C) max

Figure 25 – The ΔT (Tmax-Tmin) as a function of the peak temperature for NiAl and Ru determined using peak resistance and thermocouple measurements, and for RuAl using the whole waveform test.

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D. Data Analysis

After the sample failed, the voltage and current data were converted to resistance and maximum temperature data. The maximum temperature data from the first minute of the test were averaged and this value was used as the peak temperature for the test and also to calculate the ΔT for the test. As the temperature is not constant throughout the test, and for most samples continues to increase until failure, the peak temperature and ΔT will be underestimated for most samples, and can be considered as a lower bound. A few samples showed a decrease in resistance at the beginning of the test, which would result in a temperature decrease, therefore for these samples the peak and ΔT will be slightly overestimated. The reason for using the temperature calculated early in the test is to ensure that the resistance calculated from the I-V data only has contributions from Joule heating and not from annealing or damage formation and accumulation. As damage/annealing occurs in the sample, the resistance of the sample will change. This means that the calibrations of R(T) performed prior to testing will no longer be valid and will lead to inaccuracies in the calculation of temperature. The error in using the initial R(T) calibration was evaluated by stopping a test after a 15% increase in resistance was observed and then re-measuring the resistance as a function of temperature. The results are shown in Figure 26, which shows that a 15% increase in resistance corresponds to a 6.5% increase in resistivity once the sample has been returned to room temperature. The remainder of the increase in resistance is due to Joule heating during testing. Due to the changed sample resistivity, the temperature would be overestimated by ~25-45°C by using the initial calibrations depending on the testing temperature. Prior to a 15% increase in resistance (during testing) the errors will be smaller than this, with larger errors as the resistance increases beyond 15%.

36 a) 50 b)

32 45 cm)

 40 28

35 Resistivity ( Resistivity

24 30 Error in Temperature (°C) Temperature in Error before thermal cycling 25 after cycling to a 15% increase in R 20 0 50 100 150 200 250 100 200 300 400 500 600 700 Temperature (°C) Predicted Temperature (°C) Figure 26 – (a) Resistivity as a function of temperature calibrations performed prior to testing and after TMF testing to a 15% increase in resistance. (b) The error in temperature obtained after a 15% increase in resistance has been observed if the initial resistivity as a function of temperature curve is used to calculate the temperature. The above tests were performed on a Ni/Al multilayer sample annealed at 800°C. Error bars are within the size of the symbols.

The curves of resistance as a function of testing time were typically plotted as the instantaneous resistance divided by the value of the first measured resistance (R/R1) so that tests with varying temperature could all be viewed in the same plot. All curves are plotted using a five point

130 adjacent averaging (each data point is replaced by the average of 5 data points), and were curve fit using an equation that was developed that can describe all of the different trends noted in the data. The equation consists of exponential, power, and linear terms and will be discussed in more detail in the Results and Discussion.

III. RESULTS AND DISCUSSION

This section details the results of TMF tests including R(t) curves, the time to failure as a function of current density and temperature, and SEM observations of the samples after failure. The majority of the R(t) curves are plotted as the relative resistance R/R1, where R1 is the first value of resistance recorded during the AC TMF test. Each material will be discussed in turn followed by a comparison of the results from the different materials tested. Due to different values of the CTE for the different materials tested, equal values of ΔT will result in different values of stress. All of the tested materials have CTE’s (NiAl ~13×10-6/°C, Au 14.2×10-6/°C, RuAl 5.5×10-6/°C, and Ru 6.4×10-6/°C) [33-35] that are larger than Si (2.6×10-6/°C) [33], so the stresses will become more compressive as the temperature increases. The difference in CTE between RuAl/Si (112%) is much smaller than the difference in CTE between NiAl/Si (400%), so more stresses are likely to be generated in NiAl than in RuAl. The section will conclude with a detailed analysis of the shape of R(t) curves and a comparison between AC and DC results obtained on samples of RuAl. The value of Tmax plotted in the figures is the average of the temperature obtained from voltage measurements recorded during the first minute of testing. The error in this value (the standard deviation) was typically ± 8-10°C (sometimes as low as ± 3°C), but for some samples was as high as ± 18-24°C (Ru deposited at 7 mTorr). All values given for the resistivity are the average of the six measurements obtained at room temperature from probe station measurements on individual patterned lines. The standard deviation of the resistivity for each line was typically less than 0.1 μΩcm, with a higher line-to-line variability. Values for the thickness correspond to the average of at least 8 measurements with a standard deviation of 2-3 nm.

A. Ruthenium

Prior to the testing of the NiAl and RuAl intermetallics, several samples of sputtered ruthenium lines were examined. The deposition conditions, sample dimensions, and testing conditions are summarized in Table 1. All of the samples were annealed at 400°C for 4 h.

131

Table 1 – Details of the ruthenium samples tested and the testing conditions.

Deposition AC current Thickness Line width Resistivity AC current Time to Material pressure density T (°C) Comments (nm) (μm) (μΩcm) (mA) max failure (h) (mTorr) (MA/cm2) Ruthenium 1.5 195 3.2 18.1 74.7 11.97 497 .03 1.5 195 3.2 18.6 72.4 11.60 477 .27 1.5 195 3.2 19.2 70.0 11.22 446 1.7 1.5 195 3.2 19.7 66.4 10.64 427 2.4 1.5 195 3.2 20.1 47.4 7.60 216 337.2 no failure Ruthenium 7 287 2.8 52.1 59.9 7.45 × 0.006 7 287 2.8 53.6 62.9 7.83 273 0.02 7 287 2.8 53.0 65.8 8.19 260 0.067 7 287 2.8 53.7 59.9 7.45 305 0.14 7 287 2.8 48.9 59.3 7.38 310 3.1 7 287 2.8 48.8 58.7 7.30 328 13.6 7 287 2.8 53.9 56.9 7.08 331 30.5

The resistance as a function of time curves are shown in Figure 27a for the samples deposited at 1.5 mTorr and in Figure 27b for the samples deposited at 7 mTorr. As is shown in the figures, the two ruthenium sample sets show distinctly different R(t) curves: the low pressure samples show only an increase in resistance until failure, whereas the high pressure samples show an initial decrease in resistance followed by an increasing resistance until failure. The different behavior of the two samples is believed to be due to microstructural differences due to changing the sputtering pressure. Decreasing the sputtering pressure typically results in a film with a denser microstructure, which was clearly shown in the study by Thomas [36] for platinum films that had a density ranging from 9-20.8 g/cm3 depending on the sputtering pressure. A microstructure with fewer voids at the grain boundaries would explain the difference in resistivity of the ruthenium films, where the sample deposited at high pressure has a significantly higher resistivity (~2.5 times) than the low pressure sample, indicating that the high pressure sample has more sites that scatter electrons. It is possible that the temperatures during cycling are sufficient to impart additional annealing to the more resistive sample, thus decreasing its resistivity/resistance. Once damage begins to form, the decrease in resistance slows until the effects of damage formation outweigh those from annealing and the resistance begins to increase. The resistivity of the samples deposited at low pressure is initially fairly low, so any changes in resistance due to annealing are expected to be small and are outweighed by heating and damage formation.

132

210 1.08 a) 47.4 mA b) 56.9 mA 66.4 mA 58.7 mA 200 70.0 mA 1.04 59.3 mA 72.4 mA 59.9 mA 74.7 mA 62.9 mA 190

) 1.00 65.8 mA 

180

1 0.96

R/R 170

0.92 Resistance ( Resistance

132 130 0.88 128 126 0.84 0 1 2 3 0.01 0.1 1 10 Time (h) Time (h) Figure 27 – Resistance as a function of time for ruthenium sputtered at (a) 1.5 mTorr and (b) 7 mTorr.

The ruthenium samples with the lower resistivity (low pressure deposition) can sustain higher current densities than the samples with a higher resistivity (see Figure 28a). This is a result of decreased Joule heating for the lower resistivity samples. The low pressure samples were also able to withstand higher temperatures than the higher pressure samples (Figure 28b). Reasons for this could be due to microstructural differences as discussed above, or due to the reduced thickness of the 1.5 mTorr samples, which would tend to increase the strength of the sample.

14 Ru - 1.5 mTorr Ru - 1.5 mTorr

Ru - 7 mTorr 500 Ru - 7 mTorr

) 2 12

400

10 (°C)

max T

300

8 Current Density (MA/cm Density Current 200 a) b) 6 1E-3 0.01 0.1 1 10 100 1000 0.01 0.1 1 10 100 1000 Time to Failure (h) Time to Failure (h) Figure 28 – Comparison of the lifetime data for the two sets of ruthenium samples as a function of (a) current density and (b) the initial peak temperature during AC thermal fatigue testing.

SEM images of failed lines of ruthenium are shown in Figure 29 for samples deposited at 1.5 mTorr and in Figure 30 and Figure 31 for samples deposited at 7 mTorr. The damage is localized to the region of failure as was seen for the majority of the samples tested; however, the damage is more severe than was noted for most of the other samples tested in this work. Most of the ruthenium samples showed cracking near the failure location as well as surface wrinkling. Cracking was rarely noted in any of the other samples tested in this work as well as in the

133 literature. Thinning was also noticed in one sample (Figure 31b) as was some voiding/dimpling (Figure 30c). The failures all tended to be violent, with significant melting and loss of material in the vicinity of the failure.

a) b)

10 μm

Figure 29 – SEM images of a ruthenium line deposited at 1.5 mTorr and tested at 66.4 mA (10.64 MA/cm2) until failure, which occurred after 1,728,000 thermal cycles. (a) Low magnification image of the region where failure occurred. (b) High magnification image of the bottom of the failure region.

134

a) b) c)

2 μm 1 μm

Figure 30 – SEM images of a ruthenium line deposited at 7 mTorr and tested at 56.9 mA (7.08 MA/cm2) until failure, which occurred after 21,960,000 thermal cycles. (a) and (c) show high magnification images of regions near the bottom and top of the failure, respectively. (b) Low magnification image of the region where failure occurred.

a) b)

2 μm

Figure 31 – SEM images of a ruthenium line deposited at 7 mTorr and tested at 58.7 mA (7.30 MA/cm2) until failure, which occurred after 9,792,000 thermal cycles. (a) Low magnification image of the region where failure occurred. (b) A high magnification image of the bottom side of the failure region.

135

B. NiAl

1. NiAl from annealed e-beam deposited 20 nm wavelength multilayers – Samples of NiAl fabricated from annealed Ni/Al multilayers were stressed via AC. Details on the deposition and testing conditions are given in Table 2. Two different annealing conditions were used: the standard 4 h anneal at 400°C, and an anneal that included 10 min at 700°C after the 400°C anneal. This second anneal was added in an effort to decrease the resistivity in order to decrease the Joule heating and in turn increase the current density. The second anneal was successful in decreasing the resistivity of the NiAl lines by 17%, and it was expected that this would increase the current density of the lines due to decreased Joule heating. However, this was not the case as the higher resistivity lines were able to withstand a higher current density for similar times to failure. Reasons for this will be discussed below.

Table 2 – Details on the multilayers deposited and tested. NF refers to samples that did not fail. AC current Thickness Line width Resistivity AC current Time to density T (°C) (nm) (μm) (μΩcm) (mA) max failure (h) (MA/cm2) annealed at 400°C for 4 h 293 3.1 24.4 85.1 9.68 390 0.73 293 3.1 24.0 82.2 9.35 411 1.73 293 3.1 23.6 74.6 8.49 350 259.7 (NF) 293 3.1 24.6 60.9 6.93 253 306.9 (NF) annealed at 400°C for 4 h followed by 700°C for 10 min 293 2.9 19.7 76.0 8.65 512 1.05 293 2.9 20.3 71.1 8.09 489 16.8

Curves showing the resistance change as a function of testing time are shown in Figure 32a for the samples annealed with the standard low temperature anneal and for the samples including the additional high temperature anneal. The two sets of samples show distinctly different R(t) curves. The samples annealed at low temperature show a region of decreasing resistance followed by increasing resistance until failure, whereas the samples annealed at high temperature show only an increasing resistance until failure. The samples annealed at low temperature were tested near or above the annealing temperature, so the region of decreasing resistance is likely due to additional annealing. The samples annealed at high temperature were tested at temperatures significantly lower than the annealing temperature, so no region of decreasing resistance occurred.

136

1.20 10.0 600 a) 74.6 mA b) 82.2 mA 4 h 400°C 1.15 85.1 mA 500

+ ) 9.6 2 71.1 mA 10 min 700°C (°C)

76.0 mA max 400 1.10 T

1.15 9.2 300 0.1 1 10 100 1000 1 1.10 Time to Failure (h)

1.05 1.05

R/R

1.00 8.8 1.00 0.95 0.90 0.0 0.5 1.0 1.5 2.0 2.5 8.4 0.95 (MA/cm Density Current 400°C anneal 400°C + 700°C anneal 0.90 8.0 0 2 4 6 8 10 12 14 16 0.1 1 10 100 1000 Time (h) Time to Failure (h) Figure 32 – (a) Relative resistance as a function of time for the e-beam multilayer NiAl. The inset shows the same data, but zoomed in to highlight tests with shorter lifetimes (axes and units are the same in the main graph and the inset). (b) Lifetime plot as a function of current density. Temperature data are included in the inset.

Figure 32b shows lifetime plots with respect to current density and temperature (in the inset). The samples annealed at low temperature show the ability to carry an increased current density in comparison to the samples annealed at high temperature. This is opposite to what would be expected as the samples annealed at high temperature have a lower resistivity and thus decreased Joule heating. Two possible reasons for this observation are: 1) a higher tensile residual stress in the samples annealed at high temperature, which can be seen in SEM images as a partial loss of adhesion and cracking in the contact pads (see Figure 33), and 2) the annealing that occurs during testing in the samples annealed at low temperature. As the samples anneal, the resistance decreases and the temperature is reduced due to less Joule heating, which in turn reduces the cyclic stresses experienced by the metal line. In contrast, for the samples annealed at high temperature, the increasing resistance leads to increased Joule heating and a temperature (and stress) that increases throughout the test. The data on the time to failure as a function of temperature indicate that the samples annealed at low temperature may not be able to withstand temperatures as high as the samples annealed at high temperature. However, as will be shown later, the scatter in the thermal fatigue data is fairly large, and the apparently large difference between the temperature data for the two samples (inset to Figure 32b) may be an artifact of the small sample set. It was also difficult to run the samples annealed at low temperature at temperatures in excess of 400°C as the samples immediately began to anneal as soon as the temperature exceeded 400°C, thus decreasing the temperature.

An SEM image of a failed line is shown in Figure 33 for one of the samples annealed at high temperature. Samples from the low temperature anneal looked similar except there was no edge curling in these samples. The edge curling is due to high tensile stresses induced during cooling from the annealing temperature due to the difference in thermal expansion coefficient between NiAl and the substrate. Not much damage was noted along the lines after AC testing. The only damage was noted in the region surrounding the failure location and included a few small depressions. Another feature noticed along the edges of the lines is what appears to be grain growth. Similar features have been noted in many of the other NiAl samples.

137

a) b)

100 μm 2 μm

Figure 33 – SEM images of a failed NiAl TMF sample tested at 76.0 mA (8.60 MA/cm2). Failure occurred after 756,000 thermal cycles. The sample was annealed at 400°C followed by a 700°C anneal. (a) Low magnification image of the entire sample, and (b) a higher magnification image of the failure location.

2. NiAl from annealed sputter deposited 30 nm wavelength multilayers: deposition 1 – The results for annealed 30 nm wavelength NiAl multilayers are shown in Table 3 for two annealing conditions: the standard low temperature anneal (400°C for 4 h) and the standard anneal followed by a high temperature anneal (800°C for 30 s).

Table 3 – Sample details, testing conditions, and results for 30 nm wavelength annealed Ni/Al multilayers. AC current Thickness Line width Resistivity AC current Time to density T (°C) (nm) (μm) (μΩcm) (mA) max failure (h) (MA/cm2) annealed at 400°C for 4 h 252 3.3 23.3 74.7 8.98 464 0.1 annealed at 400°C for 4 h followed by 800°C for 30 s 256 3.8 25.8 88.4 9.10 591 0.02 256 3.8 24.9 85.4 8.79 488 0.12 256 3.8 20.4 83.6 8.61 385 8.1 256 3.8 22.1 82.5 8.49 407 9.1

138

Although only 1 sample was available for testing from the low temperature anneal, no significant difference between this sample and those annealed at higher temperature was noted between the resistivity, time to failure as a function current density, time to failure as a function of temperature, R(t) curves, or damage morphology of the failed lines. The resistance increase as a function of testing time is shown in Figure 34a for all of the samples tested from this sample set. All of the samples show an increasing resistance with increasing time until failure. Figure 34b-d shows SEM images of the sample annealed at low temperature as well as two of the samples annealed at high temperature. Figure 34b shows a magnified view of a region near the failure location showing a series of ridges that lie towards the center of the line. This feature has been seen in the majority of NiAl samples, and is also present in the sample shown in Figure 34d, which presents another common feature of failed NiAl samples: voids/regions of surface depression. A very uncommon occurrence is shown in the image in Figure 34c, which shows two cracks in the area adjacent to the failure. Cracking was present in only one other NiAl sample. All three images show the feature that could be grain growth along the line edges with Figure 34c showing enlargement of these features as the failure location is approached. This is not surprising as the temperature of the line will increase and approach the melting temperature of NiAl at the failure location.

139

1.30

1.16 a) b) 1 μm 1.25 1.12

1.08

1.20 1.04

1.00

0.00 0.05 0.10 0.15 0.20 1

1.15

R/R

1.10 74.7 mA 4 h 400°C 82.5 mA 74.7 mA 1.05 + 83.6 mA 9.0 MA/cm2 85.4 mA 30 s 800°C 88.4 mA 72,000 cycles 1.00 0 1 2 3 4 5 6 7 8 9 Time (h) c) 85.4 mA 82.5 mA d) 8.8 MA/cm2 8.5 MA/cm2 86,400 cycles 6,552,000 cycles

2 μm 2 μm

Figure 34 – (a) Relative resistance as a function of AC testing time for annealed 30 nm wavelength Ni/Al multilayers with the low temperature and high temperature anneal. The inset shows the same data, but zoomed in to highlight the tests with shorter times to failure (axes and units are the same in the main graph and the inset). (b) SEM image of a region near the failure site for the sample annealed at low temperature. (c) and (d) SEM images of 2 of the samples annealed at high temperature after TMF testing.

3. NiAl from annealed sputter deposited 30 nm wavelength multilayers: deposition 2 – Another set of 30 nm wavelength Ni/Al multilayers were deposited and tested. These samples are being discussed separately from the previous 30 nm wavelength samples as the line width is larger, and the samples were all annealed at higher temperatures (800-1000°C). Multiple annealing profiles were used with the ultimate goals being: 1) using higher temperatures to reduce the resistivity and increase current density, 2) keeping the time at high temperature low to minimize reactions with the substrate, which has been seen to occur at annealing temperatures of 1000°C, and 3) more slowly cooling the samples in an effort to decrease film stresses. Details of the samples, annealing profiles, testing conditions, and results are given in Table 4.

140

Table 4 – Sample details for the second set of 30 nm wavelength Ni/Al multilayers, the testing conditions, and TMF test results. All heating ramp rates were 50°C/s. AC current Thickness Line width Resistivity AC current Time to density T (°C) (nm) (μm) (μΩcm) (mA) max failure (h) (MA/cm2) anneal 1: 800°C(30 s) – 10°C/s – 500°C(30s) 259 4.6 23.4 94.0 7.89 495 0.57 259 4.6 23.4 91.6 7.69 504 0.8 259 4.6 22.0 89.2 7.49 387 155.88 259 4.6 21.3 90.5 7.60 367 360.95 anneal 2: 1000°C(30 s) – 10°C/s – 500°C(30 s) 258 4.5 48.8 78.9 6.80 606 0.01 258 4.5 44.9 77.7 6.69 504 0.2 258 4.5 46.2 74.2 6.39 469 2.4 258 4.5 40.7 76.5 6.59 429 6.7 anneal 3: 400°C(5 min) – 1000°C(30 s) – 10°C/s – 500°C(30 s) 256 4.8 44.8 82.5 6.71 510 0.01 256 4.8 43.9 81.3 6.62 510 0.27 256 4.8 45.1 78.9 6.42 470 0.47 256 4.8 42.9 80.1 6.52 465 1.6 256 4.8 46.2 76.5 6.23 476 3.75 anneal 4: 1000°C(30 s) – 5°C/s – 900°C(15 s) – 2°C/s – 800°C(15 s) – 2°C/s – 600°C(30s) 260 4.4 45.7 74.2 6.79 471 0.08 260 4.4 55.7 73.0 6.38 561 0.4 260 4.4 45.9 71.8 6.00 438 10.6

The data in Table 4 shows that anneal 1 results in a resistivity much lower than anneals 2-4, and that the variability in the resistivity is much lower for anneal 1. The decreased resistivity results in decreased Joule heating and an increased current density for anneal 1 compared to the other three anneals, which all show very similar results (see Figure 35a). However, the time to failure as a function of temperature does not appear to be affected by the different anneals (Figure 35b). The increased resistivity (decreased current density) is likely due to reaction with the substrate (see details in Chapter 2) in the samples annealed at 1000°C (anneals 2-4); however, this does not seem to have a deleterious effect on the time to failure as a function of temperature.

141

8.0 650 a) b) anneal 1 anneal 2

7.6 600 anneal 3 ) 2 anneal 4

7.2 550

6.8 (°C) 500

max T 6.4 450

anneal 1

Current Density (MA/cm Density Current 6.0 anneal 2 400 anneal 3 anneal 4 5.6 350 1E-3 0.01 0.1 1 10 100 1000 1E-3 0.01 0.1 1 10 100 1000 Time to Failure (h) Time to Failure (h) Figure 35 – Lifetime data plotted as a function of (a) current density and (b) temperature for the samples from anneals 1-4.

The curves for the resistance as a function of time are all very similar: the resistance in all cases increases with increasing time until failure. Examples are shown in Figure 36a for the samples from anneal 1. These data were used to plot the time necessary to reach a given amount of resistance increase (Figure 36b).

1.35 103 89.2 mA a) b) 15% increase in R 1.30 89.3 mA 10% increase in R 90.5 mA 6.8 % increase in R 91.6 mA 2 10 1.25 94.0 mA

1.20 1 1 10

1.3

R/R

1.15 Time (h) Time

1.2

1.10 0

10 1.1 1.05

1.0 0 1 2 -1 1.00 10 0 100 200 300 400 350 400 450 500 550 Time (h) Initial Temperature (°C) Figure 36 – (a) Relative resistance as a function of testing time (the inset shows a zoomed in view of the data at short times) and (b) the time necessary to reach a given amount of resistance change as a function of the peak temperature at the start of the test. Data are for samples from anneal 1.

An example of a failed sample from anneal 1 is shown in Figure 37, which shows two of the typical features of NiAl TMF samples: ridges along the middle of the line, and voids. Both are only noticed near the failure and extend up to ~20 μm in either direction. The lateral extent of the ridged features seems to increase as the site of the open circuit is approached.

142

89.3 mA 7.5 MA/cm2 576,000 cycles

5 μm

Figure 37 – SEM image of the failure region in a sample from anneal 1.

SEM images of samples from anneals 2 and 4 are shown in Figure 38, which show some similarities and differences to images of the samples from anneal 1. Ridged features, surface depressions, and voids are noted in samples from all of the anneals. However, the number of surface depressions does appear to be larger in the samples from anneals 2-4. One feature noted in samples from anneals 2-4 that was not present in the samples from anneal 1 is the serration along the edges of the lines. This is believed to be due to stresses generated during annealing and not due to the thermal cycling, as this feature was noted around the edges of the bond pads as well as along the line edges. Remarkably, even with the poor line morphology and the incorporation of silicon in the samples from anneals 2-4, the time to failure as a function of temperature is the same as the samples from anneal 1.

143

a) 74.2 mA 6.8 MA/cm2 c) 59,760 cycles

1 μm 1 μm 71.8 mA b) 77.7 mA 6.0 MA/cm2 6.7 MA/cm2 10,476,000 cycles 144,000 cycles

2 μm 2 μm

Figure 38 – SEM images of failed samples from (a) anneal 4, (b) anneal 2, and (c) anneal 4, where the top image is an enlarged region to the right of the failure and the bottom image shows the failure location.

4. NiAl from annealed sputter deposited 30 nm wavelength multilayers: doped with 0.5% Ag – A series of experiments were conducted on NiAl doped with 0.5% Ag in an effort to increase the strength without significantly increasing the resistivity of the films [37]. Blanket films of NiAl- 0.5Ag after annealing at 400°C for 4 h had a resistivity of 15.7 μΩcm in comparison to 15.5 μΩcm for un-doped films. From annealing experiments it was expected that the doped films were higher in strength than the un-doped films. The silver doped samples did not crack when annealed and rapidly cooled from 1000°C; whereas the un-doped films were severely cracked after identical annealing conditions (see Figure 39).

144

a) b)

100 μm 100 μm

Figure 39 – Nomarski light microscope images of the current pads of annealed Ni/Al multilayer TMF samples (a) un-doped and (b) doped with 0.5% Ag. The samples were annealed at 1000°C for 30 s.

Samples of NiAl-0.5Ag were fabricated from annealed 30 nm wavelength Ni/Al multilayers. Details of the deposited samples are given in Table 5 – Details of the NiAl-0.5Ag samples tested and the testing results. All of the samples were annealed at 400°C for 4 h, which resulted in a sample resistivity of 22.7 ± 2.0 μΩcm. While not the lowest obtained for NiAl films, it is one of the lower values of resistivity. It is also notable that there is some variability in the line width, which ranges from 3.4-4.0 μm. As will be demonstrated in this and the subsequent sample set, the line width significantly affects the current density by altering the I(T) behavior.

Table 5 – Details of the NiAl-0.5Ag samples tested and the testing results. AC current Thickness Line width Resistivity AC current Time to density T (°C) Comments (nm) (μm) (μΩcm) (mA) max failure (h) (MA/cm2) 258 3.8 22.5 89.6 9.16 × ~0.000001 immediate failure 266 4.0 26.8 90.2 8.48 × ~0.000001 immediate failure 261 3.4 20.3 86.0 9.69 × 0.00028 narrower line 261 3.4 20.4 81.9 9.81 428 0.1 narrower line 263 4.0 22.0 91.9 8.75 438 0.38 258 3.8 21.2 87.8 8.97 473 0.88 263 4.0 21.6 89.0 8.47 420 0.97 261 3.4 21.1 84.2 9.49 444 1.27 narrower line 263 4.0 23.3 87.8 8.36 431 1.73 266 4.0 26.3 90.8 8.54 452 1.97 266 4.0 23.4 90.2 8.48 448 4.1 266 4.0 22.2 89.0 8.37 418 5.4 258 3.8 21.9 86.6 8.85 433 5.65 258 3.8 25.1 84.8 8.67 433 140.4 no failure

145

Results from TMF testing of the NiAl-0.5Ag samples are shown in Figure 40. A first examination of the data showed a significant amount of scatter in the time to failure, but when the changing line width is taken into account the data can be separated into three regions. The narrowest lines are able to withstand a higher current density than the wider lines (Figure 40a). The reason for this has to do with the evolution of temperature with increasing current density, and how this changes with changing line width. Figure 40b shows that to obtain an equal temperature during TMF testing, a higher current density is required for narrower lines. This is why the narrower lines are able to carry a higher current. However, if this is plotted as the applied current the situation is reversed: the narrower lines are hotter at equal currents and therefore have a shorter time to failure. The above discussion seems counterintuitive until you examine the relationship between ΔT and the applied power, which shows that ΔT is proportional

to the applied power: , where A1 and A2 are constants (see

Figure 4 in the “Introduction” to this chapter). This shows that as the line width decreases the temperature increases if the current is kept constant. If this equation is written in terms of the current density j, then Eq. 6 is obtained, which shows that as the line width decreases the temperature decreases if the current density is kept constant.

Eq. 6

When the data is plotted as time to failure versus temperature (inset to Figure 40a), no effect of the changing line width is observed. The above discussion shows that depending on the metric used to compare the data the effect of decreasing the line width can be beneficial (increased current density), detrimental (decreased current), or has no effect (temperature behavior unaffected). This shows the importance of careful sample preparation and in keeping as many variables constant as possible in order to make comparisons between different samples. This also makes it difficult to determine the effect of silver doping on the TMF behavior of NiAl as the un-doped NiAl samples have different values of thickness, line width, and resistivity. The closest match to the NiAl-0.5Ag samples is included in the data of Figure 40a. Unfortunately, only 1 sample is available for comparison, so more tests are needed to determine the effects of silver doping on NiAl, but silver doping does not appear to have affected the performance in terms of the cyclic temperature.

146

10.0 500 a) NiAl-0.5Ag b) NiAl

) 3.4

2 9.6 480

9.2 460

(°C) 3.8 4.0 m NiAl-0.5Ag  3.8 m NiAl-0.5Ag max

3.4 m NiAl-0.5Ag T 8.8 500 440 3.3 m NiAl 3.8 4.0 3.4 475 4.0

(°C) 450

8.4 max 420

Current Density (MA/cm Density Current

T

425

400 0.01 0.1 1 10 100 1000 8.0 400 1E-5 1E-3 0.1 10 1000 8.0 8.5 9.0 9.5 10.0 2 Time to Failure (h) Current Density (MA/cm ) Figure 40 – Results from the NiAl-0.5Ag TMF tests. The numbers within the circled regions refer to the line width of the samples in μm. (a) The time to failure as a function of the current density. The data plotted as a function of temperature is shown in the inset. The data from a 3.3 μm wide un-doped NiAl sample is also included. (b) The temperature during TMF testing as a function of the applied current density. The circled data point in the 3.8 μm region corresponds to one of the 4.0 μm samples.

Several curves showing the resistance as a function of time are presented in Figure 41a, which shows that the trend in R/R1 depends on the peak temperature at the start of the test. If the peak temperature is within ~30°C of the annealing temperature then the resistance increases until failure. When the testing temperature is more than ~30°C higher than the annealing temperature then there is an initial region of annealing (decreasing resistance) prior to the resistance increasing until failure. An increasing resistance may not be observed if the sample fails outside of the region encompassed by the voltage pads. Figure 41b-d shows images of samples of NiAl- 0.5Ag after TMF testing until failure. The doped samples show the same typical surface morphology noted in most of the un-doped NiAl samples: voiding, ridged details, and grain growth along the line edges.

147

1.20 a) 118 b) 117 Ti = 433 C 1.15 Ti = 418 C 116

115 114 1.10 113 112 0 1 2 3 4 5 6

1.05

1

R/R 1.00

Ti = 448 C 0.95 86.6 mA 89.0 mA 0.90 Ti = 452 C 2 90.2 mA 2 μm 8.85 MA/cm 90.8 mA 4,068,000 cycles 0.85 0 1 2 3 4 5 6 Time (h) c) 84.2 mA d) 128 9.49 MA/cm2 124

120 914,400 cycles 116 112 Ti = 420 C 0.0 0.5 1.0

131 130 Ti = 444 C 129 128 89.0 mA 2 127 8.47 MA/cm 2 μm 2 μm 126 698,400 cycles 0.0 0.4 0.8 1.2 Figure 41 – (a) Relative resistance as a function of testing time for 3 samples of NiAl-0.5Ag. (b-d) SEM images of samples tested until failure. The insets show the resistance (in Ω) as a function of time (in h) for the samples in the images. Ti in (a-d) refers to the peak temperature measured at the start of the test.

5. Co-sputtered NiAl – Several samples of NiAl co-sputtered at 1.5 mTorr were tested with AC until failure. All samples were annealed at 400°C for 4 h prior to testing. The line width was varied to study the effects of changing geometry on the AC thermal lifetime. Table 6 shows details on the NiAl co-sputtered samples and the TMF results.

148

Table 6 – Overview of the co-sputtered NiAl samples and the TMF results. AC current Thickness Line width Resistivity AC current Time to density T (°C) (nm) (μm) (μΩcm) (mA) max failure (h) (MA/cm2) 299 3.4 41.35 73.4 7.34 547 0.33 299 3.3 39.45 72.4 7.24 426 4.15 299 3.3 39.89 70.6 7.06 421 24.83 298 5.2 38.12 103.2 6.67 475 0.27 298 5.2 38.19 106.2 6.86 501 1.32 298 5.2 39.07 100.3 6.48 479 33.63 298 6.4 37.86 121.0 6.40 461 1.18 298 6.4 37.42 124.0 6.56 488 5.9 298 6.4 37.31 118.1 6.25 425 137.47

Representative curves showing the resistance as a function of time for the co-sputtered samples are shown in Figure 42a, which shows that the resistance increases until sample failure. This occurred for all of the samples even though the testing temperature was higher than the annealing temperature. This means that 1) the samples are not annealing during the test, or 2) the effects of damage formation outweigh the effects of annealing. To determine whether or not any grain growth is occurring during testing, microstructural studies are needed.

The time to failure of the co-sputtered samples as a function of current density and temperature are shown in Figure 42b. The co-sputtered samples have a higher resistivity than the Ni/Al multilayers annealed at 400°C, and this leads to a lower current density for these samples. It can also be seen that as with the NiAl-0.5Ag samples, the co-sputtered samples show an increasing current density as the line width decreases and no change in the temperature response.

149

1.25 7.5 a) b) 560

520

)

1.20 2 7.2

(°C) 480 max

T 440

1.15 6.9 400 0.1 1 10 100 1000

1

R/R 1.10 6.6

1.05 (MA/cm Density Current 6.3 70.6 mA - 3.3 m 3.3 m wide 100.3 mA - 5.2 m 5.2 m wide 124.0 mA - 6.4 m 6.4 m wide 1.00 6.0 0 5 10 15 20 25 30 35 0.1 1 10 100 1000 Time (h) Time to Failure (h) Figure 42 – (a) Relative resistance as a function of time curves for co-sputtered NiAl of varying line width, and (b) the lifetime data plotted against the applied current density with the inset showing the lifetime as a function of temperature (x-axis is the same in the main graph and the inset).

SEM images of two of the co-sputtered samples are shown in Figure 43. Some similar features to the multilayer samples are seen: damage localized to the region immediately adjacent to failure, voids/surface depressions, and surface ridges extending along the middle of the line. One feature that was not observed in the multilayer samples is seen along the line edges in Figure 43b and consists of what looks like small surface extrusions.

a) b) 103.2 mA 6.67 MA/cm2 192,240 cycles

70.6 mA 7.06 MA/cm2 2 μm 1 μm 17,877,600 cycles

Figure 43 – SEM images of samples of co-sputtered NiAl with line widths of (a) 3.3 μm and (b) 5.2 μm that were cycled until failure.

6. Summary of the NiAl data – This section summarizes all of the NiAl data discussed in the preceding sections. The sample sets depicted in the subsequent plots are numbered from 1-12, with 1 having the lowest resistivity and 12 the highest resistivity. Sample sets 1-6 have a “low” resistivity (20-24 μΩcm) and samples 7-12 have a “high” resistivity (38-49 μΩcm).

150

The results for the lifetime as a function of the peak temperature are compiled in Figure 44a. No difference is seen between the 12 data sets in terms of the time to failure as a function of temperature. By changing the deposition conditions (co-sputtering vs. multilayers), annealing profiles (high temp vs. low temp), line width, and resistivity, no increase or decrease in the operating temperature is obtained. Each of the 12 data sets lies within the scatter of the other 11 data sets. When comparing the NiAl temperature data against the other materials tested (Ru, Au, RuAl), the 12 data sets will be treated as one. This has been done with the solid line plotted in Figure 44a, which was obtained from a curve fit to the entire data set.

700 600 b) 1 2 3 a) 10 4 5 6

7 8 9 )

600 2 10 11 12 500 9

500

400 8

T (°C) T

400 

300 7 1 2 3 Initial Temperature (°C) Temperature Initial 300 4 5 6 (MA/cm Density Current 7 8 9 6 10 11 12 200 200 1E-3 0.01 0.1 1 10 100 1000 1E-3 0.01 0.1 1 10 100 1000 Time to Failure (h) Time to Failure (h) Figure 44 – Summary of the NiAl TMF data. The numbers in the plots refer to individual sample sets and are ranked according to the average resistivity of the sample set. Samples 1-6 have an average resistivity in the range of 20-24 μΩcm and samples 7-12 have an average resistivity of 38-49 μΩcm. The open symbols (7-9) are for the co- sputtered samples, and the closed symbols are for the annealed multilayer samples. Arrows indicate sample that did not fail. (a) Time to failure plotted against the initial peak temperature (left) and ΔT (right). The solid line is a linear fit to the entire data set plotted as log(t) vs. T. (b) Time to failure plotted as a function of the applied current density.

In contrast to the temperature data, a difference is noted in the current density data for the 12 data sets. This is shown in Figure 44b, where each of the individual data sets does not lie within the scatter of the other data sets. The plot can be divided into two regions: higher current density for the “low” resistivity samples and lower current density for the “high” resistivity samples. Sample set 2 lies between the two regions and this will be discussed below.

The data in Figure 44b were re-plotted as the ln(tf) versus current density and linear fits to the data with 3 or more data points were obtained. From the curve fits, the current density required for failure in 10 h was estimated and this data is plotted in Figure 45 as a function of the average resistivity of the sample set. The data show that a major factor in determining the allowable current density for failure in a given amount of time is the resistivity of the metal line, and that the current density may be increased further if the resistivity of the NiAl lines can be reduced. Though not specifically plotted in Figure 45, the line width also affects the current density, as was discussed previously in relation to the co-sputtered and NiAl-0.5Ag data. The data in the plot can be grouped into “narrow” and “wide” lines. The “narrow” lines are samples 3, 5, and 9 and have line widths ranging from 3.3-3.8 μm. Samples 2, 7, 8, 10, 11, and 12 have line widths ranging from 4.4-6.4 μm, and can carry a lower current density than the “narrow” line samples.

151

This is the reason that sample 2 in Figure 44b does not line up with the rest of the “low” resistivity samples. The larger line width requires a smaller current density (although the current is higher) to achieve an equivalent temperature and thus the maximum current density is lower. Similar effects are expected to be seen with variations in film thickness; however, the samples studied have a much lower variance in film thickness than in line width so the effects are expected to be small in comparison. When comparing the current density of NiAl with the Au, Ru, and RuAl data, the results from samples 1, 3, 4, 5, and 6 will be used. These are the samples that have a “low” resistivity (20-24 μΩcm) and a “narrow” line width (2.9-3.8 μm) and represent the best results in terms of current density for the NiAl samples.

10

) 9 2 3 5

8 2

7 9 8 7 11 10

6 12 Current Density (MA/cm Density Current

5 10 20 30 40 50 60 Resistivity ( cm) Figure 45 – The current density resulting in failure in 10 h as a function of the average resistivity. The numbers in the plot refer to the same numbers as are seen in Figure 44.

C. RuAl

1. Annealed 25 nm wavelength Ru/Al multilayers – Ru/Al multilayers were deposited and exposed to a variety of annealing profiles as was done for Ni/Al multilayers. Annealing temperatures were higher than for Ni/Al due to the higher melting temperature of RuAl. Details of the deposited and tested samples are shown in Table 7. The resistivity of the RuAl samples are higher than the previously discussed NiAl samples, and the annealing is a balance between using high enough temperatures (and times) to allow for grain growth, but not so high that significant reactions with the substrate occur. Anneal 1 gives the best results in terms of resistivity, with anneals 2 and 4 having a higher resistivity likely due to the higher annealing temperature leading to a more significant reaction with the substrate (TEM has shown reaction with the substrate at 1000°C: see Chapter 3 for details). The samples from anneal 3 also have a higher resistivity than those from anneal 1, which could be due to incomplete annealing occurring at the lower annealing temperature used in anneal 3. As will be shown below, the

152 samples from anneal 3 show a region of decreasing resistance indicating that further annealing is occurring in these samples during testing, whereas this does not occur for the samples annealed at higher temperatures (anneals 1, 2, and 4).

Table 7 – Sample details and results for annealed Ru/Al 25 nm wavelength multilayers tested using AC TMF. NF refers to a sample that did not fail in the time indicated. AC current Thickness Line width Resistivity AC current Time to density T (°C) (nm) (μm) (μΩcm) (mA) max failure (h) (MA/cm2) anneal 1: 1000°C(30 s) – 10°C/s – 600°C(30s) 249 4.7 27.5 91.9 7.87 572 0.12 249 4.7 26.3 89.6 7.67 507 0.87 249 4.7 26.2 87.2 7.47 472 4.40 249 4.7 25.9 84.8 7.26 391 21.80 anneal 2: 1100°C(30 s) – 5°C/s – 900°C(15 s) – 2°C/s – 800°C(15 s) – 2°C/s – 600°C(30s) 246 5.0 30.0 94.3 7.67 642 0.07 246 5.0 31.9 89.0 7.23 558 0.28 246 5.0 30.8 92.0 7.48 593 0.47 anneal 3: 900°C(30 s) – 10°C/s – 600°C(30 s) 245 4.9 31.0 95.0 7.91 658 0.03 245 4.9 31.8 95.0 7.91 679 2.28 245 4.9 31.4 93.9 7.82 659 9.80 245 4.9 31.6 92.2 7.68 616 154.15 245 4.9 30.0 86.3 7.19 466 838.40 (NF) anneal 4: 600°C(60 s) – 900°C(30 s) – 1100°C(30 s) – 10°C/s – 600°C(30s) 247 4.9 36.0 83.6 6.99 582 0.17 247 4.9 33.6 77.7 6.49 403 33.60

The results from TMF tests of anneals 1-4 are shown in Figure 46 as the time to failure as a function of current density and temperature. The applied current density for failure in a given amount of time for anneals 1, 2, and 4 go in reverse order of the resistivity: anneal 1 has the lowest resistivity and the highest current density, and anneal 4 has the highest resistivity and the lowest current density. This trend is the same as was seen for the NiAl samples. The current density for anneal 3 does not line up with the trend seen in the rest of the data. The resistivity of anneal 3 is similar to anneal 2, so it was expected that the two would show similar performance in terms of current density, but anneal 3 resulted in the highest current density for the annealed Ru/Al multilayers. A possible explanation is that the lower annealing temperature for anneal 3 resulted in less reaction with the substrate, and it is possible that the incorporation of Si into

153

RuAl is detrimental to the fatigue performance of the films. Auger and TEM analysis of annealed Ru/Al multilayers indicated no noticeable reaction with the substrate at 800°C and a reaction with the substrate at 1000°C (see Chapter 3). Because of the higher current density for anneal 3, this sample exhibits higher testing temperatures that the other three sample sets, which all show very similar times to failure as a function of temperature. Therefore, anneal 3 is able to withstand higher cyclic stresses and likely has higher strength than the other three samples. The most likely cause is silicon incorporation, which has been shown via TEM analysis to occur at 1000°C and details are given in Chapter 3 (analysis has not been done for lower annealing temperatures). It is also notable that the time to failure increases more rapidly with decreasing temperature for anneal 3 than for anneals 1, 2, and 4.

8.4 800 a) b) 8.0

) 700 2

7.6 600

7.2 (°C)

max T

500 6.8

anneal 1 400 anneal 1 Current Density (MA/cm Density Current 6.4 anneal 2 anneal 2 anneal 3 anneal 3 anneal 4 anneal 4 6.0 300 0.01 0.1 1 10 100 1000 0.01 0.1 1 10 100 1000 Time to Failure (h) Time to Failure (h) Figure 46 – Time to failure for the annealed Ru/Al multilayers as a function of (a) current density and (b) temperature. The solid line is a fit to anneals 1, 2, and 4, and the dashed line is a fit to the data from anneal 3.

Curves showing the resistance as a function of time for the samples from anneal 1 are shown in Figure 47a. Curves are not shown for anneals 2 and 4 as the results are similar. The resistance for samples from anneals 1, 2, and 4 increases with increasing time until failure. One feature from anneal 1 that is not seen from anneals 2 and 4 can be seen in the 89.6 mA and 87.2 mA curves. These curves do not show the typical behavior of 3 regions: the first of increasing resistance with a decreasing slope, the second of increasing resistance with a nearly constant slope, and a third region of increasing resistance with an increasing slope. Two of the curves shown in Figure 47 appear to have 4 and 5 regions for the 87.2 mA and 89.6 mA curves, respectively. It is possible that the changes in slope occur due to the initiation of other failure sites although no difference was seen in the SEM images of samples showing 3 regions and the two samples showing more than 3 regions.

The data from Figure 47a are used to plot the time to reach a certain increase in resistance as a function of the test temperature, and this is shown in Figure 47b. The trends in the data are very similar to those discussed on NiAl.

154

2 210 10 210 a) b) 15% increase in R 195 10% increase in R 195 6.8% increase in R 180 1

10

) 180 165 

150 165 100

0 1 2 3 4

Time (h) Time

Resistance ( Resistance 150

10-1 84.8 mA 135 87.2 mA 89.6 mA 91.9 mA 120 10-2 0 5 10 15 20 350 400 450 500 550 600 Time (h) Initial Temperature (°C) Figure 47 – Results from anneal 1. (a) Resistance as a function of time (the inset shows a zoomed in view of the data at short times) and (b) the time to a given increase in resistance as a function of the peak temperature at the beginning of the test.

The resistance as a function of time curves obtained from the samples from anneal 3 are significantly different to those from anneals 1, 2, and 4. Rather than showing an increasing resistance until failure, there is a region of decreasing resistance between two regions of increasing resistance. The region of decreasing resistance shows that annealing is occurring in these samples even though the test temperature is significantly lower than the annealing temperature. This annealing behavior may also be enhancing the current density for these samples. As the resistance of the samples decrease, so does the Joule heating, allowing for higher current densities.

195

190

) 

185

Resistance ( Resistance 180

92.2 mA 175 93.9 mA 95.0 mA

0.01 0.1 1 10 100 Time (h) Figure 48 – Resistance as a function of time curves for samples from anneal 3.

155

SEM images of tested samples from anneals 1-4 are shown in Figure 49a-d, respectively. There are some features that are noted in samples from all of the annealing conditions, although they are not necessarily present in every sample. These are small voids along the line edges and a speckled surface appearance. The speckled surface may be due to the effects of oxidation as the surface does look similar to the surface of blanket films oxidized at temperatures of 550°C and above. Another feature that was exhibited by samples from anneals 1 and 2 was noted along the center of the lines and is shown in the images of Figure 49a. It consists of raised features, which in some cases appear to be composed of very fine surface wrinkles. These features are often surrounded by what resemble cracks; however, they do not reach from the surface of the sample to the substrate. Only 1 sample showed additional cracking and this was a sample from anneal 3 (Figure 49c). An enlarged region containing these cracks shows what could be the initiation of melting at a secondary location (bottom of the image).

a) 84.8 mA b) 92.0 mA 7.26 MA/cm2 7.48 MA/cm2 15,696,000 cycles 338,400 cycles

1 μm 1 μm 1 μm

c) 1 μm d)

95.0 mA 77.7 mA 7.91 MA/cm2 6.49 MA/cm2 2 μm 1 μm 24,192,000 cycles 1,641,600 cycles Figure 49 – SEM images of the annealed Ru/Al multilayers from (a) anneal 1 (both sides of the failure), (b) anneal 2, (c) anneal 3 (2 different magnifications), and (d) anneal 4.

2. Co-sputtered RuAl – Samples of RuAl were co-sputtered at 1.5 mTorr Ar pressure and were tested using AC TMF. Samples were annealed at 900°C or 1000°C for 1 min in Ar, and at 1000°C for 30 s or 1 min in N2. The majority of the tests were conducted on samples annealed at 1000°C for 1 min in Ar and these will be the only tests discussed. No significant differences in current density or peak temperature were noted for the Ar annealed samples due to changing the annealing temperature from 1000°C to 900°C, although the samples annealed at 1000°C had a resistivity ~4.5 μΩcm lower than the samples annealed at 900°C. For the samples annealed in

156

N2, no difference was noted in the current density, peak temperature, or resistivity when the annealing time at 1000°C was increased from 30 s to 1 min, and no appreciable difference was observed between the samples annealed in N2 or Ar.

The co-sputtered RuAl samples had a line width of 4.1 μm and a thickness in the range of 249- 267 nm. The average resistivity was 31.4 μΩcm for the samples annealed at 1000°C, which is similar to the values from anneals 2 and 3 from the annealed Ru/Al multilayers. To analyze the repeatability of TMF tests, tests were run multiple times at the same current on samples annealed for 1 min at 1000°C: 12 tests at 82.0-82.1 mA, 10 at 84.7-84.9 mA, 8 at 87.6-87.8 mA, and 8 at 89.4-89.5 mA. Results from these tests are shown in Figure 50, which shows a fairly large range in failure time for a constant stressing current that increases as the current decreases. Thus, to get an accurate picture of the lifetime multiple tests should be run at each of several values of current.

Cycles to Failure (1000's) Cycles to Failure (1000's) 100 1000 10000 100 1000 10000 100 25% failure probability a) 90 b) 50% failure probability 80 75% failure probability

88

60

86

40 Current (mA) Current 84

Failure Probability (%) Probability Failure 20 82.1 mA 84.8 mA 87.7 mA 82 89.4 mA 0 0.1 1 10 0.1 1 10 Time to Failure (h) Time to Failure (h) Figure 50 – Results from the repeatability tests on RuAl co-sputtered at 1.5 mTorr and annealed at 1000°C for 1 min. (a) Plot of the failure probability for the four different currents tested and (b) the time to failure as a function of the testing current for failure probabilities of 25%, 50%, and 75%.

SEM images from the tests shown in Figure 50 are shown in Figure 51a-e for tests run at 82.1 mA, 84.8 mA, 84.8 mA, 87.7 mA, and 89.5 mA, respectively. For the samples tested at the lowest current (82.0-82.1 mA) only minor damage was observed near the failure location and only for tests with the higher cycles to failure. As the current increased so did the severity of damage, the size of the damaged patches, and the likelihood of finding damage in regions other than the failure location. An example showing a damaged patch not associated with the electrical open is shown in Figure 51d. In all samples, the damaged patches tended to have a semi-elliptical shape and were located at the edges of the lines. In all samples where damage was observed, failure occurred through the center of one of these damaged patches. Around the periphery of the damaged region was often a depression resembling a crack, but that did not extend all the way through the thickness of the sample. In some cases the damaged patches appear to be somewhat porous (Figure 51c) and in others the morphology appears coarser than

157 the surrounding area (see Figure 51e). Higher magnification images of two of the damaged patches are shown in Figure 52.

a) 82.1 mA b) c) 1 μm 12,528,000 cycles

84.8 mA 2 μm 554,400 cycles 2 μm

87.7 mA d) 89.5 mA 93,600 cycles e) 86,400 cycles

1 μm 1 μm

Figure 51 – Select images from the failure region (a, b, c, and e) and from a region far from the failure (d) for co- sputtered RuAl annealed at 1000°C for 1 min and then thermally cycled until failure. The inset in (a) shows an enlarged view of the region close to the failure. The boxed region in (d) highlights the damaged patch.

The damage noted in the co-sputtered RuAl is different from that noticed in the other samples tested, although there are some similarities to the annealed Ru/Al multilayers. Both the annealed Ru/Al multilayers and the co-sputtered RuAl showed features that resembled cracks, but were not through-thickness; however, these were more numerous in the case of the co-sputtered RuAl. Also, both types of RuAl showed a “speckled” surface appearance, which could be due to the initial stages of oxidation (where changes in surface morphology are noted for blanket films, but without changes in resistance – see Chapter 4 on oxidation). The damage observed in the co- sputtered RuAl is along the line edges, whereas in most of the other samples tested the damage was primarily noted along the center of the lines, with the exception being Ru, which had damage extending along the entire width of the line. The co-sputtered RuAl was also the only sample that showed damage formation in more than 1 location (i.e., locations other than the 1 causing failure).

158

a) 0.5 μm

b)

1 μm

Figure 52 – Close-up views of the damaged patches observed in RuAl co-sputtered samples tested at (a) 89.4 mA (0.2 h to failure) and (b) 89.5 mA (0.117 h to failure).

All of the co-sputtered RuAl films tested showed similar curves of resistance as a function of time. Examples are shown in Figure 53 for samples annealed for 1 min at 1000°C that were tested at 82.0-82.1 mA and failed between the voltage pads. The only difference between the curves shown and for those where the samples failed outside the voltage pads is the absence of the third region. For all samples there was an initial region of increasing resistivity with a decreasing slope and a second region of increasing resistivity with a nearly constant slope. For samples that failed between the voltage pads there was a third region of increasing resistivity with an increasing slope. This is the same behavior exhibited by nearly all of the materials tested, and will be examined in more detail below.

159

1.08 1.20 a) b)

1.06 1.15

1 1

1.04 1.10

R/R R/R

1.02 1.05

1.00 1.00 0 1 2 3 4 0 1 2 3 4 Time (h) Time (h) Figure 53 – Relative resistance as a function of testing time for the 6 RuAl co-sputtered samples tested at 82.0-82.1 mA that failed between the voltage pads. For clarity, the curves in (a) are re-plotted in (b) with an offset of 0.016 in the y-direction between each curve.

A summary of all of the results obtained from the co-sputtered RuAl are shown in Figure 54. Figure 54a shows the time to failure as a function of the applied current density and the time to failure as a function of the peak temperature is plotted in Figure 54b.

8.8 700 a) tests to failure b) tests to failure

no failure no failure

) 2 8.4 600

8.0 (°C) 500

max T

7.6 400 Current Density (MA/cm Density Current

7.2 300 0.01 0.1 1 10 100 1000 0.01 0.1 1 10 100 1000 Time to Failure (h) Time to Failure (h) Figure 54 – Results from all of the co-sputtered RuAl samples. (a) Time to failure as a function of the current density and (b) the time to failure as a function of the peak temperature.

3. Summary of the RuAl TMF data – A summary of the RuAl thermal fatigue results is given in the plots shown in Figure 55a and Figure 55b for the time to failure as a function of temperature and current density, respectively. The multilayers from anneal 3 (lower temperature 900°C anneal) result in the best performance in terms of temperature; they can withstand higher temperatures than the other samples for equal times to failure. The RuAl from the multilayers of anneal 3 also show a more rapidly increasing time to failure as the test temperature is lowered. The co-sputtered RuAl and the other multilayer samples result in similar times to failure for

160 equivalent temperatures. The co-sputtered RuAl may result in a longer lifetime at lower temperatures; however, this could be an artifact due to the small number of multilayer samples from anneals 1, 2, and 4.

The highest current density is obtained from the co-sputtered samples and the multilayers annealed using anneal 3, which result in similar times to failure for equal current densities. However, as the multilayer samples have a larger line width it is likely that the samples from anneal 3 would result in a higher current density than the co-sputtered RuAl if the line widths were the same.

Co-sputter a) 8.5 b) 700 Multilayer: anneal 1

600 Multilayer: anneal 2 ) 2 Multilayer: anneal 3 Multilayer: anneal 4 8.0 600 500

(°C) 7.5

T (°C) T

max

 T 500 7.0 400

Co-sputter RuAl (MA/cm Density Current 400 Multilayer: anneals 1, 2, 4 6.5 Multilayer: anneal 3

0.01 0.1 1 10 100 1000 0.01 0.1 1 10 100 1000 Time to Failure (h) Time to Failure (h) Figure 55 – Summary of the results from TMF tests on RuAl. (a) Time to failure as a function of the peak temperature (left) and the ΔT (right) experienced during testing and (b) time to failure as a function of the AC current density. The lines in (a) are linear fits to the data obtained from plots of log(t) vs. T.

D. Gold

Several tests were run on sputtered gold samples for comparison to the intermetallics and ruthenium data. Details of the gold samples and the test results are given in Table 8. All of the samples were annealed at 400°C for 4 h.

Table 8 – Details of the gold samples, testing conditions, and results of thermal fatigue testing.

Deposition AC current Thickness Line width Resistivity AC current Time to Material pressure density T (°C) Comments (nm) (μm) (μΩcm) (mA) max failure (h) (mTorr) (MA/cm2) Gold 7 254 3.4 4.34 134.3 15.55 298 0.4 7 254 3.4 4.22 131.1 15.18 311 13.6 7 254 3.4 4.20 135.9 15.74 309 17.5 Gold 4 340 3.5 4.65 141.5 11.89 × ~0.000278 immediate failure 4 340 3.5 4.58 137.2 11.53 249 51.4 4 340 3.5 4.67 134.2 11.28 245 59.1

161

The resistance during AC testing as a function of test time is shown in Figure 56 for the gold samples tested. All 5 tests show a sharp increase in resistance in the early stages of testing, which is believed to be mainly due to sample heating. After this initial sharp increase in resistance 4 of the 5 samples showed a decrease in resistance, which is most likely due to the effects of annealing. Even though all of the samples failed outside of the voltage pads, some of the tests showed an increase in resistance prior to failure indicating that some damage was forming between the voltage pads as well as in the region that ultimately caused the open circuit failure.

20.0 29.0 134.2 mA 137.2 mA 28.5

19.5 )

) 28.0

 

19.0 27.5

Resistance ( Resistance 27.0 Resistance ( Resistance 18.5 26.5 131.1 mA 134.3 mA a) 135.9 mA b) 18.0 26.0 0 10 20 30 40 50 60 0 3 6 9 12 15 18 Time (h) Time (h) Figure 56 – Resistance as a function of time curves for gold sputtered at (a) 4 mTorr and (b) 7 mTorr Ar pressure. The sharp changes in resistance as indicated by the regions between the arrows shown in (b) are due to a slightly fluctuating current.

Although the sample set is small, referring to the data listed in Table 1 shows that the sample deposited at 7 mTorr was able to withstand a higher current density than the sample deposited at 4 mTorr, although there was not much difference in the performance of the two gold samples as a function of the testing temperature. The higher current density is most likely due to the lower resistivity of the sample deposited at 7 mTorr, which would enable higher current densities for equal temperatures.

Figure 57 shows SEM images of the failure location for two of the gold samples, one at each deposition pressure. The two different samples show distinctly different damage morphologies from one another and to the results from Keller et al. (see for example ref. [12]). Unlike the results from Keller et al. on Al and Cu films, the gold films tested in this study showed no surface wrinkling. Instead, the lines showed regions of surface depression (4 mTorr samples, Figure 57a), or what looks like grain growth near the line edges combined with changes in the line width (7 mTorr samples, Figure 57b). The 4 mTorr samples showed damage that was highly localized (within a few μm of the failure site), whereas the 7 mTorr samples showed signs of damage in regions not immediately adjacent to the failure location. This observation explains why a several hour period of increasing resistance was noted in the 7 mTorr samples even though failure occurred outside of the voltage pads. Differences between these results and those of Keller et al. are likely due to differences in the grain size. While the grain size was not measured

162 in the samples tested, they are expected to have a grain size significantly less than 1 μm as AFM indicated that blanket gold films deposited at 1.5 mTorr and 7 mTorr had a grain size on the order of 50 nm (see Figure 58). The grain size in the samples tested by Keller et al. were on the order of 1 μm (1.4 μm for Al-1Si [11] and 0.3-1.5 μm for Cu [3]).

Figure 57 – ESEM images of gold TMF samples cycled until failure. (a) Sample deposited at 4 mTorr and tested at 137.2 mA (11.53 MA/cm2). Failure occurred after ~37,008,000 thermal cycles. (b) Sample deposited at 7 mTorr Ar pressure and tested at 135.9 mA (15.74 MA/cm2). Failure occurred after ~12,600,000 thermal cycles. (c) The same sample shown in (b) showing a larger region of the sample.

100 nm

Figure 58 – AFM image of a blanket gold film sputtered at 1.5 mTorr Ar pressure.

163

E. Comparison of the TMF Results for the Different Materials Tested

A summary of the results from all of the TMF tests are shown in Figure 59 for the data as a function of temperature and in Figure 60 as a function of current density. Both the time to failure (bottom axis) and the cycles to failure (top axis) are plotted. For high temperatures (T > ~500°C) the RuAl samples have a longer time to failure than all of the other samples followed by NiAl and Ru-1.5 mTorr, which have similar times to failure. The gold samples and the ruthenium sputtered at 7 mTorr did not attain such high temperatures during testing. As the testing temperature decreases the time to failure for RuAl (with the exception of Ru/Al anneal 3) approaches that of NiAl, and as the NiAl samples show a faster increase in time to failure with

decreasing temperature,Cycles toit isFailure likely that NiAl will attain longer times to failure as the temperature

is further4 decreased.5 The6 annealed7 Ru/Al8 multilayer from anneal 3 shows much higher times to 10 10 10 10 10 700 failure at high temperature than the other RuAl600 samples, which could be due to the lower annealing temperature (900°C compared to 1000-1100°C) and a decreased reaction with the substrate. No similar effect of processing was seen for NiAl, which was annealed at 600 500 temperatures ranging from 400-1000°C with no appreciable change noted in the time to failure as a function of the test temperature. The enhanced lifetime for all temperatures tested of RuAl and

NiAl500 compared to gold indicates the higher strength400 of the intermetallics. (°C)

(°C) T

max 

T 400 Cycles to Failure 300 104 105 106 107 108 700 600 RuAl RuAl* 200 NiAl Au 300 Ru-1.5 mTorr Ru-7.0 mTorr 600 500 100 200 0.01 0.1 1 10 100

500 Time to Failure (h) 400

(°C)

T (°C) T

max 

T 400 300

RuAl RuAl* 200 NiAl Au 300 Ru-1.5 mTorr Ru-7.0Au mTorr

100 200 0.01 0.1 1 10 100 Time to Failure (h) Figure 59 – Summary of the results from TMF testing in terms of the peak temperature and ΔT experienced during thermal cycling. (a) Plot of Tmax (and ΔT) versus the time to failure and (b) a schematic of the same data showing approximate boundaries to the various data sets. RuAl* refers to the annealed multilayer from anneal 3.

The results for current density are opposite to those for the testing temperature; gold results in significantly higher current densities than the other materials. This is due in large part to the much lower resistivity and hence much lower operating temperatures than the other samples. For NiAl samples it was shown that the current density depended on the resistivity, and increases in the current density are expected for NiAl if the resistivity can be further reduced. This may be possible for RuAl as well if reactions with the substrate can be minimized.

164

16 RuAl NiAl

Au-7.0 mTorr )

2 14 Au-4.0 mTorr Thickness Line Width Resistivity Material Ru-1.5 mTorr (nm) (μm) (μΩcm) Ru-7.0 mTorr 12 Au: 7 mTorr 254 3.4 4.2-4.3 Au: 4 mTorr 340 3.5 4.6-4.7 Ru: 1.5 mTorr 165 3.2 18.1-20.1 10 NiAl 252-293 2.9-4.6 19.7-26.8 RuAl 249-267 4.1 29.5-37.6

Current Density (MA/cm Density Current 8

Ru: 7 mTorr 287 2.8 48.8-53.9

6 1E-3 0.01 0.1 1 10 100 1000 Time to Failure (h) Figure 60 – Summary of the current density results from TMF tests of the different materials tested. The table shows details of the samples represented in the plot. The data points circled in the plot are for tests that were stopped prior to failure (indicated by the arrow).

F. Analysis of the Resistance as a Function of Time Curves

As mentioned in the “Materials and Methods” section of this chapter, an equation was developed that can be used to describe all of the trends seen in the R(t) curves. The equation consists of exponential, power, and linear terms as is shown in Eq. 7, where C1 through C9 are fit parameters with no physical meaning. The constants C7 and C8 were equal to zero for most samples.

Eq. 7

The three types of curves obtained during testing are shown in Figure 61. Two are partitioned into 3 regions (I, II, and III) and the other into 4 regions (Ia, Ib, II, and III). Regions I and Ia are described by the first exponential term, region Ib by the second exponential term, region II by the linear term, and region III is described by the power term. Due to the geometry of the samples and the location of the voltage pads, not all of the samples were curve fit using Eq. 7. Samples that failed outside the region encompassed by the voltage pads were not curve fit as the increase in resistance near the onset of failure (region III) was not captured. The first two regions could have been curve fit using the above equation, but with the constants C4 and C5 set to zero. Of the 55 data sets analyzed more than half had a correlation coefficient (r2) value greater than 0.98, and all had an r2 value greater than 0.93 with the exception of 1 sample with a value of 0.84. Values for the constants from the curve fits are given at the end of the chapter in Appendix II.

165

R  t   t   C exp   C expC    C t  C  C t C5 R 1  C  7  8 C  3 6 4 Ib 1  2   9 

II III

I/Ia Ib II III Ia 1.20 data curve fit III

I II 1.15

(au)

1

R/R 1

1.10 R/R

I II III 1.05

1.00 0.0 0.2 0.4 0.6 0.8 1.0 0 1 2 3 4 5 6 Time (t/t ) Time (h) f Figure 61 – Schematics showing the 3 different types of curves obtained during TMF testing, the equation used to fit the data, and an example of a curve fit to an experimental data set.

After curve fitting the data sets using Eq. 7, the length of time in each region was recorded. This was done by first finding the center of curvature of the fit to the dataset: for curves with a region of decreasing resistance this was the point with the lowest curvature and for curves with only increasing resistance this was the point of zero curvature. Once this point was located, which is in the center of region II, increasing sections of the curve were fit with a linear fit (starting from the center of curvature and moving in a positive and negative direction simultaneously). The linear fits were increased by 19-21 data points for each subsequent fit, with the amount in a positive and negative direction depending on the location of the center of curvature in the dataset. It was determined that incrementing by unequal amounts in the positive and negative x- directions rather than equal amounts (10 points in each direction) resulted in a better determination of the 3 regions. The last point at which the r2 value was above 0.999 determined the start and end points of region II and thus the end point of region I and the start of region III, respectively. The procedure used for determining the locations of regions I, II, and III is shown schematically in Figure 62.

166

1.3 1.20

1.2 1.15

1

1

R/R R/R 1.1 1.10

1.0 1.05 0.24

0.04

0.16 Slope Slope 0.08 0.02

0.00

0.0 0.00 -0.2 -0.04

-0.4

Curvature Curvature -0.6 a) -0.08 b) -0.8 0 2 4 6 8 10 2 4 6 Time (h) Time (h) Figure 62 – An example of the procedure used to determine the transitions between regions I, II, and III. (a) The entire data set and (b) the region encompassed by the lines in (a). The thick red line shows the location of the center of curvature for this dataset. The dashed black lines show the last point where the r2 value is greater than 0.999 when the size of the curve fit is increased by 10 data points in each direction. The solid black lines show the last point where r2 is greater than 0.999 when the curve fits are increased by 5 data points in the negative x-direction and 14 data points in the positive x-direction.

The three different trends that were obtained are shown in Figure 63 along with the calculated curve fits. In all cases the curve fits describe the data very well, and transitions between the three regions can be easily determined. Transitions between the three regions are shown in the plots of Figure 63 as the vertical dashed lines.

167

1.00 1.08 data data 0.98 curve fit curve fit

1.06 0.96 I II III

0.94

1 1

1.04 R/R 0.92 R/R

0.90 1.02 I II III

0.88

0.86 1.00 0 2 4 6 8 10 12 14 0.0 0.1 0.2 0.3 Time (h) Time (h) 1.04

data 1.03 curve fit

1.02 R  t   t  Ib II III C5 1.01  C1 exp   C7 expC8    C3t  C6  C4t 1     R1  C2   C9 

R/R 1.00

0.99 Ia I/Ia Ib II III

0.98

0.97 0 2 4 6 8 10 Time (h) Figure 63 – Example data sets and curve fits showing the three different trends noted in the R(t) curves. Also shown is the equation used to obtain the curve fits as well as lines showing the obtained transition points between the three regions.

The first region noted in the R(t) curves, region I, is characterized by a region of increasing (or decreasing) resistance with a decreasing slope. For most samples this region shows an increasing resistance, which is primarily due to sample heating. In some cases the resistance decreases initially with increasing time and this is due to annealing, which results in a decreasing testing temperature. A third trend in region I is a combination of the two aforementioned cases: initial increasing resistance from heating, followed by a decreasing resistance due to annealing. After a certain amount of time, all three types of R(t) curves show a transition to region II, which is characterized by an increasing resistance with a slope that is nearly constant. In region II, damage is beginning to form that leads to an increase in the sample resistance. The damage that forms in this region is likely due to point defects and dislocations. Eventually the sample transitions to region III, which is characterized by an increasing resistance with an increasing slope. The damage forming in this region is believed to be more severe, i.e., voids and cracks. The trends seen in regions II and III are similar to what has been observed in the literature for fatigue studies conducted on bulk metals at room temperature: i.e., a region of steadily increasing resistance that was due to the formation of dislocations and point defects, followed by another

168 region where the resistance began to increase more rapidly due to the onset of cracking [25-27]. More specifics on the trends seen in each of the three regions are contained in the discussion below.

All of the samples that showed an R(t) curve exhibiting region III behavior (all samples that failed between the voltage pads) were fit using the equation shown in Figure 63, and the time spent in each region was recorded as well as the slope at the beginning of region I, the slope in region II, and the slope at the end of region III. The slope at the beginning of region I was calculated as the instantaneous slope at 1 s, the slope in region II as the slope of the linear fit to this region, and the slope at the end of region III as the instantaneous slope at tf (the failure time). The results from this analysis are presented below for samples that showed 3 regions of increasing resistance. There were very few samples showing the other 2 behaviors, making trends and conclusions difficult to ascertain.

An analysis of the results from region I are shown in Figure 64 for the time and slope in region I as a function of the maximum temperature (a and c) and the time to failure (b and d). All of the NiAl samples that showed 3 regions of increasing resistance and the co-sputtered RuAl annealed for 1 min in Ar at 1000°C are included. The same is true for all of the following graphs. There is more scatter to the data plotted as a function of temperature (this is also true for the graphs for regions II and III) than for the data plotted as a function of the time to failure. This is due to the scatter in the fatigue data; equal testing temperatures may result in a range of failure times, but samples with equal times to failure all have very similar R(t) curves. As the testing temperature increases, the time in region I decreases and the slope increases. With higher testing temperatures, the temperature increases more rapidly and damage begins to form more rapidly, decreasing the time for the transition to region II. At the higher testing temperatures there is a larger ΔT, and due to the difference in the CTE between the film and substrate there is a larger change in stress during the test. The larger stress results in more rapid damage formation. Trends as a function of the current density are the same as for the test temperature since increasing the current density increases the temperature. Again, there is some scatter in the data, which is due in part to the different sample sets used which had different values of resistivity, line width, and film thickness, all of which will affect the temperature resulting from a certain current density. The trends as a function of the time to failure are the opposite to those of temperature and current density, i.e., as the time to failure increases, the time in region I increases and the slope at the beginning of region I decreases.

Comparisons between the results from NiAl and RuAl in terms of temperature are difficult as the RuAl samples were run at higher temperatures than the NiAl samples and the overlap in the datasets is small. The trends in terms of time to failure are easier to ascertain as there is more overlap in the data, at least for shorter times to failure (tf < 10 h). For tests with a shorter time to failure, the time and slope in region I are similar for the NiAl and RuAl samples.

169

100 100 0.8 a) b)

0.4 10 10

0.0 7.5 8.0 8.5 2

1 Current Density (MA/cm ) 1

(h)

(h)

1

1

t t

0.1 0.1

NiAl NiAl RuAl RuAl 0.01 0.01 350 400 450 500 550 600 650 700 0.01 0.1 1 10 100 1000 T (°C) max Time to Failure (h) 10 10 c) NiAl d) RuAl

1 1

0.1 1.5 0.1

1.0

0.01 0.5 0.01

Slope in Region 1 (1/dt) 1 Region in Slope Slope in Region 1 (1/dt) 1 Region in Slope

0.0 NiAl 7.5 8.0 8.5 2 RuAl Current Density (MA/cm ) 1E-3 1E-3 350 400 450 500 550 600 650 700 0.01 0.1 1 10 100 1000 T (°C) Time to Failure (h) max Figure 64 – Results from the analysis of curve fits to the NiAl and RuAl data showing three regions of increasing resistivity. (a) The time in region I as a function of the maximum test temperature, (b) the time in region I as a function of the time to failure, (c) the slope in region I as a function of the maximum test temperature, and (d) the slope in region I as a function of the time to failure. The insets in (a) and (c) show the time in region I and the slope in region I as a function of the current density, respectively (the y-axes have the same units in the insets and the main graphs).

A similar analysis to the above was completed on the data from region II, i.e. the region with an almost constant slope. The results are displayed in Figure 65 for the time and slope in region II as a function of the temperature (a and c) and as a function of the time to failure (b and d). The insets in a and c show the trends as a function of current density for the RuAl tests. The trends for region II are the same as those seen for region I. As the testing temperature increases the films are subjected to larger stresses, so damage accumulates faster, leading to a faster increase in resistance (slope) and to a shorter transition to region III. NiAl and RuAl show the same trend for the time in region II as a function of time to failure, i.e., they spend approximately the same amount of time in region II, but the trends for the slope in region II are slightly different. Both NiAl and RuAl show a decreasing slope with increasing time to failure; however, NiAl shows a higher slope in region II than does RuAl. This indicates that damage in NiAl is accumulating more rapidly than in RuAl, and as the time in region II is similar for both, more damage

170 accumulates in region II in NiAl than in RuAl. By multiplying the slope in region II by the time in region II it was determined that the NiAl samples showed an increase in resistance of 3-11% in region II compared to 0.3-1.3% for RuAl. The values for RuAl agree well with the data of Zaharia et al. [26], who determined that notched tool steel samples showed an increase in resistance of 0.2-1.6% due to the accumulation of point defects and dislocations introduced during fatigue. The values for NiAl are closer to the data of Kleinert et al. for the fatigue of copper, which resulted in a 5% increase in resistivity that was due to the formation of dislocations and vacancies. The increased amount of damage formation in NiAl could be due to the different levels of stress experienced by the two materials. NiAl has a CTE that is ~400% higher than the substrate, whereas RuAl has a CTE that is ~112% higher than the substrate. This will lead to the NiAl samples experiencing a larger change in stress during testing. As the stresses are lower in the case of RuAl, damage accumulates more slowly; however, less damage is accumulated in RuAl before crack formation and failure (region III).

1000 1000 100 a) b) 10

100 1 100

0.1

10 0.01 10 7.5 8.0 8.5

Current Density (MA/cm2)

(h)

(h)

2

2

t t 1 1

0.1 0.1

NiAl NiAl RuAl RuAl 0.01 0.01 350 400 450 500 550 600 650 700 0.01 0.1 1 10 100 1000 T (°C) Time to Failure (h) max 1 1 c) d)

0.1 0.1

0.01

0.01

1E-3 0.3

0.2

1E-3 Slope in Region 2 (1/dt) 2 Region in Slope 1E-4 0.1 (1/dt) 2 Region in Slope NiAl NiAl 0.0

RuAl 7.5 8.0 8.5 RuAl 2 1E-5 Current Density (MA/cm ) 1E-4 350 400 450 500 550 600 650 700 0.01 0.1 1 10 100 1000 T (°C) Time to Failure (h) max Figure 65 – Results from the analysis of curve fits to the NiAl and RuAl data showing three regions of increasing resistivity. (a) The time in region II as a function of the maximum test temperature, (b) the time in region II as a function of the time to failure, (c) the slope in region II as a function of the maximum test temperature, and (d) the slope in region II as a function of the time to failure. The insets in (a) and (c) show the time in region II and the slope in region II as a function of the current density, respectively (the y-axes have the same units in the insets and the main graphs).

171

Results from the analysis of the curve fits in region III are shown in Figure 66 for the time and slope in region III as a function of temperature (a and c) and as a function of the time to failure (b and d). The trends as a function of temperature are the same as those seen for regions I and II: the time in region III decreases and the slope increases as the test temperature increases. This is consistent with the higher stress associated with higher temperatures, leading to more rapid crack growth and failure. The trends as a function of the time to failure are also analogous to what was observed for regions I and II. The trends seen in the plots of Figure 66 are very similar for NiAl and RuAl, with the possibility of a slightly shorter time in region III for RuAl compared to NiAl tests of equal time.

1000 1000 a) b) 1 100 100

0.1

10 10 0.01

7.5 8.0 8.5 Current Density (MA/cm2)

1 1

(h)

(h)

3

3

t t

0.1 0.1

0.01 0.01 NiAl NiAl RuAl RuAl 1E-3 1E-3 350 400 450 500 550 600 650 700 0.01 0.1 1 10 100 1000 T (°C) Time to Failure (h) max 10 10 c) d)

1 1

0.1

0.1 3 0.01 2

1 0.01 Slope in Region 3 (1/dt) 3 Region in Slope 1E-3 (1/dt) 3 Region in Slope 0 NiAl NiAl 7.5 8.0 8.5 RuAl Current Density (MA/cm2) RuAl 1E-4 1E-3 350 400 450 500 550 600 650 700 0.01 0.1 1 10 100 1000 T (°C) Time to Failure (h) max Figure 66 – Results from the analysis of curve fits to the NiAl and RuAl data showing three regions of increasing resistivity. (a) The time in region III as a function of the maximum test temperature, (b) the time in region III as a function of the time to failure, (c) the slope in region III as a function of the maximum test temperature, and (d) the slope in region III as a function of the time to failure. The insets in (a) and (c) show the time in region III and the slope in region III as a function of the current density, respectively (the y-axes have the same units in the insets and the main graphs).

172

By using Johnson’s equation for crack growth and the resistance increase in region III, we estimated the length of the cracks present just prior to failure. Two different values were used for the resistance increase in region III: one assuming that the entire resistance increase in region III was due to cracking, and the second assuming that some of the resistance increase was due to the formation of dislocations and point defects (region II). These two different values are shown

schematically in Figure 67. The crack length estimated from the values of - ranged

from 0.35-0.94 μm for NiAl and from 0.25-0.48 μm for RuAl. The crack lengths estimated from

were ~40% higher than those estimated from III-II for NiAl and ~10% higher for

RuAl. The crack lengths estimated from - represent ~13% (±4) of the line width for

NiAl and ~9% (±2) of the line width for RuAl. To determine whether these values are realistic it would be necessary to stop the test at various positions along the curve in region III to measure the crack length. Now that the shape of the curve has been characterized it should be possible in the future to estimate the time to reach a given point on the R(t) curve (see discussion below) so that crack growth may be observed.

1.201.20 data data I curve curve fit fit II III R Δ III- II R R R 1.151.15  mt  b 1 Δ III R

R1 1

1 1

1.101.10 R R/R R/R Δ II R1

1.051.05

1.001.00 0 0 1 1 2 2 3 3 4 4 5 5 6 6 tI/II tII/III tf TimeTime (h) (h)

Figure 67 – Figure showing the two different values of the resistance due to crack growth ( and

- ), which were used to calculate the crack length using Johnson’s equation.

The relationships between the slopes in region I and II, and between the slopes in region I and III, are shown in Figure 68a and b, respectively. For all of the RuAl and NiAl samples tested the slope in region I is higher than the slope in region II, which means that the resistance increase due to heating is larger than the resistance increase due to the initial stages of damage formation. When comparing the slopes in regions I and III a transition is noted where for shorter tests the slope in region III is higher than in region I, and for longer tests the slope in region I is higher than in region III.

173

1 10 NiAl NiAl RuAl RuAl slope 1 = slope 2 slope 1 = slope 3 0.1 1

↓T, ↑tf

0.01 0.1

1E-3 0.01

Slope in Region 2 (1/dt) 2 Region in Slope Slope in Region 3 (1/dt) 3 Region in Slope ↑T, ↓tf a) b) 1E-4 1E-3 1E-3 0.01 0.1 1 10 1E-3 0.01 0.1 1 10 Slope in Region 1 (1/dt) Slope in Region 1 (1/dt) Figure 68 – Analysis of the curve fits to all of the NiAl and RuAl data showing 3 regions of increasing resistance. (a) The relationship between the slope in regions I and II, and (b) the relationship between the slopes in region I and III. The solid lines are guides to the eye and represent the case where the slopes are equal.

The analysis of the curve fits of the NiAl and RuAl samples has shown that the curves have a very distinctive shape, which can be described by a combination of power, exponential, and linear terms. If the slope in the first second of the test is measured, then the times and slopes in all of the regions can be estimated (Figure 64a, Figure 68a-b, and Figure 69a-b). This will allow for automated testing to be set up so that tests can be stopped at a certain point along the curve just by comparing the initial slope to the already obtained relationships between the various parameters. Stopping the test at specific points along the curve will allow for periodic examination of the surface morphology of the line and the progress of damage formation can be monitored.

100 1000 a) b)

100 10

10

1

(h)

(h)

2

1

t t 1

0.1 0.1

NiAl NiAl RuAl RuAl 0.01 0.01 1E-3 0.01 0.1 1 10 1E-3 0.01 0.1 1 10 Slope in Region 1 (1/dt) Slope in Region 1 (1/dt) Figure 69 – The relationships between the slope in region I and (a) the time in region I and (b) between the slope in region I and the time in region II.

174

G. Comparison between AC and DC Tests

AC and DC tests were run on co-sputtered RuAl samples that were annealed for 1 min at 1000°C. The samples for the AC and DC tests were all from the same deposition so as to minimize differences between the samples. An example resistance curve is shown in Figure 70a, which shows an AC R(t) curve for comparison. The tests have similar times to failure, but the DC curve shows a significantly larger increase in resistance, indicating that more damage occurs during DC stressing than during AC stressing. This is also shown in the plot of Figure 70b, which shows the resistance increase as a function of the time to failure for AC and DC tests (only samples showing region III). Not only do the DC tests have a larger increase in resistance for equivalent failure times, but they have larger increases in resistance than all of the AC tests. Although the DC tests have a larger increase in resistance, the R(t) curves have a similar shape to the AC curves: heating (region I), slow damage formation (region II), and rapid damage accumulation and failure (region III).

1.16 16 a) b) AC tests 14 DC tests T = 511 C 1.12

) 12 1

10 1.08 8

Resistance (R/R Resistance 6 1.04 T = 609 C

ΔT = 536 C (%) Increase Resistance

2 4 DC 7.20 MA/cm AC 8.13 MA/cm2 peak (5.75 MA/cm2 rms) 1.00 2 0.00 0.06 0.12 0.18 0.24 0.30 0.36 0.01 0.1 1 10 100 Time (h) Time to Failure (h) Figure 70 – Comparison between AC and DC tests conducted on RuAl co-sputtered at 1.5 mTorr and then annealed at 1000°C for 1 min. (a) An example of a DC R(t) curve compared to an AC curve with a similar time to failure, and (b) the increase in resistance observed during electrical testing as a function of the time to failure.

Comparisons between the time and slopes of each region are complicated by the fact that there are only 2 data points showing the three region behavior for the DC tests; however, the time for each region is very similar for the AC and DC tests (as a function of time to failure). It appears that the slopes of each region are slightly higher in the case of the DC tests. For region I (heating), the samples increase in temperature more rapidly as the current is constant for the DC tests, and varying between 0 and Imax in the case of AC tests. Higher slopes in the other two regions would indicate that damage accumulates more rapidly in the case of DC stressing in comparison to AC stressing and at failure more damage should be present. This is confirmed in SEM images of the failed lines, which are shown in Figure 71 and Figure 72. Images of the failed lines show some similarities and some differences between the AC and DC tests. Both types of tests show damaged patches at the line edges with a semi-elliptical shape/semi-circular shape (Figure 72). In the failure region these appear almost porous, with the failure occurring through the center of one of these patches. Another similarity is the appearance of crack-like

175 features, which in the case of AC do not reach through the thickness of the film, but in the case of DC there are regions where these do extend through to the substrate. These features occurred in all of the DC tests, but only in the AC tests with the higher currents. The width of these features near the failure location was ~1 μm for the DC tests. Only one of the AC tests on RuAl showed a second crack near the failure, but more than half of the DC samples had at least 1 crack in the region near the failure site. Additional features that were seen in DC tests, but not in AC tests, were noticeable grain growth (bottom right of Figure 71a) and surface wrinkling (Figure 71c). Due to the increased rate of damage formation in the DC tests, lifetimes were lower when samples were stressed with DC than with AC (see Figure 73). This is similar to what has been seen in the literature as the role of electromigration decreases as the frequency increases.

176

a) 76.7 mA – 0.55 h b) 74.4 mA – 0.07 h

2 μm

2 μm

2 μm 1 μm

c) 76.7 mA d) 2 μm 0.20 h

76.7 mA 0.32 h 2 μm 1 μm

Figure 71 – SEM images of 4 co-sputtered RuAl samples tested under DC conditions until sample failure.

177

a) 87.7 mA b) 82.6 mA 93,600 cycles 0.03 h

1 μm 1 μm

Figure 72 – SEM images of RuAl TMF samples in regions away from the failure location for samples tested until failure using (a) AC and (b) DC stressing conditions.

9.0 AC to failure AC no failure

8.5 DC to failure

) 2

8.0

7.5

7.0

Current Density (MA/cm Density Current 6.5 Plotted as peak AC current density 6.0 0.1 1 10 100 Time to Failure (h) Figure 73 – Comparison of the time to failure under AC and DC stressing conditions for co-sputtered RuAl.

IV. CONCLUSIONS

Samples of gold, ruthenium, NiAl, and RuAl have been tested using an alternating current, which due to Joule heating induces cyclic heating, which in turn causes cyclic stresses in the metallic line structures (due to CTE differences between the line and substrate). The results from tests conducted on NiAl showed minimal effects of sample preparation on the time to failure as a

178 function of ΔT. The ΔT for tests on NiAl ranged from ~250-520°C (Tmax: 360-600°C), with a time to failure of 100’s of hours when ΔT was near 300°C (Tmax: ~400°C). In contrast to the time to failure as a function of temperature, processing did affect the time to failure as a function of the current density. In general, processing that resulted in a lower resistivity led to higher current densities for equal times to failure due to reduced Joule heating for the lower resistivity samples. Ruthenium aluminide was more sensitive to the processing conditions than NiAl, and a lower resistivity did not always increase the current density of the sample. It is believed that reactions with the substrate may play an important role in the failure of RuAl resulting in a decrease in the time to failure for annealing temperatures of 1000°C and higher. RuAl had a longer time to failure than NiAl at high testing temperatures (ΔT > ~350°C), but trends in the data indicate that the time to failure at lower temperature may be higher for NiAl. One of the RuAl data sets had significantly better high temperature performance than the rest of the RuAl data sets or the NiAl samples. For example, for a time to failure of 100 h this data set had a ΔT of ~525°C whereas the rest of the RuAl samples and NiAl had a ΔT of ~300°C. Ruthenium samples with a low resistivity were able to withstand similar levels of ΔT as NiAl while gold resulted in a significantly lower ΔT (100-200°C) than the rest of the samples. In terms of current density, gold was superior to the other materials, which was most likely due to the significantly lower resistivity of the gold samples and hence reduced Joule heating. The gold samples had a resistivity that was 5-8 times lower than the resistivity of the NiAl and RuAl samples. NiAl was able to withstand a slightly higher current density than RuAl again due to the lower resistivity of the NiAl samples in comparison to the RuAl samples.

The resistance as a function of time was plotted for all of the samples during the course of the AC tests. Three different trends were obtained, but the majority of the samples showed very similar R(t) curves: an initial region of increasing R with a decreasing slope, a second region of increasing R with an almost constant slope, and a third region of increasing R with an increasing slope. These three regions have been ascribed to heating, the production of defects (dislocations and vacancies), and crack growth (and possibly voids). It was also determined that by measuring the slope of the R(t) curve at the beginning of the test, the time to failure and the length of each of the three regions could be estimated. This will enable future tests to be paused at specific locations along the R(t) curves so that damage formation and crack growth may be studied with electron microscopy.

179

APPENDIX I: RESISTANCE AS A FUNCTION OF TEMPERATURE

The figures below contain the results of the resistance as a function of temperature measurements obtained prior to TMF testing and used to calculate the temperature during TMF tests. The resistance as a function of temperature results are presented in the same order as in the main body of the chapter, with the exception of the RuAl samples used for DC tests, which are presented along with the rest of the co-sputtered RuAl data. Each data point in the plots represents the average of 6 measurements obtained on a single TMF line. The errors bars (standard deviation) are within the symbol size and are omitted for clarity. One linear fit was obtained for each TMF chip (set of 6 TMF lines) and the results from the curve fits are also plotted for each data set. The slope of the linear fit (TCR) was used to determine the temperature at the beginning of each TMF test. One line was fit to the data from all 6 lines on a TMF chip to obtain an average TCR for that sample. Alternately, each TMF line could have been fit and the slopes averaged to obtain the average TCR for the 6 mm chip. Both methods produced the same value for the TCR, but the r2 value was higher for the individual linear fits than for the single curve fit to all the data from the 6 TMF lines. An example of the procedure used to determine the TCR is shown in Figure A1, which plots the temperature as a function of resistance for 5 TMF lines contained within one 6 mm chip. A linear fit to each TMF line is plotted as well as a linear fit to the entire data set. The TCR obtained by averaging the 5 slopes, or by using the slope from the single fit are identical. Additionally, for the majority of the 6 mm chips the difference between the slope of one of the single curve fits and the average slope was less than 3%.

200 R = 0.1931T + 125.53 r2 = 0.6286 180

) R = 0.1864T + 115.54

 r2 = 0.9999 160

R = 0.1890T + 115.21 r2 = 0.9999

140 R = 0.2050T + 140.84 R = 0.1959T + 123.95

all data linear fit r2 = 0.9998 r2 = 1.000 Resistance ( Resistance line 2 line 3 R = 0.1892T + 122.10 120 line 4 r2 = 0.9999 line 5 line 6 100 0 50 100 150 200 250 Temperature (°C) Figure A1 – Comparison of two different methods of determining the TCR for a set of samples. The graph on the left shows the resistance as a function of temperature for 5 different TMF lines labeled lines 2-5. The line number in the graph corresponds to the number of rectangles located either directly above or below the TMF line on the 6 mm chip (schematic shown on the right). The curve fit to the data from all 5 TMF lines is shown as the thick black line in the plot, with the corresponding equation given in the upper right of the graph. Equations obtained for linear fits to each TMF line are shown adjacent to the corresponding line in the schematic diagram to the right. Taking an average of the slopes from the 5 linear fits results in a value of 0.1931, which is the same as the value obtained for the single curve fit. The data in this figure are the same as anneal 4 shown in the left hand graph of Figure A4.

180

275 120 sample 1 sample 1 sample 2 sample 2 250 sample 3

110

)

)

  225

100 R = 0.25818T + 184.92696 R = 0.16758T + 91.22508

200 r2 = 0.81 r2 = 0.70

Resistance ( Resistance Resistance ( Resistance 90 175

a) b) 150 80 0 50 100 150 200 250 0 25 50 75 100 125 Temperature (°C) Temperature (°C) Figure A2 – Resistance as a function of temperature for (a) ruthenium samples deposited at 7 mTorr and (b) ruthenium samples deposited at 1.5 mTorr.

110 140 sample 1 - 700°C anneal sample 1 - 400°C anneal sample 1 - 400°C anneal sample 1 - 800°C anneal 100 sample 2 - 400°C anneal 120

R = 0.11691T + 72.76378 R = 0.14391T + 76.29154

) )

2 2  r = 0.96  r = 0.61 90 100

80 R = 0.13940T + 65.51129 80 Resistance ( Resistance r2 = 0.92 ( Resistance

70 60 R = 0.15468T + 74.36506 r2 = 0.71 a) b) 60 40 0 50 100 150 200 250 0 50 100 150 200 250 Temperature (°C) Temperature (°C) Figure A3 – Resistance as a function of temperature for (a) annealed Ni/Al multilayers with a wavelength of 20 nm deposited by e-beam and (b) annealed Ni/Al multilayers with a wavelength of 30 nm (deposition 1) deposited by sputtering.

181

225 110 R = 0.13554T + 65.25183 anneal 1 3.8 m wide 2 200 anneal 2 r = 0.79 R = 0.19310T + 123.52852 4 m wide, 263 nm thick anneal 3 100 3.4 m wide r2 = 0.63

175 anneal 4 4 m wide, 266 nm thick )

) 90 R = 0.12325T + 67.43804   150 r2 = 0.79 125 80 R = 0.18699T + 109.86069 r2 = 0.71 100 R = 0.11888T + 65.42381

R = 0.17936T + 105.55429 2 Resistance ( Resistance Resistance ( Resistance 70 r = 0.70 r2 = 0.89 75 R = 0.12245T + 61.17707 2 R = 0.14203T + 53.33891 60 r = 0.95 50 r2 = 0.94 a) b) 25 50 0 50 100 150 200 250 0 50 100 150 200 250 Temperature (°C) Temperature (°C) Figure A4 – Resistance as a function of temperature for (a) annealed Ni/Al multilayers with a wavelength of 30 nm (deposition 2) deposited by sputtering and (b) annealed Ni/Al multilayers with a wavelength of 30 nm, doped with 0.5% Ag, and deposited by sputtering.

175 140 6.4 m lines a) anneal 4 b) 5.2 m lines anneal 3 R = 0.12230T + 119.21009 150 3.3 m lines-1 anneal 2 2 R = 0.20982T + 80.10914 3.3 m lines-2 r = 0.93 120 anneal 1

r2 = 0.97

)

)

  125 R = 0.11831T + 114.71914 r2 = 0.84 100

100 R = 0.07504T + 72.7232 Resistance ( Resistance r2 = 0.96 ( Resistance R = 0.18206T + 73.15236 80 r2 = 0.69 75 R = 0.17257T + 72.61528 r2 = 0.98 R = 0.06043T + 57.78265 R = 0.17358T + 64.73028 2 r2 = 0.96 r = 0.97 50 60 0 50 100 150 200 250 0 50 100 150 200 250 Temperature (°C) Temperature (°C) Figure A5 – Resistance as a function of temperature for (a) co-sputtered NiAl and (b) annealed Ru/Al multilayers with a wavelength of 25 nm deposited by sputtering.

182

130 130 256 nm 266 nm 258 nm - 1 264 nm 120 255 nm 120 267 nm - 2 258 nm - 2 263 nm

267 nm - 1 262 nm

)

)

  110 110

100 100

Resistance ( Resistance Resistance ( Resistance

90 90

a) b) 80 80 0 50 100 150 200 250 0 50 100 150 200 250 Temperature (°C) Temperature (°C) 130 2 260 nm - DC Sample Slope Y-intercept r 264 nm - DC 256 nm 0.15105 84.78260 0.98 120 258 nm – 1 0.14265 81.52072 0.98

) 255 nm 0.15382 85.90134 0.94  110 258 nm – 2 0.15046 82.99244 0.97 267 nm – 1 0.14640 82.15409 0.97

100 266 nm 0.14811 82.36334 0.96

Resistance ( Resistance 264 nm 0.14583 83.85583 0.96 267 nm – 2 0.14806 82.41446 0.94 90 263 nm 0.15002 83.90443 0.97 c) 262 nm 0.15035 84.75060 0.94 80 260 nm – DC 0.15248 84.42227 0.97 0 50 100 150 200 250 Temperature (°C) 264 nm – DC 0.15565 84.71199 0.97

Figure A6 – Resistance as a function of temperature for the co-sputtered RuAl samples tested with (a-b) AC and (b) DC conditions. The table presents the results from the curve fits.

24 20 sample 1 sample 2

22 18 )

) R = 0.03985T + 13.79132 R = 0.03239T + 10.83359   2 20 r2 = 0.97 16 r = 0.99

18 14

Resistance ( Resistance Resistance ( Resistance

16 12

a) b) 14 10 0 50 100 150 200 250 0 50 100 150 200 250 Temperature (°C) Temperature (°C) Figure A7 – Resistance as a function of temperature for gold deposited at (a) 7 mTorr and (b) 4 mTorr.

183

APPENDIX II: CURVE FIT RESULTS

The tables below contain the results from the current fitting of the R(t) curves discussed in this chapter. Each table represents the data from a separate sample set.

Table A1 – Results of the curve fits for NiAl samples fabricated from annealed Ni/Al multilayers with a wavelength of 20 nm deposited by electron-beam evaporation. Time to Current Failure C C C C C C R2 (mA) 1 2 3 4 5 6 (h) annealed at 400°C(4 h) 85.1 0.73 0.07580 0.02631 0.01974 4.31642 10.12888 0.90749 0.98768 82.2 1.73 0.05852 0.05851 0.03665 0.00576 6.36001 0.93082 0.99388 annealed at 400°C(4 h) – 700°C(10 min) 71.1 8.09 -0.02009 0.28901 0.00722 1.7576E-09 5.9624 1.02338 0.99981

Table A2 – Results of the curve fits for NiAl samples fabricated from annealed Ni/Al multilayers with a wavelength of 30 nm deposited by sputtering – deposition 1. Time to Current Failure C C C C C C R2 (mA) 1 2 3 4 5 6 (h) annealed at 400°C(4 h) – 800°C(30 s) 83.6 8.10 -0.05180 0.31409 0.01189 1.0193E-12 11.16829 1.06479 0.99650 82.5 9.10 -0.04886 0.24131 0.01708 6.4880E-09 7.15606 1.05906 0.99883

Table A3 – Results of the curve fits for NiAl samples fabricated from annealed Ni/Al multilayers with a wavelength of 30 nm deposited by sputtering – deposition 2. Time to Current Failure C C C C C C R2 (mA) 1 2 3 4 5 6 (h) anneal 1: 800°C(30 s) – 10°C/s – 500°C(30 s) 91.6 0.80 -0.03997 0.09555 0.20204 0.22227 6.9756 1.04428 0.99844 89.2 155.88 -0.05065 4.14378 0.00093 8.3954E-20 8.19731 1.0969 0.99765 90.5 360.95 -0.04261 8.29858 0.0004 3.6339E-18 6.31497 1.09265 0.99854 anneal 2: 1000°C(30 s) – 10°C/s – 500°C(30 s) 77.7 0.20 -0.01052 0.01807 0.1733 1.4777E+42 66.74786 1.01078 0.98248 anneal 3: 400°C(5 min) – 1000°C(30 s) – 10°C/s – 500°C(30 s) 80.1 1.60 -0.01902 0.09216 0.03949 0.00039 6.69549 1.02337 0.99271 76.5 3.75 -0.02382 0.11624 0.02178 5.8375E-06 6.23399 1.02199 0.99676 anneal 4: 1000°C(30 s) – 5°C/s – 900°C(15 s) – 2°C/s – 800°C(15 s) – 2°C/s – 600°C(30 s) 71.8 14.55 -0.02429 0.18006 0.00436 1.1697E-09 5.97206 1.0301 0.99826

184

Table A4 – Results of the curve fits for NiAl samples doped with 0.5% Ag fabricated from annealed Ni/Al multilayers with a wavelength of 30 nm deposited by sputtering. Samples were annealed at 400°C for 4 h. Time to Current Failure C C C C C C R2 (mA) 1 2 3 4 5 6 (h) 84.2 1.27 0.02995 0.15703 0.00474 0.00476 5.49182 0.97000 0.93733 90.2 4.10 0.03161 0.32541 0.00164 0.00002 4.85532 0.97352 0.98376 89.0 5.40 -0.05246 0.20631 0.01501 1.115E-6 6.23912 1.06085 0.99690

Table A5 – Results of the curve fits for co-sputtered NiAl films annealed at 400°C for 4 h. Time to Current Failure C C C C C C R2 (mA) 1 2 3 4 5 6 (h) 3.3 μm line width 70.6 24.83 -0.08737 0.81670 0.00290 1.3867E-07 3.91025 1.08140 0.99809 5.2 μm line width 100.3 33.63 -0.06940 1.35995 0.00348 3.36E-22 12.49934 1.08778 0.99813

185

2

R

0.98600 0.98319 0.96749

0.99503

9

0 0

C

0.71281 4.51427

8

0 0

C

2.51433 2.37235 -

-

600°C(30s)

–

7 0

0

C 0.80539

0.70303

2°C/s 2°C/s

–

6

C

1.04779 1.02079 0.98801 0.97027 C(30s)

C(30s)

°

°

5

600

800°C(15 s) 800°C(15s)

600

C

–

–

–

6.26703 2.44739 3.01939

10.55508

C/s

C/s °

°

2°C/s

10 10

–

–

–

8 -

s)

10

-

4

C (15

C

0.00017

C(30 s) C(30s)

3.84E

°

C(30 s) C(30s)

1.3031E °

°

00

9

00

17846604948.99407

–

3: 9 3:

C/s

3

5°

C

–

0.00334 0.00009

anneal

0.00384 0.36015

- -

anneal 1: 1000 anneal

0

2 C(30 s) C(30s)

C

°

0.01621 0.03552 0.04035

0.1021

00

1

2: 11 2:

C

0.02004 0.04288 0.04801

0.04505

- - -

-

the forfits curve RuAl fabricatedfrom annealed Ru/Almultilayers a wavelength with of 25 nm.

anneal anneal

(h)

0.07 9.80

21.80

154.15 Failure Failure

Time to to Time

Results ofResults

–

6

A

84.8 94.3 93.9 92.2

(mA)

Current Current

Table

186

Table A7 – Results of the curve fits for co-sputtered RuAl. Time to Current Failure C C C C C C R2 (mA) 1 2 3 4 5 6 (h) 1000°C(60 s) – 15°C/s – 500°C(30 s) 88.0 0.08 -0.03865 0.04553 0 1.888E+16 16.91440 1.03800 0.96773 89.4 0.12 -0.04022 0.04222 0 16172.70552 6.20590 1.04162 0.98736 89.4 0.13 -0.03117 0.03321 0.16828 1.1964E+23 28.81626 1.03292 0.98618 89.4 0.20 -0.02457 0.04014 0.08410 116857626298.16447 18.79989 1.02659 0.97354 87.7 0.30 -0.03874 0.06800 0.03996 913178448.98537 20.62566 1.04255 0.98200 87.7 0.38 -0.04011 0.06491 0.04872 2171938509.08727 26.76635 1.04211 0.98623 87.7 0.38 -0.03310 0.06598 0.04489 6.1118E+12 35.69387 1.03652 0.97857 82.1 0.42 -0.02559 0.06025 0.05089 373601372.56642 27.62372 1.02758 0.97760 82.1 0.78 -0.03388 0.08175 0.02049 3663077.72848 82.98335 1.03696 0.97550 87.6 0.87 -0.04589 0.09248 0.01935 2294.57694 97.17256 1.05073 0.98143 84.7 1.12 -0.04082 0.11707 0.00703 1.0566E-07 90.09999 1.05928 0.97146 82.1 2.38 -0.03890 0.12731 0.00330 1.0205E-13 28.28646 1.04423 0.97225 82.1 2.93 -0.03567 0.12067 0.00313 1.678E-15 27.28844 1.04223 0.95627 84.8 3.28 -0.04283 0.133353 0.00188 4.79E-35 62.54987 1.05092 0.96043 82.1 3.58 -0.04062 0.13134 0.00173 1.2408E-16 24.80353 1.04712 0.95988 82.0 3.68 -0.03968 0.15651 0.00234 1.5843E-29 47.41407 1.04557 0.96281 84.9 19.77 -0.04142 0.12814 0.00016 3.4051E-156 118.17640 1.04774 0.84037 N2 anneal: 1000°C(30 s) – 10°C/s – 500°C(30 s) 84.7 0.12 -0.01677 0.02170 0.13243 101.4693 4.06732 1.01714 0.98836 84.8 0.22 -0.01270 0.01911 0.17941 80658.51329 10.25778 1.01268 0.99452 84.7 0.23 -0.01084 0.02759 0.17671 842.19051 7.12992 1.01089 0.99554 N2 anneal: 1000°C(60 s) – 10°C/s – 500°C(30 s) 84.9 0.22 -0.01711 0.02831 0.12508 6.5635E+15 27.29717 1.01749 0.99144 84.8 0.35 -0.02440 0.04500 0.07524 7.1522E+24 57.46514 1.02675 0.99230 84.7 0.37 -0.02337 0.04482 0.08489 2.67E+34 82.36987 1.02627 0.98657 84.9 0.45 -0.02549 0.05537 0.05769 37006070750.34821 33.57531 1.02547 0.99091 85.0 1.03 -0.03605 0.10584 0.01270 0.00050 75.46031 1.04051 0.98437 84.9 1.43 -0.03800 0.11979 0.00844 7.35E-09 39.09495 1.04176 0.98749 81.9 1.75 -0.03641 0.12616 0.00534 9.7159E-38 144.99985 1.04000 0.98359 81.9 6.25 -0.03858 0.16200 0.00098 3.3059E-153 189.53989 1.04448 0.97308 DC: 1000°C(60 s) – 15°C/s – 500°C(30 s) 76.8 0.33 -0.04695 0.04265 0.15248 4.567E+15 35.27187 1.0499 0.97293 76.8 0.23 -0.05892 0.04018 0.16625 7.8763E+33 57.35003 1.05599 0.96521

187

Table A8 – Results of the curve fits to ruthenium samples deposited by sputtering and annealed at 400°C for 4 h. Time to Current Failure C C C C C C R2 (mA) 1 2 3 4 5 6 (h) deposited at 1.5 mTorr 72.4 0.27 -2.17074 2.72704 0 2575092.99826 13.63961 3.16789 0.99769 70.0 1.70 -0.11293 0.46976 0.02266 0.00555 4.50734 1.11724 0.99772 deposited at 7.0 mTorr 58.7 13.60 0.06032 0.37106 0.00217 1.9585E-67 57 0.87773 0.93610 56.9 30.50 0.07066 0.40405 0.00004 7.7164E-09 4.33792 0.86928 0.95646

188

REFERENCES

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[37] J.T. Guo, J. Zhou, G.S. Li, and Y.H. Qi, Effect of Ag alloying on microstructure, mechanical and electrical properties of NiAl intermetallic compound. Materials Science Forum, 2003. 426-432: p. 1631-1636.

191

CHAPTER 6

NICKEL ALUMINIDE AND RUTHENIUM ALUMINIDE FILMS FOR MICROELECTROMECHANICAL SYSTEMS

192

I. INTRODUCTION

Structural films for MEMS require a combination of a low electrical resistivity, low residual stress and stress gradients, and sufficient thermal and mechanical stability so that the resulting devices can be driven electrically and respond predictably. When we add the requirement of superior damage tolerance (i.e., a room temperature fracture toughness greater than the ~1 MPa m of silicon-based films) and limit the upper processing temperature to insure compatibility with CMOS processing (i.e., processing temperatures < 450ºC) [1], the list of materials systems that are amenable to micromachining becomes very limited. While SiGe- based films with Ge sacrificial layers can be deposited under conditions that allow for direct integration with CMOS electronics, their damage tolerance is essentially the same as micromachined Si films with SiOx sacrificial layers. In contrast, pure metals such as Au, Al, and Ni and their alloys provide low resistivity, controllable residual strains, high damage tolerance, and low processing temperatures. Unfortunately, they lack the mechanical stability required for many MEMS applications, especially when deposited under conditions that lead to nanoscale grain morphologies [1]. In principle, metal-metal (intermetallic) compound-based structural films should have the requisite combination of physical properties for a wide range of applications and service environments. In this work we demonstrate both electrically- and thermally-actuated MEMS micromachined from Ni-Al and Ru-Al intermetallic thin films using processing conditions that are amenable to co-fabrication with integrated circuits.

Ordered intermetallic thin films, like their bulk counterparts, exhibit a desirable combination of properties including, but not limited to, high tensile strength at room and elevated temperatures, good oxidation and corrosion resistance, and metallic electrical conductivity [2]. While Texas Instruments, Inc. has successfully demonstrated the use of intermetallic films for micromirrors [3] and amorphous TiAl3 for hinges in their digital micromirror devices (DMD) [4], there has been limited development of other micromachined intermetallic films. Structural films based on the Ni-Al binary system are particularly attractive because of their low electrical resistivity, potential for low as-fabricated residual strains on silicon substrates, and a range of thin film processing routes that can be used. Although there are few studies on thin film RuAl, structural layers based on Ru-Al should have similar physical properties to Ni-Al with the added benefit of a lower thermal expansion mismatch with silicon substrates.

Intermetallic thin films have been fabricated for a variety of binary systems using co-deposition [5-8] and multilayer reaction methods [9-12]. The specific case of NiAl has received considerable attention in the literature. NiAl thin films have been fabricated at temperatures between 225ºC and 500ºC by the annealing of multilayers [13-16], at 200ºC and above by co- evaporation [17], and close to room temperature by co-sputtering (unheated substrates, 77ºC < T < 150ºC) [18-20]. NiAl thin films are more conductive (resistivity: 24-54 cm) [20-23] than both Si (bulk resistivity: 100 mcm) [24] and SiGe (resistivity: 0.5-75 mcm) [25-26], and bulk NiAl is more oxidation resistant than SiGe and is described as having good corrosion resistance. While this presents a challenge from the perspective of micromachining, it also makes them potentially useful in severe environments (i.e., high temperature and/or corroding atmospheres). Also, NiAl has a room temperature fracture toughness of 4-14 MPa m , which increases to 10- 50 as the temperature increases to 350-400ºC [27-28], compared with a room

193 temperature fracture toughness of 0.8-1.1 MPa m for Si [29-32] that shows a moderate increase to 3.3 at 925ºC. There are no values cited in the literature for SiGe, but the fracture toughness should fall between the values of Si and Ge (Ge fracture toughness: 0.60-0.64 [33-34]). Even the relatively brittle Al-rich NiAl has a slightly higher fracture toughness (1.1-1.6 ) than Si and SiGe [35].

Although less well studied, there are a few literature reports on the fabrication of RuAl thin films. RuAl has been formed in regions of an Ru/Al bi-layer using laser irradiation [36], and RuAl films have been co-sputtered at 100-350ºC [37] for use as an underlayer for growing epitaxial FePt. A more recent study [38] fabricated nearly single phase RuAl from 2 m thick Ru/Al multilayers with a bi-layer thickness of 88 nm. The authors determined that annealing temperatures below 350ºC resulted in the formation of a solid solution, RuAl6 began to form at 350ºC, and no further phase formation was noted until ~550ºC, when RuAl started to form. After annealing at 630ºC the film was mostly RuAl, but some unreacted Ru was also present. In addition, we have fabricated RuAl films using co-sputtering and annealed Ru/Al multilayers (see Chapter 3 for additional details). Bulk RuAl also has a lower resistivity (13-65 cm) [39-40] than Si and SiGe; is resistant to attack from many chemicals such as aqua regia, FeCl3, NaOH, and mixtures of HNO3 and HF [41]; and is expected to have a higher fracture toughness than NiAl, with an estimated value of ~10 [42].

While countless intermetallic compounds can be fabricated, this work focuses on NiAl and RuAl ordered intermetallic structural films because of their potential combination of low resistivity, high strength, moderate toughness, good oxidation and corrosion resistance, and in the case of NiAl, demonstrated low temperature thin film processing. Additionally, both should exhibit some room temperature ductility (i.e., damage tolerance), unlike the majority of intermetallic compounds. In this work we explore co-deposition and multilayer reaction growth modes that can be used to fabricate relatively low residual strain, high conductivity NiAl and RuAl films. We then explore two potential sacrificial layers by determining the resistance of the ordered intermetallic films to dissolution in HF and XeF2. Finally, we use a single mask liftoff process and a XeF2 release procedure to demonstrate electrostatically actuated torsional resonators [43- 44] and bent beam thermal actuators [45] that are micromachined from the NiAl and RuAl structural films.

II. EXPERIMENTAL PROCEDURE

The development of the intermetallic micromachining strategy was divided into three phases: blanket film deposition, film patterning and release, and device demonstration. The blanket films were used to establish the processing conditions required to create the ordered intermetallic structural films and the physical properties relevant to electromechanical applications (crystal structure, grain morphology, and electrical resistivity). The samples studied were primarily NiAl and RuAl intermetallic thin films, although some films were a combination of phases. The Ni-Al films were fabricated using annealed multilayers of Ni/Al that were deposited with electron

194 beam evaporation (e-beam) or sputtering, as well as co-sputtering of elemental targets1. Ru-Al films were deposited using sputtered Ru/Al multilayers and co-sputtering of Ru and Al targets. Source materials for e-beam deposition were aluminum (99.9999% purity) and nickel pellets from Alfa Aesar (99.995% purity). Targets for sputtering were nickel (99.99%), aluminum (99.9995%) and ruthenium (99.95%) elemental targets from Kurt J. Lesker (Al, Ni) and Williams Advanced Materials (Ru). All samples studied were deposited in a vacuum chamber with a base pressure < 10-7 Torr. During deposition the sample was rotated to increase film uniformity. No substrate heating was applied, although the temperature of the substrate rose to ~40-70ºC by the end of the deposition, based on a thermocouple placed near the substrate (sputtered samples). After deposition, some of the samples were annealed in ultra-high purity (UHP) Ar in a rapid thermal annealing (RTA) furnace for short anneals (1-10 min), or in a tube furnace for longer anneals (4 h). After annealing, the resistivity and film thickness were measured using the 4-point probe method (Keithley 2400 SourceMeter, 9 measurements, 10 mA current) and mechanical profilometry (Tencor P-10, at least 7 measurements), respectively. Additionally, some of the films were studied with grazing angle x-ray diffraction (XRD; Philips MPD), tapping mode atomic force microscopy (AFM; Digital Instruments Dimension 3100), and transmission electron microscopy (TEM; JEOL 2010F). Cross-sectional samples for TEM were prepared using a Fischione Ion Mill using conventional techniques.

The patterning and etching phase of the project examined potential sacrificial layer materials and the selectivity of the etching chemistry for both 50% HF and XeF2 release strategies. Two different sets of samples were used: The first set were blanket films (143-247 ± 2 nm thick) deposited onto oxidized (100) silicon wafers and then immersed in HF (49%) for 5 or 10 min to determine whether HF etching could be used to remove the sacrificial SiO2 layer without damaging the intermetallic thin film. The second set of specimens to determine the viability of release were fixed-fixed beams (~35 μm long × 10 μm wide, 165-229 nm thick) that were patterned on (100) Si wafers using a conventional liftoff process with a contact mask and photoresist (SIPR 3251-2.0, Shin-Etsu MicroSi), which was removed after deposition with acetone. Based on the results of the blanket layer studies, the following films were patterned on Si: as-deposited films of NiAl co-sputtered at 1.5 and 7.0 mTorr, annealed and as-deposited RuAl co-sputtered at 1.5 mTorr, and annealed Ru/Al multilayers deposited by sputtering. The underlying silicon substrates were then removed with XeF2 (Xactix, 7-11 min etch time) to create free-standing, fixed-fixed beams. The blanket and patterned intermetallic structural films were examined for cracking or other evidence of damage after exposure to the various etching environments using light (Olympus BX60M, Sony DXC-960MD CCD camera) and scanning electron microscopy (SEM: FEI Quanta 200). In the cases where the compressive residual strains in the film were large enough to induce buckling along the edges of the undercut anchor regions, white light interferometry (Wyko NT1100) was used to characterize the shape of the deformed film. The compressive stress, , in these buckled regions was then calculated using the following equation [46-48]

2 E E  πA  σ  ε    Eq. 1 1 υ 1 υ  λ 

1 In this manuscript the hyphen between two elements (M) indicates a binary material system (i.e., M1-M2), and a slash denotes alternating layers of the metals (i.e., M1/M2). 195

where E is the film elastic modulus,  is the Poisson’s ratio of the film,  is the strain in the film, A is the amplitude of the buckled film, and  is the wavelength of the buckled film.

Finally, electrostatically actuated, torsional resonators and bent beam thermal actuators (~600 nm thick) were fabricated on (111) Si wafers and released with XeF2. Arrays of 18 resonators consisting of a toroidal shaped mass (inner radius ~30 μm, outer radius ~300 μm, spanning an arc of 60) attached to a cantilever beam spring (with and without notches located at the center of rotation) were fabricated. The cantilever beams were ~40 μm long with various widths and notch configurations (3 with 10 μm wide cantilever beams, 6 with 20 μm wide cantilever beams, and 3 groups of 3 cantilever beams 20 m in width containing V-shaped notches that extended approximately 25, 50, and 75% across the width of the beam). The comb drive actuators consisted of 10, ~5 μm wide fingers spanning a ~10 arc with 5 of overlap. The bent beam thermal actuators consisted of two sets of 10 beams in a V configuration with an angle of ~170 between the two set of beams. There were 3 different sizes of thermal actuators: 4 μm  300 μm, 5 μm  375 μm, and 6 μm  450 μm (width  length of the beams). The resulting MEMS were inspected using light and scanning electron microscopy to establish process yield and the quality of the fabrication process. The electrostatically actuated resonators were then driven at their first (out-of-plane) bending mode in laboratory air using an arbitrary waveform generator (Agilent 33120A) and a high voltage amplifier (Apex PA85). The waveform was sinusoidal with no direct current (DC) offset with amplitudes ranging from 19-95 V. The bent beam thermal actuators were driven with a constant current ranging from 0 to 105 mA. The motion of the electrostatically and thermally actuated MEMS was viewed on a manual probe station (Suss Microtek PM5) and recorded with a digital camera (PixeLINK PL-A776).

III. DEPOSITION AND ANNEALING OF INTERMETALLIC STRUCTURAL FILMS

A variety of processing conditions have been used to create blanket Ni-Al and Ru-Al films that were subsequently annealed to control the phases present, electrical resistivity (shown in Table 1), grain morphologies, and residual stresses. Ni-Al films formed from the reaction of multilayers had wavelengths2 of 25 and 30 nm. After annealing for 4 h, TEM on the 30 nm wavelength film confirmed that this sample was single phase NiAl with grains ranging in size from ~20-60 nm (Figure 1a). Ru-Al films formed from the reaction of Ru/Al multilayers had wavelengths from 6.7 to 25 nm, and required higher temperature anneals to become single phase RuAl (in comparison to NiAl). A Ru/Al multilayer with a wavelength of 25 nm was a combination of Ru and RuAl after annealing for 193 h at 400ºC (Figure 1b), whereas a 30 nm wavelength Ni/Al multilayer was completely reacted to form NiAl after 4 h at 400ºC. XRD indicated less Ru and more RuAl as the wavelength decreased.

Co-sputtered NiAl films were deposited using Ar pressures of 1.5 mTorr and 7.0 mTorr. Films deposited at 1.5 mTorr were determined from XRD and TEM analysis to be the single phase B2 NiAl in the as-deposited and annealed conditions. The films exhibited columnar grains with a grain size (~10-25 nm) less than the annealed multilayers, as shown in Figure 1c. The grain size

2 The wavelength of a sample refers to the bi-layer thickness of the multilayer (1 layer of Al and 1 layer of Ru or Ni). 196 of the annealed co-sputtered NiAl (~15-25 nm) was similar to the as-deposited co-sputtered film. The film co-sputtered at 7.0 mTorr was determined from XRD to be a combination of NiAl and Ni2Al3.

RuAl co-sputtered films were only deposited at 1.5 mTorr Ar pressure, and were determined by TEM to be single phase and polycrystalline as-deposited with a grain size of ~5-10 nm. The grain size was not measured after annealing, but grain growth is expected to be limited due to the relatively low homologous temperature during annealing (T/Tm = 0.3 for a 400ºC anneal, where T and Tm are the annealing and melting temperatures in K). Also, the resistivity of the film is significantly higher than the bulk value [39], ~8.6 times higher, compared to the 2.4 times higher NiAl thin film value compared to bulk NiAl [49-50].

Table 1 – Resistivity of Ni-Al and Ru-Al Films after Annealing at 400ºC for up to 193 h. Resistivity (cm) – Annealed at 400ºC† Material Deposition Conditions 0 h 4 h 40 h 193 h Ni-Al co-sputter – 1.5 mTorr 45.5 (1.5) 24.1 (0.1) multilayer 25 nm wavelength 18.3 (0.1) multilayer 30 nm wavelength 15.5 (0.0) Ru-Al co-sputter – 1.5 mTorr 157 (4) 103 (0) multilayer – 6.7 nm wavelength 121 (0) multilayer – 9.6 nm wavelength 70.5 (0.1) multilayer – 25 nm wavelength 125 (0) 72.7 (0.1) 58.3 (0.1) †Numbers in parentheses are the standard deviation of 9 measurements.

197

a)

100 110 111

200 210 10 nm 310 211 220 221

b) c)

300/202 211/103 105 210/110 200 111 100 110/002 110 100 310 20 nm 310 25 nm 111 220/004 200 211 220 210 Figure 1 – Cross-sectional TEM images of films formed from annealed multilayers and co-sputtering. (a) Ni/Al 30 nm wavelength multilayer deposited by sputtering and annealed at 400ºC for 4 h. (b) Ru/Al 25 nm wavelength multilayer deposited by sputtering and annealed at 400ºC for 193 h. (c) As-deposited NiAl co-sputtered at 1.5 mTorr. Indexed diffraction patterns are shown in the insets of the figures (blue-RuAl, yellow-Ru, white-NiAl), and the scale bar in each is 5 nm-1. The diffraction patterns in (a) and (c) only have diffraction rings corresponding to NiAl (white), whereas many of the rings in (b) can be indexed as either RuAl (first number-blue) or Ru (second number-yellow). Rings that can be clearly identified as RuAl are the (100), (111), (200), and (310) planes, whereas there is only 1 ring that can be clearly identified as belonging solely to Ru, i.e. the (105) plane.

IV. PATTERNING AND RELEASE OF INTERMETALLIC STRUCTURAL FILMS

In this section we describe how blanket films were exposed to aqueous HF and the gas phase XeF2, and then fixed-fixed beams were fabricated from NiAl and RuAl by removing the sacrificial Si with XeF2. The purpose was to determine an appropriate combination of structural

198 layer, sacrificial layer, and etchant that produced low-stress, crack-free, free-standing films. Excessive stress in the films may cause deformation, fracture, buckling, delamination, as well as microstructural changes [51], some of which may be exacerbated in the presence of an aggressive environment, especially if the stress is tensile. In general, sputtered films have tensile stresses at high sputtering pressures, which decrease as the sputtering pressure decreases until the stress becomes compressive [52-54]. Some of the intrinsic stress may be relieved by annealing; however, additional extrinsic stresses can occur during cooling (from the deposition temperature or after annealing) and are due to the different coefficients of thermal expansion (CTE) of the film and substrate. These extrinsic stresses can be estimated using the following equation, assuming the stresses are elastic [55],

  ET Eq. 2 where  is the change in stress, E is the Young’s modulus of the film,  is the difference in thermal expansion coefficients of the film and substrate, and T is the temperature change. -6 -6 Because NiAl and RuAl have higher CTEs than Si (Si = 2.610 /ºC [24], NiAl = ~1310 /ºC -6 [56], RuAl = 5.510 /ºC [57]), the stresses formed due to cooling will be tensile. Because of the closer match in CTE between RuAl and Si, the magnitude of the extrinsic stresses will be lower (by ~42%) in the case of RuAl for equivalent annealing temperatures. Due to intrinsic film stress and the stress created/relieved during annealing, NiAl and RuAl films have been fabricated with stresses ranging from compressive to tensile, and with further improvements in processing it is expected that stress-free films can be realized.

A. Resistance of Blanket Films to HF

While both bulk RuAl and NiAl are resistant to HF attack [41], thin films cracked and buckled in the presence of HF. The cracking is presumed to arise from environmentally-assisted mechanisms in HF, and the density of cracks should be correlated with the magnitude of the tensile residual stress in the film. Similarly, buckling of the films arises due to decohesion of the film-substrate interface after ingress of HF via a crack or other flaw, and the extent scales with the magnitude of the compressive residual stress in the film. Although etching with HF was not successful in creating crack-free, free-standing films, the results yield valuable information on the stress-state in the deposited films, as well as the effect of annealing on the stress-state. The cracking that was observed in optical images of the annealed and HF etched Ni/Al multilayers deposited by e-beam evaporation indicated the presence of tensile stresses in the films. In an effort to reduce the stresses in the film and to decrease the film cracking, different annealing profiles were used on sections of a Ni/Al multilayer from the same deposition. Changing the annealing conditions did not stop the films from cracking after exposure to HF, but did change the fraction of the surface covered in cracks. The best annealing condition was found to be an initial anneal at 400ºC followed by incremental cooling (cooled to 350ºC and held for 2 min, cooled to 300ºC and held for 2 min, and then cooled to 40ºC). An image of a section of a sample cooled incrementally and then etched in HF is shown in Figure 2. The spiral cracking shown in Figure 2 was present in most of the etched films and is an indication of tensile stresses.

199

75 μm

Figure 2 – Light microscope image of a NiAl film immersed in HF for 10 min. The film was deposited as a Ni/Al multilayer with a 25 nm wavelength using e-beam evaporation. After deposition the film was annealed in an RTA for 10 min at 400ºC, cooled to 350ºC and held for 2 min, cooled to 300ºC and held for 2 min, and then cooled to 40ºC.

The cracking and tearing of the films after HF etching indicates that the annealed multilayer films deposited by e-beam evaporation have tensile residual stresses, which occur as the film is cooled after the reaction of Ni with Al to form NiAl. Assuming the stresses in the film cause only elastic and not plastic strains, the stress created on cooling the NiAl from 400ºC to room temperature (~25ºC) can be approximated by Eq. 2. This approximation yields a change in stress of 733 MPa (tensile), which is higher than several values quoted for both the yield and fracture stress of bulk NiAl [58-59]. To reduce the tensile stresses that occur during cooling, lower annealing/reaction temperatures are needed, which will require a decrease in the wavelength of the film in order to fully react the Ni/Al multilayers and create NiAl. The literature shows that NiAl can be formed at lower temperatures as the film wavelength is decreased, and we have noted that a Ni/Al 30 nm wavelength film annealed at 200ºC for 4 h is a combination of NiAl3 and Ni, and therefore not completely reacted.

In addition to multilayers deposited by evaporation, samples were fabricated via co-sputtering, and the results are shown in Figure 3 for films sputtered at 1.5 mTorr and then exposed to HF for 10 min. As-deposited films show a considerable amount of compressive stress as is evidenced by the buckling of the film (Figure 3a-b) including some regions of periodic, tortuous (“telephone-cord”) buckling. There are also some cracks present around the edges of the raised features. The film is more intact (less cracked) than the annealed multilayers, but the considerable buckling when etched with HF makes this material/etchant combination undesirable for MEMS applications. When the film is annealed, the stresses appear to change from compressive to tensile as is shown by the change from primarily buckling to cracking when annealing at 400ºC (Figure 3c) and then film tearing after annealing at 600ºC (Figure 3d). The

200 change from compressive to tensile stresses after annealing indicates that much of the compressive stress has been relaxed during the short anneals at 400ºC and 600ºC; the tensile stresses resulting from cooling the NiAl back to room temperature as discussed earlier. The amount of stress produced during cooling from 600ºC to room temperature would be approximately 1.12 GPa. It is possible that using lower annealing temperatures may result in films with lower stress. This approach will be used on the fabricated resonators. The tensile stress in the annealed films discussed in this section demonstrates that they are unsuitable for MEMS based on an HF sacrificial layer release process, as the devices would have a tendency to crack.

a) b)

50 μm 10 μm

c) d)

50 μm 100 μm

Figure 3 – NiAl film co-sputtered at 1.5 mTorr Ar pressure and then exposed to HF for 10 min. The images represent different sections of the same sample: (a) as-deposited material, (b) higher magnification of the as- deposited NiAl, (c) annealed at 400ºC for 10 min, and (d) annealed at 600ºC for 1 min. In (c) and (d), the lighter gray represents areas of exposed substrate.

Multilayers of Ru/Al were also fabricated, annealed and then exposed to HF for 10 min. The result for an annealed and etched 6.7 nm wavelength multilayer, which is primarily RuAl is shown in Figure 4a, where the buckling of the film indicates compressive stresses rather than the tensile stresses observed in annealed Ni/Al multilayers (similar results to Figure 4a were obtained for 9.6 nm and 15.4 nm wavelength Ru/Al multilayers).

201

a) b)

50 μm 50 μm

c) d)

50 μm 5 μm

Figure 4 – (a) Ru/Al multilayer film annealed at 400ºC and then immersed in HF for 10 min. The film is a 6.7 nm wavelength multilayer annealed for 29 h. RuAl films co-sputtered at 1.5 mTorr Ar pressure and then exposed to HF for 10 min: (b) as-deposited film, (c) film annealed at 400ºC for 30 min, and (d) high magnification image of the film annealed at 600ºC for 1min.

RuAl films co-sputtered at 1.5 mTorr Ar pressure were also exposed to HF for 10 min (Figure 4b-d). As with the co-sputtered NiAl films, as-deposited RuAl is under compressive stress, but less damage (e.g., cracking and/or delamination) was observed after HF etching in the case of RuAl than NiAl. Another difference in the two materials is that annealing does not significantly affect the stress state in the RuAl film: the as-deposited (Figure 4b) and annealed films (Figure 4c-d) all show buckling, indicating compressive stresses in the material. As with the as- deposited NiAl film, the RuAl films under compressive stress all show cracking around some of the raised areas of the film. Since all of the films etched with HF exhibited some amount of cracking regardless of the processing conditions, a different combination of etchant/sacrificial layer is required for NiAl and RuAl MEMS. The next section explores the use of the gas-phase etchant XeF2 for removal of a Si sacrificial layer.

202

B. Resistance of Patterned Films to XeF2

Another possibility explored for sacrificial layer etching was XeF2. To check the compatibility of XeF2 as a sacrificial layer etchant for NiAl on Si, a section of a blanket co-sputtered NiAl film was etched for 4 min in XeF2. Both the as-deposited and annealed (400ºC) conditions of the NiAl films were unaffected by the XeF2. Similar to what was seen in the HF etching, the as-deposited film is under a compressive stress, whereas the annealed film is under a tensile stress. This was seen as edge wrinkling (compressive) and curling under (tensile) where the substrate had been etched away from the film. These results are promising for the application of NiAl as MEMS devices as the free-standing sections of the films did not show any cracking, in contrast to the samples that were exposed to HF.

The results have shown that it is possible to obtain free-standing films of NiAl and RuAl by etching the silicon substrate with XeF2. Images of free-standing NiAl films fabricated in this manner are shown in Figure 5a and b. The best result obtained with the NiAl films was for those deposited at 1.5 mTorr Ar pressure (Figure 5a). The films deposited at 1.5 mTorr show some slight buckling indicating compressive stresses, but no cracks were observed in the samples. On the other hand, the sample deposited at 7.0 mTorr Ar pressure had tensile stresses and showed significant cracking in both the as-deposited (Figure 5b) and annealed samples. The difference in stress state between the 1.5 mTorr and 7.0 mTorr samples was expected, as a higher working pressure during sputtering tends to favor tensile stresses in the film, which decrease as the pressure decreases until the stresses become compressive.

203

a) 25 μm b) 25 μm

c) d)

25 μm 25 μm

Figure 5 – Patterned co-sputtered NiAl and RuAl films on silicon after exposure to XeF2. The SEM images are for (a) 1.5 mTorr NiAl as-deposited, (b) 7.0 mTorr NiAl as-deposited, (c) as-deposited RuAl co-sputtered at 1.5 mTorr Ar pressure, and (d) RuAl deposited at 1.5 mTorr and annealed for 4 h at 400ºC.

The results for RuAl films deposited at 1.5 mTorr and then exposed to XeF2 are shown in the images of Figure 5c and d for the as-deposited and annealed (400ºC-4 h) films, respectively. As for the NiAl films, the as-deposited RuAl is under a compressive stress and shows the absence of cracking. Unlike NiAl, the annealed RuAl also appears to be under a compressive stress (confirming the results of the HF exposure experiments), and no cracking was observed. The as- deposited NiAl sample contains significantly less wrinkling than the as-deposited RuAl, indicating that the NiAl film is under less compressive stress and therefore may yield better MEMS in terms of curvature and flatness. This is opposite to what one might predict from the HF exposure, as the RuAl co-sputtered films were more resistant to decohesion and buckling after immersion in HF. This inconsistency could be due to an increased resistance of RuAl to attack by HF, and the increased cracking seen in HF etched NiAl in comparison to RuAl would then be chemistry related rather than stress related. Thus, great care must be taken when

204 extrapolating the behavior from one sacrificial release system to another as the behavior may be chemistry-dependent.

Annealed (400ºC-11 h), sputtered Ru/Al multilayers were also exposed to XeF2, and these results are shown in Figure 6 for wavelengths of 9.6 (Figure 6a) and 25 nm (Figure 6b). Approximately half of the beams cracked in the shortest wavelength sample (9.6 nm), but the two larger wavelength samples (15.4 and 25 nm) exhibited no cracking after etching. The amount of deformation of the films decreased with increasing wavelength, indicating either a reduction in the tensile stresses in the films or an increase in the stiffness (the higher wavelength films have more Ru and less RuAl). The 25 nm wavelength Ru/Al shows very little curling around the edges, and would therefore be the better of the three films for MEMS applications. In fact, the deformation exhibited by the 25 nm wavelength Ru/Al was the least of all the fixed-fixed beam samples.

a) b)

25 μm 25 μm

Figure 6 – Annealed Ru/Al multilayer films deposited by sputtering and then exposed to XeF2. Both samples were annealed for 11 h at 400ºC. The images shown are for (a) a 9.6 nm wavelength film and (b) a 25 nm wavelength film.

The stress state noted after XeF2 etching of Ru/Al multilayers (tensile) is opposite to that noted for the HF etching of Ru/Al multilayers (compressive). The films analyzed in the two sets of experiments were from different depositions, and those etched with XeF2 were found to contain oxygen contamination. This contamination was from an insufficient pre-sputter duration and was eliminated in the HF-etched samples by an additional pre-sputter prior to sample deposition. It appears that oxygen incorporation changes the film stresses after annealing from compressive to slightly tensile.

Film stress was calculated for samples under a compressive stress by measuring the amplitude and period of the edge buckling by optical interferometry (representative data is shown in Figure 7). This was done for the 1.5 mTorr as-deposited co-sputtered NiAl and RuAl films and for Ru/Al multilayer films annealed at 200ºC for 4 h. The data were used to calculate the stress in the films, and a value of 1.53 GPa ± 0.04 (compressive) was obtained for as-deposited RuAl co- sputtered films, in comparison to 0.83 GPa ± 0.01 (compressive) for NiAl co-sputtered films.

205

The Ru/Al multilayer annealed at 200ºC for 4 h had a compressive stress of 0.82 ± 0.30 GPa. The larger error for the multilayer sample is due to uncertainty in the elastic modulus used to calculate the stress, as the film is not homogeneous, but a layered structure of Al and Ru. The stress calculated for the annealed Ru/Al multilayer is less than that for the co-sputtered RuAl film, and results in less cracking of fabricated devices (see below). The as-deposited co- sputtered films both contain significant compressive stress, which can likely be reduced by simply increasing the sputtering pressure.

Figure 7 – Surface profile obtained from optical interferometry of an as-deposited RuAl film co-sputtered at 1.5 mTorr and then etched with XeF2. The stress in the film is calculated as ~1.5 GPa compressive.

Free-standing films of RuAl and NiAl have been successfully fabricated with both tensile and compressive stresses. Small compressive stresses are preferable over small tensile stresses, as tensile stresses in a free-standing structure are more likely to cause cracking and failure. This would make the as-deposited NiAl co-sputtered at 1.5 mTorr the best candidate for MEMS applications.

V. ELECTROSTATICALLY AND THERMALLY ACTUATED MEMS

The previous section determined that using HF to etch the sacrificial SiO2 layer, in combination with the film stresses, resulted in cracking, and/or buckling of the NiAl and RuAl films. However, it was concluded that free-standing, fixed-fixed beams that were crack-free and relatively flat could be obtained by using XeF2 to etch the sacrificial Si layer from beneath the intermetallic film. This section uses this latter procedure to fabricate two types of MEMS:

206 electrostatically actuated resonators and bent beam thermal actuators. Devices that survived the etching process without cracking were run at their first out of plane bending (mode I) resonant frequency, or were thermally actuated.

The results of resonator fabrication for the co-sputtered NiAl and multilayer Ni/Al are shown in Figure 8. The best results were obtained for as-deposited, co-sputtered NiAl. Figure 8a shows that some curvature to the free-standing devices was noted for the as-deposited samples, but none of the 18 resonators were cracked after etching. Even with the slight curvature of the resonators shown in Figure 8a, the devices are not touching the substrate, and resonance was achieved and was clearly observed with an optical microscope. The as-deposited NiAl co- sputtered resonators showed mode I resonance between 2 and 84 Hz for the wide beam resonators. Optimizing the deposition, annealing, and etching parameters (samples in many cases were over-etched) should eliminate the curvature noted in the as-deposited samples. Co- sputtered devices annealed at 100ºC for 4 h (Figure 8b) had comb drives that were less curved than the as-deposited resonators, but the comb drives were tilted to one side, which was resting on the substrate. None of the resonators annealed at 100ºC cracked, but annealing at 200ºC for 4 h (Figure 8c) resulted in cracking in 8 out of 18 devices. The stresses created during annealing at 400ºC (Figure 8d) were sufficient to cause the patterned resonators to decohere from the substrate prior to release.

As-Deposited 100 C – 4 h 200 C – 4 h 400 C – 4 h

a) 50 μm b) c) d)

NiAl

prior to etching

e) f) g)

Ni/Al

Figure 8 – Representative light microscope images of co-sputtered NiAl (a-d) and Ni/Al multilayer (e-g) resonators in the as-deposited condition (a) and after annealing for 4 h at 100ºC (b and e), 200ºC (c and f), and 400ºC (d and g). The co-sputtered resonators annealed at 400ºC were not etched as the resonators cracked and peeled after annealing. The scale bar in (a) applies to all of the images.

Resonators were fabricated from Ni/Al multilayers with a wavelength of 30 nm that were annealed at 100ºC, 200ºC, and 400ºC for 4 h (Figure 8e-g). Results from the resonators annealed at 100ºC (Figure 8e) and 200ºC (Figure 8f) were better than those annealed at 400ºC (Figure 8g)

207 as they showed fewer cracked devices, but the results were not as good as the co-sputtered NiAl devices, which had no device failure during etching. The results showed that annealing the Ni/Al multilayer devices at 100ºC for 4 h resulted in 18 of the 18 devices being crack-free, albeit with significant curvature (the majority of the devices resting on the substrate surface). Increasing the annealing temperature to 200ºC resulted in devices with less curvature, but only 10 of the 18 devices survived the etching process without cracking. Although close to half of the devices cracked, the only cracked devices were those with notches. All of the un-notched samples survived the etching process and showed mode I resonance with a large amplitude of motion (resonance from 2 to 76 Hz). From the cracking observed after HF etching it was concluded that the multilayers annealed at 400ºC were in a state of tension, and this resulted in 18 of the 18 resonators cracking in the beam region during etching.

To obtain crack-free, free-standing films, the best condition appears to be the as-deposited co- sputtered state. One possibility to reduce the curvature of the as-deposited devices, and to eliminate the need for post-deposition annealing, is to increase the sputtering pressure slightly to reduce the magnitude of the compressive stresses present in the as-deposited films. Another avenue to investigate is controlling the deposition temperature as the curvature of the 1.5 mTorr as-deposited devices could be modified by slightly changing the maximum temperature experienced during deposition. Examples are shown in Figure 9, where the resonator shown in Figure 9a experienced a higher temperature during deposition than the resonator shown in Figure 9b. The resonators from the deposition that experienced a higher temperature showed decreased curvature, and all curved upwards (away from the substrate) rather than downwards as with the lower temperature deposition. This shows the importance of temperature control during deposition, and may be another approach to control the stresses rather than annealing after deposition.

a) b)

50 μm 50 μm

Figure 9 – Nomarski optical images of NiAl resonators from 2 different depositions conducted at 1.5 mTorr. Both samples are as-deposited, but the film in (a) was deposited from 2 sequential depositions totaling 34 min and (b) was deposited from 3 sequential depositions for a total of 33.5 min. The samples were kept in vacuum for ~30 min in between each sequential deposition.

For the multilayer films, the curvature of the samples annealed at 200ºC was the least observed for the annealed Ni/Al films (and in fact for all of the etched resonators), but the yield was not as good as the as-deposited co-sputtered NiAl. Annealing the multilayers at 150ºC may result in

208 free-standing crack-free devices as the 100ºC anneal resulted in a 100% yield of crack-free devices, which were unfortunately resting on the substrate, whereas the 200ºC anneal resulted in flat, free-standing devices, but with a lower yield of crack-free resonators (56%). The 200ºC resonators showed less curvature than the as-deposited co-sputtered NiAl resonators, indicating lower stresses, yet fewer resonators survived the etching process. The samples annealed at 200ºC are actually a combination of Ni and NiAl3 rather than NiAl, so a possible explanation for the cracking in these resonators is that NiAl3 may not be as strong as NiAl, leading to fracture at lower stress levels. With modifications to the annealing and deposition conditions it should be possible to further reduce the stress to create planar, free-standing NiAl MEMS devices from multilayers or co-deposited material.

Resonators were also fabricated from RuAl formed from reacting multilayers as well as from co- sputtered films. As with the NiAl devices, more work is needed to optimize the processing conditions, but the results obtained so far are promising. Figure 10 shows representative images of devices fabricated from co-sputtered RuAl (Figure 10a-c) and Ru/Al multilayers (Figure 10d- f). Figure 10a shows a resonator fabricated from as-deposited co-sputtered RuAl. The films are under a significant compressive stress (calculated to be ~1.53 GPa), which is indicated by the large amount of wrinkling of the free-standing regions. Even with the high level of stress, a large number of devices survived etching with minimal cracking, suggesting that RuAl may be a good candidate for MEMS applications as the films are able to withstand a large amount of stress. The devices also showed mode I resonance in the frequency range of 20-24 Hz. The co- sputtered films were annealed in an effort to decrease the stress. By annealing the co-sputtered films for 4 h at 400ºC (Figure 10b) prior to etching, we were able to somewhat decrease the stresses in the film. This resulted in a higher yield (14 of 18 survived rather than 11 of 18), and further improvements are expected with further optimization of the processing parameters. Increasing the annealing temperature to 600ºC for 1 min (Figure 10c) resulted in fewer resonators surviving the etching process. An alternative method to decrease the stresses in co- sputtered RuAl may be to increase the sputtering pressure so that a low stress as-deposited film is formed. The current processing conditions were optimized for electrical properties, not residual stress, but the results are promising, as the majority of the films are able to withstand ~1.5 GPa compressive stress without failure.

209

As-Deposited 200 C – 4 h 400 C – 4 h 600 C – 1 min a) b) c)

RuAl

50 μm

d) e) f)

Ru/Al

Figure 10 – Representative images of resonators fabricated from co-sputtered RuAl (a-c) and Ru/Al multilayers (d- f). The resonators are shown in the as-deposited condition (a), after annealing for 4 h at 200ºC (d) and 400ºC (b and e), and after annealing for 1 min at 600ºC (c and f). The scale bar in (a) applies to all of the images.

The resonators shown in Figure 10d-f were fabricated from annealed 25 nm wavelength Ru/Al multilayers using the deposition procedure that minimized the oxygen content. The resonators are likely not single phase RuAl as the anneals were carried out at low temperature (200-600ºC), and TEM has shown a layered structure of Ru and Al for films annealed at 200ºC, and a layered film of Ru and RuAl after annealing at 400ºC. The image in Figure 10d shows that the multilayers annealed at 200ºC are in a state of compression (calculated: 0.82 ± 0.3 GPa, from edge buckling), and all 18 resonators survived etching without cracking. All of the resonators show a significant amount of curvature, with the devices curling up (away from the substrate). This could indicate that a stress gradient is present in these films. When the annealing temperature was increased to 400ºC (Figure 10e) and reaction between the multilayers began, there was a decrease in the amount of compressive stress in the film. However, the majority of the devices are resting on the surface of the substrate again indicating that a stress gradient is present. Increasing the annealing temperature also decreased the number of surviving resonators from 18 of 18 to 14 of 18. A further increase in the annealing temperature to 600ºC (Figure 10f) resulted in only 1 resonator surviving the etching process. By comparing the results in Figure 10d and Figure 10e, it appears that an optimal annealing temperature for 25 nm wavelength Ru/Al multilayers may be found between 200ºC and 400ºC, although the films will not be single phase RuAl. Decreasing the wavelength of the Ru/Al multilayers to create fully reacted RuAl films may also increase the device yield. A summary of the results from all of the resonators fabricated is shown in Table 2.

210

Table 2 - Summary of the results from the ruthenium and nickel aluminide resonator fabrication. All of the anneals were for 4 h except the 600ºC anneal, which was for 1 min. Percent Survival Material Deposition Method Annealing Conditions All Devices Un-Notched Devices NiAl co-sputter as-deposited 100 100 100ºC 100 100 200ºC 56 100 400ºC 0 0 multilayer 100ºC 100 100 200ºC 56 100 400ºC 0 0 RuAl co-sputter as-deposited 61 67 400ºC 78 100 600ºC 61 83 multilayer 200ºC 100 100 400ºC 78 100 600ºC 6 0

Thermal actuators were also fabricated from the intermetallic thin films. The results of a test on an as-deposited co-sputtered NiAl (1.5 mTorr) thermal actuator are shown in Figure 11, which shows a fairly large range of motion as the current is increased. The actuators fabricated were all free-standing and showed minimal curvature.

70 mA 60 mA 50 mA 40 mA 0 mA

18 μm 11 μm 6 μm 2 μm

Figure 11 – Light microscope images showing the range of motion of a thermal actuator as the current is ramped down from 70 mA to 0 mA. The actuator was fabricated from an as-deposited NiAl film that was co-sputtered with an Ar pressure of 1.5 mTorr.

VI. CONCLUSIONS

In this paper we have investigated the use of Ni- and Ru-aluminide films combined with 2 different etchants for the fabrication of MEMS. Using a conventional wet etch for SiO2 sacrificial layers (HF) led to cracking and/or buckling depending on the stress state in the film (tensile/compressive). However, using a gas phase etchant for silicon sacrificial layers, XeF2, free-standing regions could be formed that were crack free. Resonators were fabricated and first

211 out-of-plane bending mode resonance was observed by using XeF2 etching, and the best results were obtained for as-deposited NiAl that was co-sputtered at 1.5 mTorr and was under a compressive stress of ~0.83 GPa. While the devices were not completely flat, they were free- standing, and improvements are expected by decreasing the stress in the co-sputtered films by increasing the sputtering pressure. Ni/Al multilayers annealed at 200ºC for 4 h (Ni-NiAl3) produced devices with the least amount of curvature, but only 56% of the devices survived etching (100% of the un-notched), in comparison to 100% of the as-deposited co-sputtered NiAl. The results on multilayer Ni/Al should be improved by using smaller wavelength Ni/Al multilayers to form NiAl at temperatures lower than 400ºC to decrease the tensile stresses (~0.733 GPa, estimated due to CTE mismatch) that occurred during cooling and lead to 100% device failure. The results for RuAl resonators are also promising as the films can withstand high compressive stresses (~1.5 GPa, calculated from edge buckling), and improved performance from co-sputtered RuAl is expected by increasing the sputtering pressure in order to decrease the film stress. Thermal actuators fabricated from as-deposited co-sputtered NiAl showed ~18 m of motion with an applied current of 70 mA, and very little curvature/buckling was noted. The effect of the high compressive stress is less pronounced in the doubly supported actuators than in the singly supported resonators. Improvements in processing should lead to improvements in the device yield and curvature, but the results presented plus the high strength, high electrical conductivity, good oxidation/corrosion resistance, and moderate toughness make NiAl and RuAl suitable materials for MEMS.

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CHAPTER 7

CONCLUSIONS AND SUGGESTIONS FOR FUTURE WORK

217

I. SUMMARY

In this study B2 aluminide films (NiAl and RuAl) have been successfully deposited and characterized. Films were fabricated by the annealing of multilayers and co-sputtering, and NiAl films were found to have a lower resistivity than RuAl films. NiAl films fabricated from annealed multilayers were fully reacted at lower temperatures than RuAl films. NiAl multilayers with a wavelength of 30 nm and below were fully reacted to form NiAl after annealing for 2 h at 400°C. These films had a resistivity of 15.5-26.7 μΩcm (wavelengths from 15.4-30 nm) after a 4 h anneal at 400°C, and the lower values of resistivity correspond to films with larger wavelengths. Increasing the annealing temperature to above 400°C resulted in a decreased resistivity (lowest: 11.0 ± 0.1 μΩcm after 20 min at 800°C), but reactions with the substrate were observed by TEM after annealing at 1000°C. A reaction with the substrate likely occurred at annealing temperatures as low as 600°C and was noted as an increase in resistivity after sufficiently long annealing times. The time after which an increase in resistivity occurred was ~1-2 min for anneals at 600°C, ~7-12 min at 800°C, and ~0.5-1 min at 1000°C. For RuAl films fabricated from annealed multilayers, shorter wavelength films (6.7 and 9.6 nm) were likely fully reacted after annealing at 400ºC for 40 h with a resistivity of 121 ± 0 cm and 70.5 ± 0.1 cm, respectively. Longer wavelength films (15.4 and 25.0 nm) were not fully reacted at 400ºC, and consisted of Ru and RuAl. These films are likely single phase after annealing at 500ºC and 900-1000ºC for the 15.4 nm and 25 nm wavelength samples, respectively. The lowest resistivity in RuAl films was somewhat higher than in similar NiAl films (19.3 ± 0.1 cm in comparison to 11.0 ± 0.1 cm), and required a much higher annealing temperature (1100°C for RuAl and 800°C for NiAl). RuAl also reacted with the substrate after annealing at 1000°C, but no increase in resistivity was observed due to this reaction. Both NiAl and RuAl films obtained the lowest resistivity from the largest wavelength studied for each material.

As-deposited co-sputtered films of NiAl and RuAl were crystalline as-deposited, with small columnar grains. Co-sputtering at 7 mTorr led to NiAl films with tensile stresses, a structure with small voids, and a high resistivity. Decreasing the sputtering pressure to 1.5 mTorr resulted in films with compressive stresses, a dense microstructure, and a lower resistivity. The film as- deposited at 1.5 mTorr had a resistivity of 45.5 ± 1.5 μΩcm, which was not much higher than the film deposited at 7 mTorr and then annealed at 1000°C for 30 s. After annealing the film deposited at 1.5 mTorr at 1000°C, the resistivity was 18.7 ± 0.0 μΩcm, which is higher than the annealed multilayer films. RuAl films co-sputtered at 1.5 mTorr Ar pressure have an as- deposited resistivity of 157 ± 4 cm, which decreased to a minimum of 19.6 ± 0.1 cm after annealing at 1100ºC for 1 min. This value is very similar to the annealed Ru/Al multilayers. Co- sputtering is an attractive method to fabricate these intermetallic films as no annealing is required to acquire single phase crystalline films. However, annealing is required in order to decrease the resistivity of the films, and RuAl requires higher temperature anneals due to the higher melting temperature. After annealing, NiAl co-sputtered films have a slightly lower resistivity than RuAl co-sputtered films. However, the best results in terms of resistivity were for annealed Ni/Al multilayers, which obtained a minimum resistivity similar to the value for bulk NiAl.

The oxidation resistance of the NiAl and RuAl films fabricated by co-sputtering and the annealing of multilayers was studied by annealing the films in flowing oxidation for 1 h at various temperatures. The intermetallic thin films were compared to Al, Ni, and Ru thin films,

218 and were found to be superior to the pure metal films in terms of maintaining a smooth surface before and after oxidation and the electrical resistance after oxidation. The intermetallic films showed no observable surface changes (light microscopy) up to 500ºC in flowing oxygen and were conductive to higher oxidation temperatures (800°C for RuAl, 850ºC for NiAl) than Ni (500ºC), Al (600ºC), and Ru (800ºC), although vaporization may have begun at ~700ºC in the case of the Ru film. After oxidation, the RuAl film still had a very reflective surface, but a distinct haziness was noted on all of the Ni-Al (NiAl and Ni2Al3) films. Both intermetallics after being oxidized to 850ºC had surfaces composed solely of aluminum and oxygen, which was most likely Al2O3. Below this Al-O layer, RuAl had an oxide scale composed of alternating Ru- rich and alumina layers, with no evidence of a reaction with the substrate or any remaining RuAl. The oxide on NiAl had a more complex and heterogeneous structure due to reactions with the substrate. Although some surface roughening was observed in the RuAl film, the surface was still relatively smooth, reflective, and homogeneous in comparison to NiAl, and showed little change in surface morphology from the onset of oxidation (500ºC) up to 800ºC. In contrast, NiAl showed the appearance of discrete surface oxides at ~550ºC, with a more uniform oxide forming between these particles, which began to noticeably thicken at 850ºC. In terms of the surface quality RuAl outperformed NiAl, whereas NiAl had a lower resistivity than RuAl after oxidation.

The intermetallics NiAl and RuAl along with Ru and Au have been patterned into thin line structures and then tested using an alternating current to induce thermomechanical fatigue. RuAl samples were able to withstand higher cyclic values of ΔT than NiAl for comparable times to failure. Both NiAl and RuAl were able to withstand higher values of ΔT than gold. The ΔT for tests on NiAl ranged from ~300-520°C (Tmax: 400-600°C), with a time to failure of 100’s of hours when ΔT was near 300°C (Tmax: ~400°C). Samples of NiAl that had a lower resistivity were able to withstand higher current densities due to reduced Joule heating. RuAl had a longer time to failure than NiAl at high testing temperatures (ΔT > ~350°C), but trends in the data indicate that the time to failure at lower temperature may be higher for NiAl. NiAl was able to withstand a slightly higher current density than RuAl due to the lower resistivity of the NiAl samples in comparison to the RuAl samples. If the resistivity of the intermetallic samples can be reduced, then increases in the current density are expected.

A method to monitor the accumulation of damage and to estimate the time to failure during AC tests has been developed, and only involves monitoring the change in resistance during the test. Curves of the resistance as a function of time were plotted for all of the samples during the course of the AC tests, and three different trends were obtained, with the majority of the samples displaying very similar R(t) curves: an initial region of increasing R with a decreasing slope, a second region of increasing R with an almost constant slope, and a third region of increasing R with an increasing slope. These three regions have been ascribed to heating, the production of defects (dislocations and vacancies), and crack growth (and possibly voids). It was determined that by measuring the slope of the R(t) curve at the beginning of the test, the time to failure and the length of time spent in each of the three regions could be estimated. This will enable future tests to be paused at specific locations along the R(t) curves so that damage formation and crack growth may be studied.

219

We have also investigated the use of Ni- and Ru-aluminide films combined with 2 different etchants for the fabrication of MEMS. Using a conventional wet etch to remove the SiO2 sacrificial layer (HF) led to cracking and/or buckling of the intermetallic films depending on the stress state in the film (tensile/compressive). However, when a gas phase etchant was used for etching a silicon sacrificial layer, XeF2, free-standing regions could be formed that were crack free. Resonators were fabricated using XeF2 as an etchant, and first out-of-plane bending mode resonance was observed for RuAl and NiAl devices. The best results were obtained for as- deposited NiAl that was co-sputtered at 1.5 mTorr, which resulted in a compressive stress of ~0.83 GPa. While the devices were not completely flat, they were free-standing, and improvements can be expected by decreasing the stress in the co-sputtered films by increasing the sputtering pressure. Ni/Al multilayers annealed at 200ºC for 4 h (Ni-NiAl3) produced devices with the least curvature, but only 56% of the devices survived etching (100% of the un- notched), in comparison to 100% of the as-deposited co-sputtered NiAl. The results on multilayer Ni/Al should be improved by using smaller wavelength Ni/Al multilayers to form NiAl at temperatures lower than 400ºC to decrease the tensile stresses (~0.733 GPa) that occurred during cooling from 400°C and led to 100% device failure. The results for RuAl resonators are also promising as the films can withstand high compressive stresses (~1.5 GPa), and improved performance from co-sputtered RuAl is expected by increasing the sputtering pressure in order to decrease the film stress. Thermal actuators fabricated from as-deposited co- sputtered NiAl showed ~18 μm of motion with an applied current of 70 mA, and very little curvature/buckling was noted. The effect of the high compressive stress is less pronounced in the doubly supported actuators than in the singly supported resonators. Improvements in processing should lead to improvements in the device yield and curvature, but the results presented plus the high strength, high electrical conductivity, good oxidation/corrosion resistance, and moderate toughness make NiAl and RuAl interesting materials for further development of MEMS.

The results presented on NiAl and RuAl thin films show that these intermetallics are promising materials for thin film applications. Both NiAl and RuAl thin films can be fabricated at low temperature, and the films can withstand a large amount of stress without cracking. The thin films have a low resistivity, and can be heated to 500°C in oxygen (for at least 1 h) without any changes in surface morphology. NiAl and RuAl can withstand 100’s of hours in an environment where the temperature is fluctuating by as much as 300°C (200 Hz heating frequency on Si substrates). Finally, it has been shown that free-standing structures of NiAl and RuAl can be fabricated for applications such as MEMS.

220

VITA

Jane A. Howell

Education

Ph.D., Materials Science and Engineering (December 2010) The Pennsylvania State University, University Park, PA 16802 THESIS: “Deposition, Characterization, and Thermomechanical Fatigue of Nickel Aluminide and Ruthenium Aluminide Thin Films” ADVISORS: Suzanne Mohney, Chris Muhlstein

Master of Science, Materials Science and Engineering (May 2005) The Pennsylvania State University, University Park, PA 16802 THESIS: “Surface Properties of Glass Evaluated Using Nanoindentation” ADVISORS: John Hellmann, Chris Muhlstein

Bachelor of Science, Materials Science and Mechanics (May 2001) Michigan State University, East Lansing, MI

Publications

J.V. Ryan, P. Columbo, J.A. Howell, and C.G. Pantano. “Tribology-structure relationships in silicon oxycarbide thin films.” International Journal of Applied Ceramic Technology. (2009)

J.A. Howell, J.R. Hellmann, and C.L. Muhlstein. "Nanomechanical properties of commercial float glass." Journal of Non-Crystalline Solids 354(17): 1891-1899. (2008)

J.A. Howell, J.R. Hellmann, and C.L. Muhlstein. "Correlations between free volume and pile-up behavior in nanoindentation reference glasses." Materials Letters 62(14): 2144-2146. (2008)

J.A. Howell, A. Telang, J.G. Lee, S. Choi, and K.S. Subramanian. "Surface damage accumulation in Sn-Ag solder joints under large reversed strains." Journal of Materials Science: Materials in Electronics 13(6): 335-344. (2002)