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Alternative nitride diffusion barriers on silicon and germanium for copper metalization

University of Florida Institutional Repository
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ALTERNATIVE NITRID E DIFFUSION BARRIERS ON SILICON AND GERMANIUM FOR COPPER METALLIZATION By SEEMANT RAWAL A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVERSITY OF FLOR IDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 2006

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Copyright 2006 by Seemant Rawal

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Dedicated to my mother Rasilaben Ra wal and my dear wife Purvi Rawal.

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iv ACKNOWLEDGMENTS I would like to immensely thank my advi sor Dr. David P. Norton. His enthusiasm and knowledge on different subjects of scie nce are astounding. Dr. Norton’s exemplary dedication towards his work and students along with his zeal to be the best in the world in his area of research motivated me greatly to pursue and conduct supe rior quality research. I would like to thank Prof. Rajiv K. Singh for providing me guidance and support during my initial years in USA. I would like to thank Prof. Tim Anderson and Prof. Lisa McElwee-White for their guidance and help with my research. They have helped me to think and analyze my research critically. I would like to thank my mother, Rasilaben Rawal, for her passion towards education. Her undying zeal towards learning sowe d the first seeds in me to respect and acquire knowledge. Many thanks go to my hi gh school chemistry teacher, Mr. Dubey and physics teacher Mr. Parmar for providing support throught my 12th standard. Special thanks go to Mahendra Pandya and Geetaben Pandya, for supporting me throughout my undergraduate years. I would like to thank my brother, Rakesh Rawal, and his wife, Padmaja Rawal, for their love and support. Mo st importantly, I would like to thank my wife, Purvi Rawal, for her unconditional l ove and support. Her immense understanding and perceptive nature have made my jour ney cherishing. I would like to thank all my friends and well-wishers for th eir support and encouragement. Last but not least, I thank GOD for giving me this lif e and opportunity to serve HIS world.

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v TABLE OF CONTENTS page ACKNOWLEDGMENTS.................................................................................................iv TABLES......................................................................................................................... ..vii LIST OF FIGURES.........................................................................................................viii ABSTRACT....................................................................................................................... xi 1 INTRODUCTION........................................................................................................1 2 LITERATURE REVIEW.............................................................................................7 Refractory Metals as Diffusion Barrier........................................................................7 Ti Diffusion Barrier...............................................................................................7 Ta Diffusion Barrier..............................................................................................8 Cr Diffusion Barrier..............................................................................................8 W Diffusion Barrier...............................................................................................8 Binary Diffusion Barrier...............................................................................................9 Refractory Intermetallics.......................................................................................9 Refractory Carbides.............................................................................................10 Refractory Nitrides..............................................................................................12 Ternary Diffusion Barriers.........................................................................................15 3 EXPERIMENTAL DETAILS AND CHARACTERIZATION.................................18 Sputtering....................................................................................................................1 8 DC Magnetron Sputtering...................................................................................19 RF Magnetron Sputtering....................................................................................20 Characterization..........................................................................................................21 X-ray Diffraction.................................................................................................21 Auger Electron Spectroscopy..............................................................................22 X-ray Photoelectron Spectroscopy......................................................................23 Scanning Electron Microscopy and En ergy Dispersive Spectroscopy................24 Atomic Force Microscopy...................................................................................26 Transmission Electron Microscopy.....................................................................26 Focused Ion Beam...............................................................................................28 Van der Pauw Measurement and Four Point Probe.............................................28

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vi 4 PROPERTIES OF W-Ge-N AS A DIFFUSION BARRIER MATERIAL FOR COPPER.....................................................................................................................36 Introduction.................................................................................................................36 Experimental Details..................................................................................................37 Results and Discussion...............................................................................................38 Conclusion..................................................................................................................41 5 INVESTIGATION OF W-Ge-N DEPO SITED ON Ge AS A DIFFUSION BARRIER FOR Cu METALLIZATION...................................................................46 Introduction.................................................................................................................46 Experimental Details..................................................................................................47 Results and Discussion...............................................................................................48 Conclusion..................................................................................................................51 6 COMPARATIVE STUDY OF HfNX AND Hf-Ge-N DI FFUSION BARRIERS ON Ge.........................................................................................................................5 9 Introduction.................................................................................................................59 Experimental Details..................................................................................................60 Results and Discussion...............................................................................................62 Conclusion..................................................................................................................66 7 EFFECT OF Ge OVERLAYER ON TH E DIFFUSION BARRI ER PROPERTIES OF HfNX.....................................................................................................................75 Introduction.................................................................................................................75 Experimental Details..................................................................................................76 Results and Discussion...............................................................................................78 Conclusion..................................................................................................................80 8 PROPERTIES OF Ta-Ge-N AS A DIFFUSION BARRIER FOR Cu ON Si...........87 Introduction.................................................................................................................87 Experimental Details..................................................................................................88 Results and Discussion...............................................................................................89 Conclusion..................................................................................................................92 9 CONCLUSIONS........................................................................................................98 LIST OF REFERENCES.................................................................................................102 BIOGRAPHICAL SKETCH...........................................................................................110

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vii TABLES Table page 3-1 Sputtering yields of different ions............................................................................30

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viii LIST OF FIGURES Figure page 1-1 RC delay vs. technology nodes..................................................................................6 1-2 Void and Hillock formation in Al interconnects........................................................6 3-1 Schematic diagram of DC sputtering system with parallel plate discharge.............31 3-2 Schematic diagram of RF sputtering system with a capacitive, parallel plate discharge...................................................................................................................31 3-3 Schematic diagram of X-ray diffraction set-up........................................................32 3-4 Schematic set-up principle of an atomic force microscope......................................32 3-5 Interaction output between a high-en ergy electron beam and thin specimen..........33 3-6 Sample geometries for Van der pauw measurements..............................................33 3-7 Schematic diagram of Van der pauw configuration for measurement of RA and RB resistances, respectively......................................................................................34 3-8 Schematic diagram of four point probe measurement..............................................35 4-1 X-ray diffraction patterns of as-depos ited and annealed films at different temperature of (a) WNx films (b) W-Ge-N films.....................................................42 4-2 AES depth profiles of Cu/WNx/SiO2/Si (a) as-deposited (b ) annealed at 600 C/1 hr (c) annealed at 800 C/1 hr..................................................................................43 4-3 AES depth profiles of Cu/W-Ge-N/SiO2/S i (a) as-deposited (b) annealed at 600 C/1 hr (c) annealed at 800 C/1 hr..........................................................................44 4-4 Resistivity vs. annealing te mperature (a) for W-Ge-N and WNx films. Also shown (b) is the resistivity as a function of sputter target power for the Ge and W targets..................................................................................................................45 5-1 X-ray diffraction patterns of as-depos ited and annealed films at different temperature of (a) WNx films (b) W-Ge-N films.....................................................53

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ix 5-2 AES depth profiles of Cu/WNx/Ge (a) as-deposited (b) annealed at 400 C/1 hr (c) annealed at 500 C/1 hr.......................................................................................54 5-3 AES depth profiles of Cu/W-Ge-N/Ge (a) as-deposited (b) annealed at 400 C/1 hr (c) annealed at 500 C/1 hr..................................................................................55 5-4 EDS depth profile of Cu/WNx/Ge annealed at 500 C/1 hr.....................................56 5-5 EDS depth profile of Cu /W-Ge-N/Ge annealed at 500 C/1 hr...............................57 5-6 X-TEM images of (a) Cu/WNx/Ge annealed at 500 C/1 hr (b) Cu/W-Ge-N/Ge annealed at 500 C/1 hr............................................................................................58 6-1 X-ray diffraction patterns of Cu/HfNx/Ge films in as-deposited and annealed conditions for varying thickness of (a) 300 nm, (b) 50 nm, and (c) 15 nm.............68 6-2 X-ray diffraction patterns of Cu/Hf-Ge-N /Ge films in as-deposited and annealed conditions for varying thickness of (a) 300 nm, (b) 50 nm, and (c) 15 nm.............69 6-3 FE-SEM images of 50 nm films annealed at 500 C for 1 hr: (a) Cu/HfNx/Ge and (b) Cu/Hf-Ge-N/Ge...........................................................................................70 6-4 FE-SEM images of 50 nm films annealed at 600 C for 1 hr: (a) Cu/HfNx/Ge and (b) Cu/Hf-Ge-N/Ge...........................................................................................70 6-5 X-TEM images of 50 nm films annealed at 600 C for 1 hr: (a) Cu/HfNx/Ge and (b) Cu/Hf-Ge-N/Ge..................................................................................................71 6-6 EDS depth profile of 50 nm thick Cu/HfNx/Ge annealed at 600 C for 1 hr...........72 6-7 EDS depth profile of 50 nm thick Cu/Hf-Ge-N/Ge annealed at 600 C for 1 hr.....73 6-8 XPS chemical state data of Cu 2p pe ak at various sputte ring times for 50 nm thick film of Cu/Hf-Ge-N/Ge (a) as deposited and (b) annealed at 600 C for 1 hr.74 7-1 X-ray diffraction patterns of as-dep osited films and annealed ones at the temperature shown in the figure...............................................................................82 7-2 EDS depth profile of Cu/Ge(25nm)/HfNx(7nm)/Si for a) as-deposited: b) annealed at 500 C for 1 hr: and c) annealed at 500 C for 3 h................................84 7-3 AFM images of Cu/Ge(25nm)/HfNx(7nm)/Si for a) as-deposited: b) annealed at 500 C for 1 hr: and c) annealed at 500 C for 3 hr.................................................85 7-4 HRTEM images of Cu/Ge(25nm)/HfNx(7nm)/Si for a) as-deposited: b) annealed at 500 C for 1 hr: and c) annealed at 500 C for 3 hr.............................................86

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x 8-1 X-ray diffraction patterns of as-deposite d and annealed at the temperature shown in figure of a) Cu/TaN/Si (001) for 1 hr: b) Cu/Ta-Ge-N/Si (001) for 1 hr.............93 8-2 Field-emission SEM images after annealing at 400 oC for 1 hr for a) Cu/TaN/Si (001) and b) Cu/Ta-Ge-N/Si (001)...........................................................................94 8-3 AES depth profile of Cu/TaN/Si (001) fo r a) as-deposited and b) annealed at 400 C for 1 hr.................................................................................................................95 8-4 AES depth profile of Cu/Ta-Ge-N/Si ( 001) for a) as-deposited and b) annealed at 400 C for 1 hr......................................................................................................96 8-5 Sheet resistance of Cu vs. anneal ing temperature for TaN and Ta-Ge-N................97

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xi Abstract of Dissertation Pres ented to the Graduate School of the University of Florida in Partial Fulfillment of the Requirements for the Degree of Doctor of Philosophy ALTERNATIVE NITRID E DIFFUSION BARRIERS ON SILICON AND GERMANIUM FOR COPPER METALLIZATION By Seemant Rawal December 2006 Chair: David P. Norton Major Department: Materials Science and Engineering As device dimensions shrink, there is an ur gent need to replac e conventionally used Al interconnects to achieve increased current density requirements and better performance at future technology node. Copper is a viable candidate due to its lower resistivity and higher resistan ce to electromigration. Howeve r, Cu has its own problems in its integration. It di ffuses rapidly through SiO2, Si and Ge into the active regions of the device, thereby deteriorating device performance. There is therefore a need to find a diffusion barrier for Cu. Amorphous ternary nitrides are investigated as a candidate diffusion barrier for Cu metallization on single crystal Si and Ge substrates. Ge was chosen as a third element in the binary matrix of WNx, HfNx and TaN due to its chemical similarity to Si and possible integration in SiGe or Ge based devices. The addition of Ge helps in amorphization of the binary matrix. It hinders grain boundary form ation and increases the recrystallization temperature compared to their respectiv e binary nitrides. Cu was deposited in-situ after

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xii nitride deposition. In the case of W-Ge-N on Si, the recrys tallization temperature was raised by 400 oC, while for Ta-Ge-N it was raised by 100 oC indicating better diffusion barrier properties than th eir corresponding binaries. A bilayer approach for diffusion barrier was also studied. Ge/HfNx bilayer was deposited on Si followed by in-situ depositi on of Cu. Copper reacts with Ge and forms Cu3Ge which has low resistivity, is less reacti ve with oxygen, and is a diffusion barrier for Cu. By combining the material properties of Cu3Ge and HfNx, an excellent diffusion barrier performance was demonstrat ed under stringent test conditions.

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1 CHAPTER 1 INTRODUCTION The integrated circuit industry doubles the number of transistors in a chip every 18 months. With more transistors built into a chip, minimum feature size decreases, leading to faster devices. At present, we are moving towa rds the 45 nm node technology and beyond which is guided by Internationa l Technology Roadmap for Semiconductors (2005)2. For many years, gate length determined the device speed and was the bottleneck for future generation devices. However, with decreasing dimensions, delay in interconnects now plays a major role in infl uencing device speed. Th e interconnect delay, also called resistance-capacitance (RC) tim e constant delay, is mostly caused by a increase in resistance at smaller dimensions The RC delay is expressed by the following equation: ILD M ILDt t RC *2 where is resistivity, is interconnection length, Mtis interconnect thickness, ILD and ILDt are permittivity and thickness of interlevel dielectric (ILD). Also, the current density increases as width of the interconnect decr eases. Figure 1-1 shows the increase in RC delay time in interconnects with decreasing f eature size which domina tes overall delay at sub micron technology nodes.1 Aluminum has been the interconnect materi al of choice for many generations with SiO2 ( = 3.9) as the dielectric material. This results in a tremendously high RC delay as dimensions decrease. One way to reduce the RC delay is to decrease the dielectric

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2 constant of SiO2. Fluorinated silicon dioxide (FSG) was used with a dielectric constant of 3.7 for the 180 nm technology node.2 Not until recently was low-k material ( = 2.7 – 3.0) being integrated in devices. At high curre nt densities, electromigration performance of Al is severely degraded leading to voids and hillock fo rmation. Creation of voids leads to discontinuous circu it causing failure of devices. Fi gure 1-2 shows void and hillock formation in Al interconnect. Alloying Al with Cu showed an increase in electromigration resistance, but it reached its usability limits as subsequent milestones in IC industry were achieved. With increased demands on performance, an alternative interconnect material is eventually desired. The logical choice is a transition to Cu interconnects. There are several advantages in using Cu as an interconnect material. Bulk resistivity of Cu is 1.67 -cm, which is almost 40% less than Al ( =2.65 -cm).3 Resistance to electromigration property of Cu is highly s uperior to Al, which leads to better device reliability. Replacing the Al/SiO2 gate stack with Cu/lowreduces the RC delay significantly. However, Cu has its own share of problems. It diffuses rapidly through Si, Ge and SiO2 and forms parasitic silicides (Cu3Si) and hinders the device performance. It forms deep and shallow level traps in Si and Ge, respectively. Cu reacts with dopants and forms Cu-D complexes (D is dopant atom) deterior ating device performance. Cu exhibits poor adhesion on SiO2 and other lowdielectrics. In addition to the above difficulties, solutions also need to be found for anis otropic etching, poor corrosion and oxidation properties of Cu. In order to get the desi red benefits by switching from Al to Cu interconnect, a viable diffusion barrier is required to prevent Cu diffusion.

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3 The primary purpose of a diffusion barrier is to prevent intermixing of chemical species with each other. This can be achieve d via various mechanisms thereby classifying the diffusion barriers based on the method. i.e., passive, sacrificial, stuffed and amorphous diffusion barriers. A pa ssive diffusion barrier is an ideal barrier which does not react with any of the layers it separates. A sacrificial barrier would react with either or both layers that it separate s and get consumed. For a sacr ificial barrier, the reaction rate between the barrier and the layers is very important. It should be slow enough so that the barrier can perform to acceptable levels for the useful lifetime of the device. For polycrystalline or nanocrystal line thin films, rapid diffusion occurs via grain boundaries, dislocations and surface defects.4 Diffusion is highly influenced by temperature. The temperature dependence of diffusion coefficient (D) is given by the following equation: kT Q D Dd oexp whereoD is temperature-independe nt pre-exponential factor, dQ is the activation energy, k is the Boltzmann constant and T is the temperature. The activation energy for grain boundary diffusion is approximately half the activation energy required fo r lattice diffusion. As a result, grain boundary diffusion dominates at lower temperatures.4 For polycrystalline thin fi lms, diffusion through grain boundaries and dislocations is the fastest. This mode of diffusion can be stopped using two approaches. The grain boundaries and/or dislocations can be “stuffed” by impurity atoms which would hinder diffusion through th em. Another approach is to lower the amount or eliminate the grain boundaries al together. This can be achieved by low temperature or room temperature deposition of nanocrystalline or a amorphous diffusion barrier where the short-circu it pathways are lowered or eliminated. This is highly

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4 desirable because of low/no addition to the thermal budget. This may not be the best approach as films deposited at lower temp eratures can have undesired modified properties during subsequent high temperature processing steps. Apart from the primary function of pr eventing diffusion of Cu, the diffusion barrier should meet other stringent specifications. Some of them are: The diffusion barrier should be thermodynamically stable with Cu and underlying substrate under standard operating conditions. It should not react with Cu or the substrate under thermal, mechanical or electrical stresses encountered during other processing steps. The density of diffusion barrier should be close to its bulk density in order to avoid any defects, voids or disl ocations which can compromise its integrity. The diffusion barrier should have reasonable thermal and electrical conductivity to avoid an y parasitic capacitan ce effects and unwanted heating effects. The contact resistance of diffusion barri er with Cu and substrate should be minimal. Diffusion barrier should behave well under the applied mechanical and electrical stresses durin g subsequent processing. Microstructure of the diffusion barri er should ideally be amorphous at room temperature and remain amorphous after thermal treatment at higher temperatures. In general, diffu sion barriers with higher melting temperatures have high recr ystallization temperatures. The barrier should have good conformality especially over high aspect ratio structures. Diffusion barrier should adhere we ll with Cu and the surrounding material. Deposition of diffusion barrier should be compatible with existing processing facilities and infrastructures.

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5 Chapter 2 will present background and a lite rature review of various diffusion barriers, method of deposition and barrier prop erties. In Chapter 3, experimental methods used to deposit diffusion barrier films for this study will be explained. Various quantitative characterization techniques used to evaluate the diffusion barrier performance under extreme thermal treatment wi ll also be described. Recrystallization properties and diffusion ba rrier performance of WGe-N deposited on Si/SiO2 are evaluated in Chapter 4. The behavior and performance of W-Ge-N diffusion barrier deposited on Ge are investigated in Chap ter 5. A comparison of barrier properties and performance of HfNx and Hf-Ge-N deposited on Ge is described in Chapter 6. Chapter 7 will describe the excellent diffusion barrier properties of a Ge/HfNx bi-layer diffusion barrier deposited on Si. Propertie s of Ta-Ge-N diffusion barrier on Si will be evaluated in Chapter 8 followed by conclusion in Chapter 9.

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6 Figure 1-1. RC delay vs. technology nodes. Figure 1-2. Void and Hillock formation in Al interconnects.

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7 CHAPTER 2 LITERATURE REVIEW Due to the unique diffusion property of Cu in Si, Ge, SiO2 and other lowdielectrics, there is an immi nent need for a diffusion barrier for Cu. There has been a plethora of research on differe nt diffusion barriers. In genera l, metals or compounds with high melting temperatures are suitable for the purpose because of less chance of having grain boundaries, which are the fastest diffu sion pathways. Therefore, refractory based materials seem to be the best choice due to their high melting temperature and decent electrical conductivity. A considerable amount of research has been conducted to find a viable diffusion barrier for Cu based on refract ory metals. Most of the research can be classified in three categories. Refractory metals as diffusion barriers Binary diffusion barriers based on refractory metals Ternary diffusion barriers based on refractory metals This chapter will cover a comprehensive review of the above classifications. Refractory Metals as Diffusion Barrier As it became evident that switching from Al based interconnects to Cu based interconnects was necessary, initial research focused on refractory metals as candidate diffusion barriers. Primary interest was in Ti Ta, Cr, and W as Cu diffusion barriers. Ti Diffusion Barrier Ti based diffusion barriers were extensivel y used for Al interconnects. A natural choice was to extend its functi onality in acting as a Cu di ffusion barrier. However, a Ti diffusion barrier fails at 350 oC by Ti-Cu compound formation.5 A study conducted by

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8 Ohta et al.6 also revealed failure of Ti diffusion barrier at 400 oC indicated by an increase in resistivity after annealing. The resistivit y increase was attributed to Cu diffusion through Ti and subsequent reaction w ith Si to form Cu silicides. Ta Diffusion Barrier The two major advantages of using Ta as a diffusion barrier in comparison to Ti are its very high melting temperature (3017 oC compared to 1668 oC for Ti) and thermodynamical (interface and solubility) stability with Cu for very high temperatures. The Cu-Ta phase diagram shows that both Cu an d Ta are insoluble in each other even at high temperature. Failure of Ta diffusion barrier occurs due to Cu diffusion through grain boundaries formed in Ta after high temperature annealing followed by Cu3Si formation at Ta/Si interface.7 Ta also reacts with Si and forms silicides thus rendering itself unusable as a Cu diffusion barrier. Cr Diffusion Barrier Cr was studied as a diffusion barrier for Cu because of its good corrosion resistance properties and ex cellent adhesion promoter.8 However, Cu/Cr/Si multilayer structure showed a huge increase in resistivity afte r annealing at temperatures higher than 400 oC. This was attributed to Cu-Cr binary phase formation. The formation of such a phase compromises the integrity of the Cu/C r/Si multilayer structure resulting in Cu silicide formation indicati ng diffusion barrier failure.6 W Diffusion Barrier W is a better diffusion barrier compared to Ti and Cr. It is chemically and thermodynamically stable with Cu.9 Initial study showed that Cu/W/Si structures behaved well even after annealing at 600 oC. However, it failed after annealing at 700 oC.6 Cu intermixes with W even at low temperatures of 260 oC, thus rendering it non-viable as a

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9 solution to Cu diffusion.10 Preferentially oriented W(110) failed after an nealing at 690 oC for 1 hr due to consumption of W by a uniform silicidation reaction.11 Selective chemical vapor deposited W layer on p+-n junction diodes led to fa ilure after annealing at temperatures above 650 oC.12 Binary Diffusion Barrier In order to overcome the drawbacks of lower recrystallization temperatures and formation of grain boundaries in refractory metals used as diffusion barriers, their respective binaries were explored. This cl ass of diffusion barriers can generally be classified in three categories. Refractory intermetallics Refractory carbides Refractory nitrides These individual classifications will be e xplored further in the following sections. Refractory Intermetallics Several of the refractory intermetallics were studied as Cu diffusion barriers. TiW was used extensively in Al-based interconne cts and so was also tried as a diffusion barrier for Cu interconnects.13 The study showed that TiW (30 at. % Ti) not exposed to air before deposition of Cu failed at 775 oC when RTA annealed in N2 gas for 30 seconds. The authors also summarized results of a va riety of Cu/diffusion barrier/Si systems. Table-1 in reference 13 shows a list of diffusion barrier inve stigated between Si and Cu.13 Amorphous Ni60Nb40 diffusion barrier failed at 600 oC after annealing for 60 min. A different composition of Ni57Nb42 failed at a lower temperature of 500 oC after 60 min.14 The authors also investigat ed the feasibility of Ni60Mo40 as a candidate diffusion barrier for copper. It failed at 500 oC after annealing for 60 min. Amorphous Ir-Ta alloy was investigated for Cu diffu sion barrier application.15 Sandwich structure of Si/a-Ir45Ta55/Cu/

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10 a-Ir45Ta55 failed after annealing at temperatures higher than 700 oC. Failure occurred at 750 oC by interdiffusion of Cu and Si as detected by Rutherford backscattering spectrometry. It was also observed that th e recrystallization temperature of a-Ir45Ta55 was lowered from 900 oC to 750 oC due to the presence of Cu. Alloyed Ta-Co16 was shown to act as a Cu diffusion barrier deposited by e-b eam evaporation. However, an intermetallic phase of Co2Ta and metal-silicide phase of Co2Si formed after annealing at 500 oC. Recently, a study conducted by J.S. Fang et al.17 explored the possibility of using sputtered Ta-TM (TM = Fe, Co) as diffusion barriers for Cu. Ta0.5Fe0.5 diffusion barriers failed at 650 oC while Ta0.5Co0.5 failed at 700 oC after annealing, indicating better performance of Ta0.5Co0.5. Refractory Carbides Interest in refractory carbides to se rve the purpose of copper diffusion barriers was due to two reasons. First, they have a high melting temperature, and second, they have low resistivity indica ting higher thermal and chemical stability. Among the refractory carbides studied, primary interest has been in Ti, Ta and W based carbides. A detailed study of TiCx as copper diffusion barriers was carried out by S.J. Wang et al.18 However, the Cu/TiCx/Si structure was unstable after a nnealing at temperatures higher than 600 oC for 30 min. A more sensitive electric al characterization done by measuring leakage current across Cu/TiCx/p+ n-Si diode structure revealed an early failure just after annealing at 500 oC for 30 min. As many as 60% of the diodes measured for leakage current after annealing at 530 oC had the leakage current in the range of 10-5 A/cm2, which is three orders of ma gnitude higher than the leakag e current measured at room temperature. Approximately 80% of the diode s had a leakage current in the range of 10-3

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11 A/cm2 after annealing at 550 oC, indicating a failure temp erature somewhere around 500 oC. TaC has a melting temperature of 3985 oC19 suggesting good thermal stability,21 and room temperature resistivity of 27 -cm,20 thus making it a potential candidate as a Cu diffusion barrier. In this study, Junji Imahori et al.20 compared films with different carbon concentrations, namely Ta53C47, Ta40C60, and Ta20C80. Of these, Ta53C47 films performed better than others. However, it t oo failed after annealing at temperatures higher than 600 oC for 30 minutes. Films with 60 and 80% C concentrations failed primarily due to diffusion of Cu through th e amorphous C phase and lesser diffusion of Cu occurred through grain boundaries of TaC. Ho wever, the prime reason for failure of Ta53C47 was due to the Cu diffusion through th e grain boundaries with Cu activation energy of 0.9 eV. A study of the interfacial reactions in the Cu/TaC/Si before and after annealing was conducted to determine the failure mechanism of the barrier layer.22 The failure temperature was a function of thic kness as 7, 35 and 70 nm TaC films failed at 600 oC, 650 oC and 750 oC respectively. The thicker TaC film integrity was compromised at a lower temperature of 600 oC due to the formation of amorphous Ta(O,C)x layer at the Cu/TaC interface. WCx was studied as a candidate diffusion ba rrier for Cu due to its high melting temperature of around 2785 oC23 and low electrical resistivity.24 A hexagonal closed pack W2C phase was achieved by chemical vapor de position at a growth temperature of 600oC. This growth temperature is fairly high keeping in mind future technology requirements where incorporation of lowmaterials will put an even more stringent temperature constraint.25 Low temperature CVD process to deposit W2C was demonstrated by Y. M.

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12 Sun et.al.26 TEM studies revealed 5-6 nm W2C crystallites in an amorphous matrix with a W/C ratio of 2:1 until 450 oC. The Cu/W2C/SiO2/Si stack remained intact after 400 oC annealing for 8-9 hours. Unfortunately, higher temperature annealing was not done in the study which would have revealed capacity of the diffusion barrier to withstand higher thermal stress. Room temp erature sputter deposited WCx was tested as a candidate diffusion barrier for Cu on Si.24 Fifty nanometer WCx failed according to analytical results after annealing at temperatures higher than 650 oC for 30 min. Electrical measurements, however, lowered the failure temperature to 550 oC. At 700 oC, the metallurgical stability of the WCx/Si interface is compromised. W reacts with Si to form W5Si3 at the WCx/Si interface. Cu diffuses through th e barrier layer and possibly through defects in the barrier layer into the Si substrate to form Cu3Si phase. In as-deposited condition, 60% of the measured diodes fail with a leak age current of 10-9 A/cm2. Also, 60% of the diodes fail after anne aling at a temperature of 550 oC with a leakage current of 10-7 A/cm2. Refractory Nitrides The activation energy for Cu to diffuse through grain boundaries is low making this the fastest diffusion pathway and killer of devices. A method to stop Cu diffusion is to block these super highways by “stuffing” th em. Oxygen was used as a stuffing agent in TiN diffusion barriers used for Al metallization. Al reacts with O2 and forms an aluminum oxide phase that hinders Cu from further diffusion. However, the same concept could not be applied to Cu.3 Nitrogen is a good stuffing agent that can solve the purpose. Excess nitrogen in the film moves out to the grain boundaries and stuffs them. Cu diffusing through the grain boundaries experien ces a repulsive force from the nitrogen thus stopping it from diffusi ng through the barrier film.27 Refractory nitrides are

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13 interesting candidates for Cu diffusion barrier application due to their lower resistivity,28 high melting temperatures, lower heat of fo rmation indicating better stability and the ability to block Cu diffusion by stuffing th e grain boundaries by nitr ogen. Some of the binary refractory nitrides investigated for Cu diffusion barrier app lications are Ti, Ta, W and Hf based which will be discussed further in detail. TiN was extensively used as a diffusion barrier for Al based metallization. Considerable effort was made to utilize the existing knowledge and processing methods for Cu based metallization as well. Howe ver, TiN deposited by both PVD and CVD methods resulted in a columnar grain struct ure in which grain boundaries run through the entire thickness of the barrier film.29-32 These columnar grain boundaries provide easy pathways for Cu diffusion and subsequent reaction with Si to form silicides. The properties of TiN films greatly depend on the deposition conditions which affect the microstructure, density and other relevant ba rrier properties of the film. The resistivity values range from 20 – 2000 -cm and density ranges from 3.2 – 5.0 g/cm3 33 A comparative study of different deposition me thods and conditions to deposit TiN films was done by Park et al.33 Porous films have lower densities making them susceptible for impurities like oxygen which alter the diffusion barrier properties. The TiN films failed in the temperature range of 500 – 750 oC depending on the f ilm properties, 500 oC being for CVD deposited films having lowest density and 750 oC being for sputter deposited films having highest density. Atomic layer deposi tion techniques were also tried for TiN barrier deposition with focus on the influence of microstructure, resistivity and impurity content on film barrier properties.34 However, the barrier film failed after annealing at 500 oC for 1 hr. The major drawbacks of CVD based techniques are unwanted C and O

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14 impurities, particulate generation, high deposition temperatures, and uncontrolled thickness variations across the surface. WNx is also an attractive candidate for Cu diffusion barrier applications due to its low resistivity and the ability to be deposit ed as an amorphous phase. Due to the absence of grain boundaries, Cu diffusion is hindere d or slowed. However, the phase of WNx mainly governs the barrier pr operties. A tungsten rich WNx (x 0.5) phase tends to dissociate at temperatures as low as 450 oC into W and W2N leading to barrier failure by diffusion of Cu through the grain boundari es due to lower recrystallization temperatures.35 A nitrogen rich WNx phase (x = 1) tends to be amorphous and remain amorphous for higher temperatures, but the gain is traded off with higher resistivity. Uekubo et al demonstrated the feasibility of W2N as a diffusion barrier, as compared to W and WN, and showed that 8 nm W2N was able to stop Cu diffusion in Si until 600 oC for 30 min.36 Failure of the diffusion barrier was due to recrystallization and grain boundary formation. An array of deposition methods have been tried to deposit WNx like sputtering,37 inorganic CVD,38 plasma-enhanced CVD,39 metal-organic CVD,40 and ALD.41,42 TaN is being currently used for Cu diffu sion barrier applications due to its high melting temperature and thermal stabilit y. Among the various phases of Ta-N, stoichiometric TaN has a melting temperature of 3087 oC and heat of formation ( Hf = -120 kJ/mol)43 making it more stable than Ta2N which has a melting temperature of 2050 oC and heat of formation of Hf = -98 kJ/mol. Chemical vapor deposition of tantalum nitride results in an insulating tetragonal Ta3N5 phase which is not suitable for diffusion barrier applications.44,45 A bi-layer structure of Ta/T aN is currently being used

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15 to overcome the obstacle of adhesion problem of TaN. TaN does not adhere well to Cu46 but it does adhere with SiO2.47 In case of Ta, it adheres we ll to TaN but not with SiO2 and is not an effective diffusion barrier by itself.47 Also, TaN helps nucleate Ta in the -Ta phase, which has a BCC structure and a resistivity of ~15-30 -cm, as compared to when Ta is directly deposited on SiO2 where it nucleates in the -Ta phase, with a resistivity of ~150-220 -cm.47 However, with device dimensions shrinking, conformality of high aspect ra tio is very critical. PVD48 has been used until now for deposition of liner and Cu seed layer, with ionized PVD (I-PVD)49 being the latest in the technology that has been able to extend the functionality to lower dimensions. However, with future technology constraints, the liner thickness should be less than 10 nm.50 Since the overall liner-Cu seed layer thickness re quirement decreases, either some major modifications has to be done to the I-PVD process or a switch to a better conformal process like atomic layer deposit ion (ALD) has to be adopted. Ternary Diffusion Barriers The activation energy required for Cu diffusion through the grain boundaries is low. Binary diffusion barriers ha ve been investigated in wh ich one of the elements, like nitrogen, “stuffs” the grain boundaries and di sallows Cu diffusion by blocking. However, at higher temperatures the binary compound recrystallizes, forming parasitic grain boundaries that should be avoided. Another way to solve the problem of creating a diffusion barrier for Cu is to form an amor phous matrix that acts as a diffusion barrier and remains amorphous (not recrystallized) wh en annealed at higher temperatures. This can be achieved by adding a third element into the binary matrix. This addition frustrates the binary lattice structure and delays or avoids the recrysta llization process when it is annealed at higher temperatures. Consider able research has been done on amorphous

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16 ternary nitride diffusion barriers for Cu metalli zation. Most of the research work has been done by incorporating Si, B or C as the third el ement in the binary matrix of refractory nitrides. Some of the work will be discussed below. TiSixNy were initially studied as Ti based li ners and used as diffusion barriers for Al metallization. Amorphous TiSixNy is attractive due to the absence of grain boundaries. The properties of diffusion barriers depend on the method of depositi on and its chemical composition. Chemical vapor deposited TiSixNy films (25 nm thick) were studied as candidate diffusion barriers for Cu.51 The barrier failed after a nnealing at temperatures higher than 700 oC for 30 min. Unfortunately, the barr ier resistivity is too high (800 cm) for applications. MOCVD deposited TiSiN f ilms were also investigated as diffusion barriers for Cu.52 The deposition steps involved a H2/N2 plasma treatment to increase the density of the film. Comparison of films that were plasma treated and not plasma treated was done. Films were characterized by sec ondary ion mass spectroscopy (SIMS) and Secco etch method for diffusion barrier failure. Secco etch revealed that failure occurred in films that were not plasma treated after annealing at 550 oC for 1 hr. For plasma treated films, failure occurred after annealing at 600 oC for 1 hr. SIMS detected Cu diffusion in the sample that was not plasma treated at 450 oC, whereas Secco etch revealed etch pits in the same sample after annealing at 550 oC. TaSixNy (x=1.4, y=2.5) films with different thickness (t= 5-40 nm) were sputter deposited on Si by Lin et al53 to evaluate their diffusion barrier characteristics. Electrical measurements on a p+n diode structure revealed failure of 5 nm TaSixNy after annealing at 500 oC for 30 min. The TaSixNy structure remained amorphous even after annealing at 800 oC, which is a good indicator that addition of Si to the TaN matrix was effective in

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17 keeping the lattice st ructure amorphous. Failure of the diffusion barrier was due to Cu diffusion through localized defects into Si. Th e nitrogen content plays an important role in determining the diffusion barrier characteristics. With the increase in the nitrogen content, the diffusion barrier pr operties are enhanced as the nitrogen helps in preventing formation of TaSi2 phase after annealing at higher temperatures in Ta-Si-N films.54 There has been some research on WSixNy diffusion barriers for Cu metallization. LPCVD deposited WSixNy shows promise as a diffusion barrier compared to Ta and Ti counterparts.55,56 Efforts are also made to take ad vantage of good conformal coverage obtained using the CVD deposition process while achieving superi or diffusion barrier properties similar to those achieved by PVD methods.56 Researchers have worked on using boron or carbon as a third el ement in the binary matrix.57 W-B-N seems to be a better diffusion barrier compared to W-Si-N due to its lower resistivity while having the same performance as W-Si-N.57

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18 CHAPTER 3 EXPERIMENTAL DETAILS AND CHARACTERIZATION Sputtering Sputtering is a physical vapor deposition t echnique where incident ions remove atoms (sputter) from the target surface by mo mentum transfer process. Sputtering was discovered in the 19th century by Grove and Pulker a nd reported by Wright in 1877. The number of atoms removed from the target su rface by a single incident ion is called the sputtering yield. The sputtering yield depends on the mass of th e atoms (target) and ions, bombardment energy of the inci dent ions, angle of incide nce and binding energy of the atoms in the target material. Table 3-1 shows the sputter yields of different ions when bombarded on the target materials as indica ted. Different classifi cations are used to define a particular sputtering technique depe nding on the type of sputtering configuration and also on the type of a reac tive species inside the depositio n chamber. As a result, there are different sputtering techni ques like DC sputtering, RF sp uttering, triode sputtering, magnetron sputtering and “unbalanced” magne tron sputtering. Depending on the absence or presence of a reactive species, it can be called non-reactive or reactive sputtering, respectively. The energy transfer taking place due to collision between two hard spheres can be given as: 2 24t i i t i tM M Cos M M E E

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19 wheretE iE are the energy of the target and incident particle respectively, i tM M ,are the mass of the target and incident particle and is the angle of incidence with respect to the surface normal of the target. Currently, almost any target material can be deposited by sputtering due to the latest advancement in technology. Co-sput tering can be carried out to deposit a compound on the substrate by sputtering two ta rgets simultaneously or by sputtering a single target in a reactive atmosphere or by using a compound target in an inert atmosphere. Some of the advantages of the sputtering deposition process are given below. Ability to coat large areas with uniform thickness Low-temperature deposition process Flexibility of target materials choice including insulators and semiconductors Target composition is replicated in the composition of deposited film. Good adhesive property obtaine d in the deposited films Sputtering can be up, down or sideways Reactive sputte ring possible Reproducibility Environmentally frie ndly process technology Can be easily scale up for commercial production The two major kind of sputtering configura tions conventionally used are DC and RF magnetron sputtering. They are described in detail below. DC Magnetron Sputtering In a DC sputtering process, the target is the cathode and either the substrate or the chamber walls are the anode. A very high negati ve DC voltage (several kV) is applied to the cathode. Ar gas is supplied in the chamber to create plasma. Ionized Ar+ ions are accelerated to the cathode, bombarding the su rface and sputtering atoms from the cathode surface. They also generate secondary el ectrons that ionize the gas atoms creating increased amounts of Ar+ ions. This increases the probability of collisions at the cathode surface. A circular magnet (magnetron) below the target traps the electrons emitted from

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20 the cathode surface to remain near the cathode. These electrons hop in a cycloid fashion on the target surface due to (E x B) forces ac ting on it. As the electrons are trapped near the surface region, the probabi lity of creating more Ar+ ions increases which increases the sputtering yield. Another advantage of us ing the magnetron source is that low gas pressures can be used inside the chamber to maintain a stable plasma. Since the gas pressure is low, internal gas collisions decr ease leading to higher yields of sputtering target material. As there is less chance of in ternal gas collisions, the sputtered materials impact the anode surface with a high kinetic energy. Figur e 3-1 shows a schematic DC sputtering chamber set-up. The disadvantage of DC sputtering is that the cathode material needs to be conductive. Insulating material s cause a charge build-up near the surface quenching the plasma. RF Magnetron Sputtering The major difference between DC and RF magnetron sputtering is the source. In radio frequency (RF) sputtering, an RF s ource is used, typically 13.56 MHz. A matching network is used to optimize pow er transfer from the RF s ource to the discharge and a blocking capacitor is used in the circuit to crea te a DC bias. With the use of an RF source, sputtering of insulating targets has been made possible. Figure 3-2 shows a schematic setup of a RF sputtering. The sputtering pro cess is similar to DC sputtering. Low gas pressures can be used for RF sputtering. One consideration during RF sputtering is to use target materials with high thermal conduc tivity. Low thermally conductive materials develop a high enough thermal gradient due to the sputtering proce ss to cause brittle fracture of the target.

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21 Characterization Characterization is an importa nt part of research. It he lps in identifying physical, electrical, magnetic, optical and chemical properties of the samples that helps in understanding the fundamental behavior of the material being studied. Some of the characterization techniques used in this study are discussed below. X-ray Diffraction X-ray diffraction (XRD) is one of the most versatil e non-destructive technique used todate to identify the crystalline phases and crystallinity of the sample. Advances in X-ray diffraction and understanding the basic physics between the in teraction of X-ray’s and the sample have given access to a plethora of information. Currently, information regarding strain, in-plane epitaxy, defect st ructure, film thickness etc. can easily be obtained from the constructive and destructive interference of X-rays with the sample. Figure 3-3 shows a basic X-ra y diffraction set up. Cu K radiation is generated and impinged on the sample. This radiati on undergoes constructive or destructive interference after reflecting from the sample depending on the path difference between the incident and reflected X-rays. Constr uctive interference will cause a peak at a particular 2 angle according to Bragg’s rule as given below: Sin d n 2 where is the wavelength of X-ray (Cu K radiation 1.54 ), d is the distance between two consecutive (hkl) planes and is the angle between the incident X-ray and the sample surface. In a cubic system the d-spacing is given as: 2 2 2l k h a do hkl whereoa is the lattice constant.

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22 For a single crystal, there are only specifi c orientations that satisfy the Bragg’s law. Diffraction peaks from these planes appear in the diffraction pattern. However, for a polycrystalline film having differently orie nted grains, diffraction peaks appear when those grains meet the diffraction conditions. The full width at half maximum of an X-ray diffraction peak gives information about the grain size. The relationshi p is called the Scherrer equation: BCos t 9 0 where t is the grain size, is the wavelength of x-ray, B is the full width at half maximum and is the Bragg angle. Auger Electron Spectroscopy Auger electron spectroscopy is an excelle nt technique for surface and sub-surface analysis. It is a highly sensitive technique as it probes only from few angstroms to few nanometers in depth from the surface. Usually an electron beam is used to probe and excite electrons in the atoms of the sample. As the core shell electron leaves the atom, it is in an excited state. A higher shell electr on fills the core shell vacancy and the energy emitted by this transition is transferred in ex citing another electron to emit from the atom which is called the Auger electron. Depe nding on the kinetic en ergy of the electron measured, the binding energy of the emitted Auger electron is calculated by the following equation: L L KE E E E K where K.E is the kinetic energy of the Auger electron, and EK EL and EL are the energies of the K and L shells, respectively.

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23 A total of three electrons are needed in order to fulfill an Auger transition. As a result, elements like H and He are undetectab le with AES. Suppose an electron from the K-shell of an element is removed by photons or electrons. The electron from the L shell fills up this vacancy with enough energy left to emit an electron from the L shell. This is called a KLL transition. Similarly, now highe r shells would fill up the L shell and would result in a LMM type of transition. The Auger equipment can be combined with an ion gun, which would sputter the sample and expose a new surface to be examined. A group of all those data points would give the depth profile character istics of the sample studied. This is therefore a very good method to study interfaces and diffusion profiles of elements inside the sample. The following in formation can be obtained from an Auger profile. Type of elements Amount percentage of the elements Chemical state of elements Valence band density of states X-ray Photoelectron Spectroscopy As the name suggests, a high energy phot on source (X-ray) is bombarded on the surface of the sample. It ejects a core shell electron in th e atom transferring it to an excited state. The atom returns back to ground state as the higher sh ell electron fills up the core shell vacancy, emitting the excess ener gy in the form of photon or by emitting an Auger electron. The kinetic energy of the em itted core shell electron is detected which gives a plethora of information about th e sample. The expression of kinetic energy measured is as below: E B h E K .

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24 X-ray photoelectron spectroscopy is a ve ry sensitive tec hnique for surface analysis as it probes only few angstroms fr om the surface. This is because only few electrons near the surface are able to escape the sample while most of the electron excited by the x-ray lose their energy by collisions with atoms and are not able to exit the surface. The flux of electrons able to exit the surf ace of the sample without being scattered ( Id) is given by the following expression: Sin d I Ie o dexp where Io is the original flux of el ectrons generated at depth d, e is the inelastic mean free path of electrons, is the angle of electron emission. As the binding energy of an electron is a characteristic of an atom, elemental information can by obtained by XPS. Also, a ge neral trend is that if the charge on an atom increases, so does the binding energy of th e electron. As a result, the valence state of the atom can be known. Amount percentage of the element presen t can be determined by peak intensity. The valence electrons pa rticipate in the bondi ng process and thus compound information can also be known from XPS data analysis. A depth profile is possible. However, it is very time consuming. Scanning Electron Microscopy and Energy Dispersive Spectroscopy Scanning electron microscopy is used fo r surface and cross-section imaging, topographical information and compositional in formation of a binary alloy system. An electron beam strikes the sample surface a nd various kinds of elastic or inelastic interactions occur resulting in emission of electrons. The electrons emitted out of the sample are of generally two types Secondary electrons

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25 Back-scattered electrons When the incident primary beam of elect rons impinge on the sample surface, they can undergo inelastic scattering by colliding w ith electrons of the sample, transferring some of its energy to them which exit the su rface of the sample depending on the amount of energy transferred. If sufficient energy is transferred, then electrons from the sample exit the surface and are collected by the detector. Electrons with energy of less than 50 eV are classified as secondary electrons. The primary beam can also undergo elastic scattering with the nucleus of the atom where transfer of energy is null or very small. These electrons have high energy and can ex it the surface easily. The amount of backscattered electrons depends on the atomic nu mber of the atom. A higher atomic number (Z) yields a higher amount of back-scattere d electrons which results in increased brightness of the image. However, for a cons tant Z, the back scattering yield remains unchanged if the primary beam energy is above 5 keV. The secondary electron yield does not depend much on the Z. The backscatteri ng and secondary electron images can be used to complement each other in seeking information. Since the primary beam has high energy, it can remove core electrons from the atom thus exciting it. The atom goes back into the ground state by filling the core vacancy by an electron from a higher orbital. The excess energy can either be emitted in the form of a photon or an Auger electron. The emitted photon can be detected and used for further study depending on its energy. This method is called as energy dispersive spectroscopy (EDS). Since the emitted photon is characteristic of the atom from which it emits, elemental information of the sample can be known. An elemental map of the surface can be created with use of today’s in struments. Quantification of the amount of a

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26 particular element can be carried out de pending on the peak intensity and using ZAF correctional factors, where Z is atomic numbe r, A is absorption and Z is secondary x-ray fluorescence. Atomic Force Microscopy The topography and roughness of thin films can be determined using atomic force microscope. The probe in an AFM is a sharp tip created at the end of a cantilever beam having a spring constant of 0.1-1.0 N/m. As the tip is brought near the surface of the sample, van der Waals forces become activ e between the tip and sample surface atoms resulting in the deflection of the tip. This de flection is monitored and measured by a laser striking the back of the cantile ver beam which plots the surface of the sample as the tip is moved across the sample. A general schematic of the AFM set up is show in Figure 3-4. The tip can be rastered over the sample surface and a 3-D plot can be obtained with nanometer spatial resolution. Lower spatial re solution can be obtained if the scanned area is reduced and slow scans are used. AFM is operated in two modes namely contact mode, where the tip is in contact with the sample surface, and tapping mode, where the tip is vibrated over the sample surface. The tapping m ode is particularly useful for analyzing soft samples like polymers. Transmission Electron Microscopy The wave nature of electron was first theo rized by Louis de Broglie in 1925. The term “electron microscope” was first used by Knoll and Ruska in 1932 when they build the first electron microscope and obtained im ages. TEM is an excellent tool to obtain high resolution lattice images and diffraction patterns of electron transparent samples. The high spatial resolution is due to the fact that high energy electrons are used to probe

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27 the sample which have extremely small wave length (on the order of ). The wavelength of the electron depends on the en ergy as per the expression: E 22 1 where is wavelength of electrons in nm E is energy of electrons in eV. A typical TEM operates at 200 to 400 keV. A highly coherent beam of monochromatic electrons is fo cused on the sample using a series of electromagnetic lenses. Some of the possible interactions be tween sample and the electrons are shown in Figure 3-5. For transmission electron micros copy, information is collected below the sample from the deflected and un-deflected electrons on a fluorescent screen to form images or diffraction pattern depending on wh ere it’s focused below the sample. Unlike scanning electron microscope, the entire sample is in focus all the time as along as its electron transparent. Some of the drawbacks of TEM are as follows Limited sampling volume Complex data interpretation Beam damage Sample preparation Electron diffraction is a very important quantitative characterizing technique to analyze a sample. Information from a partic ular area of a sample can be obtained by using selected area aperture (SAD) which fo cuses only on the area of interests. Other information, like defects (line and point def ects), burger’s vectors, 3-D loops, crystal orientation, orientation relationship between film and substrate or multilayer films and sample crystallinity can be obtained.

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28 Focused Ion Beam The samples required for TEM in this st udy were prepared by focused ion beam (FIB) technique. Focused ion beam uses an electron beam for imaging and a gallium ion source for imaging as well as milling. The reason for using Ga+ ions for milling is because the mass of Ga is approx. 127000 times that of electron and so provides a huge momentum transfer. The Ga source is heated to a liquid on a tip and Ga+ ions are extracted and focused on the sample by applyi ng a bias. The ion source is also used for milling purposes to prepare an electron transparent sample. Platinum is used to avoid milling the area of interest. Once the sample is electron transparent it is set free and transferred using glass rods to a grid for TEM analysis. Some of the advantages of using FIB as compared to conventional TEM sample preparations are Selected feature can be prepared for analysis Excellent control over the cross-section to be prepared Insulating samples can also be used No mechanical damage in the area of interest Van der Pauw Measurement and Four Point Probe Van der Pauw technique is used to measur e the resistivity of the sample in the semiconductor industry due to its ease and si mplicity. The technique doesn’t depend on the shape of sample. Four ohmic contacts are pr epared on the four corn ers or periphery of a sample. An acceptable version of sample geometry for this measurement is shown in Figure 3-6. As shown in the Figure3-7, a DC current is applied between contacts 1 and 2 (I12) and voltage is measured along contacts 3 and 4 (V43). This resistance RA is measured as follows: 12 43I V RA

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29 This is followed by another measurement by applying DC current between contacts 1 and 4 (I14) and voltage is measured between contacts 2 and 3 (V23). This gives resistance RB as given by the following equation: 23 14V I RB The sheet resistance RS is related to RA and RB through the following equation: 1 exp exp S B S AR R R R The bulk resistivity can be calculated as per the expression: d RS where d is thickness of the film The sheet resistance of coppe r before and after annea ling was measured by four point probe. A schematic set-up of four point probe is as shown in Figure 3-8. The DC current (I) is supplied through the extrem e probes and corresponding voltage (V) is measure by the inside probes. The sheet resistivity of a thin film of thickness (d) is calculated by using the following equation: I V d 2 ln And the sheet resistance (RS) is given by I V k RS where k is the geometric factor and its va lue is 4.53 for semi-infinite thin film.

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30 Table 3-1. Sputtering yi elds of different ions. Sputtering yield by 500 eV ions Be (9) Al (27) Si (28) Cu (64) Ag (106) W (184) Au (197) He+ (4 amu) 0.24 0.16 0.13 0.24 0.2 0.01 0.07 Ne+ (20 amu) 0.42 0.73 0.48 1.8 1.7 0.28 1.08 Ar+ (40 amu) 0.51 1.05 0.5 2.35 2.4-3.1 0.57 2.4 Kr+ (84 amu) 0.48 0.96 0.5 2.35 3.1 0.9 3.06 Xe+ (131 amu) 0.35 0.82 0.42 2.05 3.3 1.0 3.01

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31 Figure 3-1. Schematic diagram of DC sputte ring system with parallel plate discharge. Figure 3-2. Schematic diagram of RF sputte ring system with a capacitive, parallel plate discharge.

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32 Figure 3-3. Schematic diagra m of X-ray diffraction set-up. Figure 3-4. Schematic set-up princi ple of an atomic force microscope.

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33 Figure 3-5. Interaction out put between a high-energy elect ron beam and thin specimen. Figure 3-6. Sample geometries for Van der Pauw measurements.

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34 Figure 3-7. Schematic diagram of Van de r Pauw configuration for measurement of RA and RB resistances, respectively.

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35 Figure 3-8. Schematic diagram of four point probe measurement.

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36 CHAPTER 4 PROPERTIES OF W-Ge-N AS A DIFFU SION BARRIER MATERIAL FOR COPPER Introduction For present and future integrated elec tronic technologies, the use of barrier materials to enable materials integration is becoming increasingly impor tant. In current Si technology, the push for higher circuit densit y and low RC time delays has made copper the material of choice for interconnects due to its higher resistance to electromigration and lower resistivity as compared to Al Unfortunately, Cu is known to show poor adhesion to most dielectric materi als and rapidly di ffuses into SiO2 and Si. This obviously degrades the electr ical properties of devices58, creating the need for intermediate layers that provide a barrier to Cu diffusion. When considering the role of microstr ucture in diffusion processes, amorphous materials are generally bette r suited than polycrystallin e phases, as grain boundaries provide high diffusivity pathways for Cu diffusion through the barrier material. Among the material systems currently being deve loped, binary nitrides, such as TaN59,60 and WNx 61,62 are receiving significant attention. Fo r Ta-based barriers, the Ta-Cu phase diagram indicates that Ta and Cu are eff ectively immiscible even at their melting temperature. Unfortunately, recrystallizati on of TaN films occurs at approximately 600 oC, which is relatively low for diffusion barrier applications. WNx is also an interesting candidate as it is rela tively easy to synthesize as an amorphous film. In this case, the nitrogen content in WNx films has a significant infl uence on its diffusion barrier properties. For WNx films with low nitrogen content, the recrystallization temperature

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37 can be on the order of 450 oC.3 Higher nitrogen content yiel ds higher recrystallization temperature. WNx films grown by physical vapor depos ition have been reported to exhibit a recrystallization temperature as high as 600 oC. Yet, the primary mode of failure remains diffusion through grain boundaries that form during heat treatments. One approach to achieve higher recrysta llization temperature is to consider ternary compositions. The additional element added to the refractory metal-nitride composition frustrates the recrystallizati on behavior, rendering a stable amorphous mixture at high temperatures and thus mini mizing grain boundary diffusion. While failure of ternary diffusion barrier s still occurs through grain boundaries formed due to decomposition and recrystallization of the f ilm, this process gene rally takes place at higher processing temperature. For this reason, there is significant interest in ternary nitride alloys, such as Ta-Si-N,63-65 W-Si-N,56,66 W-B-N,67 and Ta-W-N68 due to their high recrystallization temperatur e as compared to the binaries. In this study, we report on the diffusion barrier properties of W-Ge-N thin films for Cu metallization. The W-Ge-N alloy is chemically similar to W-Si-N, should be more resistant to recrystallization than WNx, and may prove attractive for integr ation with SiGe or Ge devices. Experimental Details The W-Ge-N films were deposited at the rate of 8.64 nm/min on thermally grown SiO2 (630 )/n-type (100) Si substrates by re active sputter deposition. For comparison, WNx films were also deposited under similar co nditions. The substrates were sequentially cleaned with trichloroethylene, acetone, and methanol for 5 min each in an ultrasonic bath. The substrates were then loaded into the multi-target R.F. sputter deposition system via a load-lock. The base pressure of the sputtering chamber was 7 x 10-6 Torr. Typical forward sputtering power for the W and Ge targets was 200 W and 100 W, respectively.

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38 Nitrogen was incorporated into the fi lms by leaking a mixture of Ar and N2 at the ratio of 1 : 0.9 into the chamber at a fixed chamber pr essure of 11.5 mTorr. The thickness of the films was measured using a st ylus profilometer. In the experiments reported here, film thickness was maintained in the range 300 to 360 nm. All targets were pre-sputtered before deposition to remove any contamin ant present on the target surface. To assess the compatibility and diffusion properties of W-Ge-N with respect to Cu metallization, the nitride layer deposition was followed by in situ deposition of a Cu film 90 nm thick. During Cu deposition, Ar ga s was used as the sputter deposition gas at a fixed chamber pressure of 15 mTorr. Afte r deposition, the individual samples were annealed in a separate vacuum cham ber with a base pressure of 4 x 10-5 Torr at 400, 600 and 800 oC for 1 hr to study and compare the diffu sion barrier properties of the films. The crystallinity of films and formation of a ny intermetallic compounds by annealing were characterized by X-ray diffraction (XRD) m easurements. Resistivity was measured by the Van der Pauw method, while Auger elect ron spectroscopy was us ed to characterize the Cu diffusion profile in the nitride films. Energy dispersive spectrometry (EDS) was used to determine the composition of the films. Results and Discussion Initial studies focused on the crystallinity of the W-Ge-N films, both as-deposited and after high temperature annealing. Figure 4-1 shows XRD spectra of WNx and W-GeN films, both as-deposited and annealed at 400, 600, and 800 oC. The data show that films of both compositions are amorphous in the as-deposited condition. For the WNx films [Fig.4-1(a)], recrystallization is clea rly evident in the XRD pattern for annealing temperature of 400 oC and higher. The onset of grain structure will provide undesired diffusion paths via the grain boundaries. In c ontrast, the W-Ge-N film shows no evidence

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39 of recrystallization upon annealing at 400 or 600 oC. The XRD data [Fig.4-1(b)] for WGe-N films annealed at 400 and 600 C show no peaks related to the nitride material. Only at 800 oC do polycrystalline peaks appear. The addition of Ge to the W-N solid presumably frustrates crystallization, thus rendering the films amorphous for more severe annealing conditions relative to WNx. It should also be noted that the W-Ge-N film was far less susceptible to oxidation via ambient atmosphere exposure as compared to WNx. This may prove advantageous in terms of device processing. To assess the behavior of these films as diffusion barriers to Cu, the chemical profile of the annealed structures was dete rmined by Auger electron spectroscopy. Figure 4-2 shows the depth profiles for the WNx film (a) as-deposited and upon annealing at (b) 600 oC, and (c) 800 oC. For the as-deposited diffusion barr ier layer in Figure 4-2(a), there is a well-defined interface between Cu and WNx and the SiO2 buffer layer is evident. In contrast, after annealing at 600 oC, the Cu signal is seen th roughout the barrier layer and into the SiO2/Si. This is in agreement with the XRD data, which show the formation of grain structure upon anneali ng at a temperature of 600 oC. Annealing at 800 oC [Fig. 42(c)] yields a Cu di ffusion profile similar to that for the 600 oC anneal. There is also some apparent intermixing of W and SiO2 seen at the WNx/SiO2 interface in the Auger depth profiles of the annealed samples, which is perhaps related to a change in surface roughness. The intensity of the nitrogen profile decreases sl ightly as we increase the annealing temperature to 800 oC. This may reflect the decomposition of WNx and liberation of N2 from the films under those conditions This is not unexpected as Affolter et al. 69 have shown that nitrogen is liberated from W-N alloy thin films when annealed at temperatures above 700 oC.

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40 The chemical composition of Cu/W-Ge-N/SiO2/Si structures was also examined with Auger electron spectrometry. Figure 4-3 sh ows the chemical depth profiles in (a) asdeposited W-Ge-N and after annealing at (b) 600 oC and (c) 800 oC. In Figure 3(a), distinct interfaces at th e Cu/W-Ge-N and W-Ge-N/SiO2 boundaries show that there is no Cu diffusion during growth. At an annealing temperature of 600 oC, the interfaces and layers remain distinct, consistent with no or minimal diffusion of Cu through the barrier layer. These results suggest that W-Ge-N films possess superior diffusion barrier properties as compared to WNx. At an annealing temperature of 800 oC, Cu is observed in the diffusion barrier as seen for the WNx film. Again, Cu diffusion correlates with the appearance of grain structure in th e film (i.e., (111) reflection of -W2N in the XRD pattern). This is also consistent with N loss suggested by the AE S sputter profile, and thus its ability to stuff the diffusion pathways. The resistivity of the W-Ge-N films was measured using the Van der Pauw method. In general, the resistance of the as-deposited W-Ge-N films is higher than WNx. As shown in Figure 4-4(a), the resistivity of WNx and W-Ge-N decreases as the annealing temperature increases While the resistivity of WNx decreases upon annealing at 400 oC, the resistivity for W-Ge-N remains relatively unchanged after annealing at 400 oC. There is a progressive decrease in resis tivity for both materials with further increase in annealing temperature. The change in re sistivity correlates with onset of grain structure, suggesting electron transport ac ross grain boundaries. The decrease may also be due to the loss of nitrogen from the films. The resistivity of WNx is two orders of magnitude lower than W-Ge-N at 800 oC reflecting the robustness of W-Ge-N film, for which increased decomposition temperature slows nitrogen evolut ion. Figure 4-4(b)

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41 shows the resistivity of W-Ge-N films vs. spu ttering power to the Ge target. Clearly, the resistivity scales with Ge content in the film. Conclusion In conclusion, the diffusion barrier propert ies of W-Ge-N thin films have been investigated. X-ray diffraction shows recrysta llization of WNx films at an annealing temperature of 400 oC and higher, while W-Ge-N films show recrystallization peaks only at an annealing temperature of 800 oC, suggesting that the addition of Ge frustrates the recrystallization behavior of WNx. The AES data show complete Cu diffusion across the WNx layer for an annealing temperature of 600 oC, while for W-Ge-N films the Cu/WGe-N and W-Ge-N/SiO2 interfaces remain dis tinct at that temperature, indicating that WGe-N has better diffusion barrier properties. Th e resistivity of both films decreases with increasing anneal temperature, with the resistivity of WNx two orders of magnitude lower than W-Ge-N after annealing at 800 oC. This behavior is consistent with enhanced stability of W-Ge-N with respect to f ilm decomposition and subsequent nitrogen evolution.

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42 Figure 4-1. X-ray diffracti on patterns of as-deposited and annealed films at different temperature of (a) WNx films (b) W-Ge-N films.

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43 Figure 4-2. AES depth profiles of Cu/WNx/SiO2/Si (a) as-deposited (b) annealed at 600 C/1 hr (c) annealed at 800 C/1 hr.

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44 Figure 4-3. AES depth profiles of Cu/W-GeN/SiO2/Si (a) as-deposited (b) annealed at 600 C/1 hr (c) annealed at 800 C/1 hr.

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45 Figure 4-4. Resistivity vs. anneal ing temperature (a) for W-Ge-N and WNx films. Also shown (b) is the resistivity as a function of sputter target power for the Ge and W targets.

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46 CHAPTER 5 INVESTIGATION OF W-Ge-N DEPOSITED ON Ge AS A DI FFUSION BARRIER FOR Cu METALLIZATION Introduction The progression to ever-decreasing semiconduc tor device dimensions brings with it new challenges for material integration. As SiGe-based microelectronic technology moves toward a 45 nm technology node, ther e is an imminent need to replace Al interconnects. Cu is an attractive candidate because of its low resistivity and high resistance to electromigration as compared to Al.70,71 Replacing Al/SiO2 interconnect technology with Cu/low-k dielectrics can yiel d a large reduction in RC time constant delay3, providing for a significant motivation to ca rry out the replacement search for Al. Unfortunately, Cu has known problems with adhe sion to low-k materials. It also has a high diffusion rate in silicon a nd germanium creating deep level traps in Si and a shallow level in Ge located at 0.04 eV near the valence band.72-74 Diffusion barriers are needed to achieve in tegration of Cu with Si-Ge and Ge. In diffusion barrier materials the diffusion pathwa ys are interstices, vacancies and grain boundaries. Diffusion through grai n boundaries is fastest and more prominent. By “stuffing” the grain boundaries with selected dopants, diffusion can often be hindered. In general, amorphous materials are highly prefe rred due to the absence of grain boundaries. Common diffusion barrier materi als studied for Cu metallization are refractory metalnitrides, which include TaN, TiN, HfN and WNx.42,59,75-79 However, these have limited utility as viable diffusion barriers becaus e of their relatively low recrystallization

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47 temperatures. Recrystallization of amorphous diffusion barriers can often be inhibited by the addition of a third element to the matrix. Previous results show higher recrystallization temperature fo r W-Ge-N as compared to WNx as the introduction of Ge effectively frustrates the matrix.80 SiGe devices have higher mobility as compared to Si devices and the flexibility of band-gap engineering, thus bei ng applicable in high speed el ectronics. The properties of the metal/SixGe1-x contact layer is important for semiconductor device applications. Considerable research has been done on the chemical reactivity of metals, such as Co, Ti,72,82 Pt,83 Pd,84 and Cu85 on SiGe. In most cases, it results in the formation of metal-Si, metal-Ge, or metal (SixGe1-x) alloys in the temp erature range of 400-600 oC. In particular, direct deposition of Cu on Si1-xGex results in the formation of unstable Cu3(Si1-xGex) in the temperature range of 250-400 oC.85 One solution could be to grow a metal-rich Cu3Si or Cu3Ge phase directly on Si1-xGex. Unfortunately the Cu3Si phase reacts with oxygen when exposed to air. The Cu3Ge phase is more stable and less reactive with oxygen.86 In general, the need to identify viable barriers of Cu integration with Ge-based structures persists. In this work, we report on the diffusi on barrier properties of W-Ge-N thin films deposited on Ge for Cu metallization. The properties are compared to WNx deposited on Ge under similar conditions to evaluate its su itability and properties Due to the ternary nature of this diffusion barrier, W-Ge-N is expected to have be tter recrystallization properties as compared to WNx. Experimental Details W-Ge-N films were deposited on p-type Ge (001) by reactive sputtering. The Ge substrates were cleaned by a sta ndard procedure reported elsewhere77 to remove any

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48 organic residue or impurity on the surface. Th e substrates were loaded in the reactive sputter chamber with a ba se pressure of 3 x 10-6 Torr via a load-lock. Nitrogen was incorporated in the film by flowing a mixture of N2 and Ar at a ratio of 1:3 at a fixed pressure of 10 mTorr. Prior to deposition, a ll targets were cleaned by pre-sputtering by flowing Ar+N2 in the chamber. Forward sputtering power for the W and Ge targets were 200 W and 100 W respectively. The deposition rate for W-Ge-N film on Ge was 10.2 nm/min under the above mentioned conditi ons. Thickness was measured by a stylus profilometer and was kept in the range of 50 to 300 nm. Cu metallization was carried out in-situ after depositing W-Ge-N thin films to determine suitability as a diffusion barrier. S putter deposition of Cu was carried out by flowing Ar gas at a fixed chamber pressure of 5 mTorr. After depositing the film stack, individual samples were separate ly annealed in the range of 400 oC – 700 oC in a tube furnace. Ar gas was flowed through the t ube furnace at 65 sccm for at least 10 hours before starting the annealing to remove a ny residual air mixture. Typical annealing experiment was carried out for 1 hr to study the diffusion barrier prop erties of the film. X-ray diffraction was used to identify any inte rmetallic phases and access the crystallinity of the films after annealing. Cu diffusion profile through the film was determined by Auger electron spectroscopy (AES) and Energy dispersive spectrometry (EDS). Interface properties were determined by cross-secti on transmission electron microscopy (X-TEM). Results and Discussion Crystallinity and high temperature phase formation of the Cu/W-Ge-N/Ge film structure before and after annealing was determined by XRD. The XRD spectra for Cu/WNx/Ge and Cu/W-Ge-N/Cu film structure is shown in Figure 5-1, both as-deposited

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49 and annealed in the temperature range of 400 – 700 oC. The WNx diffusion barrier shows little crystallization in as-deposited conditi on as evident by (111) peak, whereas W-Ge-N film is amorphous. The (111) peak intensity increases at 500 oC indicating further crystallization as evident from Figure 5-1(a). This leads to increase in the formation of undesired grain boundaries that provide fast diffusion paths. Cu diffuses through the grain boundaries and reacts with the underlying Ge substrate re sulting in the formation of a Cu3Ge phase that is clearly evident from Fi gure. 5-1(a). In comparison, there is no recrystallization of W-Ge -N films at 500 and 600 oC. However, at 600 oC, Cu reacts with Ge in the W-Ge-N layer, and subsequently with Ge substrate to form the Cu3Ge phase which is evident from Figure 5-1(b). This re sults in depletion of Ge from the W-Ge-N film. Upon further high temperature annealing, recrystallization of the Ge-depleted WGe-N takes place leading to barrier failure. Th is result indicates that adding Ge to W-N alloy effectively hampers the recrystall ization process even at high annealing temperatures as compared to WNx. The Cu diffusion profile through the diffusion barrier in as-deposited and annealed structures was determined by A uger electron spectroscopy. The film thickness for WNx and W-Ge-N was 50 nm for the AES chem ical profile study. Figure 5-2 shows the chemical profile of Cu for Cu/WNx/Ge structure (a) as-depos ited and after annealing at (b) 400 oC and (c) 500 oC. The Auger profile shows negligible Cu diffusion in WNx at 400 oC. However, upon annealing to 500 oC, Cu is seen to ra pidly diffuse through the barrier film and to the substrate. This is in agreement with XRD data at 500 oC where it shows increased crys tallization of WNx, thus increasing the grain boundaries, i.e., diffusion pathways, resulting in the formation of Cu3Ge phase. The nitrogen signal

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50 intensity steadily decreases at subsequent hi gh temperature annealing compared to the asdeposited sample indicating decomposition of WNx by liberating N2 at higher temperatures. The chemical profile of Cu/W-Ge-N/Ge structures was also measured by Auger electron spectroscopy as shown in Figure 5-3( a) as-deposited and af ter annealing at (b) 400 oC and (c) 500 oC. As is evident, there is little or no Cu diffusion through the barrier upon annealing at temperatures less than 500 oC. The Cu/W-Ge-N and W-Ge-N/Ge interfaces are distinct in as-deposited and 400 oC annealed samples. At 500 oC, Cu starts consuming the Ge in the film. This is also supported by XRD data [Fig. 5-1(b)], where a small intensity (-111) Cu3Ge peak appears for annealing temperature of 500 oC. Upon annealing at higher temperature, Cu complete ly consumes the Ge in the film, thereby depleting the W-Ge-N matrix. This results in the recrystallization of WNx as is evident from Figure 5-1(b) at 700 oC resulting in rapid Cu diffusion. The result above suggests that W-Ge-N hinders recrystallization by frustr ating recrystallization of the matrix and is a better diffusion barrier as compared to WNx. The high oxygen cont ent noticed in the chemical profiles is due to background oxygen during sputteri ng. The negative enthalpies of formation of W-O and Ge-O bond are 672 kJ/mol and 659.4 kJ/mol respectively indicating the stability of th e compound after formation. Longer pre-sputtering time of target could help in reducing the oxy gen content in th e deposited film. Figure 5-4 and Figure 5-5 shows th e Cu diffusion profile in WNx and W-Ge-N film annealed at 500 oC, respectively, which was m easured by Energy dispersive spectroscopy (EDS) attached to a crosssection transmission electron microscope (XTEM) system. As seen in Figure 5-4, Cu signal is s een throughout the WNx film and

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51 into the Ge substrate. However, very little (half counts as compared to WNx) Cu signal is seen coming from the W-Ge-N (Fig. 5-5) di ffusion barrier. This result corroborates the above mentioned XRD and AES data suggesti ng excellent diffusion ba rrier properties for W-Ge-N as compared to WNx. The interface properties were determined by XTEM. Figure 5-6 shows XTEM images of (a) Cu/WNx/Ge structure and (b) Cu/W-Ge-N/Ge structure both annealed at 500 oC. The WNx/Ge and W-Ge-N/Ge in terfaces are abrupt, indicating no intermixing or reactions between them even after annealing. However, the Cu/WNx and Cu/W-Ge-N interface is rough, su ggesting some intermixing and Cu diffusion in WNx films. XTEM images shows well-de fined (111) grain structure for Cu films deposited on W-Ge-N [Fig. 6(b)] as compared to WNx films [Fig. 6(a)]. This may be important as (111) oriented Cu films has high resistance to electromigration. Conclusion In conclusion, the diffusion barrier propert ies of W-Ge-N thin films deposited on p-Ge (001) substrates were investigated. W-Ge-N films have a higher recrystallization temperature (700 oC) as compared to WNx films as shown by X-ray diffraction. The failure of W-Ge-N diffusion barrier films at high temperatures o ccurs by Cu diffusing through the grain boundaries form ed by recrystallization of Ge depleted W-Ge-N. This Ge depletion is caused by the consumption of Ge in the barrier film by Cu at high temperatures, thereby forming the Cu3Ge phase. Removal of Ge from the film makes the W-Ge-N barrier more likely to recrysta llize. In contrast, failure of WNx diffusion barrier at high temperature takes pl ace by diffusion of Cu through WNx grain boundaries. Nitrogen evolution from the film at hi gh temperature causes decomposition of the diffusion barrier, thus enhanc ing crystallization of WNx film. The AES data clearly

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52 shows complete Cu diffusion throughout the WNx layer at 500 oC annealing temperature, whereas there is little or negligible Cu diffusion through W-Ge-N films. This suggests that W-Ge-N is a better diffusion barrier. This is substantiated by th e Cu profile measured by EDS. The Cu films deposited on W-Ge-N ha ve better orientati on as compared to WNx after annealing at high temperat ures, thereby increasing its resi stance to electromigration.

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53 Figure 5-1. X-ray diffracti on patterns of as-deposited and annealed films at different temperature of (a) WNx films (b) W-Ge-N films.

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54 Figure 5-2. AES depth profiles of Cu/WNx/Ge (a) as-deposited (b) annealed at 400 C/1 hr (c) annealed at 500 C/1 hr.

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55 Figure 5-3. AES depth profiles of Cu/W-GeN/Ge (a) as-deposited (b) annealed at 400 C/1 hr (c) annealed at 500 C/1 hr.

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56 Figure 5-4. EDS depth profile of Cu/WNx/Ge annealed at 500 C/1 hr.

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57 Figure 5-5. EDS depth profile of Cu/W-Ge-N/Ge annealed at 500 C/1 hr.

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58 Figure 5-6. X-TEM images of (a) Cu/WNx/Ge annealed at 500 C/1 hr (b) Cu/W-GeN/Ge annealed at 500 C/1 hr.

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59 CHAPTER 6 COMPARATIVE STUDY OF HfNX AND Hf-Ge-N DIFFUSI ON BARRIERS ON Ge Introduction For many years, aluminum has been the primary interconnect metal for Si-based integrated circuits. However, with devi ce dimensions shrinking to sub-45 nm and demands for high current density increasi ng, the conductivity and electromigration properties of Al become limitations to perf ormance. In response, Cu is beginning to replace conventional Al interconnects given its better electromig ration resistance and lower electrical resistance.70,71,87 The use of low resistivity Cu compared to Al significantly reduces the circuit time constant delay to make the circuit faster. As the need for high speed electronics grows, there is also a renewed interest in Ge and SiGe based devices because of inherent adva ntages of Ge over Si, i.e., smaller Eg, higher mobility of charge carriers a nd lower dopant activation energy.88-90 Si1-xGex -based devices are also of interest because of the innate flexibil ity to tailor the bandgap through the alloy composition.91-93 These factors provide sufficient impetus to investigate barrier layer materials needed in inco rporating Cu interconnects in Si1-xGex and Ge-based devices. For interconnect applications, copper cannot be deposited directly on Si-Ge since it diffuses rapidly in Si and Ge creating deep level traps.73,94 For the case of Si, it forms copper silicides at saturation. It al so passivates dopants by forming Cu-D (D is dopant atom) covalent pairs thus alte ring the intended doping levels.95 Copper is also known to diffuse rapidly in Ge with an av erage diffusion coefficient of 3 x 10-5 cm2s-1 in the 700 to

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60 900 oC temperature range.96 Copper introduces three accep tor levels in Ge, two at Ev+0.04 and Ev+0.32 eV and another at Ec-0.26 eV respectively.74 Direct deposition of Cu on Si1-xGex results in the formation of Cu3(Si1-xGex) and passivation of the dopants.97 In addition to the above issues, Cu also exhi bits poor adhesion to dielectrics commonly used in Si device structures.7 Considerable work has focused on id entifying viable Cu diffusion barrier materials on Si. Since amorphous materials l ack grain boundaries that are fast diffusion pathways, they are ideally suited for applic ation as a diffusion barrier. Recent material systems that have been studied as possible Cu diffusion barriers for Si include refractory metal nitrides, such as TaN, TiN, WNx.59,61,75,97 These binary nitrid es, however, tend to recrystallize at moderate temperature, thus becoming susceptible to rapid Cu diffusion. There is significant interest in identify ing diffusion barrier materials that remain amorphous at high processing te mperature and effectively bl ock Cu diffusion. Increasing the temperature necessary for crystallization can often be achieved by the addition of a third element to a binary matrix material. Some of the ternary materials systems that have been studied include Ta -Si-N, W-Si-N, W-Ge-N.98-100 In this paper, we report on the recrystallization of HfNx and Hf-Ge-N thin films deposit ed on Ge and their diffusion barrier properties for Cu metallization. Experimental Details HfNx and Hf-Ge-N thin films with varyi ng thickness (15, 50 and 300 nm) were deposited on p-Ge (001) si ngle crystal substrates by r eactive sputtering at room temperature. Prior to deposition, the substrates were cleaned with trichloroethylene, acetone and methanol in an ultra-sonic bath for 5 min each to remove any organic residue from the surface. The substrates were introdu ced in a reactive sputter deposition chamber

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61 with a base pressure of 3 x 10-7 Torr via a load-lock. Nitr ogen was incorporated in the film by flowing Ar and N2 in the chamber at a ratio of 3:1. The total chamber pressure during deposition was 10 mTorr. Prior to depos ition, the targets were cleaned in-situ by pre-sputtering with Ar+N2 at a fixed chamber pressure of 15 mTorr. The forward sputtering power for Hf and Ge was 200 and 100 W, respectively. The typical deposition rate for HfNx and Hf-Ge-N films was 1.8 and 6.23 nm/min, respectively. Identical thickness was achieved for both films by vary ing the deposition time. Film thickness was measured by a stylus profilometer. Nitride film deposition was followed by in-situ deposition of Cu films. The forward power used for Cu deposition was 200 W. The Cu thickness was maintained constant at 300 nm for all films. The deposit ion was carried out by fl owing Ar inside the chamber at a fixed chamber pressure of 5 mTorr. Individual film stacks were then separately annealed in a tube furnace in the temperature range of 400 to 700 oC for 1 hr. Before starting the annealing process, the tube was purged by flowing Ar gas at 65 sccm for at least 10 hrs. The film crystallinity before and after annealing was determined by Xray diffraction (XRD) while the film su rface morphology and roughness after annealing were determined by field emission-sca nning electron microscopy (FE-SEM). The chemical depth profile of Cu diffusion th rough the diffusion barrier was determined by energy dispersive spectroscopy (EDS). The chemical state analysis of Cu and intermetallic compound formation with Ge was investigated by X-ray photoelectron spectroscopy (XPS). Interface reactions and prop erties were determined by cross-section transmission electron microscopy (X-TEM).

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62 Results and Discussion The HfNx films were amorphous in the as-d eposited condition and showed no signs of recrystallization for any film thickne ss even after high temperature annealing. A lack of crystallization upon a nnealing is desirable as form ation of grain boundaries leads to rapid Cu diffusion. The HfNx diffusion barrier properties are expected to be attractive based on its high melting temperature (3330 oC).101 Materials that have a high melting temperature also generally show a have hi gh recrystallization temperature since both process involve bond breaking. HfN films have been shown to be stable to thermal decomposition up to 1000 oC.102 Fig. 1 shows the X-ra y diffraction patterns for Cu/HfNx/Ge as-deposited and after high temperatur e annealing in the range of 400 to 700 oC in an Ar atmosphere. For 300 and 50 nm th ick HfNx diffusion barrier films that were annealed at a temperature of 600 oC or greater, the Cu films e xhibit a shift in the Cu (111) peak towards smaller 2 values. This may indicat e a reaction with the HfNx film. It is noted, however, that for the 300 nm thick HfNx barrier, no Cu3Ge phase was formed even after annealing at 700 oC. For the 50 nm thick HfNx film (Fig 1b), as aforementioned, there is a definitive shift of Cu (111) peak at 700 oC annealing temperature which may be due to reaction of Cu with the underlying HfNx layer. Also evident for the 50 nm thick HfNx sample is the formation of non-stoichiometric Cu3-xGe phase after annealing at 700 oC. This indicates Cu diffusion through the HfNx diffusion barrier and reaction with the underlying Ge substrate. Cu3-xGe phase could also have b een formed by possible Ge outdiffusion through the diffusion barrier and subseq uent reaction with Cu For the ultra-thin HfNx diffusion barrier film (15 nm), the barrie r fails at lower temperature as evident from Fig. 1c that shows Cu3Ge phase formation after annealing at 500 oC and above. The 400 oC anneal pattern, however, does not re veal evidence of barrier failure.

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63 The properties of Cu/Hf-Ge-N/Ge multilaye rs were then examined and compared to the Cu/HfNx/Ge samples. Fig. 2 shows X-ray di ffraction patterns for Cu/Hf-Ge-N/Ge as-deposited and after high temperature annealing in the range of 400 to 700 oC in Ar atmosphere. Based on the behavior of the th ickest film, the Hf-Ge-N films remained amorphous after annealing at a temperature as high as 700C for all film thicknesses. For the 300 nm thick Hf-Ge-N film (Fig. 2a), there is little or no shift in the Cu (111) peak even after annealing at 700 oC. For the 50 nm thick Hf-Ge-N film thickness (Fig 2b), Cu3xGe phase formation is evid ent after annealing at 600 oC indicating that Cu has diffused through the barrier film to reac t with the underlying Ge substrat e. It is also noted for the 600 oC annealed sample that the Cu (200) and (111) peaks are no longer present, suggesting significant loss of Cu to the underlying material. At 700 oC annealing temperature, sufficient Cu diffuses through th e barrier layer to fo rm stoichiometric Cu3Ge phase, again indicati ng barrier failure. These data suggests that while Hf-Ge-N and HfNx have similar recrystallization behavior, the diffusion ba rrier properties of HfNx are superior. In particular, the absence of the Cu (200) peak for the film on 50 nm thick Hf-Ge-N annealed at 600 C suggests significant diffusion as compared to HfNx. One possible factor in determining the properties of the two materials relates to the relative percentages of the elements pr esent in the diffusion barrier. As mentioned before, HfNx and Hf-Ge-N films were deposited at the rate of 1.8 and 6.23 nm/min, respectively. As the forwarding power to Hf was kept constant during deposition of both films, the total Hf content in the HfNx sample is about 3.5 times greater than that for the corresponding Hf-Ge-N film of the same thickness. This results in the deposition of a Ge-

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64 rich Hf-Ge-N film. Cu is known to react r eadily with Ge. For example, at room temperature, a 20 nm Cu3Ge reaction layer will form at a Cu/Ge interface in 24 hrs in a binary reaction couple103 and the reaction rate should increase with increased anneal temperature. The atomic percentages of each element present in Hf-Ge-N film as measured by Auger electron spectroscopy were 3 at. % nitrogen, 41.5 at. % oxygen, 28.8 at. % germanium and 26.7 at. % hafnium, respectively, in as-deposited condition. The surface morphology of the barrier mate rials was revealed by field emissionscanning electron microscopy. Fig. 3 shows a comparison of FE-SEM micrographs for (a) Cu/HfNx/Ge and (b) Cu/Hf-Ge-N/Ge annealed at 500 oC; samples that retained barrier integrity as evidenced by the XRD patterns shown in Fig. 1b and 2b. The thickness of each HfNx and Hf-Ge-N layers was 50 nm. The grain structure observed in the micrographs is that of the Cu. Note that th ere is no evidence of delamination. Fig. 4 shows the FE-SEM micrographs for 50 nm thick (a) Cu/HfNx/Ge and (b) Cu/Hf-Ge-N/Ge films annealed at 600 oC. After annealing at 600 oC, the surface morphology is significantly different fo r the copper films on HfNx as compared to that on Hf-Ge-N films. For Cu on the HfNx, the Cu films are continuous w ith a roughness similar to that seen for the 500 C anneal. For the Cu on Hf-Ge-N, however, significant Cu segregation is observed. This is consistent with th e suppression of the (002) Cu peak for this structure and annealing temperature. R eaction with the Hf-Ge-N and possible Cu diffusion through the Hf-Ge-N film leads to depletion of Cu from the surface and segregation of Cu islands. This apparent Cu loss on the surface is in agreement with the XRD data showing the appearance of Cu3Ge peaks for Cu films on Hf-Ge-N barriers that

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65 are annealed at 600oC. No such peaks are detected for the comparison HfNx film suggesting improved diffusion barrier quality of the latter film. The interface properties a nd reactions were examined by cross-section transmission electron microscopy. Fig 5 show s the X-TEM images of the 50 nm thick (a) Cu/HfNx/Ge and (b) Cu/Hf-Ge-N/Ge f ilms after annealing at 600 oC for 1 hr. Cu diffusion is clearly seen in the Hf-Ge-N film with the formation of Cu3Ge phase formed below the diffusion barrier f ilm. The image of the HfNx film, however, shows a negligible amount of Cu diffusion as indicate d by a continuous Cu film on the surface and no indication of formation of the Cu3Ge phase. The discontinuous layer at the HfNx/Ge interface is due to delamination of HfNx. This could be due to TEM sample preparation prepared by focused ion beam (F IB). The chemical diffusion profile of Cu was determined by energy dispersive spectroscopy attached to the cross-section TEM. Figs. 6 and Fig. 7 show the chemical diffu sion profile of Cu after annealing at 600 oC for 1 hr for the HfNx and Hf-Ge-N films respectively. A Cu signal is present in the Hf-Ge-N barrier layer and Cu3Ge has clearly formed by transport through the barrier film to the Ge substrate. In contrast, HfNx shows little Cu signal is seen from the Ge substrate indicating that the HfNx barrier layer prevented Cu diffusion to the Ge substrate. The EDS peak positions of Cu and Hf overlap each other. As a result, a bump in the Cu intensity profile is observed in the HfNx layer. Also, a strong Hf EDS intensity is seen from the entire region of the Cu layer due to the EDS detector’s inability to differentiate between Cu and Hf. The chemical state of Cu and intermetal lic phase formation after annealing was determined by X-ray photoelectron spectro scopy (XPS). Fig. 8 compares the Cu 2p3/2

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66 peak shifts in Cu/Hf-Ge-N/Ge film for differe nt sputtering times in the (a) as deposited material (b) after anne aling the film at 600 oC for 1 hr and Ge 2p3/2 peak shifts in (c) Cu/Hf-Ge-N/Ge for different sputte ring times after annealing at 600 oC. The Cu surface of the as-deposited film is clearly oxidized fo rming a CuO layer. This is evident from the characteristics satellite peaks formed for Cu+2. After sputtering, however, the peak shifts and matches with pure Cu (932.8 eV). As s een in Fig. 8b, after annealing the 50nm Cu/Hf-Ge-N film at 600 oC, the Cu 2p3/2 peak in the as-received condition forms at 934.8 eV indicating its reaction with Ge and formation of Cu3-xGe. This is also consistent with the XRD and X-TEM data which show the formation of Cu3-xGe. The peak intensity increases with sputtering time as more Cu participation in the Cu-Ge bond is revealed. Apparently, there is slig ht shift in the Cu 2p3/2 peak to 933.3 eV after 60 minutes sputtering indicating that Cu might react with oxygen a nd form a Cu-O compound. The Ge 2p3/2 peak appears at 1221.9 eV after annealing at 600 oC. The Ge 2p3/2 peak intensity increases with sputtering time indicating in creased Ge participation in the Cu-Ge bond formation. There is a shift in the Ge 2p3/2 peak position to 1221 eV after sputtering for 60 minutes. This might be due to reaction with oxygen and formation of Ge-O bond. Conclusion In conclusion, a comparative study of th e diffusion barrier properties of Hf-Ge-N and HfNx deposited on (001) Ge single crysta l wafers was conducted. The FE-SEM images show almost identical surface mo rphology of Cu films af ter annealing at 500 oC. Annealing at 600 oC, however, results in considerable extent of diffusion across the HfGe-N films leaving a discontinuous Cu film on the surface. Furthermore, sufficient Cu transport occurs to form Cu3Ge which is evident from XRD da ta. In contrast, little or no diffusion takes place for HfNx films of the same 50 nm th ickness and annealing condition

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67 leaving Cu films continuous and smoother. Th is is also substantia ted by cross-sectional TEM images which clearly show the formation of a Cu3Ge below the Hf-Ge-N diffusion barrier, but no such phase is formed in the comparison HfNx film after annealing at 600 oC. The chemical valence state was determined by XPS and the results point to Cu-Ge bond formation after high temperature ann ealing. The chemical diffusion profile measured by EDS shows Cu signal emanating from the Hf-Ge-N diffusion barrier and the underlying Cu3Ge phase formed after annealing at 600 oC. There is little or no Cu signal however, observed in the HfNx diffusion barrier and underlyi ng Ge substrate, indicating that the HfNx diffusion barrier was successful in prev enting Cu diffusion to the substrate. It is thus concluded that HfNx is an attractive diffusion barrier for Cu on Ge, while HfGe-N demonstrates limited utility.

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68 Figure 6-1. X-ray diffr action patterns of Cu/HfNx/Ge films in as-deposited and annealed conditions for varying thickness of (a) 300 nm, (b) 50 nm, and (c) 15 nm.

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69 Figure 6-2. X-ray diffraction patterns of Cu/Hf-Ge-N/Ge films in as-deposited and annealed conditions for varying thickne ss of (a) 300 nm, (b) 50 nm, and (c) 15 nm.

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70 Figure 6-3. FE-SEM images of 50 nm films annealed at 500 C for 1 hr: (a) Cu/HfNx/Ge and (b) Cu/Hf-Ge-N/Ge. Figure 6-4. FE-SEM images of 50 nm films annealed at 600 C for 1 hr: (a) Cu/HfNx/Ge and (b) Cu/Hf-Ge-N/Ge.

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71 Figure 6-5. X-TEM images of 50 nm films annealed at 600 C for 1 hr: (a) Cu/HfNx/Ge and (b) Cu/Hf-Ge-N/Ge.

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72 Figure 6-6. EDS depth profile of 50 nm thick Cu/HfNx/Ge annealed at 600 C for 1 hr.

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73 Figure 6-7. EDS depth profile of 50 nm thick Cu/Hf-Ge-N/Ge annealed at 600 C for 1 hr.

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74 Figure 6-8. XPS chemical state data of Cu 2p peak at various sputtering times for 50 nm thick film of Cu/Hf-Ge-N/Ge (a) as deposited and (b) annealed at 600C for 1 hr.

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75 CHAPTER 7 EFFECT OF Ge OVERLAYER ON THE DI FFUSION BARRIER PROPERTIES OF HfNX Introduction Cu is gradually replacing Al as an interc onnect material in integrated circuits due to its lower resistivity and higher electromigration resistance.104 The limitations on materials properties will become more stringent with shrinking device dimensions as we move towards the 45 nm device node technology.105 For example, the current density and the associated heat generation increase rapidly as more devices are packaged on a single chip. The properties and behavior of in terconnects become the limiting factor in determining the circuit speed under these extreme circumstances. Although Cu is preferred in metallization schemes for resistiv ity reasons, it rapidly diffuses in Si creating deep level traps and reacts with dopant atoms, deteri orating device performance.106-108 There is thus a need to find a diffusion barrier material for Cu metallization. Ideally, a single crystal diffusion barrier wi th no defects is highly desi rable due to lack of grain boundaries. Growth of single crysta l barriers, however, is not practical due to the process restrictions (e.g. thermal budget) and materi al properties (e.g. lattice parameter and thermal expansion coefficient mismatches betw een Cu and the underlying substrate (Si or low material). An amorphous system is then preferred due to lack of grain boundaries, which provide facile pathways for diffusi on of Cu. Avoiding recr ystallization and the formation of grain boundaries is thus necessa ry for the diffusion barrier to function. Refractory metal nitrides are being studied in tensely as diffusion barriers for Cu because

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76 of their high melting temperatures. Some of th e refractory metal nitrides being considered are WNx,42,109,110 TaN,111-113 and TiN.114-117 Diffusion barriers of these materials fail at relatively low temperature (500 to 600 oC) due to recrystalliza tion. Ternary diffusion barriers such as W-Si-N,55 Ti-Si-N,118-119 W-B-N,120 and W-Ge-N80 have better diffusion barrier properties owing to higher recrysta llization temperatures compared to their respective binary systems. Recently HfNx 78,121 has received significan t interest as a diffu sion barrier material due to its high melting temperature (3330 oC),101 which translates to higher recrystallization temperatures., Failure of di ffusion barriers typically takes place due to recrystallization of the barrier film giving rise to para sitic grain boundaries. The integrity of the diffusion barrier can be enhanced in several ways, for example, by adding a third element into the binary matrix,122 choosing a refractory nitride system having very high melting temperature, or by combination of mate rials that collectively hinder Cu diffusion. In this paper it is demonstrated that the third approach is a viable one to formulating a diffusion barrier for Cu metallization. Specifi cally, a novel bilayer diffusion barrier of Ge (25 nm)/HfNx (7 nm thick) on Si is tested as a barrier for Cu. The candidate diffusion barrier is compared with a simple HfNx layer deposited under identical conditions. The results indicate that the integrity of th e diffusion barrier was maintained even under strenuous test conditions. Experimental Details Ge/HfNx diffusion barrier films were deposited on p-Si (001) si ngle crystal wafers by a reactive sputtering process. The native ox ide on Si was first etched by dipping the substrate in 7:1 buffered oxide etchant and rinsing with D.I. wa ter. After drying the substrates with ultra high purity N2 gas, the substrates were loaded into the deposition

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77 chamber via a load-lock. The chamber base pressure was maintained at 3 x 10-7 Torr. The sputtering targets were pre-cleaned before deposition by flowing Ar gas inside the chamber at a fixed chamber pressure of 15 mT orr. The forward sputtering power used for Hf and Ge targets was 200 and 100 W, respectively. HfNx diffusion barrier films were deposited by flowing Ar + N2 gas inside the chamber at a fixed pressure of 10 mTorr. Once the desired thickness of HfNx was achieved, the N2 flow was stopped and a Ge over layer was deposited via Ar ion sputte ring. The thicknesses of the HfNx and Ge films were maintained at 7 nm and 25 nm, respectively, by varying deposition time. Cu thin films were deposited in-situ on the Ge layer at a fixed pressure of 5 mTorr without breaking vacuum. Forward sputte ring power used for Cu was 200 W. The thickness for the Cu film was 300 nm for all sa mples. The substrate was rotated at 20 rpm throughout each deposition process to ensu re film uniformity. Subsequent to Cu metallization, individual diffusion barrier samples were annealed separately in a tube furnace in the temperature range 400 to 700 oC for 1 hr to test the integrity of the diffusion barriers. Ultra high pur ity Ar gas was used to purge the tube furnace at a flow rate of 65 sccm for at least 10 hr. Indi vidual samples were also annealed at 500 oC for 1 to 3 hr. The film crystallinity and intermeta llic phase formation were determined using a Phillips APD 3720 X-ray diffractometer (X RD) while the Cu depth profile was determined using a JEOL Superprobe 733 energy-dispersive spectrometer (EDS). The integrity of the interface wa s assessed by examining cros s-sections on a JEOL 2010F high-resolution transmission electron microsc ope (HRTEM). Samples were prepared using Dual beam strata DB235 focused ion beam (FIB) and the surface roughness of Cu was determined using atomic force microscopy (AFM) on a SPM/AFM Dimension 3100.

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78 Results and Discussion Thick HfNx films (thickness 300 nm) without a Ge layer were first grown with an apparent amorphous structure in the as-de posited condition since no diffraction peak was observed that was assignable to HfNx and the underlying Si (4 00) reflection at 69.155 was evident (not shown). The HfNx (thickness 300 nm) remained amorphous even after annealing at high te mperature (400 – 700 oC) (not shown). Figure 1 shows the X-ray diffraction patterns of thinner HfNx (7 nm) diffusion barrier films with (Fig. 1a) and without Ge (Fig. 1b). The sample with only a 7 nm thick HfNx film, however, failed after annealing at 400 oC for 1 hr as evident from the formation of Cu3Si (Figure 1b). By providing an over layer of Ge, the efficacy of HfNx was significantly enhanced (Figure 1a). The integrity of the Ge/HfNx stack remains intact up to 600 oC annealing temperature indicating a significant increase in its perfor mance as a diffusion barrier. Only after annealing at 700 oC are copper silicide p eaks evident, indicating improved performance of the bilayer diffusion barrier until 600 oC. For Ge/HfNx films annealed at 500 oC in the range of 1 to 3 hrs (Figure 1c), the diffusion barrier stack does not show any indication of Cu transport across it. The integrity of the Ge/HfNx stack at these severe test conditions is very encouraging. This excellent performance can be attributed to a combined effect of both Ge and HfNx. Ge was deposited as an over layer because of its rapid reactivity with Cu to form Cu3Ge, which also has a hi gh electrical conductivity.23,103,123 Moreover, the Cu3Ge is less reactive with oxygen86 than Cu3Si and itself is a good diffusion barrier for Cu.95 Using Ge as a stand alone diffusion barrier for Cu, however, fails at 500 oC as evident by formation of copper silicide peaks (Fig. 1d). When Ge is used as an over layer, it reacts with Cu to form Cu3Ge, which hinders Cu diffusion. After passing through the Cu3Ge layer, Cu

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79 encounters the amorphous HfNx layer, which further limits its diffusion. This is an excellent example of a synergistic effect of two different materials in achieving a common goal. The Cu intensity peak ratio be tween Cu (111) and (200) increases as the temperature is increased. It also increases when annealed at constant temperature but with increasing anneal time. This apparent recr ystallization of Cu to yield a preferred orientation of Cu along the [111] direction is advantageous as Cu has superior resistance to electromigration in the [111] direction. Figure 2 shows the atomic depth profiles in Cu/Ge/HfNx/Si film stacks before and after annealing at 500 oC. As clearly seen in Figure 2a for the as-deposited stack, there is a noticeable overlap of the Cu intensity profile with the Ge over layer. This supports the hypothesis that Cu reacts with Ge even at r oom temperature. The EDS peak positions of Cu and Hf overlap each other. As a result, a bump in the Cu intensity profile is observed in the HfNx layer. Also, a strong Hf EDS intensity is seen from the entire region of the Cu layer due to the EDS detector ’s inability to differentiate between Cu and Hf. The Cu intensity profile declines sharply as it m oves inside the Si substrate indicating no Cu diffusion into the Si at room temperature. For samples annealed at 500 oC for 1 hr (Fig. 2b) and 3 hr (Fig. 2c), the Cu intensity profile again decreas es sharply as the scan moves into the Si substrate. This is consistent with the XRD data, which also show no sign of Cu diffusion and formation of copper silicides. The oxygen content as measured by auger electron spectroscopy in HfNx films was approx. 20 at. %. Th is can be rectified by using lower base pressure and longer target pre-sputtering times. The surface roughness was quantitatively m easured by atomic force microscopy. Figure 3 shows AFM plots of the same Cu/Ge/HfNx/Si stacks profiled in Figure 2. As

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80 evident in these images, there is no signi ficant change in the rms roughness values between the as-deposited samples (5. 182 nm) and samples annealed at 500 oC for 1 hr (9.373 nm) and 3 hr (6.476 nm). The data suggest that the Cu layer is fairly smooth with the small amount of roughness likely produced by the reaction of the Cu with Ge. Subsequent annealing does not promote fu rther reaction or surface atom migration. The film interface integrity and abrupt ness were determined by high-resolution TEM for the same 3 samples analyzed by EDS depth profiling and AFM. As shown in Figure 4 the HfNx/Si film interface appears to be very abrupt for all three samples, which suggests no intermixing between the layers. This is consistent with the XRD data, which show no evidence of copper silicid e formation for annealing at 600 oC for 1 hr and at 500 oC for 3 hr. Conclusion In summary, the effectiven ess of a novel Ge(25 nm)/HfNx(7 nm) bilayer to serve as a diffusion barrier for Cu was investigated XRD data suggests that the barrier stack does not fail even after annealing at 600 oC for 1 hr and shows no si gns of copper silicide formation even after annealing at 500 oC for 3 hr, indicating excellent diffusion barrier quality of the Ge/HfNx bilayer structure. The as-deposited Cu/Ge/HfNx/Si (001) stack as well as ones annealed at 500 oC for 1 and 3 hr were further examined by EDS, AFM, and HRTEM. Elemental EDS profiles show a sharp decline in the Cu intensity profile at the HfNx/Si interface, indicating no Cu diffusion into the Si substrate. The results of these measurements also suggest th at Cu reacts at room temperature with the Ge to form Cu3Ge. The Cu film surface was smooth in the as-deposited condition with no significant changes after annealing at 500 oC for 3 hr. This is benefici al as smoother films exhibit lower contact resistance. Lattice images of the HfNx/Si interface reveal atomic

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81 abruptness, indicating excellent stability of the diffusion barrier and showing no signs of intermixing. Annea ling of a single HfNx layer barrier, however, showed Cu silicide formation at 400 oC. Overall, the 32 nm Ge/HfNx bilayer stack has excellent diffusion barrier properties for Cu metallization due to synergistic behavior of two different material systems.

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82 Figure 7-1. X-ray diffracti on patterns of as-deposited film s and annealed ones at the temperature shown in the figure: a) Cu/Ge(25nm)/HfNx(7nm)/Si (001) for 1 hr: b) Cu/HfNx(7nm)/Si (001) for 1 hr: c) Cu/Ge(25nm)/HfNx(7nm)/Si annealed at 500 oC for increasing tim e durations and d) Cu/Ge(25nm)/Si (001) for 1 hr.

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83 Figure 7-1. (cotd.).

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84 Figure 7-2. EDS depth profile of Cu/Ge(25nm)/HfNx(7nm)/Si for a) as-deposited: b) annealed at 500 C for 1 hr: and c) annealed at 500 C for 3 hr.

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85 Figure 7-3. AFM images of Cu/Ge(25nm)/HfNx(7nm)/Si for a) as-deposited: b) annealed at 500 C for 1 hr: and c) annealed at 500 C for 3 hr.

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86 Figure 7-4. HRTEM images of Cu/Ge(25nm)/HfNx(7nm)/Si for a) as-deposited: b) annealed at 500 C for 1 hr: and c) annealed at 500 C for 3 hr.

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87 CHAPTER 8 PROPERTIES OF Ta-Ge-N AS A DI FFUSION BARRIER FOR Cu ON Si Introduction As the minimum feature size in integrated circuits continues to decrease, there is a need to replace Al in interconnects du e to the limited conductivity and poor electromigration performance of Al under hi gh current densities. Cu is a viable alternative for replacing Al due to its 1.7 tim es lower resistivity compared to Al and better electromigration propertie s under high current densities.71,124 In addition, the RC time delay achieved by an Al/SiO2 stack needs to be lowered significantly to increase the device speed. The RC time constant can be si gnificantly lowered by combining Cu with a low-k material thus enhancing the performance of the device.3 Although the aforementioned properties of Cu are advantageous for device manufacture, Cu diffuses very effectively through Si, SiO2 and Ge, degrading the electrical properties of the device.104 It is an acceptor in Ge introducing traps at Ev + 0.04 eV, Ev + 0.32 eV and Ec – 0.26 eV.74 Cu is a donor in Si and creates traps from 0.2 to 0.5 eV above the valence band.125 It also reacts with Si to fo rm parasitic copper silicides at the interface. In addition, Cu reacts with dopant s in Si to form complexes that affect the device characteristics.95 Thus, use of Cu as an interconnect for future technology nodes will require a viable diffusion barrier. There ha s been a significant interest in refractory nitrides such as WNx, TiN, HfNx, and TaN61,110,114-116,121,126-128 as diffusion barrier candidates for Cu metallization. These diffu sion barriers, however, typically fail at moderate temperature (400-600 oC) limiting their use. The primary mode of failure for

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88 these diffusion barriers is by Cu diffus ion through grain boundaries formed by recrystallization of the ba rrier material upon annea ling. By increasing the recrystallization temperature, grain boundary formation can be delayed, thereby increasing the robustness of the diffusion ba rrier. Addition of a third element to the binary matrix induces amorphization at room temperature and also delays the recrystallization process, ma king ternary solutions intere sting as diffusion barrier candidates. Some of the ternary nitride di ffusion barriers being st udied include W-Si-N, Ta-Si-N, and W-Ge-N.99,81,129 In this paper, we report on the barrier layer properties when Ge is added to TaN. Ge is of interest because it di splays chemical behavior simila r to that of its congener Si and might be compatible with future Ge and SiGe based devices. Diffusion barrier properties of Ta-Ge-N were compared with TaN deposited under identical conditions. The results indicate that a Ta-Ge-N diffusion barrier fails at a higher temperature than TaN, suggesting superior di ffusion barrier properties. Experimental Details Ta-Ge-N diffusion barriers were depos ited on p-Si (001) wafers by reactive sputtering process at room temperature. Prior to deposition, the wafer was etched in 7:1 buffered oxide etch to remove its native oxide and then rinsed with deionized water. The substrate was loaded in the sputtering chamber that is maintained at 3 x 10-7 Torr base pressure. Sputtering targets were pre-sputtered before deposition at an Ar pressure of 15 mTorr to remove any contamination on the su rface. The forward power used for Ta and Ge was 200 W and 100 W, respectively. The di ffusion barrier films were then deposited by flowing Ar and N2 at a chamber pressure of 10 mTorr. For comparison, TaN diffusion

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89 barriers were deposited under identical ci rcumstances. The diffusion barrier film thickness was maintained at 50 nm in all cases. Cu was then deposited on the nitride in-situ at room temperature without breaking vacuum. Pre-sputtering of the Cu target was ca rried out to remove any contaminants prior to deposition. The forward sputtering power for Cu was 200 W. Cu was deposited on the diffusion barrier by flowing UHP Ar gas at a chamber pressure of 5 mTorr. The thickness of Cu was maintained at 300 nm in all cases. Substrates were rotate d at 20 rpm during all depositions to maintain film uniformity. Individual samples were then annealed separately in a tube furnace in the temperature range 400 to 700 oC for 1 hr. Before annealing, the furnace tube was purged by flowi ng Ar gas at 65 sccm for at least 10 hrs. The diffusion barrier analysis was co nducted using a Phill ips APD 3720 X-ray diffractometer (XRD) to detect any intermet allic phase formation between Cu and Si, which would indicate bulk Cu diffusion in th e substrate leading to barrier failure. The chemical depth profile of Cu through the diffusion barrier was determined by AES Perkin-Elmer PHI 660 Scanning Auger Multip robe (AES). Film surface morphology was evaluated by using a JEOL JSM 6335F fieldemission scanning electron microscope (FESEM). Sheet resistance of Cu wa s measured by four-point probe. Results and Discussion Figures 1 shows the X-ray diffraction pa tterns for Cu/TaN/Si and Cu/Ta-Ge-N/Si both as-deposited and afte r annealing in the temp erature range 400 to 700 oC for 1 hr. As evident from Figure 1, both TaN and Ta-G e-N films are amorphous in as-deposited condition with only Cu (111) and (200) peaks present. Upon annealing at 400 oC in Ar atmosphere for 1 hr, sufficient Cu diffuses through TaN into the Si substrate to form copper silicide (Fig. 1a), indicating diffusion barrier failure. Cu film crystallinity

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90 increases after 400 oC annealing, as evidenced by the sharpening of the Cu (111) and (200) peaks. Further annealing at higher te mperature induces bulk Cu diffusion into the Si substrate as evidenced by a decrease in the intensity of the Cu peaks after 500 oC annealing. After 600 oC, the Cu peak finally disappear s with a corresponding increase in copper silicide peak intensity indicating complete Cu diffu sion through the TaN diffusion barrier. In comparison, for the Ta-Ge-N diffusi on barrier there is no formation of copper silicide peaks after annealing at 400 oC indicating superior di ffusion barrier properties (Fig. 1b). Cu film crystallinity increases as evident by an increase in sharpness of Cu (111) and (200) peaks. Only after anneali ng the Ta-Ge-N diffusion barrier films at 500 oC, does the observation of copper silicide peaks indicate barrier failure. Subsequent annealing at higher temperature re sults in behavior similar to th at of TaN. The addition of Ge to the binary TaN matrix, however, raises the diffusion barrier failure temperature by at least 100 oC, resulting in better diffusion barri er properties of Ta-Ge-N films. The film morphology was examined by field-emission scanning electron microscopy. The film surface morphology was f eatureless and identical in as-deposited condition (not shown). After annealing at 400 oC, however, delamina tion of Cu films on TaN is evident as seen in Figure 2a. The delamination was sufficiently severe that Cu peeling occurred from the TaN diffusion barrie r even by soft nitrogen blowing. In the case of Ta-Ge-N (Fig. 2b), there was no dela mination of Cu consistent with better adhesion properties. Thus, addition of Ge to TaN could eliminate the need for the Ta layer that is normally used to prevent de lamination of Cu from TaN, thus saving a process step.

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91 The chemical depth profile of Cu through the diffusion barrier was determined by Auger electron spectroscopy. Fi gure 3 shows the Auger depth profile for Cu/TaN/Si asdeposited and after annealing at 400 oC for 1 hr. The as-deposited depth profile shows distinct interfaces at Cu-TaN and TaN-Si layers indicating no Cu diffusion. After annealing at 400 oC, however, Cu signal is seen th roughout the diffusion barrier and into the substrate suggesting signifi cant Cu diffusion and thus TaN barrier failure. This supports the XRD data in which copper silic ide peak formation is evident after 400 oC annealing (Fig. 1a). Figure 4 shows the Auger depth profile for the Cu/Ta-Ge-N/Si structure as-deposited an d after annealing at 400 oC for 1 hr. In contrast to TaN, the Cu/Ta-Ge-N and Ta-Ge-N/Si interfaces are qui te distinct both in as-deposited condition and after annealing at 400 oC for 1 hr. The Cu profile is identical for Figures 4a and 4b indicating no Cu diffusion through the diffusion barrier. The lack of evidence for copper diffusion in the depth profile is consistent with the XRD data that show no copper silicide peak formation even after annealing at 400 oC. Cu sheet resistance was measured by fou r-point probe. Figure 5 shows the sheet resistance of Cu for TaN and Ta-Ge-N in as-deposited and after high temperature annealing conditions. The sheet resistance of Cu decreases af ter annealing until 300 oC or higher for both TaN and Ta-Ge-N. This coul d be due to grain growth and increased crystallization of Cu after annealing. Incr ease in grain size causes decrease in grain boundaries, which contributes to lower el ectron scattering. Upon annealing at 500 oC (Figure 5), the sheet resistance of Cu increase s drastically. This increase is consistent with Cu diffusion and formation of coppe r silicides as eviden t by X-ray diffraction patterns.

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92 Conclusion In summary, the properties of Ta-Ge-N thin films as diffusion barriers for Cu have been investigated. X-ray diffraction patterns showed that the addition of Ge to the binary TaN matrix causes the diffusion barrier failure temperature to increase by at least 100 oC as compared to TaN. Delamination doe s not occur for Cu deposited on Ta-Ge-N while significant delamination is observed for Cu deposited on TaN as evidenced by FESEM. The AES depth profile showed signi ficant Cu diffusion through TaN diffusion barriers and into the Si subs trate after annealing at 400 oC, while no Cu diffusion occurs at similar temperature th rough the Ta-Ge-N diffusion ba rrier suggesting superior diffusion barrier properties of Ta-Ge-N thin films. The sheet resistance of Cu initially decreased with increase in annealing temperature for bot h TaN and Ta-Ge-N suggesting Cu grain growth. With further temperature increase, Cu diffusion eventually occurs through both diffusion barriers leading to form ation of highly resis tive copper silicides with a corresponding increase in the Cu sheet resistance.

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93 Figure 8-1. X-ray diffracti on patterns of as-deposited and annealed at the temperature shown in figure of a) Cu/TaN/Si (001) for 1 hr: b) Cu/Ta-Ge-N/Si (001) for 1 hr.

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94 Figure 8-2. Field-emission SEM images after annealing at 400 oC for 1 hr for a) Cu/TaN/Si (001) and b) Cu/Ta-Ge-N/Si (001).

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95 Figure 8-3. AES depth profile of Cu/TaN/Si (0 01) for a) as-deposited and b) annealed at 400 C for 1 hr.

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96 Figure 8-4. AES depth profile of Cu/TaGe-N/Si (001) for a) as-deposited and b) annealed at 400 C for 1 hr.

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97 Figure 8-5. Sheet resistance of Cu vs. annealing temperature for TaN and Ta-Ge-N.

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98 CHAPTER 9 CONCLUSIONS There is an imminent need for a viable diffusion barrier for copper metallization for future technology nodes. In general, ternar y refractory nitride diffusion barriers are superior to their respective binary nitrides for diffusion barrier applications. The major reason for increased functionality of ternary diffusion barriers is because of their increased recrystallization temperatures. By adding a third element in the binary matrix, the crystal lattice is disrupted causing an in crease in recrystallization temperatures. The initial part of the research was focused on i nvestigating the recrysta llization behavior of W-Ge-N diffusion barriers deposited on Si/SiO2 substrates and comparison of relevant diffusion barrier properties with WNx films. By adding Ge to the binary matrix of WNx, the crystal structure was di srupted causing the recrysta llization temperature of WNx to increase from room temperature to 600 oC after annealing for 1 hr. as seen from the X-ray diffraction results. The chemical depth profile of Cu as determined by auger electron spectroscopy supported the superior behavior of W-Ge-N films for barrier applications as compared to WNx. The decrease in resistiv ity of both W-Ge-N and WNx films can be attributed to grain growth with increase in annealing temperature and possible nitrogen evolution at higher temperat ures. The failure of ternar y diffusion barrier occurs by diffusion of copper through grain boundaries formed at higher temperatures. A similar study was conducted using pGe (001) substrates for possible integration in Ge or SiGe based devices The results again suggested that the recrystallization behavior of W-Ge-N films was increase d to higher temperatures by

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99 addition of Ge to WNx matrix. The depth profile as m easured using energy dispersive spectroscopy on cross-secti on TEM samples and auger el ectron spectroscopy suggests increased copper diffusion in WNx films as compared to W-Ge-N films after annealing at 500 oC for 1 hr indicating better diffusion barrie r properties of W-Ge-N films. The high oxygen content in the films could be due the background oxygen in the chamber. A comparative study of HfNx and Hf-Ge-N deposited on single crystal Ge substrates with varying thickness was carried out to evaluate their barrier properties and performance. HfNx was chosen for this study due to its high melting temperature. In general, high melting temperat ure relates to high er recrystallization temperatures. The study indicated similar recrys tallization behavior for bot h films although the diffusion barrier performance of HfNx was superior to Hf-Ge-N films. The reason can be attributed to Ge content in Hf-Ge-N films. As the Hf -Ge-N films were Ge-rich, copper diffuses and reacts with Ge very fast th rough the barrier and s ubsequently with the substrate causing barrier failure. The energy dispersive sp ectroscopy depth profile conducted on crosssection TEM samples suggested significan t copper diffusion in Hf-Ge-N films as compared to HfNx. The oxygen content in the films was high. Bonding information of Cu-Ge as investigated by X-ray photoelectron spectroscopy revealed Cu3-xGe formation as supported by X-ray diffraction patterns. The above studies were carried out by co-sput tering Ge target w ith either W or Hf targets. A study was conducted by forming a continuous overlayer of Ge on HfNx deposited on p-Si (001) substrates. C opper reacts readily with Ge to form Cu3Ge, which has good copper diffusion barrier properties a nd also has low resi stivity. The Ge/HfNx bilayer showed superior performance for copper diffusion barrier applications as

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100 compared to stand alone HfNx film when annealed at higher temperatures in Ar atmosphere for 1 hr. The Ge/HfNx film stack showed no signs of failure even after annealing at 500 oC for 3 hrs indicating its superior properties. The surface morphology and roughness of copper films showed no change after annealing. The Cu (111) to (200) XRD peak ratio increased with increasing annealing time at 500 oC. This is beneficial as Cu has excellent electromigration resistance in the [111] direction. The interface study as determined using high resolution transmission electron microscope revealed abrupt Sidiffusion barrier interface ev en after annealing at 500 oC for 3 hrs. A concluding study was conducte d by adding Ge to TaN binary matrix to evaluate the diffusion barrier performance. The recr ystallization temperat ure was increased by 100 oC as suggested by X-ray diffraction pa tterns. The surface morphology showed delamination of copper deposited on TaN but not when it is deposited on Ta-Ge-N films. The chemical depth profile suggested si gnificant copper diffusion through TaN as compared to Ta-Ge-N at similar temperat ures suggesting the robustness of Ta-Ge-N films. The sheet resistance of copper deposit ed on both films decreases with increased annealing temperatures suggesting grain gr owth at higher temperature. The sheet resistance increases rapidly when anneal ed at higher temperatures suggesting intermetallic Cu-Si phase formation. Based on the above work, the following conclusions can be drawn. 1. Addition of a third element in a binary refractory nitride matrix effectively frustrates the matrix and promotes amor phization at higher temperatures. This is evident from the X-ray diffraction pattern s of W-Ge-N films deposited on Si/SiO2 and Ge substrates. The results show that the recrystallization temperature of WNx films deposited on Si/SiO2 was increased from room temperature to 600 oC. The WNx films deposited on Ge were crystalline at room temperature. However, addition of Ge to the matrix caused amorphization and increased the recrystallization te mperature to 600 oC.

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101 2. Addition of Ge in the binary refracto ry nitride system improved the diffusion barrier performance in some materials su ch as W-Ge-N and Ta-Ge-N films. This is evident from the X-ray diffraction pattern in the said cases which shows failure of ternary diffusion barrier at higher temp eratures than their respective binary nitrides. The chemical depth profile s uggests significant copper diffusion through binary nitrides indicating barrier failur e. However, little or no copper diffusion occurs through ternary nitrides. In the case of HfNx films, addition of Ge did not improve the barrier performance as sugge sted by the X-ray diffraction patterns, energy dispersive spectroscopic analysis and cross-section transmission electron spectroscopic analysis. This was because of formation of Ge-rich Hf-Ge-N films which caused rapid reaction and diffu sion of copper through the film and subsequent failure. 3. The bilayer stack of Ge/HfNx is an attractive candidate for copper diffusion barrier applications. By formation of a Ge overlayer, a continuous film of Cu3Ge can be formed by its reaction with Cu on HfNx. The bilayer stack shows no signs of failure suggesting superior perfor mance even after undergoing strenuous barrier test conditions as evident from the X-ray diffraction, energy dispersive spectroscopy and cross-section TEM results. The selective orientation of Cu in the [111] direction shows promise due to its superior electrom igration resistance properties in that direction.

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PAGE 122

110 BIOGRAPHICAL SKETCH Seemant Rawal was born in Vadnagar, Gujarat, India, on 4th May 1975. He did his high school till 10th standard in Mumbai, India. Later, he moved to Mehsana, Gujarat, and completed his 11th and 12th standard education. He got admitted to the Metallurgical Engineering Department, M.S. University Baroda, from where he completed his undergraduate bachelor’s degree (B.E) in metallurgical engin eering in 1997 with distinction. After his B.E., he worked as application e ngineer in Electrotherm (I) Ltd. and later in Indabrator Ltd., two medium sized co mpanies manufacturing M.F. induction melting furnace and shot blasting machines, respectively. He was admitted to Materials Science a nd Engineering Department, University of Florida, in August 2001 and joined Dr. Rajiv Singh’s group. He worked on UV effect on nitrogen incorporation in high-k gate di electrics and non-melt laser annealing of Cu(InGa)Se2 solar cells. He joined Dr. Davi d Norton’s group in September 2003 and worked on transition metal doped perovskite. Later he changed his research topic and worked on ternary nitride diffusion barrier for Cu interconnects. The diffusion barrier project was supported by the National Science Foundation. Mr. Rawal and his advisor devised a novel diffusion barrier for Cu interconnects on Si using a bi-layer approach. The diffusion ba rrier was deposited by reactive sputtering in the nanofabrication laboratory at the University of Florida.


Permanent Link: http://ufdc.ufl.edu/UFE0017418/00001

Material Information

Title: Alternative nitride diffusion barriers on silicon and germanium for copper metalization
Physical Description: Mixed Material
Language: English
Creator: Rawal, Seemant ( Dissertant )
Norton, David P. ( Thesis advisor )
Singh, Rajiv K. ( Reviewer )
Anderson, Tim ( Reviewer )
McElwee-White, Lisa ( Reviewer )
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2006
Copyright Date: 2006

Subjects

Subjects / Keywords: Materials Science and Engineering thesis, Ph.D
Dissertations, Academic -- UF -- Materials Science and Engineering

Notes

Abstract: As device dimensions shrink, there is an urgent need to replace conventionally used Al interconnects to achieve increased current density requirements and better performance at future technology node. Copper is a viable candidate due to its lower resistivity and higher resistance to electromigration. However, Cu has its own problems in its integration. It diffuses rapidly through SiO2, Si and Ge into the active regions of the device, thereby deteriorating device performance. There is therefore a need to find a diffusion barrier for Cu. Amorphous ternary nitrides are investigated as a candidate diffusion barrier for Cu metallization on single crystal Si and Ge substrates. Ge was chosen as a third element in the binary matrix of WNx, HfNx and TaN due to its chemical similarity to Si and possible integration in SiGe or Ge based devices. The addition of Ge helps in amorphization of the binary matrix. It hinders grain boundary formation and increases the recrystallization temperature compared to their respective binary nitrides. Cu was deposited in-situ after nitride deposition. In the case of W-Ge-N on Si, the recrystallization temperature was raised by 400 oC, while for Ta-Ge-N it was raised by 100 oC indicating better diffusion barrier properties than their corresponding binaries. A bilayer approach for diffusion barrier was also studied. Ge/HfNx bilayer was deposited on Si followed by in-situ deposition of Cu. Copper reacts with Ge and forms Cu3Ge which has low resistivity, is less reactive with oxygen, and is a diffusion barrier for Cu. By combining the material properties of Cu3Ge and HfNx, an excellent diffusion barrier performance was demonstrated under stringent test conditions.
Subject: barrier, copper, diffusion, interconnects, metallization, nitrides, refractory, ternary
General Note: Title from title page of source document.
General Note: Document formatted into pages; contains 122 pages.
General Note: Includes vita.
Thesis: Thesis (Ph.D.)--University of Florida, 2006.
Bibliography: Includes bibliographical references.
General Note: Text (Electronic thesis) in PDF format.

Record Information

Source Institution: University of Florida
Holding Location: University of Florida
Rights Management: All rights reserved by the source institution and holding location.
System ID: UFE0017418:00001

Permanent Link: http://ufdc.ufl.edu/UFE0017418/00001

Material Information

Title: Alternative nitride diffusion barriers on silicon and germanium for copper metalization
Physical Description: Mixed Material
Language: English
Creator: Rawal, Seemant ( Dissertant )
Norton, David P. ( Thesis advisor )
Singh, Rajiv K. ( Reviewer )
Anderson, Tim ( Reviewer )
McElwee-White, Lisa ( Reviewer )
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2006
Copyright Date: 2006

Subjects

Subjects / Keywords: Materials Science and Engineering thesis, Ph.D
Dissertations, Academic -- UF -- Materials Science and Engineering

Notes

Abstract: As device dimensions shrink, there is an urgent need to replace conventionally used Al interconnects to achieve increased current density requirements and better performance at future technology node. Copper is a viable candidate due to its lower resistivity and higher resistance to electromigration. However, Cu has its own problems in its integration. It diffuses rapidly through SiO2, Si and Ge into the active regions of the device, thereby deteriorating device performance. There is therefore a need to find a diffusion barrier for Cu. Amorphous ternary nitrides are investigated as a candidate diffusion barrier for Cu metallization on single crystal Si and Ge substrates. Ge was chosen as a third element in the binary matrix of WNx, HfNx and TaN due to its chemical similarity to Si and possible integration in SiGe or Ge based devices. The addition of Ge helps in amorphization of the binary matrix. It hinders grain boundary formation and increases the recrystallization temperature compared to their respective binary nitrides. Cu was deposited in-situ after nitride deposition. In the case of W-Ge-N on Si, the recrystallization temperature was raised by 400 oC, while for Ta-Ge-N it was raised by 100 oC indicating better diffusion barrier properties than their corresponding binaries. A bilayer approach for diffusion barrier was also studied. Ge/HfNx bilayer was deposited on Si followed by in-situ deposition of Cu. Copper reacts with Ge and forms Cu3Ge which has low resistivity, is less reactive with oxygen, and is a diffusion barrier for Cu. By combining the material properties of Cu3Ge and HfNx, an excellent diffusion barrier performance was demonstrated under stringent test conditions.
Subject: barrier, copper, diffusion, interconnects, metallization, nitrides, refractory, ternary
General Note: Title from title page of source document.
General Note: Document formatted into pages; contains 122 pages.
General Note: Includes vita.
Thesis: Thesis (Ph.D.)--University of Florida, 2006.
Bibliography: Includes bibliographical references.
General Note: Text (Electronic thesis) in PDF format.

Record Information

Source Institution: University of Florida
Holding Location: University of Florida
Rights Management: All rights reserved by the source institution and holding location.
System ID: UFE0017418:00001


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ALTERNATIVE NITRIDE DIFFUSION BARRIERS ON SILICON AND
GERMANIUM FOR COPPER METALLIZATION














By

SEEMANT RAWAL


A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL
OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT
OF THE REQUIREMENTS FOR THE DEGREE OF
DOCTOR OF PHILOSOPHY

UNIVERSITY OF FLORIDA


2006

































Copyright 2006

by

Seemant Rawal

































Dedicated to my mother Rasilaben Rawal and my dear wife Purvi Rawal.
















ACKNOWLEDGMENTS

I would like to immensely thank my advisor Dr. David P. Norton. His enthusiasm

and knowledge on different subj ects of science are astounding. Dr. Norton' s exemplary

dedication towards his work and students along with his zeal to be the best in the world in

his area of research motivated me greatly to pursue and conduct superior quality research.

I would like to thank Prof. Rajiv K. Singh for providing me guidance and support during

my initial years in USA. I would like to thank Prof. Tim Anderson and Prof. Lisa

McElwee-White for their guidance and help with my research. They have helped me to

think and analyze my research critically.

I would like to thank my mother, Rasilaben Rawal, for her passion towards

education. Her undying zeal towards learning sowed the first seeds in me to respect and

acquire knowledge. Many thanks go to my high school chemistry teacher, Mr. Dubey and

physics teacher Mr. Parmar for providing support through my 12th standard. Special

thanks go to Mahendra Pandya and Geetaben Pandya, for supporting me throughout my

undergraduate years. I would like to thank my brother, Rakesh Rawal, and his wife,

Padmaj a Rawal, for their love and support. Most importantly, I would like to thank my

wife, Purvi Rawal, for her unconditional love and support. Her immense understanding

and perceptive nature have made my journey cherishing. I would like to thank all my

friends and well-wishers for their support and encouragement.

Last but not least, I thank GOD for giving me this life and opportunity to serve HIS

world.





















TABLE OF CONTENTS


page

ACKNOWLEDGMENT S .............. .................... iv


T ABLE S .............. .................... vii


LIST OF FIGURES ............ ....... ..............viii...


AB STRAC T ................ .............. xi


1 INTRODUCTION ................. ...............1.......... ......


2 LITERATURE REVIEW .............. ...............7.....


Refractory Metals as Diffusion Barrier .............. ...............7.....
Ti Diffusion Barrier............... ...............7.
Ta Diffusion Barrier ................. ...............8................
Cr Diffusion Barrier .............. ...............8.....
W Diffusion Barrier............... ...............8.
Binary Diffusion Barrier............... ...............9.
Refractory Intermetallics ................. ...............9.................
Refractory Carbides ................. ...............10.................
Refractory Nitrides ................. ...............12.......... .....
Ternary Diffusion Barriers .............. ...............15....


3 EXPERIMENTAL DETAILS AND CHARACTERIZATION ................. ...............18


Sputtering ............... ... .......... ...............18......
DC Magnetron Sputtering .............. ...............19....
RF Magnetron S puttering ................. ...._ ....___ ............2
Characterization ............ ..... .._ ...............21...

X-ray Diffraction ................. ...............21.................
Auger Electron Spectroscopy ................. ...............22........... ....
X-ray Photoelectron Spectroscopy ................... ............ ............... 23 .....
Scanning Electron Microscopy and Energy Dispersive Spectroscopy ................24
Atomic Force Microscopy ................. ......... ...............26......
Transmission Electron Microscopy ................. ...............26........... ....
Focused lon Beam ............... ... ............. ...............2
Van der Pauw Measurement and Four Point Probe ................. .....................28












4 PROPERTIES OF W-Ge-N AS A DIFFUSION BARRIER MATERIAL FOR
COPPER ................. ...............36.......... .....


Introducti on ................. ...............36.................

Experimental Details .............. ...............37....
Results and Discussion .............. ...............38....
Conclusion ................ ...............41.................


5 INVESTIGATION OF W-Ge-N DEPOSITED ON Ge AS A DIFFUSION
BARRIER FOR Cu METALLIZ ATION ....__ ....____ .......... .............4


Introducti on ............. ...... ._ ...............46...

Experimental Details .............. ...............47....
Results and Discussion .............. ...............48....
Conclusion ............. ...... ...............51...


6 COMPARATIVE STUDY OF Hf~Nx AND Hf-Ge-N DIFFUSION BARRIERS
ON Ge ................. ...............59.................


Introducti on ................. ...............59.................

Experimental Details .............. ...............60....
Results and Discussion .............. ...............62....
Conclusion ................ ...............66.................


7 EFFECT OF Ge OVERLAYER ON THE DIFFUSION BARRIER PROPERTIES
OF Hf~Nx .............. ...............75....


Introducti on ................. ...............75.................

Experimental Detail s .............. ...............76....
Results and Discussion .............. ...............78....
Conclusion ................. ...............80.......... .....


8 PROPERTIES OF Ta-Ge-N AS A DIFFUSION BARRIER FOR Cu ON Si ...........87


Introducti on ................. ...............87.____.......

Experimental Details .............. ...............88....
Results and Discussion .............. ...............89....
Conclusion ............. ...... __ ...............92....


9 CONCLUSIONS .............. ...............98....


LIST OF REFERENCES ............. ...... ._ ...............102...


BIOGRAPHICAL SKETCH ............. ......___ ...............110...
















TABLES

Table pg

3-1 Sputtering yields of different ions ........................... ........._ ...... 3

















LIST OF FIGURES


Figure pg

1-1 RC delay vs. technology nodes .............. ...............6.....

1-2 Void and Hillock formation in Al interconnects ................. .......... ................6

3-1 Schematic diagram of DC sputtering system with parallel plate discharge .............3 1

3-2 Schematic diagram of RF sputtering system with a capacitive, parallel plate
di schar ge ................. ...............3.. 1..............

3-3 Schematic diagram of X-ray diffraction set-up ................. .......... ................3 2

3-4 Schematic set-up principle of an atomic force microscope............... ...............3

3-5 Interaction output between a high-energy electron beam and thin specimen ..........33

3-6 Sample geometries for Van der pauw measurements .............. ....................3

3-7 Schematic diagram of Van der pauw configuration for measurement of RA and
RB TOSistances, respectively ................. ......... ...............34......

3-8 Schematic diagram of four point probe measurement ................. ............. .......3 5

4-1 X-ray diffraction patterns of as-deposited and annealed films at different
temperature of(a) WNx films (b) W-Ge-N fi1ms ........................... ...............42

4-2 AES depth profiles of Cu/WNx/SiO2/Si (a) as-deposited (b) annealed at 600 OC/1
hr (c) annealed at 800 OC/1 hr .............. ...............43....

4-3 AES depth profiles of Cu/W-Ge-N/SiO2/Si (a) as-deposited (b) annealed at 600
oC/1 hr (c) annealed at 800 OC/1 hr............... ...............44...

4-4 Resistivity vs. annealing temperature (a) for W-Ge-N and WNx films. Also
shown (b) is the resistivity as a function of sputter target power for the Ge and
W targets .............. ...............45....

5-1 X-ray diffraction patterns of as-deposited and annealed films at different
temperature of(a) WNx films (b) W-Ge-N films ........................... ...............53










5-2 AES depth profies of Cu/WNx/Ge (a) as-deposited (b) annealed at 400 OC/1 hr
(c) annealed at 500 OC/1 hr............... ...............54...

5-3 AES depth profies of Cu/W-Ge-N/Ge (a) as-deposited (b) annealed at 400 OC/1
hr (c) annealed at 500 OC/1 hr .............. ...............55....

5-4 ED S depth profie of Cu/WNx/Ge annealed at 500 OC/1 hr ................ ................. 56

5-5 ED S depth profie of Cu/W-Ge-N/Ge annealed at 500 OC/1 hr .............. .............57

5-6 X-TEM images of (a) Cu/WNx/Ge annealed at 500 OC/1 hr (b) Cu/W-Ge-N/Ge
annealed at 500 OC/1 hr ........._ ...... .___ ...............58...

6-1 X-ray diffraction patterns of Cu/Hf~Nx/Ge fi1ms in as-deposited and annealed
conditions for varying thickness of (a) 300 nm, (b) 50 nm, and (c) 15 nm .............68

6-2 X-ray diffraction patterns of Cu/Hf-Ge-N/Ge fi1ms in as-deposited and annealed
conditions for varying thickness of (a) 300 nm, (b) 50 nm, and (c) 15 nm .............69

6-3 FE-SEM images of 50 nm films annealed at 500 OC for 1 hr: (a) Cu/Hf~Nx/Ge
and (b) Cu/Hf-Ge-N/Ge .............. ...............70....

6-4 FE-SEM images of 50 nm films annealed at 600 OC for 1 hr: (a) Cu/Hf~Nx/Ge
and (b) Cu/Hf-Ge-N/Ge .............. ...............70....

6-5 X-TEM images of 50 nm films annealed at 600 OC for 1 hr: (a) Cu/HfNx/Ge and
(b) Cu/Hf-Ge-N/Ge .............. ...............71....

6-6 EDS depth profile of 50 nm thick Cu/Hf~Nx/Ge annealed at 600 OC for 1 hr...........72

6-7 EDS depth profile of 50 nm thick Cu/Hf-Ge-N/Ge annealed at 600 OC for 1 hr.....73

6-8 XPS chemical state data of Cu 2p peak at various sputtering times for 50 nm
thick film of Cu/Hf-Ge-N/Ge (a) as deposited and (b) annealed at 6000C for 1 hr.74

7-1 X-ray diffraction patterns of as-deposited films and annealed ones at the
temperature shown in the figure ....__ ......_____ .......___ ...........8

7-2 EDS depth profile of Cu/Ge(25nm)/Hf'Nx(7nm)/Si for a) as-deposited: b)
annealed at 500 OC for 1 hr: and c) annealed at 500 OC for 3 h............... ................84

7-3 AFM images of Cu/Ge(25nm)/Hf'Nx(7nm)/Si for a) as-deposited: b) annealed at
500 oC for 1 hr: and c) annealed at 500 OC for 3 hr .............. .....................8

7-4 HRTEM images of Cu/Ge(25nm)/Hf'Nx(7nm)/Si for a) as-deposited: b) annealed
at 500 oC for 1 hr: and c) annealed at 500 OC for 3 hr .............. .....................8










8-1 X-ray diffraction patterns of as-deposited and annealed at the temperature shown
in figure of a) Cu/TaN/Si (001) for 1 hr: b) Cu/Ta-Ge-N/Si (001) for 1 hr .............93

8-2 Field-emission SEM images after annealing at 400 oC for 1 hr for a) Cu/TaN/Si
(001) and b) Cu/Ta-Ge-N/Si (001)............... ...............94.

8-3 AES depth profile of Cu/TaN/Si (001) for a) as-deposited and b) annealed at 400
oC for 1 hr............... ...............95...

8-4 AES depth profile of Cu/Ta-Ge-N/Si (001) for a) as-deposited and b) annealed
at 400 oC for 1 hr............... ...............96...

8-5 Sheet resistance of Cu vs. annealing temperature for TaN and Ta-Ge-N ................97
















Abstract of Dissertation Presented to the Graduate School
of the University of Florida in Partial Fulfillment of the
Requirements for the Degree of Doctor of Philosophy

ALTERNATIVE NITRIDE DIFFUSION BARRIERS ON SILICON AND
GERMANIUM FOR COPPER METALLIZATION

By

Seemant Rawal

December 2006

Chair: David P. Norton
Major Department: Materials Science and Engineering

As device dimensions shrink, there is an urgent need to replace conventionally used

Al interconnects to achieve increased current density requirements and better

performance at future technology node. Copper is a viable candidate due to its lower

resistivity and higher resistance to electromigration. However, Cu has its own problems

in its integration. It diffuses rapidly through SiO2, Si and Ge into the active regions of the

device, thereby deteriorating device performance. There is therefore a need to find a

diffusion barrier for Cu.

Amorphous ternary nitrides are investigated as a candidate diffusion barrier for Cu

metallization on single crystal Si and Ge substrates. Ge was chosen as a third element in

the binary matrix of WNx, HfNx and TaN due to its chemical similarity to Si and possible

integration in SiGe or Ge based devices. The addition of Ge helps in amorphization of the

binary matrix. It hinders grain boundary formation and increases the recrystallization

temperature compared to their respective binary nitrides. Cu was deposited in-situ after









nitride deposition. In the case of W-Ge-N on Si, the recrystallization temperature was

raised by 400 oC, while for Ta-Ge-N it was raised by 100 oC indicating better diffusion

barrier properties than their corresponding binaries.

A bilayer approach for diffusion barrier was also studied. Ge/Hf~Nx bilayer was

deposited on Si followed by in-situ deposition of Cu. Copper reacts with Ge and forms

Cu3Ge which has low resistivity, is less reactive with oxygen, and is a diffusion barrier

for Cu. By combining the material properties of Cu3Ge and HfNx, an excellent diffusion

barrier performance was demonstrated under stringent test conditions.















CHAPTER 1
INTTRODUCTION

The integrated circuit industry doubles the number of transistors in a chip every

18 months. With more transistors built into a chip, minimum feature size decreases,

leading to faster devices. At present, we are moving towards the 45 nm node technology

and beyond which is guided by International Technology Roadmap for Semiconductors

(2005)2. For many years, gate length determined the device speed and was the bottleneck

for future generation devices. However, with decreasing dimensions, delay in

interconnects now plays a maj or role in influencing device speed. The interconnect delay,

also called resi stance-capacitance (RC) time constant delay, is mostly caused by a

increase in resistance at smaller dimensions. The RC delay is expressed by the following

equation:


RC = 2rDo
t, tRD

where p is resistivity, is interconnection length, t, is interconnect thickness, SED and

tlL are permittivity and thickness of interlevel dielectric (ILD). Also, the current density

increases as width of the interconnect decreases. Figure 1-1 shows the increase in RC

delay time in interconnects with decreasing feature size which dominates overall delay at

sub micron technology nodes.l

Aluminum has been the interconnect material of choice for many generations with

SiO2 (K = 3.9) as the dielectric material. This results in a tremendously high RC delay as

dimensions decrease. One way to reduce the RC delay is to decrease the dielectric









constant of SiO2. Fluorinated silicon dioxide (FSG) was used with a dielectric constant of

3.7 for the 180 nm technology node.2 Not until recently was low-k material (x = 2.7 -

3.0) being integrated in devices. At high current densities, electromigration performance

of Al is severely degraded leading to voids and hillock formation. Creation of voids leads

to discontinuous circuit causing failure of devices. Figure 1-2 shows void and hillock

formation in Al interconnect. Alloying Al with Cu showed an increase in

electromigration resistance, but it reached its usability limits as subsequent milestones in

IC industry were achieved.

With increased demands on performance, an alternative interconnect material is

eventually desired. The logical choice is a transition to Cu interconnects. There are

several advantages in using Cu as an interconnect material. Bulk resistivity of Cu is 1.67

CLD-cm, which is almost 40% less than Al (p=2.65 CLD-cm).3 Resistance to

electromigration property of Cu is highly superior to Al, which leads to better device

reliability. Replacing the Al/SiO2 gate stack with Cu/low-K reduces the RC delay

significantly .

However, Cu has its own share of problems. It diffuses rapidly through Si, Ge and

SiO2 and forms parasitic silicides (Cu3Si) and hinders the device performance. It forms

deep and shallow level traps in Si and Ge, respectively. Cu reacts with dopants and forms

Cu-D complexes (D is dopant atom) deteriorating device performance. Cu exhibits poor

adhesion on SiO2 and other low-K dielectrics. In addition to the above difficulties,

solutions also need to be found for anisotropic etching, poor corrosion and oxidation

properties of Cu. In order to get the desired benefits by switching from Al to Cu

interconnect, a viable diffusion barrier is required to prevent Cu diffusion.









The primary purpose of a diffusion barrier is to prevent intermixing of chemical

species with each other. This can be achieved via various mechanisms thereby classifying

the diffusion barriers based on the method. i.e., passive, sacrificial, stuffed and

amorphous diffusion barriers. A passive diffusion barrier is an ideal barrier which does

not react with any of the layers it separates. A sacrificial barrier would react with either

or both layers that it separates and get consumed. For a sacrificial barrier, the reaction

rate between the barrier and the layers is very important. It should be slow enough so that

the barrier can perform to acceptable levels for the useful lifetime of the device. For

polycrystalline or nanocrystalline thin films, rapid diffusion occurs via grain boundaries,

dislocations and surface defects.4 Diffusion is highly influenced by temperature. The

temperature dependence of diffusion coefficient (D) is given by the following equation:





whereDo is temperature-independent pre-exponential factor, Qd is the activation energy,

k is the Boltzmann constant and Tis the temperature.

The activation energy for grain boundary diffusion is approximately half the

activation energy required for lattice diffusion. As a result, grain boundary diffusion

dominates at lower temperatures.4 For polycrystalline thin films, diffusion through grain

boundaries and dislocations is the fastest. This mode of diffusion can be stopped using

two approaches. The grain boundaries and/or dislocations can be "stuffed" by impurity

atoms which would hinder diffusion through them. Another approach is to lower the

amount or eliminate the grain boundaries altogether. This can be achieved by low

temperature or room temperature deposition of nanocrystalline or a amorphous diffusion

barrier where the short-circuit pathways are lowered or eliminated. This is highly










desirable because of low/no addition to the thermal budget. This may not be the best

approach as films deposited at lower temperatures can have undesired modified

properties during subsequent high temperature processing steps.

Apart from the primary function of preventing diffusion of Cu, the diffusion

barrier should meet other stringent specifications. Some of them are:

The diffusion barrier should be thermodynamically stable with Cu and
underlying substrate under standard operating conditions. It should not
react with Cu or the substrate under thermal, mechanical or electrical
stresses encountered during other processing steps.

The density of diffusion barrier should be close to its bulk density in order
to avoid any defects, voids or dislocations which can compromise its
integrity .

The diffusion barrier should have reasonable thermal and electrical
conductivity to avoid any parasitic capacitance effects and unwanted
heating effects.

The contact resistance of diffusion barrier with Cu and substrate should be
minimal .

Diffusion barrier should behave well under the applied mechanical and
electrical stresses during subsequent processing.

Microstructure of the diffusion barrier should ideally be amorphous at
room temperature and remain amorphous after thermal treatment at higher
temperatures. In general, diffusion barriers with higher melting
temperatures have high recrystallization temperatures.

The barrier should have good conformality especially over high aspect
ratio structures.

Diffusion barrier should adhere well with Cu and the surrounding
material.

Deposition of diffusion barrier should be compatible with existing
processing facilities and infrastructures.









Chapter 2 will present background and a literature review of various diffusion

barriers, method of deposition and barrier properties. In Chapter 3, experimental methods

used to deposit diffusion barrier films for this study will be explained. Various

quantitative characterization techniques used to evaluate the diffusion barrier

performance under extreme thermal treatment will also be described. Recrystallization

properties and diffusion barrier performance of W-Ge-N deposited on Si/Sio2 aef

evaluated in Chapter 4. The behavior and performance of W-Ge-N diffusion barrier

deposited on Ge are investigated in Chapter 5. A comparison of barrier properties and

performance of HfNx and Hf-Ge-N deposited on Ge is described in Chapter 6. Chapter 7

will describe the excellent diffusion barrier properties of a Ge/Hf~Nx bi-layer diffusion

barrier deposited on Si. Properties of Ta-Ge-N diffusion barrier on Si will be evaluated in

Chapter 8 followed by conclusion in Chapter 9.











45 -

40~ nIM Cu 5m de~lly. Al nd 5502
r: r Umo de~lra Cu arid km ki2 Ij
35 O= Intezrronne t dly. alm 5102l~
SInterconned delay,.oc and lowr Ll.9.0





~~: ;F Undlow k




10 -I
Gate



650 500 320 250 180j 130 100
Genercallon mnra)


Figure 1-1. RC delay vs. technology nodes.


Figure 1-2. Void and Hillock formation in Al interconnects.















CHAPTER 2
LITERATURE REVIEW

Due to the unique diffusion property of Cu in Si, Ge, SiO2 and other low-K

dielectrics, there is an imminent need for a diffusion barrier for Cu. There has been a

plethora of research on different diffusion barriers. In general, metals or compounds with

high melting temperatures are suitable for the purpose because of less chance of having

grain boundaries, which are the fastest diffusion pathways. Therefore, refractory based

materials seem to be the best choice due to their high melting temperature and decent

electrical conductivity. A considerable amount of research has been conducted to find a

viable diffusion barrier for Cu based on refractory metals. Most of the research can be

classified in three categories.

Refractory metals as diffusion barriers
Binary diffusion barriers based on refractory metals
Ternary diffusion barriers based on refractory metals

This chapter will cover a comprehensive review of the above classifications.

Refractory Metals as Diffusion Barrier

As it became evident that switching from Al based interconnects to Cu based

interconnects was necessary, initial research focused on refractory metals as candidate

diffusion barriers. Primary interest was in Ti, Ta, Cr, and W as Cu diffusion barriers.

Ti Diffusion Barrier

Ti based diffusion barriers were extensively used for Al interconnects. A natural

choice was to extend its functionality in acting as a Cu diffusion barrier. However, a Ti

diffusion barrier fails at 350 oC by Ti-Cu compound formation.' A study conducted by









Ohta et al.6 also revealed failure of Ti diffusion barrier at 400 oC indicated by an increase

in resistivity after annealing. The resistivity increase was attributed to Cu diffusion

through Ti and subsequent reaction with Si to form Cu silicides.

Ta Diffusion Barrier

The two maj or advantages of using Ta as a diffusion barrier in comparison to Ti

are its very high melting temperature (3017 oC compared to 1668 oC for Ti) and

thermodynamical (interface and solubility) stability with Cu for very high temperatures.

The Cu-Ta phase diagram shows that both Cu and Ta are insoluble in each other even at

high temperature. Failure of Ta diffusion barrier occurs due to Cu diffusion through grain

boundaries formed in Ta after high temperature annealing followed by Cu3Si formation at

Ta/Si interface.' Ta also reacts with Si and forms silicides thus rendering itself unusable

as a Cu diffusion barrier.

Cr Diffusion Barrier

Cr was studied as a diffusion barrier for Cu because of its good corrosion

resistance properties and excellent adhesion promoter.8 However, Cu/Cr/Si multilayer

structure showed a huge increase in resistivity after annealing at temperatures higher than

400 oC. This was attributed to Cu-Cr binary phase formation. The formation of such a

phase compromises the integrity of the Cu/Cr/Si multilayer structure resulting in Cu

silicide formation indicating diffusion barrier failure.6

W Diffusion Barrier

W is a better diffusion barrier compared to Ti and Cr. It is chemically and

thermodynamically stable with Cu.9 Initial study showed that Cu/W/Si structures behaved

well even after annealing at 600 oC. However, it failed after annealing at 700 oC.6 CU

intermixes with W even at low temperatures of 260 oC, thus rendering it non-viable as a









solution to Cu diffusion.10 Preferentially oriented W(110) failed after annealing at 690 oC

for 1 hr due to consumption of W by a uniform silicidation reaction." Selective chemical

vapor deposited W layer on p -n junction diodes led to failure after annealing at

temperatures above 650 oC.12

Binary Diffusion Barrier

In order to overcome the drawbacks of lower recrystallization temperatures and

formation of grain boundaries in refractory metals used as diffusion barriers, their

respective binaries were explored. This class of diffusion barriers can generally be

classified in three categories.

Refractory intermetallics
Refractory carbides
Refractory nitrides

These individual classifications will be explored further in the following sections.

Refractory Intermetallics

Several of the refractory intermetallics were studied as Cu diffusion barriers. TiW

was used extensively in Al-based interconnects and so was also tried as a diffusion

barrier for Cu interconnects.13 The study showed that TiW (30 at. % Ti) not exposed to

air before deposition of Cu failed at 775 oC when RTA annealed in N2 gaS for 30 seconds.

The authors also summarized results of a variety of Cu/diffusion barrier/Si systems.

Table-1 in reference 13 shows a list of diffusion barrier investigated between Si and Cu.13

Amorphous Ni60N T40 diffusion barrier failed at 600 oC after annealing for 60 min. A

different composition of Ni57r\b42 failed at a lower temperature of 500 oC after 60 min.14

The authors also investigated the feasibility of Ni60Mo40 aS a candidate diffusion barrier

for copper. It failed at 500 oC after annealing for 60 min. Amorphous Ir-Ta alloy was

investigated for Cu diffusion barrier application." Sandwich structure of Si/a-Ir45Tass/Cu/









a-Ir45Tass failed after annealing at temperatures higher than 700 oC. Failure occurred at

750 oC by interdiffusion of Cu and Si as detected by Rutherford backscattering

spectrometry. It was also ob served that the recrystallization temperature of a-Ir45Tass was

lowered from 900 oC to 750 oC due to the presence of Cu. Alloyed Ta-Col6 was shown to

act as a Cu diffusion barrier deposited by e-beam evaporation. However, an intermetallic

phase of Co2Ta and metal-silicide phase of Co2Si formed after annealing at 500 oC.

Recently, a study conducted by J. S. Fang et al.17 explored the possibility of using

sputtered Ta-TM (TM = Fe, Co) as diffusion barriers for Cu. Tao~sFeo~s diffusion barriers

failed at 650 oC while Tao~sCoo~s failed at 700 oC after annealing, indicating better

performance of Tao~sCoo~s.

Refractory Carbides

Interest in refractory carbides to serve the purpose of copper diffusion barriers

was due to two reasons. First, they have a high melting temperature, and second, they

have low resistivity indicating higher thermal and chemical stability. Among the

refractory carbides studied, primary interest has been in Ti, Ta and W based carbides. A

detailed study of TiGx as copper diffusion barriers was carried out by S.J. Wang et al.l

However, the Cu/TiCx/Si structure was unstable after annealing at temperatures higher

than 600 oC for 30 min. A more sensitive electrical characterization done by measuring

leakage current across Cu/TiCx/p+ n-Si diode structure revealed an early failure just after

annealing at 500 oC for 30 min. As many as 60% of the diodes measured for leakage

current after annealing at 530 oC had the leakage current in the range of 10-5 A/cm2,

which is three orders of magnitude higher than the leakage current measured at room

temperature. Approximately 80% of the diodes had a leakage current in the range of 10-3









A/cm2 after annealing at 550 oC, indicating a failure temperature somewhere around 500

oC.

TaC has a melting temperature of 3985 oC19 Suggesting good thermal stability,21

and room temperature resistivity of 27 CLD-cm,20 thus making it a potential candidate as a

Cu diffusion barrier. In this study, Junji Imahori et al.20 COmpared films with different

carbon concentrations, namely Ta53 47, Ta40 60, and Ta20 80. Of these, Ta53 47 f11ms

performed better than others. However, it too failed after annealing at temperatures

higher than 600 oC for 30 minutes. Films with 60 and 80% C concentrations failed

primarily due to diffusion of Cu through the amorphous C phase and lesser diffusion of

Cu occurred through grain boundaries of TaC. However, the prime reason for failure of

Ta53 47 WAS due to the Cu diffusion through the grain boundaries with Cu activation

energy of 0.9 eV. A study of the interfacial reactions in the Cu/TaC/Si before and after

annealing was conducted to determine the failure mechanism of the barrier layer.22 The

failure temperature was a function of thickness as 7, 3 5 and 70 nm TaC films failed at

600 oC, 650 oC and 750 oC respectively. The thicker TaC film integrity was compromised

at a lower temperature of 600 oC due to the formation of amorphous Ta(O,C)x layer at the

Cu/TaC interface.

WCx was studied as a candidate diffusion barrier for Cu due to its high melting

temperature of around 2785 oC23 and low electrical resistivity.24 A hexagonal closed pack

W2C phase was achieved by chemical vapor deposition at a growth temperature of 600oC.

This growth temperature is fairly high keeping in mind future technology requirements

where incorporation of low-K materials will put an even more stringent temperature

constraint.25 Low temperature CVD process to deposit W2C WAS demonstrated by Y. M.









Sun et.al.26 TEM studies revealed 5-6 nm W2C CryStallites in an amorphous matrix with a

W/C ratio of 2:1 until 450 oC. The Cu/W2C/SiO2/Si stack remained intact after 400 oC

annealing for 8-9 hours. Unfortunately, higher temperature annealing was not done in the

study which would have revealed capacity of the diffusion barrier to withstand higher

thermal stress. Room temperature sputter deposited WCx was tested as a candidate

diffusion barrier for Cu on Si.24 Fifty nanometer WCx failed according to analytical

results after annealing at temperatures higher than 650 oC for 30 min. Electrical

measurements, however, lowered the failure temperature to 550 oC. At 700 oC, the

metallurgical stability of the WCx/Si interface is compromised. W reacts with Si to form

W5Si3 at the WCx/Si interface. Cu diffuses through the barrier layer and possibly through

defects in the barrier layer into the Si substrate to form Cu3Si phase. In as-deposited

condition, 60% of the measured diodes fail with a leakage current of 10-9 A/cm2. Also,

60% of the diodes fail after annealing at a temperature of 550 oC with a leakage current

of 10-' A/cm2.

Refractory Nitrides

The activation energy for Cu to diffuse through grain boundaries is low making

this the fastest diffusion pathway and killer of devices. A method to stop Cu diffusion is

to block these super highways by "stuffing" them. Oxygen was used as a stuffing agent in

TiN diffusion barriers used for Al metallization. Al reacts with 02 and forms an

aluminum oxide phase that hinders Cu from further diffusion. However, the same concept

could not be applied to Cu.3 Nitrogen is a good stuffing agent that can solve the purpose.

Excess nitrogen in the film moves out to the grain boundaries and stuffs them. Cu

diffusing through the grain boundaries experiences a repulsive force from the nitrogen

thus stopping it from diffusing through the barrier film.27 Refractory nitrides are









interesting candidates for Cu diffusion barrier application due to their lower resistivity,28

high melting temperatures, lower heat of formation indicating better stability and the

ability to block Cu diffusion by stuffing the grain boundaries by nitrogen. Some of the

binary refractory nitrides investigated for Cu diffusion barrier applications are Ti, Ta, W

and Hf based which will be discussed further in detail.

TiN was extensively used as a diffusion barrier for Al based metallization.

Considerable effort was made to utilize the existing knowledge and processing methods

for Cu based metallization as well. However, TiN deposited by both PVD and CVD

methods resulted in a columnar grain structure in which grain boundaries run through the

entire thickness of the barrier film.29-32 These columnar grain boundaries provide easy

pathways for Cu diffusion and subsequent reaction with Si to form silicides. The

properties of TiN films greatly depend on the deposition conditions which affect the

microstructure, density and other relevant barrier properties of the film. The resistivity

values range from 20 2000 CLD-cm and density ranges from 3.2 5.0 g/cm3 33 A

comparative study of different deposition methods and conditions to deposit TiN films

was done by Park et al.33 Porous films have lower densities making them susceptible for

impurities like oxygen which alter the diffusion barrier properties. The TiN films failed in

the temperature range of 500 750 oC depending on the film properties, 500 oC being for

CVD deposited films having lowest density and 750 oC being for sputter deposited films

having highest density. Atomic layer deposition techniques were also tried for TiN

barrier deposition with focus on the influence of microstructure, resistivity and impurity

content on film barrier properties.34 However, the barrier film failed after annealing at

500 oC for 1 hr. The maj or drawbacks of CVD based techniques are unwanted C and O









impurities, particulate generation, high deposition temperatures, and uncontrolled

thickness variations across the surface.

WNx is also an attractive candidate for Cu diffusion barrier applications due to its

low resistivity and the ability to be deposited as an amorphous phase. Due to the absence

of grain boundaries, Cu diffusion is hindered or slowed. However, the phase of WNx

mainly governs the barrier properties. A tungsten rich WNx (x < 0.5) phase tends to

dissociate at temperatures as low as 450 oC into W and W2N leading to barrier failure by

diffusion of Cu through the grain boundaries due to lower recrystallization

temperatures.35 A nitrogen rich WNx phase (x = > 1) tends to be amorphous and remain

amorphous for higher temperatures, but the gain is traded off with higher resistivity.

Uekubo et al demonstrated the feasibility of W2N as a diffusion barrier, as compared to

W and WN, and showed that 8 nm W2N was able to stop Cu diffusion in Si until 600 oC

for 30 min.36 Failure of the diffusion barrier was due to recrystallization and grain

boundary formation. An array of deposition methods have been tried to deposit WNx like

sputtering,37 inOrganic CVD,38 plasma-enhanced CVD,39 metal-organic CVD,40 and

ALD.41,42

TaN is being currently used for Cu diffusion barrier applications due to its high

melting temperature and thermal stability. Among the various phases of Ta-N,

stoichiometric TaN has a melting temperature of 3087 oC and heat of formation

(A~Hf= -120 kJ/mol)43 making it more stable than Ta2N which has a melting temperature

of 2050 oC and heat of formation of AHf = -98 kJ/mol. Chemical vapor deposition of

tantalum nitride results in an insulating tetragonal Ta3N5 phase which is not suitable for

diffusion barrier applications.44,45 A bi-layer structure of Ta/TaN is currently being used









to overcome the obstacle of adhesion problem of TaN. TaN does not adhere well to Cu46

but it does adhere with SiO2.47 In case of Ta, it adheres well to TaN but not with SiO2 and

is not an effective diffusion barrier by itself.47 Also, TaN helps nucleate Ta in the a-Ta

phase, which has a BCC structure and a resistivity of ~15-30 CLD-cm, as compared to

when Ta is directly deposited on Sio2 where it nucleates in the P-Ta phase, with a

resistivity of ~150-220 CpD-cm.47 However, with device dimensions shrinking,

conformality of high aspect ratio is very critical. PVD48 has been used until now for

deposition of liner and Cu seed layer, with ionized PVD (I-PVD)49 being the latest in the

technology that has been able to extend the functionality to lower dimensions. However,

with future technology constraints, the liner thickness should be less than 10 nm.5o Since

the overall liner-Cu seed layer thickness requirement decreases, either some maj or

modifications has to be done to the I-PVD process or a switch to a better conformal

process like atomic layer deposition (ALD) has to be adopted.

Ternary Diffusion Barriers

The activation energy required for Cu diffusion through the grain boundaries is

low. Binary diffusion barriers have been investigated in which one of the elements, like

nitrogen, "stuffs" the grain boundaries and disallows Cu diffusion by blocking. However,

at higher temperatures the binary compound recrystallizes, forming parasitic grain

boundaries that should be avoided. Another way to solve the problem of creating a

diffusion barrier for Cu is to form an amorphous matrix that acts as a diffusion barrier

and remains amorphous (not recrystallized) when annealed at higher temperatures. This

can be achieved by adding a third element into the binary matrix. This addition frustrates

the binary lattice structure and delays or avoids the recrystallization process when it is

annealed at higher temperatures. Considerable research has been done on amorphous









ternary nitride diffusion barriers for Cu metallization. Most of the research work has been

done by incorporating Si, B or C as the third element in the binary matrix of refractory

nitrides. Some of the work will be discussed below.

TiSixN, were initially studied as Ti based liners and used as diffusion barriers for

Al metallization. Amorphous TiSixN, is attractive due to the absence of grain boundaries.

The properties of diffusion barriers depend on the method of deposition and its chemical

composition. Chemical vapor deposited TiSixN, films (25 nm thick) were studied as

candidate diffusion barriers for Cu.51 The barrier failed after annealing at temperatures

higher than 700 oC for 30 min. Unfortunately, the barrier resistivity is too high (800 CLa-

cm) for applications. MOCVD deposited TiSiN films were also investigated as diffusion

barriers for Cu.52 The deposition steps involved a H2 N2 plaSma treatment to increase the

density of the film. Comparison of films that were plasma treated and not plasma treated

was done. Films were characterized by secondary ion mass spectroscopy (SIMS) and

Secco etch method for diffusion barrier failure. Secco etch revealed that failure occurred

in films that were not plasma treated after annealing at 550 oC for 1 hr. For plasma treated

films, failure occurred after annealing at 600 oC for 1 hr. SIMS detected Cu diffusion in

the sample that was not plasma treated at 450 oC, whereas Secco etch revealed etch pits

in the same sample after annealing at 550 oC.

TaSixN, (x=1.4, y=2.5) films with different thickness (t= 5-40 nm) were sputter

deposited on Si by Lin et al53 to evaluate their diffusion barrier characteristics. Electrical

measurements on a pn diode structure revealed failure of 5 nm TaSixN, after annealing

at 500 oC for 30 min. The TaSixN, structure remained amorphous even after annealing at

800 oC, which is a good indicator that addition of Si to the TaN matrix was effective in









keeping the lattice structure amorphous. Failure of the diffusion barrier was due to Cu

diffusion through localized defects into Si. The nitrogen content plays an important role

in determining the diffusion barrier characteristics. With the increase in the nitrogen

content, the diffusion barrier properties are enhanced as the nitrogen helps in preventing

formation of TaSi2 phase after annealing at higher temperatures in Ta-Si-N films.54

There has been some research on WSixN, diffusion barriers for Cu metallization.

LPCVD deposited WSixN, shows promise as a diffusion barrier compared to Ta and Ti

counterparts.55,56 Efforts are also made to take advantage of good conformal coverage

obtained using the CVD deposition process while achieving superior diffusion barrier

properties similar to those achieved by PVD methods.56 Researchers have worked on

using boron or carbon as a third element in the binary matrix." W-B-N seems to be a

better diffusion barrier compared to W-Si-N due to its lower resistivity while having the

same performance as W-Si-N.5














CHAPTER 3
EXPERIMENTAL DETAILS AND CHARACTERIZATION

Sputtering

Sputtering is a physical vapor deposition technique where incident ions remove

atoms (sputter) from the target surface by momentum transfer process. Sputtering was

discovered in the 19th century by Grove and Pulker and reported by Wright in 1877. The

number of atoms removed from the target surface by a single incident ion is called the

sputtering yield. The sputtering yield depends on the mass of the atoms (target) and ions,

bombardment energy of the incident ions, angle of incidence and binding energy of the

atoms in the target material. Table 3-1 shows the sputter yields of different ions when

bombarded on the target materials as indicated. Different classifications are used to

define a particular sputtering technique depending on the type of sputtering configuration

and also on the type of a reactive species inside the deposition chamber. As a result, there

are different sputtering techniques like DC sputtering, RF sputtering, triode sputtering,

magnetron sputtering and "unbalanced" magnetron sputtering. Depending on the absence

or presence of a reactive species, it can be called non-reactive or reactive sputtering,

respectively. The energy transfer taking place due to collision between two hard spheres

can be given as:

E, 4M~,MCos 6
E, (M, +M,)2










whereE,, E, are the energy of the target and incident particle respectively, M~,,M~, are

the mass of the target and incident particle and B is the angle of incidence with respect to

the surface normal of the target.

Currently, almost any target material can be deposited by sputtering due to the

latest advancement in technology. Co-sputtering can be carried out to deposit a

compound on the substrate by sputtering two targets simultaneously or by sputtering a

single target in a reactive atmosphere or by using a compound target in an inert

atmosphere. Some of the advantages of the sputtering deposition process are given below.

Ability to coat large areas with uniform thickness
Low-temperature deposition process
Flexibility of target materials choice including insulators and semiconductors
Target composition is replicated in the composition of deposited film.
Good adhesive property obtained in the deposited films
Sputtering can be up, down or sideways
Reactive sputtering possible
Reproducibility
Environmentally friendly process technology
Can be easily scale up for commercial production

The two maj or kind of sputtering configurations conventionally used are DC and RF

magnetron sputtering. They are described in detail below.

DC Magnetron Sputtering

In a DC sputtering process, the target is the cathode and either the substrate or the

chamber walls are the anode. A very high negative DC voltage (several kV) is applied to

the cathode. Ar gas is supplied in the chamber to create plasma. Ionized Ar' ions are

accelerated to the cathode, bombarding the surface and sputtering atoms from the cathode

surface. They also generate secondary electrons that ionize the gas atoms creating

increased amounts of Ar+ ions. This increases the probability of collisions at the cathode

surface. A circular magnet (magnetron) below the target traps the electrons emitted from









the cathode surface to remain near the cathode. These electrons hop in a cycloid fashion

on the target surface due to (E x B) forces acting on it. As the electrons are trapped near

the surface region, the probability of creating more Ar+ ions increases which increases the

sputtering yield. Another advantage of using the magnetron source is that low gas

pressures can be used inside the chamber to maintain a stable plasma. Since the gas

pressure is low, internal gas collisions decrease leading to higher yields of sputtering

target material. As there is less chance of internal gas collisions, the sputtered materials

impact the anode surface with a high kinetic energy. Figure 3-1 shows a schematic DC

sputtering chamber set-up. The disadvantage of DC sputtering is that the cathode material

needs to be conductive. Insulating materials cause a charge build-up near the surface

quenching the plasma.

RF Magnetron Sputtering

The maj or difference between DC and RF magnetron sputtering is the source. In

radio frequency (RF) sputtering, an RF source is used, typically 13.56 IVHz. A matching

network is used to optimize power transfer from the RF source to the discharge and a

blocking capacitor is used in the circuit to create a DC bias. With the use of an RF source,

sputtering of insulating targets has been made possible. Figure 3-2 shows a schematic set-

up of a RF sputtering. The sputtering process is similar to DC sputtering. Low gas

pressures can be used for RF sputtering. One consideration during RF sputtering is to use

target materials with high thermal conductivity. Low thermally conductive materials

develop a high enough thermal gradient due to the sputtering process to cause brittle

fracture of the target.









Characterization

Characterization is an important part of research. It helps in identifying physical,

electrical, magnetic, optical and chemical properties of the samples that helps in

understanding the fundamental behavior of the material being studied. Some of the

characterization techniques used in this study are discussed below.

X-ray Diffraction

X-ray diffraction (XRD) is one of the most versatile non-destructive technique

used todate to identify the crystalline phases and crystallinity of the sample. Advances in

X-ray diffraction and understanding the basic physics between the interaction of X-ray's

and the sample have given access to a plethora of information. Currently, information

regarding strain, in-plane epitaxy, defect structure, film thickness etc. can easily be

obtained from the constructive and destructive interference of X-rays with the sample.

Figure 3-3 shows a basic X-ray diffraction set up. Cu Ku radiation is generated

and impinged on the sample. This radiation undergoes constructive or destructive

interference after reflecting from the sample depending on the path difference between

the incident and reflected X-rays. Constructive interference will cause a peak at a

particular 26 angle according to Bragg's rule as given below:

nAl = 2dSine

where Ai is the wavelength of X-ray (Cu Ku radiation 1.54 A+), d is the distance between

two consecutive (hkl) planes and O is the angle between the incident X-ray and the

sample surface. In a cubic system the d-spacing is given as:


d,
hkih +k +1~


where ao is the lattice constant.









For a single crystal, there are only specific orientations that satisfy the Bragg's

law. Diffraction peaks from these planes appear in the diffraction pattern. However, for a

polycrystalline film having differently oriented grains, diffraction peaks appear when

those grains meet the diffraction conditions.

The full width at half maximum of an X-ray diffraction peak gives information

about the grain size. The relationship is called the Scherrer equation:

0.9il
BCose

where t is the grain size, ii is the wavelength of x-ray, B is the full width at half maximum

and O is the Bragg angle.

Auger Electron Spectroscopy

Auger electron spectroscopy is an excellent technique for surface and sub-surface

analysis. It is a highly sensitive technique as it probes only from few angstroms to few

nanometers in depth from the surface. Usually, an electron beam is used to probe and

excite electrons in the atoms of the sample. As the core shell electron leaves the atom, it

is in an excited state. A higher shell electron fills the core shell vacancy and the energy

emitted by this transition is transferred in exciting another electron to emit from the atom

which is called the Auger electron. Depending on the kinetic energy of the electron

measured, the binding energy of the emitted Auger electron is calculated by the following

equation:




where K.E is the kinetic energy of the Auger electron, and Ex- EL and EL are the energies

of the K and L shells, respectively.










A total of three electrons are needed in order to fulfill an Auger transition. As a

result, elements like H and He are undetectable with AES. Suppose an electron from the

K-shell of an element is removed by photons or electrons. The electron from the L shell

fills up this vacancy with enough energy left to emit an electron from the L shell. This is

called a KLL transition. Similarly, now higher shells would fill up the L shell and would

result in a LMM type of transition. The Auger equipment can be combined with an ion

gun, which would sputter the sample and expose a new surface to be examined. A group

of all those data points would give the depth profie characteristics of the sample studied.

This is therefore a very good method to study interfaces and diffusion profies of

elements inside the sample. The following information can be obtained from an Auger

profile.

Type of elements
Amount percentage of the elements
Chemical state of elements
Valence band density of states

X-ray Photoelectron Spectroscopy

As the name suggests, a high energy photon source (X-ray) is bombarded on the

surface of the sample. It ej ects a core shell electron in the atom transferring it to an

excited state. The atom returns back to ground state as the higher shell electron fills up

the core shell vacancy, emitting the excess energy in the form of photon or by emitting an

Auger electron. The kinetic energy of the emitted core shell electron is detected which

gives a plethora of information about the sample. The expression of kinetic energy

measured is as below:


K.E = hv B.E









X-ray photoelectron spectroscopy is a very sensitive technique for surface

analysis as it probes only few angstroms from the surface. This is because only few

electrons near the surface are able to escape the sample while most of the electron excited

by the x-ray lose their energy by collisions with atoms and are not able to exit the surface.

The flux of electrons able to exit the surface of the sample without being scattered (Id) is

given by the following expression:



,,=,,,(AlSine

where Io is the original flux of electrons generated at depth d, 3Ae is the inelastic mean free

path of electrons, 9 is the angle of electron emission.

As the binding energy of an electron is a characteristic of an atom, elemental

information can by obtained by XPS. Also, a general trend is that if the charge on an

atom increases, so does the binding energy of the electron. As a result, the valence state

of the atom can be known. Amount percentage of the element present can be determined

by peak intensity. The valence electrons participate in the bonding process and thus

compound information can also be known from XPS data analysis. A depth profile is

possible. However, it is very time consuming.

Scanning Electron Microscopy and Energy Dispersive Spectroscopy

Scanning electron microscopy is used for surface and cross-section imaging,

topographical information and compositional information of a binary alloy system. An

electron beam strikes the sample surface and various kinds of elastic or inelastic

interactions occur resulting in emission of electrons. The electrons emitted out of the

sample are of generally two types

*Secondary electrons









*Back-scattered electrons

When the incident primary beam of electrons impinge on the sample surface, they

can undergo inelastic scattering by colliding with electrons of the sample, transferring

some of its energy to them which exit the surface of the sample depending on the amount

of energy transferred. If sufficient energy is transferred, then electrons from the sample

exit the surface and are collected by the detector. Electrons with energy of less than 50

eV are classified as secondary electrons. The primary beam can also undergo elastic

scattering with the nucleus of the atom where transfer of energy is null or very small.

These electrons have high energy and can exit the surface easily. The amount of back-

scattered electrons depends on the atomic number of the atom. A higher atomic number

(Z) yields a higher amount of back-scattered electrons which results in increased

brightness of the image. However, for a constant Z, the back scattering yield remains

unchanged if the primary beam energy is above 5 keV. The secondary electron yield does

not depend much on the Z. The backscattering and secondary electron images can be

used to complement each other in seeking information.

Since the primary beam has high energy, it can remove core electrons from the

atom thus exciting it. The atom goes back into the ground state by filling the core

vacancy by an electron from a higher orbital. The excess energy can either be emitted in

the form of a photon or an Auger electron. The emitted photon can be detected and used

for further study depending on its energy. This method is called as energy dispersive

spectroscopy (EDS). Since the emitted photon is characteristic of the atom from which it

emits, elemental information of the sample can be known. An elemental map of the

surface can be created with use of today's instruments. Quantifieation of the amount of a










particular element can be carried out depending on the peak intensity and using ZAF

correctional factors, where Z is atomic number, A is absorption and Z is secondary x-ray

fluorescence .

Atomic Force Microscopy

The topography and roughness of thin films can be determined using atomic force

microscope. The probe in an AFM is a sharp tip created at the end of a cantilever beam

having a spring constant of 0. 1-1.0 N/m. As the tip is brought near the surface of the

sample, van der Waals forces become active between the tip and sample surface atoms

resulting in the deflection of the tip. This deflection is monitored and measured by a laser

striking the back of the cantilever beam which plots the surface of the sample as the tip is

moved across the sample. A general schematic of the AFM set up is show in Figure 3-4.

The tip can be rastered over the sample surface and a 3-D plot can be obtained with

nanometer spatial resolution. Lower spatial resolution can be obtained if the scanned area

is reduced and slow scans are used. AFM is operated in two modes namely contact mode,

where the tip is in contact with the sample surface, and tapping mode, where the tip is

vibrated over the sample surface. The tapping mode is particularly useful for analyzing

soft samples like polymers.

Transmission Electron Microscopy

The wave nature of electron was first theorized by Louis de Broglie in 1925. The

term "electron microscope" was first used by Knoll and Ruska in 1932 when they build

the first electron microscope and obtained images. TEM is an excellent tool to obtain

high resolution lattice images and diffraction patterns of electron transparent samples.

The high spatial resolution is due to the fact that high energy electrons are used to probe









the sample which have extremely small wavelength (on the order of A+). The wavelength

of the electron depends on the energy as per the expression:

1.22



where h is wavelength of electrons in nm, E is energy of electrons in eV.

A typical TEM operates at 200 to 400 keV. A highly coherent beam of

monochromatic electrons is focused on the sample using a series of electromagnetic

lenses. Some of the possible interactions between sample and the electrons are shown in

Figure 3-5. For transmission electron microscopy, information is collected below the

sample from the deflected and un-deflected electrons on a fluorescent screen to form

images or diffraction pattern depending on where it' s focused below the sample. Unlike

scanning electron microscope, the entire sample is in focus all the time as along as its

electron transparent. Some of the drawbacks of TEM are as follows

Limited sampling volume
Complex data interpretation
Beam damage
Sample preparation

Electron diffraction is a very important quantitative characterizing technique to

analyze a sample. Information from a particular area of a sample can be obtained by

using selected area aperture (SAD) which focuses only on the area of interests. Other

information, like defects (line and point defects), burger's vectors, 3-D loops, crystal

orientation, orientation relationship between film and substrate or multilayer films and

sample crystallinity can be obtained.










Focused lon Beam

The samples required for TEM in this study were prepared by focused ion beam

(FIB) technique. Focused ion beam uses an electron beam for imaging and a gallium ion

source for imaging as well as milling. The reason for using Ga+ ions for milling is

because the mass of Ga is approx. 127000 times that of electron and so provides a huge

momentum transfer. The Ga source is heated to a liquid on a tip and Ga' ions are

extracted and focused on the sample by applying a bias. The ion source is also used for

milling purposes to prepare an electron transparent sample. Platinum is used to avoid

milling the area of interest. Once the sample is electron transparent it is set free and

transferred using glass rods to a grid for TEM analysis. Some of the advantages of using

FIB as compared to conventional TEM sample preparations are

Selected feature can be prepared for analysis
Excellent control over the cross-section to be prepared
Insulating samples can also be used
No mechanical damage in the area of interest

Van der Pauw Measurement and Four Point Probe

Van der Pauw technique is used to measure the resistivity of the sample in the

semiconductor industry due to its ease and simplicity. The technique doesn't depend on

the shape of sample. Four ohmic contacts are prepared on the four corners or periphery of

a sample. An acceptable version of sample geometry for this measurement is shown in

Figure 3-6. As shown in the Figure3-7, a DC current is applied between contacts 1 and 2

(I12) and voltage is measured along contacts 3 and 4 (V43). This resistance RA is measured

as follows:


R 43
A,
12,









This is followed by another measurement by applying DC current between

contacts 1 and 4 (114) and voltage is measured between contacts 2 and 3 (V23). This gives

resistance RB aS given by the following equation:

I,
V23

The sheet resistance Rs is related to RA and RB through the following equation:


expR + expr =1

The bulk resistivity can be calculated as per the expression:

p = Rsd

where d is thickness of the film

The sheet resistance of copper before and after annealing was measured by four

point probe. A schematic set-up of four point probe is as shown in Figure 3-8. The DC

current (I) is supplied through the extreme probes and corresponding voltage (V) is

measure by the inside probes. The sheet resistivity p of a thin film of thickness (d) is

calculated by using the following equation:


In 2 Il

And the sheet resistance (Rs) is given by


RYI

where k is the geometric factor and its value is 4.53 for semi-infinite thin film.






30


Table 3-1. Sputtering yields of different ions.


Sputtering yield

Al (27) Si (28)

0.16 0.13


0.73 0.48


1.05 0.5


0.96 0.5


0.82 0.42


by 500 eV ions

Cu (64) Ag (106)

0.24 0.2


1.8 1.7


2.35 2.4-3.1


2.35 3.1


2.05 3.3


Be (9

0.24


0.42


0.51


0.48


0.35


W (184)

0.01


0.28


0.57


0.9


1.0


Au (197)

0.07


1.08


2.4


3.06


3.01


He+ (4 amu)


Ne+ (20 amu)


Ar+ (40 amu)


Kr+ (84 amu)


Xe+ (131 amu)


















Voc


Argon


Figure 3-1. Schematic diagram of DC sputtering system with parallel plate discharge.



















I Ivacurum Chambrer
To Pump

Figure 3-2. Schematic diagram of RF sputtering system with a capacitive, parallel plate
discharge.


To Pump











Focus


X-ray tube


Detect
diapleagm


Detecto


Sample


Glancing angle
Dilfraction angle
Aperture angle


Figure 3-3. Schematic diagram of X-ray diffraction set-up.


Photodiede


Laser


Sample Surface Canileer T


Figure 3-4. Schematic set-up principle of an atomic force microscope.









In~ident
high-kV beam


Secondaryp
elct~ronsr (SE)


B;Kckscattered
electrons (BSE)


CharcEIcristic
X-rays


/1Visible


'Absorbed'
elec~trns -


Elctrron-hoic
gars


\C;~L BRem~SStralung
X-rays


:n


Elastically
searttered Dir
electrons bea


Spccirne


yllacitsulcni
Mattered
etectrons


ect
m n


Figure 3-5. Interaction output between a high-energy electron beam and thin specimen.


Square1 or arcrtanlgle:
contracts at thze edges
or inside! the
perimetr~r


Squarle or
rec ta ngle:
contracts a~t
4 the earne~rs


Cloverle~af


Pre~fer~red


(b)
Acceptable


(c)~
Nort Reco~mmended ~


Figure 3-6. Sample geometries for Van der Pauw measurements.











I,


RA = V43 / In2


Rr,= V, /I, 3,


Figure 3-7. Schematic diagram of Van der Pauw configuration for measurement of RA
and RB TOSistances, respectively.








-I+ ~I + I











Figure 3-8. Schematic diagram of four point probe measurement.















CHAPTER 4
PROPERTIES OF W-Ge-N AS A DIFFUSION BARRIER MATERIAL FOR COPPER

Introduction

For present and future integrated electronic technologies, the use of barrier

materials to enable materials integration is becoming increasingly important. In current Si

technology, the push for higher circuit density and low RC time delays has made copper

the material of choice for interconnects due to its higher resistance to electromigration

and lower resistivity as compared to Al. Unfortunately, Cu is known to show poor

adhesion to most dielectric materials and rapidly diffuses into Sio2 and Si. This

obviously degrades the electrical properties of devices ", creating the need for

intermediate layers that provide a barrier to Cu diffusion.

When considering the role of microstructure in diffusion processes, amorphous

materials are generally better suited than polycrystalline phases, as grain boundaries

provide high diffusivity pathways for Cu diffusion through the barrier material. Among

the material systems currently being developed, binary nitrides, such as TaN59,60 and

WNx61,62 are receiving significant attention. For Ta-based barriers, the Ta-Cu phase

diagram indicates that Ta and Cu are effectively immiscible even at their melting

temperature. Unfortunately, recrystallization of TaN films occurs at approximately 600

oC, which is relatively low for diffusion barrier applications. WNx is also an interesting

candidate as it is relatively easy to synthesize as an amorphous film. In this case, the

nitrogen content in WNx films has a significant influence on its diffusion barrier

properties. For WNx films with low nitrogen content, the recrystallization temperature









can be on the order of 450 oC.3 Higher nitrogen content yields higher recrystallization

temperature. WT\x films grown by physical vapor deposition have been reported to

exhibit a recrystallization temperature as high as 600 oC. Yet, the primary mode of failure

remains diffusion through grain boundaries that form during heat treatments.

One approach to achieve higher recrystallization temperature is to consider

ternary compositions. The additional element added to the refractory metal-nitride

composition frustrates the recrystallization behavior, rendering a stable amorphous

mixture at high temperatures and thus minimizing grain boundary diffusion. While failure

of ternary diffusion barriers still occurs through grain boundaries formed due to

decomposition and recrystallization of the film, this process generally takes place at

higher processing temperature. For this reason, there is significant interest in ternary

nitride alloys, such as Ta-Si-N,63-65 W-Si-N,56,66 W-B-N,67 and Ta-W-N68 due to their

high recrystallization temperature as compared to the binaries. In this study, we report on

the diffusion barrier properties of W-Ge-N thin films for Cu metallization. The W-Ge-N

alloy is chemically similar to W-Si-N, should be more resistant to recrystallization than

WNx, and may prove attractive for integration with SiGe or Ge devices.

Experimental Details

The W-Ge-N films were deposited at the rate of 8.64 nm/min on thermally grown

SiO2 (630 A+)/n-type (100) Si substrates by reactive sputter deposition. For comparison,

WNx films were also deposited under similar conditions. The substrates were sequentially

cleaned with trichloroethylene, acetone, and methanol for 5 min each in an ultrasonic

bath. The substrates were then loaded into the multi-target R.F. sputter deposition system

via a load-lock. The base pressure of the sputtering chamber was 7 x 10-6 Torr. Typical

forward sputtering power for the W and Ge targets was 200 W and 100 W, respectively.









Nitrogen was incorporated into the films by leaking a mixture of Ar and N2 at the ratio of

1 : 0.9 into the chamber at a fixed chamber pressure of 1 1.5 mTorr. The thickness of the

films was measured using a stylus profilometer. In the experiments reported here, film

thickness was maintained in the range 300 to 360 nm. All targets were pre-sputtered

before deposition to remove any contaminant present on the target surface.

To assess the compatibility and diffusion properties of W-Ge-N with respect to

Cu metallization, the nitride layer deposition was followed by in situ deposition of a Cu

film 90 nm thick. During Cu deposition, Ar gas was used as the sputter deposition gas at

a fixed chamber pressure of 15 mTorr. After deposition, the individual samples were

annealed in a separate vacuum chamber with a base pressure of 4 x 105 Torr at 400, 600

and 800 oC for 1 hr to study and compare the diffusion barrier properties of the films. The

crystallinity of films and formation of any intermetallic compounds by annealing were

characterized by X-ray diffraction (XRD) measurements. Resistivity was measured by

the Van der Pauw method, while Auger electron spectroscopy was used to characterize

the Cu diffusion profile in the nitride films. Energy dispersive spectrometry (EDS) was

used to determine the composition of the films.

Results and Discussion

Initial studies focused on the crystallinity of the W-Ge-N films, both as-deposited

and after high temperature annealing. Figure 4-1 shows XRD spectra of WNx and W-Ge-

N films, both as-deposited and annealed at 400, 600, and 800 oC. The data show that

films of both compositions are amorphous in the as-deposited condition. For the WNx

films [Fig.4-1(a)], recrystallization is clearly evident in the XRD pattern for annealing

temperature of 400 oC and higher. The onset of grain structure will provide undesired

diffusion paths via the grain boundaries. In contrast, the W-Ge-N film shows no evidence









of recrystallization upon annealing at 400 or 600 oC. The XRD data [Fig.4-1(b)] for W-

Ge-N films annealed at 400 and 600 oC show no peaks related to the nitride material.

Only at 800 oC do polycrystalline peaks appear. The addition of Ge to the W-N solid

presumably frustrates crystallization, thus rendering the films amorphous for more severe

annealing conditions relative to WNx. It should also be noted that the W-Ge-N film was

far less susceptible to oxidation via ambient atmosphere exposure as compared to WVNx.

This may prove advantageous in terms of device processing.

To assess the behavior of these films as diffusion barriers to Cu, the chemical

profile of the annealed structures was determined by Auger electron spectroscopy. Figure

4-2 shows the depth profiles for the WNx film (a) as-deposited and upon annealing at (b)

600 oC, and (c) 800 oC. For the as-deposited diffusion barrier layer in Figure 4-2(a), there

is a well-defined interface between Cu and WNx and the SiO2 buffer layer is evident. In

contrast, after annealing at 600 oC, the Cu signal is seen throughout the barrier layer and

into the SiO2/Si. This is in agreement with the XRD data, which show the formation of

grain structure upon annealing at a temperature of 600 oC. Annealing at 800 oC [Fig. 4-

2(c)] yields a Cu diffusion profile similar to that for the 600 oC anneal. There is also some

apparent intermixing of W and Sio2 Seen at the WNx/SiO2 interface in the Auger depth

profiles of the annealed samples, which is perhaps related to a change in surface

roughness. The intensity of the nitrogen profile decreases slightly as we increase the

annealing temperature to 800 oC. This may reflect the decomposition of WNx and

liberation of N2 fOm the films under those conditions. This is not unexpected as Affolter

et at. 69 have shown that nitrogen is liberated from W-N alloy thin films when annealed at

temperatures above 700 oC.









The chemical composition of Cu/W-Ge-N/SiO2/Si structures was also examined

with Auger electron spectrometry. Figure 4-3 shows the chemical depth profiles in (a) as-

deposited W-Ge-N and after annealing at (b) 600 oC and (c) 800 oC. In Figure 3(a),

distinct interfaces at the Cu/W-Ge-N and W-Ge-N/SiO2 boundaries show that there is no

Cu diffusion during growth. At an annealing temperature of 600 oC, the interfaces and

layers remain distinct, consistent with no or minimal diffusion of Cu through the barrier

layer. These results suggest that W-Ge-N films possess superior diffusion barrier

properties as compared to WNx. At an annealing temperature of 800 oC, Cu is observed in

the diffusion barrier as seen for the WNx film. Again, Cu diffusion correlates with the

appearance of grain structure in the film (i.e., (1 11) reflection of P-W2N in the XRD

pattern). This is also consistent with N loss suggested by the AES sputter profile, and

thus its ability to stuff the diffusion pathways.

The resistivity of the W-Ge-N films was measured using the Van der Pauw

method. In general, the resistance of the as-deposited W-Ge-N films is higher than WNx.

As shown in Figure 4-4(a), the resistivity of WNx and W-Ge-N decreases as the

annealing temperature increases. While the resistivity of WNx decreases upon annealing

at 400 oC, the resistivity for W-Ge-N remains relatively unchanged after annealing at 400

oC. There is a progressive decrease in resistivity for both materials with further increase

in annealing temperature. The change in resistivity correlates with onset of grain

structure, suggesting electron transport across grain boundaries. The decrease may also

be due to the loss of nitrogen from the films. The resistivity of WNx is two orders of

magnitude lower than W-Ge-N at 800 oC reflecting the robustness of W-Ge-N film, for

which increased decomposition temperature slows nitrogen evolution. Figure 4-4(b)









shows the resistivity of W-Ge-N films vs. sputtering power to the Ge target. Clearly, the

resistivity scales with Ge content in the film.

Conclusion

In conclusion, the diffusion barrier properties of W-Ge-N thin films have been

investigated. X-ray diffraction shows recrystallization of WNx films at an annealing

temperature of 400 oC and higher, while W-Ge-N films show recrystallization peaks only

at an annealing temperature of 800 oC, suggesting that the addition of Ge frustrates the

recrystallization behavior of WNx. The AES data show complete Cu diffusion across the

WNx layer for an annealing temperature of 600 oC, while for W-Ge-N films the Cu/W-

Ge-N and W-Ge-N/SiO2 interfaces remain distinct at that temperature, indicating that W-

Ge-N has better diffusion barrier properties. The resistivity of both films decreases with

increasing anneal temperature, with the resistivity of WNx two orders of magnitude lower

than W-Ge-N after annealing at 800 oC. This behavior is consistent with enhanced

stability of W-Ge-N with respect to film decomposition and subsequent nitrogen

evolution.












P-W,N (ll )


SCu (111) 80




600 "C


4 OC


A eposited


(a)





1,
ed
Y

v,


40 45 50
26 (~degrees)


5'5 60 65


26 (degrees)l


Figure 4-1. X-ray diffraction patterns of as-deposited and annealed films at different
temperature of (a) WNx films (b) W-Ge-N films.















8-



7-



M-

o 0-
`~4-



3-



40 -

30
0-





20
5-

140

2-


5 10 15
Timne (m


mO















20 25 30 35

in)



















WI









--Y--Cu--


Timne (m


(c) 70

60

50-

40 -

S30
2-


~3cc~g:
~4c-LL


O 1 1'5 20 25 30 35 40 45
Timne (min)



Figure 4-2. AES depth profiles of Cu/WNx/SiO2/Si (a) as-deposited (b) annealed at 600
oC/1 hr (c) annealed at 800 OC/1 hr.















(a) 80

70 -

60

o~50


3-

30
1-


0-




70
c~-



,5-

J




2-


5 10 15 20 25 30 35

Timne (min)


r~


O 5 10 15 20 25 30 35 40 45 50

Timne (min)


(c) 70-


0 5 10 15 20 25 30 35

Timne (min)


Figure 4-3. AES depth profiles of Cu/W-Ge-N/SiO2/Si (a) as-deposited (b) annealed at
600 oC/1 hr (c) annealed at 800 OC/1 hr.


* W















(a) 10 -




o










103



18


I I I I I I I I


4-

E
u

Y 3-
O
~

~ 2_
v,
o
d


O1-




-0.25


0.00 0.25


0.50 0.75 1.00 1.25 1.50 1.75

H target RF: power I \'ars)


Figure 4-4. Resistivity vs. annealing temperature (a) for W-Ge-N and WNx films. Also
shown (b) is the resistivity as a function of sputter target power for the Ge and
W targets.


a a -5- W-Ge-N\



















0 100 200 300 400 500 600 700 800 900

Anniealing temp. (oC)















CHAPTER 5
INVESTIGATION OF W-Ge-N DEPOSITED ON Ge AS A DIFFUSION BARRIER
FOR Cu METALLIZATION

Introduction

The progression to ever-decreasing semiconductor device dimensions brings with

it new challenges for material integration. As SiGe-based microelectronic technology

moves toward a 45 nm technology node, there is an imminent need to replace Al

interconnects. Cu is an attractive candidate because of its low resistivity and high

resistance to electromigration as compared to Al.70,71 Replacing Al/SiO2 interconnect

technology with Cu/low-k dielectrics can yield a large reduction in RC time constant

delay, prOViding for a significant motivation to carry out the replacement search for Al.

Unfortunately, Cu has known problems with adhesion to low-k materials. It also has a

high diffusion rate in silicon and germanium creating deep level traps in Si and a shallow

level in Ge located at 0.04 eV near the valence band.72-74


Diffusion barriers are needed to achieve integration of Cu with Si-Ge and Ge. In

diffusion barrier materials the diffusion pathways are interstices, vacancies and grain

boundaries. Diffusion through grain boundaries is fastest and more prominent. By

"stuffing" the grain boundaries with selected dopants, diffusion can often be hindered. In

general, amorphous materials are highly preferred due to the absence of grain boundaries.

Common diffusion barrier materials studied for Cu metallization are refractory metal-

nitrides, which include TaN, TiN, HfN and WNx.42,59,75-79 However, these have limited

utility as viable diffusion barriers because of their relatively low recrystallization









temperatures. Recrystallization of amorphous diffusion barriers can often be inhibited by

the addition of a third element to the matrix. Previous results show higher

recrystallization temperature for W-Ge-N as compared to WNx as the introduction of Ge

effectively frustrates the matrix.so

SiGe devices have higher mobility as compared to Si devices and the flexibility of

band-gap engineering, thus being applicable in high speed electronics. The properties of

the metal/SixGel-x contact layer is important for semiconductor device applications.

Considerable research has been done on the chemical reactivity of metals, such as Co,

Ti,72,82 Pt,83 Pd,84 and CusS on SiGe. In most cases, it results in the formation of metal-Si,

metal-Ge, or metal (SixGel-x) alloys in the temperature range of 400-600 oC. In particular,

direct deposition of Cu on Sil-xGex results in the formation of unstable Cu3(Sil-xGex) in

the temperature range of 250-400 oC.8 One solution could be to grow a metal-rich Cu3Si

or Cu3Ge phase directly on Sil-xGex. Unfortunately the Cu3Si phase reacts with oxygen

when exposed to air. The Cu3Ge phase is more stable and less reactive with oxygen.86 In

general, the need to identify viable barriers of Cu integration with Ge-based structures

persists. In this work, we report on the diffusion barrier properties of W-Ge-N thin films

deposited on Ge for Cu metallization. The properties are compared to WNx deposited on

Ge under similar conditions to evaluate its suitability and properties. Due to the ternary

nature of this diffusion barrier, W-Ge-N is expected to have better recrystallization

properties as compared to WVNx.


Experimental Details

W-Ge-N films were deposited on p-type Ge (001) by reactive sputtering. The Ge

substrates were cleaned by a standard procedure reported elsewhere" to remove any










organic residue or impurity on the surface. The substrates were loaded in the reactive

sputter chamber with a base pressure of 3 x 10-6 Torr via a load-lock. Nitrogen was

incorporated in the film by flowing a mixture of N2 and Ar at a ratio of 1:3 at a fixed

pressure of 10 mTorr. Prior to deposition, all targets were cleaned by pre-sputtering by

flowing Ar+N2 in the chamber. Forward sputtering power for the W and Ge targets were

200 W and 100 W respectively. The deposition rate for W-Ge-N film on Ge was 10.2

nm/min under the above mentioned conditions. Thickness was measured by a stylus

profilometer and was kept in the range of 50 to 300 nm.

Cu metallization was carried out in-situ after depositing W-Ge-N thin films to

determine suitability as a diffusion barrier. Sputter deposition of Cu was carried out by

flowing Ar gas at a fixed chamber pressure of 5 mTorr. After depositing the film stack,

individual samples were separately annealed in the range of 400 oC 700 oC in a tube

furnace. Ar gas was flowed through the tube furnace at 65 sccm for at least 10 hours

before starting the annealing to remove any residual air mixture. Typical annealing

experiment was carried out for 1 hr to study the diffusion barrier properties of the film.

X-ray diffraction was used to identify any intermetallic phases and access the crystallinity

of the films after annealing. Cu diffusion profile through the film was determined by

Auger electron spectroscopy (AES) and Energy dispersive spectrometry (EDS). Interface

properties were determined by cross-section transmission electron microscopy (X-TEM).

Results and Discussion

Crystallinity and high temperature phase formation of the Cu/W-Ge-N/Ge film

structure before and after annealing was determined by XRD. The XRD spectra for

Cu/WNx/Ge and Cu/W-Ge-N/Cu film structure is shown in Figure 5-1, both as-deposited









and annealed in the temperature range of 400 700 oC. The WNx diffusion barrier shows

little crystallization in as-deposited condition as evident by (111) peak, whereas W-Ge-N

fi1m is amorphous. The (111) peak intensity increases at 500 oC indicating further

crystallization as evident from Figure 5-1(a). This leads to increase in the formation of

undesired grain boundaries that provide fast diffusion paths. Cu diffuses through the

grain boundaries and reacts with the underlying Ge substrate resulting in the formation of

a Cu3Ge phase that is clearly evident from Figure. 5-1(a). In comparison, there is no

recrystallization of W-Ge-N fi1ms at 500 and 600 oC. However, at 600 oC, Cu reacts with

Ge in the W-Ge-N layer, and subsequently with Ge substrate to form the Cu3Ge phase

which is evident from Figure 5-1(b). This results in depletion of Ge from the W-Ge-N

fi1m. Upon further high temperature annealing, recrystallization of the Ge-depleted W-

Ge-N takes place leading to barrier failure. This result indicates that adding Ge to W-N

alloy effectively hampers the recrystallization process even at high annealing

temperatures as compared to WVNx.

The Cu diffusion profile through the diffusion barrier in as-deposited and

annealed structures was determined by Auger electron spectroscopy. The fi1m thickness

for WNx and W-Ge-N was 50 nm for the AES chemical profile study. Figure 5-2 shows

the chemical profile of Cu for Cu/WNx/Ge structure (a) as-deposited and after annealing

at (b) 400 oC and (c) 500 oC. The Auger profile shows negligible Cu diffusion in WNx at

400 oC. However, upon annealing to 500 oC, Cu is seen to rapidly diffuse through the

barrier film and to the substrate. This is in agreement with XRD data at 500 oC where it

shows increased crystallization of WNx, thus increasing the grain boundaries, i.e.,

diffusion pathways, resulting in the formation of Cu3Ge phase. The nitrogen signal









intensity steadily decreases at subsequent high temperature annealing compared to the as-

deposited sample indicating decomposition of WNx by liberating N2 at higher

temperatures.

The chemical profile of Cu/W-Ge-N/Ge structures was also measured by Auger

electron spectroscopy as shown in Figure 5-3(a) as-deposited and after annealing at (b)

400 oC and (c) 500 oC. As is evident, there is little or no Cu diffusion through the barrier

upon annealing at temperatures less than 500 oC. The Cu/W-Ge-N and W-Ge-N/Ge

interfaces are distinct in as-deposited and 400 oC annealed samples. At 500 oC, Cu starts

consuming the Ge in the film. This is also supported by XRD data [Fig. 5-1(b)], where a

small intensity (-111) Cu3Ge peak appears for annealing temperature of 500 oC. Upon

annealing at higher temperature, Cu completely consumes the Ge in the film, thereby

depleting the W-Ge-N matrix. This results in the recrystallization of WNx as is evident

from Figure 5-1(b) at 700 oC resulting in rapid Cu diffusion. The result above suggests

that W-Ge-N hinders recrystallization by frustrating recrystallization of the matrix and is

a better diffusion barrier as compared to WNx. The high oxygen content noticed in the

chemical profiles is due to background oxygen during sputtering. The negative enthalpies

of formation of W-O and Ge-O bond are 672 kJ/mol and 659.4 kJ/mol respectively

indicating the stability of the compound after formation. Longer pre-sputtering time of

target could help in reducing the oxygen content in the deposited film.

Figure 5-4 and Figure 5-5 shows the Cu diffusion profile in WNx and W-Ge-N

film annealed at 500 oC, respectively, which was measured by Energy dispersive

spectroscopy (EDS) attached to a cross-section transmission electron microscope

(XTEM) system. As seen in Figure 5-4, Cu signal is seen throughout the WNx film and









into the Ge substrate. However, very little (half counts as compared to WNx) Cu signal is

seen coming from the W-Ge-N (Fig. 5-5) diffusion barrier. This result corroborates the

above mentioned XRD and AES data suggesting excellent diffusion barrier properties for

W-Ge-N as compared to WNx. The interface properties were determined by XTEM.

Figure 5-6 shows XTEM images of (a) Cu/WNx/Ge structure and (b) Cu/W-Ge-N/Ge

structure both annealed at 500 oC. The WNx/Ge and W-Ge-N/Ge interfaces are abrupt,

indicating no intermixing or reactions between them even after annealing. However, the

Cu/WNx and Cu/W-Ge-N interface is rough, suggesting some intermixing and Cu

diffusion in WNx films. XTEM images shows well-defined (111) grain structure for Cu

films deposited on W-Ge-N [Fig. 6(b)] as compared to WNx films [Fig. 6(a)]. This may

be important as (111) oriented Cu films has high resistance to electromigration.

Conclusion

In conclusion, the diffusion barrier properties of W-Ge-N thin films deposited on

p-Ge (001) substrates were investigated. W-Ge-N films have a higher recrystallization

temperature (700 oC) as compared to WNx films as shown by X-ray diffraction. The

failure of W-Ge-N diffusion barrier films at high temperatures occurs by Cu diffusing

through the grain boundaries formed by recrystallization of Ge depleted W-Ge-N. This

Ge depletion is caused by the consumption of Ge in the barrier film by Cu at high

temperatures, thereby forming the Cu3Ge phase. Removal of Ge from the film makes the

W-Ge-N barrier more likely to recrystallize. In contrast, failure of WNx diffusion barrier

at high temperature takes place by diffusion of Cu through WNx grain boundaries.

Nitrogen evolution from the film at high temperature causes decomposition of the

diffusion barrier, thus enhancing crystallization of WNx film. The AES data clearly









shows complete Cu diffusion throughout the WNx layer at 500 oC annealing temperature,

whereas there is little or negligible Cu diffusion through W-Ge-N films. This suggests

that W-Ge-N is a better diffusion barrier. This is substantiated by the Cu profile measured

by EDS. The Cu films deposited on W-Ge-N have better orientation as compared to WVNx

after annealing at high temperatures, thereby increasing its resistance to electromigration.





























35 40 45 50 55 60
20 demeesr


B-W N (111)


Cu Ge (020)
P-W,7N (220)99


(b)


35 40 45 50 55 60


20 degrees





Figure 5-1. X-ray diffraction patterns of as-deposited and annealed films at different
temperature of (a) WNx films (b) W-Ge-N films.












Room temperature
Ge


(a) 100


20
0-





10-






80


2 4 6 8 10 12
Timne (min)


400 "C annealed


O 3 6 9) 12
Timne (min)


15 18 21


500 oC annealed


0 10 20 30
Timne (min)


40 50


Figure 5-2. AES depth profiles of Cu/WNx/Ge (a) as-deposited (b) annealed at 400 OC/1
hr (c) annealed at 500 OC/1 hr.












Roomn temperature
Cu
(a) 100 --,------- Ge

80
600


N

40-





0 5 10 15 20 25 30 35
Timne (min)


400 of annealed
Cu G


80


60 -
~~,o







0 5 10 15 20 25 30 35 40
Timne (min)


500 of annealed
Cu
() 100

80




~40-O



20 "


0 5 10 IS 20 25 30
Timne (min)




Figure 5-3. AES depth profiles of Cu/W-Ge-N/Ge (a) as-deposited (b) annealed at 400
oC/1 hr (c) annealed at 500 OC/1 hr.





































SGe x 5
-A Cu x 5


i~40


0 50 100 150 200

Depth (nm)


Figure 5-4. EDS depth profile of Cu/WNx/Ge annealed at 500 OC/1 hr.













































80
70


S50
U 40
30
20
10
0-


O 20 40 60 80 1. 120 140 160 180

Depth (nm)


Figure 5-5. EDS depth profile of Cu/W-Ge-N/Ge annealed at 500 OC/1 hr.

























































Figure 5-6. X-TEM images of (a) Cu/WNx/Ge annealed at 500 OC/1 hr (b) Cu/W-Ge-
N/Ge annealed at 500 OC/1 hr.















CHAPTER 6
COMPARATIVE STUDY OF Hf~Nx AND Hf-Ge-N DIFFUSION BARRIERS ON Ge

Introduction

For many years, aluminum has been the primary interconnect metal for Si-based

integrated circuits. However, with device dimensions shrinking to sub-45 nm and

demands for high current density increasing, the conductivity and electromigration

properties of Al become limitations to performance. In response, Cu is beginning to

replace conventional Al interconnects given its better electromigration resistance and

lower electrical resistance.70,71,s? The use of low resistivity Cu compared to Al

significantly reduces the circuit time constant delay to make the circuit faster. As the

need for high speed electronics grows, there is also a renewed interest in Ge and SiGe

based devices because of inherent advantages of Ge over Si, i.e., smaller E,, higher

mobility of charge carriers and lower dopant activation energy.88-90 Sil-xGex -based

devices are also of interest because of the innate flexibility to tailor the bandgap through

the alloy composition.91-93 These factors provide sufficient impetus to investigate barrier

layer materials needed in incorporating Cu interconnects in Sil-xGex and Ge-based

devices.

For interconnect applications, copper cannot be deposited directly on Si-Ge since

it diffuses rapidly in Si and Ge creating deep level traps.73,94 For the case of Si, it forms

copper silicides at saturation. It also passivates dopants by forming Cu-D (D is dopant

atom) covalent pairs thus altering the intended doping levels.95 COpper is also known to

diffuse rapidly in Ge with an average diffusion coefficient of 3 x 105 c2S- nte70t









900 oC temperature range.96 COpper introduces three acceptor levels in Ge, two at

Ev+0.04 and Ev+0.32 eV and another at Ec-0.26 eV respectively.74 Direct deposition of

Cu on Sil-xGex results in the formation of Cu3(Sil-xGex) and passivation of the dopants.97

In addition to the above issues, Cu also exhibits poor adhesion to dielectrics commonly

used in Si device structures.'

Considerable work has focused on identifying viable Cu diffusion barrier

materials on Si. Since amorphous materials lack grain boundaries that are fast diffusion

pathways, they are ideally suited for application as a diffusion barrier. Recent material

systems that have been studied as possible Cu diffusion barriers for Si include refractory

metal nitrides, such as TaN, TiN, WNx.59,61,75,97 These binary nitrides, however, tend to

recrystallize at moderate temperature, thus becoming susceptible to rapid Cu diffusion.

There is significant interest in identifying diffusion barrier materials that remain

amorphous at high processing temperature and effectively block Cu diffusion. Increasing

the temperature necessary for crystallization can often be achieved by the addition of a

third element to a binary matrix material. Some of the ternary materials systems that have

been studied include Ta-Si-N, W-Si-N, W-Ge-N.98-100 In this paper, we report on the

recrystallization of HfNx and Hf-Ge-N thin films deposited on Ge and their diffusion

barrier properties for Cu metallization.

Experimental Details

HfNx and Hf-Ge-N thin films with varying thickness (15, 50 and 300 nm) were

deposited on p-Ge (001) single crystal substrates by reactive sputtering at room

temperature. Prior to deposition, the substrates were cleaned with trichloroethylene,

acetone and methanol in an ultra-sonic bath for 5 min each to remove any organic residue

from the surface. The substrates were introduced in a reactive sputter deposition chamber









with a base pressure of 3 x 107 Torr via a load-lock. Nitrogen was incorporated in the

film by flowing Ar and N2 in the chamber at a ratio of 3:1. The total chamber pressure

during deposition was 10 mTorr. Prior to deposition, the targets were cleaned in-situ by

pre-sputtering with Ar+N2 at a fixed chamber pressure of 15 mTorr. The forward

sputtering power for Hf and Ge was 200 and 100 W, respectively. The typical deposition

rate for HfNx and Hf-Ge-N films was 1.8 and 6.23 nm/min, respectively. Identical

thickness was achieved for both films by varying the deposition time. Film thickness was

measured by a stylus profilometer.

Nitride film deposition was followed by in-situ deposition of Cu films. The

forward power used for Cu deposition was 200 W. The Cu thickness was maintained

constant at 300 nm for all films. The deposition was carried out by flowing Ar inside the

chamber at a fixed chamber pressure of 5 mTorr. Individual film stacks were then

separately annealed in a tube furnace in the temperature range of 400 to 700 oC for 1 hr.

Before starting the annealing process, the tube was purged by flowing Ar gas at 65 sccm

for at least 10 hrs. The film crystallinity before and after annealing was determined by X-

ray diffraction (XRD) while the film surface morphology and roughness after annealing

were determined by field emission-scanning electron microscopy (FE-SEM). The

chemical depth profile of Cu diffusion through the diffusion barrier was determined by

energy dispersive spectroscopy (EDS). The chemical state analysis of Cu and

intermetallic compound formation with Ge was investigated by X-ray photoelectron

spectroscopy (XPS). Interface reactions and properties were determined by cross-section

transmission electron microscopy (X-TEM).









Results and Discussion

The HfNx films were amorphous in the as-deposited condition and showed no

signs of recrystallization for any film thickness even after high temperature annealing. A

lack of crystallization upon annealing is desirable as formation of grain boundaries leads

to rapid Cu diffusion. The HfNx diffusion barrier properties are expected to be attractive

based on its high melting temperature (3330 oC).101 Materials that have a high melting

temperature also generally show a have high recrystallization temperature since both

process involve bond breaking. HfN films have been shown to be stable to thermal

decomposition up to 1000 oC.102 Fig. 1 shows the X-ray diffraction patterns for

Cu/Hf~Nx/Ge as-deposited and after high temperature annealing in the range of 400 to 700

oC in an Ar atmosphere. For 300 and 50 nm thick Hf~Nx diffusion barrier films that were

annealed at a temperature of 600 oC or greater, the Cu films exhibit a shift in the Cu (1 11)

peak towards smaller 26 values. This may indicate a reaction with the HfNx film. It is

noted, however, that for the 300 nm thick HfNx barrier, no Cu3Ge phase was formed even

after annealing at 700 oC. For the 50 nm thick HfNx film (Fig lb), as aforementioned,

there is a definitive shift of Cu (1 11) peak at 700 oC annealing temperature which may be

due to reaction of Cu with the underlying HfNx layer. Also evident for the 50 nm thick

HfNx sample is the formation of non-stoichiometric Cu3-xGe phase after annealing at 700

oC. This indicates Cu diffusion through the HfNx diffusion barrier and reaction with the

underlying Ge substrate. Cu3-xGe phase could also have been formed by possible Ge out-

diffusion through the diffusion barrier and subsequent reaction with Cu. For the ultra-thin

HfNx diffusion barrier film (15 nm), the barrier fails at lower temperature as evident from

Fig. Ic that shows Cu3Ge phase formation after annealing at 500 oC and above. The 400

oC anneal pattern, however, does not reveal evidence of barrier failure.









The properties of Cu/Hf-Ge-N/Ge multilayers were then examined and compared

to the Cu/Hf~Nx/Ge samples. Fig. 2 shows X-ray diffraction patterns for Cu/Hf-Ge-N/Ge

as-deposited and after high temperature annealing in the range of 400 to 700 oC in Ar

atmosphere. Based on the behavior of the thickest film, the Hf-Ge-N films remained

amorphous after annealing at a temperature as high as 7000C for all film thicknesses. For

the 300 nm thick Hf-Ge-N film (Fig. 2a), there is little or no shift in the Cu (111) peak

even after annealing at 700 oC. For the 50 nm thick Hf-Ge-N film thickness (Fig 2b), Cu3-

xGe phase formation is evident after annealing at 600 oC indicating that Cu has diffused

through the barrier film to react with the underlying Ge substrate. It is also noted for the

600 oC annealed sample that the Cu (200) and (111) peaks are no longer present,

suggesting significant loss of Cu to the underlying material. At 700 oC annealing

temperature, sufficient Cu diffuses through the barrier layer to form stoichiometric Cu3Ge

phase, again indicating barrier failure.

These data suggests that, while Hf-Ge-N and HfNx have similar recrystallization

behavior, the diffusion barrier properties of HfNx are superior. In particular, the absence

of the Cu (200) peak for the film on 50 nm thick Hf-Ge-N annealed at 600 OC suggests

significant diffusion as compared to HfNx.

One possible factor in determining the properties of the two materials relates to

the relative percentages of the elements present in the diffusion barrier. As mentioned

before, HfNx and Hf-Ge-N films were deposited at the rate of 1.8 and 6.23 nm/min,

respectively. As the forwarding power to Hf was kept constant during deposition of both

films, the total Hf content in the HfNx sample is about 3.5 times greater than that for the

corresponding Hf-Ge-N film of the same thickness. This results in the deposition of a Ge-









rich Hf-Ge-N film. Cu is known to react readily with Ge. For example, at room

temperature, a 20 nm Cu3Ge reaction layer will form at a Cu/Ge interface in 24 hrs in a

binary reaction couplel03 and the reaction rate should increase with increased anneal

temperature. The atomic percentages of each element present in Hf-Ge-N film as

measured by Auger electron spectroscopy were 3 at. % nitrogen, 41.5 at. % oxygen, 28.8

at. % germanium and 26.7 at. % hafnium, respectively, in as-deposited condition.

The surface morphology of the barrier materials was revealed by field emission-

scanning electron microscopy. Fig. 3 shows a comparison of FE-SEM micrographs for

(a) Cu/Hf~Nx/Ge and (b) Cu/Hf-Ge-N/Ge annealed at 500 oC; samples that retained barrier

integrity as evidenced by the XRD patterns shown in Fig. lb and 2b. The thickness of

each HfNx and Hf-Ge-N layers was 50 nm. The grain structure observed in the

micrographs is that of the Cu. Note that there is no evidence of delamination. Fig. 4

shows the FE-SEM micrographs for 50 nm thick (a) Cu/Hf~Nx/Ge and (b) Cu/Hf-Ge-N/Ge

films annealed at 600 oC. After annealing at 600 oC, the surface morphology is

significantly different for the copper films on HfNx as compared to that on Hf-Ge-N

films. For Cu on the HfNx, the Cu films are continuous with a roughness similar to that

seen for the 500 oC anneal. For the Cu on Hf-Ge-N, however, significant Cu segregation

is observed. This is consistent with the suppression of the (002) Cu peak for this

structure and annealing temperature. Reaction with the Hf-Ge-N and possible Cu

diffusion through the Hf-Ge-N film leads to depletion of Cu from the surface and

segregation of Cu islands. This apparent Cu loss on the surface is in agreement with the

XRD data showing the appearance of Cu3Ge peaks for Cu films on Hf-Ge-N barriers that









are annealed at 600oC. No such peaks are detected for the comparison HfNx film

suggesting improved diffusion barrier quality of the latter film.

The interface properties and reactions were examined by cross-section

transmission electron microscopy. Fig 5 shows the X-TEM images of the 50 nm thick (a)

Cu/Hf~Nx/Ge and (b) Cu/Hf-Ge-N/Ge films after annealing at 600 oC for 1 hr. Cu

diffusion is clearly seen in the Hf-Ge-N film with the formation of Cu3Ge phase formed

below the diffusion barrier film. The image of the HfNx film, however, shows a

negligible amount of Cu diffusion as indicated by a continuous Cu film on the surface

and no indication of formation of the Cu3Ge phase. The discontinuous layer at the

Hf~Nx/Ge interface is due to delamination of Hf~Tx. This could be due to TEM sample

preparation prepared by focused ion beam (FIB). The chemical diffusion profile of Cu

was determined by energy dispersive spectroscopy attached to the cross-section TEM.

Figs. 6 and Fig. 7 show the chemical diffusion profile of Cu after annealing at 600 oC for

1 hr for the HfNx and Hf-Ge-N films respectively. A Cu signal is present in the Hf-Ge-N

barrier layer and Cu3Ge has clearly formed by transport through the barrier film to the Ge

substrate. In contrast, HfNx shows little Cu signal is seen from the Ge substrate indicating

that the HfNx barrier layer prevented Cu diffusion to the Ge substrate. The EDS peak

positions of Cu and Hf overlap each other. As a result, a bump in the Cu intensity profile

is observed in the HfNx layer. Also, a strong Hf EDS intensity is seen from the entire

region of the Cu layer due to the EDS detector' s inability to differentiate between Cu and

Hf.

The chemical state of Cu and intermetallic phase formation after annealing was

determined by X-ray photoelectron spectroscopy (XPS). Fig. 8 compares the Cu 2p3/2









peak shifts in Cu/Hf-Ge-N/Ge fi1m for different sputtering times in the (a) as deposited

material (b) after annealing the fi1m at 600 oC for 1 hr and Ge 2p3/2 peak shifts in (c)

Cu/Hf-Ge-N/Ge for different sputtering times after annealing at 600 oC. The Cu surface

of the as-deposited film is clearly oxidized forming a CuO layer. This is evident from the

characteristics satellite peaks formed for Cu+2. After sputtering, however, the peak shifts

and matches with pure Cu (932.8 eV). As seen in Fig. 8b, after annealing the 50nm

Cu/Hf-Ge-N fi1m at 600 oC, the Cu 2p3/2 peak in the as-received condition forms at 934.8

eV indicating its reaction with Ge and formation of Cu3-xGe. This is also consistent with

the XRD and X-TEM data which show the formation of Cu3-xGe. The peak intensity

increases with sputtering time as more Cu participation in the Cu-Ge bond is revealed.

Apparently, there is slight shift in the Cu 2p3/2 peak to 933.3 eV after 60 minutes

sputtering indicating that Cu might react with oxygen and form a Cu-O compound. The

Ge 2p3/2 peak appears at 1221.9 eV after annealing at 600 oC. The Ge 2p3/2 peak intensity

increases with sputtering time indicating increased Ge participation in the Cu-Ge bond

formation. There is a shift in the Ge 2p3/2 peak position to 1221 eV after sputtering for 60

minutes. This might be due to reaction with oxygen and formation of Ge-O bond.

Conclusion

In conclusion, a comparative study of the diffusion barrier properties of Hf-Ge-N

and HfNx deposited on (001) Ge single crystal wafers was conducted. The FE-SEM

images show almost identical surface morphology of Cu fi1ms after annealing at 500 oC.

Annealing at 600 oC, however, results in considerable extent of diffusion across the Hf-

Ge-N films leaving a discontinuous Cu film on the surface. Furthermore, sufficient Cu

transport occurs to form Cu3Ge which is evident from XRD data. In contrast, little or no

diffusion takes place for HfNx films of the same 50 nm thickness and annealing condition









leaving Cu films continuous and smoother. This is also substantiated by cross-sectional

TEM images which clearly show the formation of a Cu3Ge below the Hf-Ge-N diffusion

barrier, but no such phase is formed in the comparison HfNx film after annealing at 600

oC. The chemical valence state was determined by XPS and the results point to Cu-Ge

bond formation after high temperature annealing. The chemical diffusion profile

measured by EDS shows Cu signal emanating from the Hf-Ge-N diffusion barrier and the

underlying Cu3Ge phase formed after annealing at 600 oC. There is little or no Cu signal ,

however, observed in the HfNx diffusion barrier and underlying Ge substrate, indicating

that the HfNx diffusion barrier was successful in preventing Cu diffusion to the substrate.

It is thus concluded that HfNx is an attractive diffusion barrier for Cu on Ge, while Hf-

Ge-N demonstrates limited utility.



























I I I I I I I I I I


f C
la) '3





_

n_ ..~.~

__~_~


700 %

600 %

j@@"G

C~~"C

Za; d~Fo5ihd


29 (de; g.


5~ 60


28 (d~eg.)


55 60


29 (deg.)


Figure 6-1. X-ray diffraction patterns of Cu/HfNx/Ge films in as-deposited and annealed
conditions for varying thickness of (a) 300 nm, (b) 50 nm, and (c) 15 nm.





























31 5 4 4 5 5


Cn~G14-121)


2e (deg.)


28 (dg.)


Figure 6-2. X-ray diffraction patterns of Cu/Hf-Ge-N/Ge films in as-deposited and
annealed conditions for varying thickness of (a) 300 nm, (b) 50 nm, and (c) 15
nm.























Figure 6-3. FE-SEM images of 50 nm films annealed at 500 OC for 1 hr: (a) Cu/Hf~Nx/Ge
and (b) Cu/Hf-Ge-N/Ge.


Figure 6-4. FE-SEM images of 50 nm films annealed at 600 OC for 1 hr: (a) Cu/Hf~Nx/Ge
and (b) Cu/Hf-Ge-N/Ge.

























































Figure 6-5. X-TEM images of 50 nm films annealed at 600 OC for 1 hr: (a) Cu/Hf~Nx/Ge
and (b) Cu/Hf-Ge-N/Ge.




































Depth (npm


11 lf i


1)
-3


Depth(nml D~pthlnni



Figure 6-6. EDS depth profile of 50 nm thick Cu/HfNx/Ge annealed at 600 OC for 1 hr.








73























II












hr.


















-#*--Ater .C -1 so.Te~rr n

- (b)










Binding Ener~g NY.:


Bindine Energy (sk `


Bindine EngErs (sY:


Figure 6-8. XPS chemical state data of Cu 2p peak at various sputtering times for 50 nm
thick film of Cu/Hf-Ge-N/Ge (a) as deposited and (b) annealed at 6000C for 1
hr.















CHAPTER 7
EFFECT OF Ge OVERLAYER ON THE DIFFUSION BARRIER PROPERTIES OF
Hf~Nx

Introduction

Cu is gradually replacing Al as an interconnect material in integrated circuits due

to its lower resistivity and higher electromigration resistance.104 The limitations on

materials properties will become more stringent with shrinking device dimensions as we

move towards the 45 nm device node technology.'os For example, the current density and

the associated heat generation increase rapidly as more devices are packaged on a single

chip. The properties and behavior of interconnects become the limiting factor in

determining the circuit speed under these extreme circumstances. Although Cu is

preferred in metallization schemes for resistivity reasons, it rapidly diffuses in Si creating

deep level traps and reacts with dopant atoms, deteriorating device performance.106-10

There is thus a need to find a diffusion barrier material for Cu metallization. Ideally, a

single crystal diffusion barrier with no defects is highly desirable due to lack of grain

boundaries. Growth of single crystal barriers, however, is not practical due to the process

restrictions (e.g. thermal budget) and material properties (e.g. lattice parameter and

thermal expansion coefficient mismatches between Cu and the underlying substrate (Si or

low x material). An amorphous system is then preferred due to lack of grain boundaries,

which provide facile pathways for diffusion of Cu. Avoiding recrystallization and the

formation of grain boundaries is thus necessary for the diffusion barrier to function.

Refractory metal nitrides are being studied intensely as diffusion barriers for Cu because









of their high melting temperatures. Some of the refractory metal nitrides being considered

are WNx,42,109,110 TaN, 111-113 and TiN.114-117 Diffusion barriers of these materials fail at

relatively low temperature (500 to 600 oC) due to recrystallization. Ternary diffusion

barriers such as W-Si-N,55 Ti-Si-N,118-119 W-B-N,120 and W-Ge-Nso have better diffusion

barrier properties owing to higher recrystallization temperatures compared to their

respective binary systems.

Recently Hf~Nx78,121 has received significant interest as a diffusion barrier material

due to its high melting temperature (3330 oC),101 which translates to higher

recrystallization temperatures., Failure of diffusion barriers typically takes place due to

recrystallization of the barrier film giving rise to parasitic grain boundaries. The integrity

of the diffusion barrier can be enhanced in several ways, for example, by adding a third

element into the binary matrix,122 choosing a refractory nitride system having very high

melting temperature, or by combination of materials that collectively hinder Cu diffusion.

In this paper it is demonstrated that the third approach is a viable one to formulating a

diffusion barrier for Cu metallization. Specifically, a novel bilayer diffusion barrier of Ge

(25 nm)/Hf~Nx (7 nm thick) on Si is tested as a barrier for Cu. The candidate diffusion

barrier is compared with a simple HfNx layer deposited under identical conditions. The

results indicate that the integrity of the diffusion barrier was maintained even under

strenuous test conditions.

Experimental Details

Ge/Hf~Nx diffusion barrier films were deposited on p-Si (001) single crystal wafers

by a reactive sputtering process. The native oxide on Si was first etched by dipping the

substrate in 7:1 buffered oxide etchant and rinsing with D.I. water. After drying the

substrates with ultra high purity N2 gaS, the substrates were loaded into the deposition









chamber via a load-lock. The chamber base pressure was maintained at 3 x 107 Torr. The

sputtering targets were pre-cleaned before deposition by flowing Ar gas inside the

chamber at a fixed chamber pressure of 15 mTorr. The forward sputtering power used for

Hf and Ge targets was 200 and 100 W, respectively. HfNx diffusion barrier films were

deposited by flowing Ar + N2 gaS inside the chamber at a fixed pressure of 10 mTorr.

Once the desired thickness of HfNx was achieved, the N2 flOw was stopped and a Ge over

layer was deposited via Ar ion sputtering. The thicknesses of the HfNx and Ge films

were maintained at 7 nm and 25 nm, respectively, by varying deposition time.

Cu thin films were deposited in-situ on the Ge layer at a fixed pressure of 5 mTorr

without breaking vacuum. Forward sputtering power used for Cu was 200 W. The

thickness for the Cu film was 300 nm for all samples. The substrate was rotated at 20 rpm

throughout each deposition process to ensure film uniformity. Subsequent to Cu

metallization, individual diffusion barrier samples were annealed separately in a tube

furnace in the temperature range 400 to 700 oC for 1 hr to test the integrity of the

diffusion barriers. Ultra high purity Ar gas was used to purge the tube furnace at a flow

rate of 65 sccm for at least 10 hr. Individual samples were also annealed at 500 oC for 1

to 3 hr. The film crystallinity and intermetallic phase formation were determined using a

Phillips APD 3720 X-ray diffractometer (XRD) while the Cu depth profile was

determined using a JEOL Superprobe 733 energy-dispersive spectrometer (EDS). The

integrity of the interface was assessed by examining cross-sections on a JEOL 2010F

high-resolution transmission electron microscope (HRTEM). Samples were prepared

using Dual beam strata DB23 5 focused ion beam (FIB) and the surface roughness of Cu

was determined using atomic force microscopy (AFM) on a SPM/AFM Dimension 3100.









Results and Discussion

Thick HfNx films (thickness 300 nm) without a Ge layer were first grown with an

apparent amorphous structure in the as-deposited condition since no diffraction peak was

observed that was assignable to Hf`Nx and the underlying Si (400) reflection at 69. 15 50

was evident (not shown). The HfNx (thickness 300 nm) remained amorphous even after

annealing at high temperature (400 700 oC) (not shown). Figure 1 shows the X-ray

diffraction patterns of thinner HfNx (7 nm) diffusion barrier films with (Fig. la) and

without Ge (Fig. lb). The sample with only a 7 nm thick HfNx film, however, failed after

annealing at 400 oC for 1 hr as evident from the formation of Cu3Si (Figure lb). By

providing an over layer of Ge, the efficacy of HfNx was significantly enhanced (Figure

la). The integrity of the Ge/Hf~Nx stack remains intact up to 600 oC annealing temperature

indicating a significant increase in its performance as a diffusion barrier. Only after

annealing at 700 oC are copper silicide peaks evident, indicating improved performance

of the bilayer diffusion barrier until 600 oC. For Ge/Hf~Nx films annealed at 500 oC in the

range of 1 to 3 hrs (Figure Ic), the diffusion barrier stack does not show any indication of

Cu transport across it. The integrity of the Ge/Hf~Nx stack at these severe test conditions is

very encouraging.

This excellent performance can be attributed to a combined effect of both Ge and

HfNx. Ge was deposited as an over layer because of its rapid reactivity with Cu to form

Cu3Ge, which also has a high electrical conductivity.23, 103,123 Moreover, the Cu3Ge is less

reactive with oxygen86 than Cu3Si and itself is a good diffusion barrier for Cu.95 Using Ge

as a stand alone diffusion barrier for Cu, however, fails at 500 oC as evident by formation

of copper silicide peaks (Fig. Id). When Ge is used as an over layer, it reacts with Cu to

form Cu3Ge, which hinders Cu diffusion. After passing through the Cu3Ge layer, Cu









encounters the amorphous HfNx layer, which further limits its diffusion. This is an

excellent example of a synergistic effect of two different materials in achieving a

common goal. The Cu intensity peak ratio between Cu (111) and (200) increases as the

temperature is increased. It also increases when annealed at constant temperature but with

increasing anneal time. This apparent recrystallization of Cu to yield a preferred

orientation of Cu along the [1 11] direction is advantageous as Cu has superior resistance

to electromigration in the [111] direction.

Figure 2 shows the atomic depth profiles in Cu/Ge/HfNx/Si film stacks before and

after annealing at 500 oC. As clearly seen in Figure 2a for the as-deposited stack, there is

a noticeable overlap of the Cu intensity profile with the Ge over layer. This supports the

hypothesis that Cu reacts with Ge even at room temperature. The EDS peak positions of

Cu and Hf overlap each other. As a result, a bump in the Cu intensity profile is observed

in the HfNx layer. Also, a strong Hf EDS intensity is seen from the entire region of the Cu

layer due to the EDS detector' s inability to differentiate between Cu and Hf. The Cu

intensity profile declines sharply as it moves inside the Si substrate indicating no Cu

diffusion into the Si at room temperature. For samples annealed at 500 oC for 1 hr (Fig.

2b) and 3 hr (Fig. 2c), the Cu intensity profile again decreases sharply as the scan moves

into the Si substrate. This is consistent with the XRD data, which also show no sign of Cu

diffusion and formation of copper silicides. The oxygen content as measured by auger

electron spectroscopy in Hf~Nx films was approx. 20 at. %. This can be rectified by using

lower base pressure and longer target pre-sputtering times.

The surface roughness was quantitatively measured by atomic force microscopy.

Figure 3 shows AFM plots of the same Cu/Ge/Hf~Nx/Si stacks profiled in Figure 2. As









evident in these images, there is no significant change in the rms roughness values

between the as-deposited samples (5.182 nm) and samples annealed at 500 oC for 1 hr

(9.373 nm) and 3 hr (6.476 nm). The data suggest that the Cu layer is fairly smooth with

the small amount of roughness likely produced by the reaction of the Cu with Ge.

Subsequent annealing does not promote further reaction or surface atom migration.

The film interface integrity and abruptness were determined by high-resolution

TEM for the same 3 samples analyzed by EDS depth profiling and AFM. As shown in

Figure 4 the HfNx/Si film interface appears to be very abrupt for all three samples, which

suggests no intermixing between the layers. This is consistent with the XRD data, which

show no evidence of copper silicide formation for annealing at 600 oC for 1 hr and at 500

oC for 3 hr.

Conclusion

In summary, the effectiveness of a novel Ge(25 nm)/Hf~Nx(7 nm) bilayer to serve

as a diffusion barrier for Cu was investigated. XRD data suggests that the barrier stack

does not fail even after annealing at 600 oC for 1 hr and shows no signs of copper silicide

formation even after annealing at 500 oC for 3 hr, indicating excellent diffusion barrier

quality of the Ge/Hf~Nx bilayer structure. The as-deposited Cu/Ge/HfNx/Si (001) stack as

well as ones annealed at 500 oC for 1 and 3 hr were further examined by EDS, AFM, and

HRTEM. Elemental EDS profiles show a sharp decline in the Cu intensity profile at the

HfNx/Si interface, indicating no Cu diffusion into the Si substrate. The results of these

measurements also suggest that Cu reacts at room temperature with the Ge to form

Cu3Ge. The Cu film surface was smooth in the as-deposited condition with no significant

changes after annealing at 500 oC for 3 hr. This is beneficial as smoother films exhibit

lower contact resistance. Lattice images of the Hf~Nx/Si interface reveal atomic









abruptness, indicating excellent stability of the diffusion barrier and showing no signs of

intermixing. Annealing of a single HfNx layer barrier, however, showed Cu silicide

formation at 400 oC. Overall, the 32 nm Ge/Hf~Nx bilayer stack has excellent diffusion

barrier properties for Cu metallization due to synergistic behavior of two different

material systems.










Cu (111)


35 40 45 50 55 60 65


20 (deigr~ee)


Si Kpr (-100)


35 40 45 50 55 60 65


26 (degrees)


Figure 7-1. X-ray diffraction patterns of as-deposited films and annealed ones at the
temperature shown in the figure: a) Cu/Ge(25nm)/Hf'Nx(7nm)/Si (001) for 1
hr: b) Cu/Hf~Nx(7nm)/Si (001) for 1 hr: c) Cu/Ge(2 5nm)/Hf~Nx(7nm)/Si
annealed at 500 oC for increasing time durations and d) Cu/Ge(25nm)/Si
(001) for 1 hr.












































(d)Si KS; (400)

o00C, I hr

Cu-Si
600 "C, I h~r



oC0", I hr


Cu!11~ 1/ 40oC, I hr


Cu (20)A deposited

IIIIII,, ,
5 40 45 50 55 60 65

26 (degrees)


83







(c)


Cu (200)















35 40 45 50

26 (degrees)


Si KBP(400)


500 oC, 3 hr


55 60 65


h
V)
3
C
3
9
v

v,
C
Q>
ct




s


Figure 7-1. (cotd.).







84



Roomn temperature


O
U 20

10
0-


0 1 20 3~0
Depth (nm)


40o 50


500 "C, I bour


0 30 60 90
Depth (nm)


120 150


0 2 40 60
Depth (rnm)


80 100o


Figure 7-2. EDS depth profile of Cu/Ge(25nm)/Hf'Nx(7nm)/Si for a) as-deposited: b)
annealed at 500 OC for 1 hr: and c) annealed at 500 OC for 3 hr.












































ym
060101 006


1 5 0.d nmdi








U m 7 .41000 pr isd iv
O 50 ,00 nm/di



Fgre -.AMiae fC/G(5m/fx7m/ifo )a-eoie:b


aneldat50o or1h:ad )anaeda 0 C o r


































S1
1;
~..
'.l~'f~LZ


Figure 7-4. HRTEM images of Cu/Ge(25nm)/Hf'Nx(7nm)/Si for a) as-deposited: b)

annealed at 500 OC for 1 hr: and c) annealed at 500 OC for 3 hr.


.~5


-
Tn
:














CHAPTER 8
PROPERTIES OF Ta-Ge-N AS A DIFFUSION BARRIER FOR Cu ON Si

Introduction

As the minimum feature size in integrated circuits continues to decrease, there is a

need to replace Al in interconnects due to the limited conductivity and poor

electromigration performance of Al under high current densities. Cu is a viable

alternative for replacing Al due to its 1.7 times lower resistivity compared to Al and

better electromigration properties under high current densities.71,124 In addition, the RC

time delay achieved by an Al/SiO2 Stack needs to be lowered significantly to increase the

device speed. The RC time constant can be significantly lowered by combining Cu with a

low-k material thus enhancing the performance of the device.3

Although the aforementioned properties of Cu are advantageous for device

manufacture, Cu diffuses very effectively through Si, SiO2 and Ge, degrading the

electrical properties of the device.104 It is an acceptor in Ge introducing traps at E, + 0.04

eV, E, + 0.32 eV and E, 0.26 eV.74 CU is a donor in Si and creates traps from 0.2 to 0.5

eV above the valence band.125 It also reacts with Si to form parasitic copper silicides at

the interface. In addition, Cu reacts with dopants in Si to form complexes that affect the

device characteristics.95 Thus, use of Cu as an interconnect for future technology nodes

will require a viable diffusion barrier. There has been a significant interest in refractory

nitrides such as WNx, TiN, HfNx, and TaN61,110,114-1 16,121,126-128 as diffusion barrier

candidates for Cu metallization. These diffusion barriers, however, typically fail at

moderate temperature (400-600 oC) limiting their use. The primary mode of failure for









these diffusion barriers is by Cu diffusion through grain boundaries formed by

recrystallization of the barrier material upon annealing. By increasing the

recrystallization temperature, grain boundary formation can be delayed, thereby

increasing the robustness of the diffusion barrier. Addition of a third element to the

binary matrix induces amorphization at room temperature and also delays the

recrystallization process, making ternary solutions interesting as diffusion barrier

candidates. Some of the ternary nitride diffusion barriers being studied include W-Si-N,

Ta-Si-N, and W-Ge-N.99,81,129

In this paper, we report on the barrier layer properties when Ge is added to TaN.

Ge is of interest because it displays chemical behavior similar to that of its congener Si

and might be compatible with future Ge and SiGe based devices. Diffusion barrier

properties of Ta-Ge-N were compared with TaN deposited under identical conditions.

The results indicate that a Ta-Ge-N diffusion barrier fails at a higher temperature than

TaN, suggesting superior diffusion barrier properties.

Experimental Details

Ta-Ge-N diffusion barriers were deposited on p-Si (001) wafers by reactive

sputtering process at room temperature. Prior to deposition, the wafer was etched in 7:1

buffered oxide etch to remove its native oxide and then rinsed with deionized water. The

substrate was loaded in the sputtering chamber that is maintained at 3 x 107 Torr base

pressure. Sputtering targets were pre-sputtered before deposition at an Ar pressure of 15

mTorr to remove any contamination on the surface. The forward power used for Ta and

Ge was 200 W and 100 W, respectively. The diffusion barrier films were then deposited

by flowing Ar and N2 at a chamber pressure of 10 mTorr. For comparison, TaN diffusion