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Research@

U.S. Army Research Laboratory

Energy & 

Energetics

Energy & 

Energetics

Dr. Betsy M. Rice and Dr. T. Richard Jow

Editors-in-Chief
Editorial Board

Dr. T. Richard Jow

Editor-in-Chief

Dr. Betsy M. Rice

Editor-in-Chief

Advisory Board

Dr. Kwong K. Choi

Dr. Brad E. Forch

Dr. Piotr J. Franaszczuk

Dr. Shashi P. Karna

Dr. Stephen J. Lee

Dr. Tomasz R. Letowski

Dr. Joseph N. Mait

Dr. James W. McCauley

Dr. Nasser M. Nasrabadi

Dr. Peter J. Reynolds

Dr. Paul H. Shen

Dr. Ananthram Swami

Dr. Bruce J. West

Dr. Brad E. Forch Dr. Joseph N. Mait Dr. James W. McCauley

 Research@ can be accessed electronically at www.arl.army.mil/ResearchARL
Energy and Energetics • 2012

contents

Research@

Introduction to ARL ............................................................................................................................................................... 2

Foreword................................................................................................................................................................................. 3

John M. Miller, Director

Introduction to ARL Research in Energy and Energetics ................................................................................................... 4

Betsy M. Rice and T. Richard Jow, Editors-in-Chief

Improved Cycle Life of Fe-Substituted LiCoPO4 .................................................................................................................. 9

Jan L. Allen, T. Richard Jow, and Jeffrey Wolfenstine

Electrolyte Additive in Support of 5 V Li Ion Chemistry ...................................................................................................17

Arthur von Cresce and Kang Xu

Interfacing Electrolytes with Electrodes in Li Ion Batteries .............................................................................................25

Kang Xu and Arthur von Cresce

Distinguishing Li+ Charge Transfer Kinetics at NCA/Electrolyte and Graphite/Electrolyte Interfaces, and NCA/

Electrolyte and LFP/Electrolyte Interfaces in Li-Ion Cells................................................................................................43

T. Richard Jow, Michelle B. Marx, and Jan L. Allen

A New Direction for the Performance Improvement of Rechargeable Lithium/Sulfur Batteries................................53

Sheng S. Zhang and Jeffrey A. Read

Oxygen Reduction Reaction Catalyst on Lithium/Air Battery Discharge Performance................................................61

Xiaoming Ren, Sheng S. Zhang, Dat T. Tran, and Jeffrey A. Read

Effect of Substitution (Ta, Al, Ga) on the Conductivity of Li7La3Zr2O12............................................................................71

Jan L. Allen, Jeffrey Wolfenstine, Ezhiylmurugan Rangasamy, and Jeff Sakamoto

Shock-Induced Behavior of Cubic Gauche Polymeric Nitrogen ......................................................................................77

William D. Mattson and Radhakrishnan Balu

Optical Cell for In Situ Vibrational Spectroscopic Measurements at High Pressures and Shear ................................85

Jennifer A. Ciezak and Timothy A. Jenkins

The First Launch of an Autonomous Thrust-Driven Microrobot Using Nanoporous Energetic Silicon.........................91

Wayne A. Churaman, Luke J. Currano, Christopher J. Morris, Jessica E. Rajkowski, and Sarah Bergbreiter

Simulations of High-Pressure Phases in RDX.................................................................................................................101

Lynn B. Munday, Peter W. Chung, Betsy M. Rice, and Santiago D. Solares

Elastic-Plastic Behavior of Cyclotrimethylene Trinitramine Single Crystals Under Spherical Indentation: 

Modeling and Simulation..................................................................................................................................................111

J. D. Clayton and R. Becker

Particle-Based Multiscale Coarse Graining with Density-Dependent Potentials: 

Application to Molecular Crystals (hexahydro-1,3,5-trinitro-s-triazine) ........................................................................121

Sergei Izvekov, Peter W. Chung, and Betsy M. Rice

Laser-Induced Plasma Chemistry of the Explosive RDX with Various Metallic Nanoparticles ................................. 139

Jennifer L. Gottfried

Biographies of ARL Authors .............................................................................................................................................149
U.S. Army Research Laboratory

Introduction to ARL

The Army Research Laboratory of the U.S. Army Research, Development and Engineering Command (RDECOM) is the Army’s 

corporate laboratory. ARL’s research continuum focuses on basic and applied research (6.1 and 6.2) and survivability/lethality 

and human factors analysis (6.6). ARL also applies the extensive research and analysis tools developed in its direct mission 

program to support ongoing development and acquisition programs in the Army Research, Development, and Engineering 

Centers (RDECs), Program Executive Offices (PEOs)/Program Manager (PM) Offices, and Industry. ARL has consistently provided 

the enabling technologies in many of the Army’s most important weapons systems.

The Soldiers of today and tomorrow depend on us to deliver the scientific discoveries, technological advances, and the analyses 

that provide Warfighters with the capabilities with which to execute full-spectrum operations. ARL has Collaborative Technology 

Alliances in Micro Autonomous Systems and Technology, Robotics, Cognition and Neuroergonomics, and Network Science, an 

International Technology Alliance, and new Collaborative Research Alliances in Multiscale Multidisciplinary Modeling of Electronic 

Materials and Materials in Extreme Dynamic Environments. ARL’s diverse assortment of unique facilities and dedicated workforce 

of government and private sector partners make up the largest source of world class integrated research and analysis in the Army. 

ARL Mission

The mission of ARL is to “Provide the underpinning science, technology, and analysis that enable full-spectrum operations.” 

Our Vision

America’s Laboratory for the Army: Many Minds, Many Capabilities, Single Focus on the Soldier

ARL’s Organization

•Army Research Office (ARO) - Initiates the scientific and far reaching technological discoveries in extramural organizations: 

educational institutions, nonprofit organizations, and private industry.

•Computational and Information Sciences Directorate (CISD) - Scientific research and technology focused on information 

processing, network and communication sciences, information assurance, and battlespace environments, and advanced 

computing that create, exploit and harvest innovative technologies to enable knowledge superiority for the Warfighter. 

•Human Research and Engineering Directorate (HRED) - Scientific research and technology directed toward optimizing Soldier 

performance and Soldier-machine interactions to maximize battlefield effectiveness, and to ensure that Soldier performance 

requirements are adequately considered in technology development and system design.

•Sensors and Electron Devices Directorate (SEDD) - Scientific research and technology in electro-optic smart sensors, 

multifunction radio frequency (RF), autonomous sensing, power and energy, and signature management for reconnaissance, 

intelligence, surveillance, and target acquisition (RISTA), fire control, guidance, fuzing, survivability, mobility and lethality.

•Survivability/Lethality Analysis Directorate (SLAD) - Integrated survivability and lethality analysis of Army systems and 

technologies across the full spectrum of battlefield threats and environments as well as analysis tools, techniques, and 

methodologies.

•Vehicle Technology Directorate (VTD) - Scientific research and technology addressing propulsion, transmission, aeromechanics, 

structural engineering, and robotics technologies for both air and ground vehicles.

•Weapons and Materials Research Directorate (WMRD) - Scientific research and technology in the areas of weapons, protection, 

and materials to enhance the lethality and survivability of the Nation’s ground forces.

ARL Workforce in 2012

•2013 Civilians - 33 Military

•1399 S&Es

•541 (39%) S&Es hold PhDs

•13 STs / 26 ARL Fellows

Arl’s Primary Sites

•Aberdeen Proving Ground, MD

•Adelphi Laboratory Center, MD

•White Sands Missile Range, NM

•Raleigh-Durham, NC

•Orlando, FL

Visit ARL’s web site at www.arl.army.mil

Unique ARL facilities provide our scientists and 

engineers access to world-class research centers.
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As we complete two decades of excellence as the Army’s 

Corporate Research Laboratory, we look forward to creating 

even more innovative capabilities for the Army, the Department 

of Defense and our Nation. Without the dedication to both basic 

and applied research, it would be impossible for us to effectively 

focus on the future. Our dynamic workforce, in collaboration 

with our national and international partners in industry, national 

laboratories and academia, ensure that we are well positioned 

to take an active role in anticipating and fulfilling the changing 

needs of our military and Nation. Visionary ARL scientists and 

engineers will provide revolutionary technology through novel 

and innovative research, analysis, design, development and 

evaluation. We are committed to scientific excellence as we 

boldly pursue new opportunities. While we cannot predict the 

exact situations and threats we will face, we focus our research 

creatively by addressing a full spectrum of potential combat, 

counterinsurgency and humanitarian scenarios. 

In this, our inaugural issue of Research@ARL monograph series, we endeavor to share with you the excitement of 

our scientists and engineers as they investigate cutting-edge science within the Energy and Energetics research 

areas through advanced experimentation, computational chemistry and physics.

Motivated by the intense need for lightweight portable power, ARL has led the development of higher voltage 

cathode materials and high voltage electrolytes. ARL is gaining significant insight into understanding the 

interfacial chemistry and kinetics in Li-ion batteries, which is vital to enable leap-ahead technology to meet 

future needs. In the energetics arena, ultrafast spectroscopic and imaging methods are enabling us to probe 

chemical, mechanical and structural failure and decomposition in ultra-fast time regimes. Our research teams 

are exploring novel ways to convert mechanical energy to thermal energy by examining initiation mechanisms, 

multi-phase combustion, detonation and the mechanisms that lead to the release of energy. Advances in science 

are facilitating the integration of nanoenergetics and electronics to provide integrated chip-level devices for highpower, precision initiation and ignition effects for specialized weapons applications, as well as to enable the 

controlled mobility of millimeter-scale robotic platforms. 

As the Army’s underpinning science organization, ARL is uniquely positioned to advance its leadership with a 

portfolio of science and technology that has been built over the years in survivability, lethality and vulnerability 

assessment and analysis, materials sciences, information sciences, ballistics and aeromechanic sciences, and 

human sciences.

In the future issues of Research@ARL, we look forward to sharing with you further highlights from our interdisciplinary 

strategic focus areas where our scientists and engineers are leading with a single focus on the Soldier.

FOREWORD

John M. Miller

Director, U.S. Army Research Laboratory
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INTRODUCTION

The discovery of novel materials and methods to store more energy and release energy faster is a crucial area of research to 

address Army needs in both weapons development and lightweight compact power for individual Soldiers, combat vehicles and 

munitions. The U.S. Army Research Laboratory (ARL), the Army’s corporate research laboratory, is breaking existing paradigms in 

these areas by developing a rigorous science-based understanding of diverse multidisciplinary domains, integrating theoretical, 

computational and experimental research. Recently, ARL announced the establishment of Enterprise for Multiscale Research 

of Materials to create a unique capability for the design of materials that are suitable for extreme dynamic environments and 

novel electronic and electromagnetic devices. Cognizant of new research initiatives during the past five years, powered by 

breakthroughs in advanced experimentation and computational chemistry, physics, and materials science, ARL is boldly creating 

and enhancing the essential computational tools and capabilities to advance these rapidly evolving fields. 

The select papers included in this monograph are representative of the high level of theoretical and experimental research, 

discoveries and achievements of ARL scientists and engineers in energetics science and technology, and in energy storageelectrochemical power, focusing on lithium-ion, lithium-air and lithium-sulfur batteries.

ENERGY STORAGE:

Current lithium-ion battery technology is an outgrowth of early space-race research. The success of using pure lithium as the 

anode in primary cells stimulated efforts to employ this material for rechargeable and direct power applications, as well as to 

develop cathode materials and compatible electrolytes that could also endure cycling. The particular difficulty encountered 

in attempts to cycle lithium pertained to the poor morphology of lithium electrodeposits produced during cell recharge in the 

organic solvent based electrolytes of the time. This poor morphology (i.e., dendritic with poor adhesion) invariably led to very 

limited cycle life and the very serious safety problem relating to short-circuiting of cells due to cell separator penetration by 

lithium dendrites. After decades of such experimentation with pure lithium and some practical success by the Sony Corporation 

in 1991 using a lithium/graphitic carbon alloy, a consensus developed that the most promising route was to retreat from the use 

of pure lithium. Instead lithium-ion batteries employ lithium as the intercalated element in graphitic carbon. 

Finding better lithium anode intercalates and compatible electrolytes and cathode materials is the central thrust of lithium-ion 

battery development today. Current Li-ion batteries based on lithium cobalt oxide (LiCoO2) with an energy density of 200 Wh/kg 

supply compact power for commercial (e.g., smartphones, laptops, cameras and power tools) and military applications. Li-ion 

batteries stand on the brink of solving many societal power needs, e.g., affordable energy storage for hybrid and electric vehicles 

and alternative energy applications (e.g., wind, solar), and extremely high power for electric weapons. Only dedicated, focused 

research will enable the identification of new materials and technology that will increase energy and power density, lengthen 

service life, enhance user safety, lower cost, and provide the all-weather service needed by rechargeable batteries. 

In this monograph, we present a representative collection of articles that reveal ARL’s cutting-edge advances in the area of 

electrochemical energy storage on lithium and lithium-ion batteries. Particular attention is given to work on developing higher 

voltage cathode materials and high voltage electrolytes; understanding interfacial chemistry and kinetics in Li-ion batteries, and 

fabricating and testing emerging lithium batteries and the computational efforts that support such research. 

Advanced Li-ion Batteries

The present state-of-the-art cathode material for Li-ion batteries is LiCoO2, which has a theoretical energy density of 518 Wh/

kg when paired with a Li/carbon anode. To increase energy density, cathode materials that can provide a higher charge storage 

capacity, higher voltage or both (energy is the product of both voltage and capacity) are needed. Among the three cathode 

materials that have higher theoretical energy density than LiCoO2, we decided to investigate lithium cobalt phosphate (LiCoPO4,), 
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which has a potential of 4.8 V vs. a reference Li/Li+ electrode and has a theoretical energy density of 800 Wh/kg. Lithium 

cobalt phosphate has the same atomic structure (olivine) as the 3.4 V lithium iron phosphate (LiFePO4) cathode material that is 

presently regarded as providing the highest user safety in a Li-ion battery. Potentially, this should result in a Li-ion battery with 

both higher energy density and good user safety. However, there are several challenges to overcome to develop this cathode for 

Li-ion batteries. It has been known since the year 2000 that it is difficult to cycle LiCoPO4 with good capacity retention. It was not 

understood whether the difficulty in cycling this cathode material with good capacity retention was due to the structural instability 

of the electrode material itself or the instability of the experimental electrolytes at high voltage. While the lack of electrolytes 

that would allow stable and prolonged operation at voltages over 4.5 V was often cited as the main cause of that difficulty, the 

structural stability of the lithium depleted phases of a LiCoPO4 cathode including CoPO4, has also not been given much attention 

in the past when considering cathode life limitations. Two main approaches were taken at ARL to tackle these challenges. Allen et 

al. (J. Power Sources, 2011, 196, 8656) proceeded with substituting part of Co with other transition metals such as Fe to see if 

this can stabilize the structure. Cresce et al (J. Electrochem. Soc., 2011, 158, A337) engaged in exploring solvents and additives 

that would allow the operation of the cathodes at voltages approaching 5 V.

ARL has made great advances on both fronts. With Fe doping, Allen et al. succeeded in increasing the structural stability of the 

LiCoPO4 cathode. Cresce et al. succeeded in identifying tris(hexafluoro) isopropyl phosphate (HFiP) as an effective additive in the 

baseline electrolyte allowing the operation of 5 V class of cathode materials by stopping further oxidation electrolyte components 

on the high voltage cathode surfaces. The structural stability of the cathode is a necessary requirement of a long cycle life with 

sustained capacity retention. The pristine LiCoPO4 lost 50% of its charge capacity in 10 cycles even in the ARL-developed high 

voltage electrolyte, while the stabilized LiCoPO4 can maintain good capacity retention even in standard baseline electrolyte. Thus, 

the capacity fade is mainly a result of LiCoPO4/CoPO4 structural decomposition and, to a lesser degree, a result of electrolyte 

decomposition. Nonetheless, the high voltage electrolyte was still needed for long-term cycling. In short, ARL has had success 

in achieving a stabilized cathode and a compatible electrolyte for a safe Li-ion battery with a cathode that can provide both high 

energy density and good cycle and storage life. The understanding of the mechanisms of structural stability by doping Fe in 

LiCoPO4 and effectiveness of HFiP in stopping further oxidation of electrolytes are being actively pursued.

Interfacial Chemistry and Interfacial Kinetics of Li-ion Battery Electrodes

While the energy density of an electrochemical couple is determined by the electrochemical potential difference and the charge 

capacity of the electrodes chosen, all of the electrochemical reactions occur at the electrode/electrolyte interfaces and the 

electrolyte is needed to perform its function of delivering the needed Li+ to the electrodes and for completing the electrochemical 

reactions by accepting the electron transfer from the external circuit. For Li-ion batteries to achieve long storage and cycle life 

with a high charge/discharge rate, it is essential that the electrolyte is able to perform its functions with low impedance and 

without undesirable parasitic reactions. This is a very challenging requirement due to the extreme electrochemical potentials of 

the anode and the cathode materials used in this chemistry. The electrochemical potential of the lithiated graphite anode, which 

is close to Li potential, is about 3 V negative of standard hydrogen potential, is reductively very active and lies outside the stability 

window of the state-of-the-art nonaqueous electrolytes. The electrochemical potential of the lithium transition metal oxides is 

near the oxygen evolution potential in aqueous systems. State-of-the-art Li-ion batteries rely on the ability of the electrolyte to 

form protective layers (referred to as solid electrolyte interfaces [SEIs]) on the surfaces of the electrodes during initial charging. 

SEIs prevent undesirable reactions. During charge, the solvated Li ion needs to be de-solvated before entering the SEI, diffused 

into SEI, and then intercalated into graphite completing the electron transfer process. The charge of the Li ion is balanced by the 

electron coming from the external circuit near or at the electrode/electrolyte interface. The same process happens at the cathode 

during discharge. These processes involve the steps of solvated Li-ion transport in the bulk, desolvation of solvents before 

entering the SEI, transport through the SEI and intercalation into the electrode, and the electron charge transfer. However, the 

SEI can impede the electrode processes if it is too thick, poorly conducting or unstable The properties of the SEI for any particular 

electrode material are determined by the composition of the electrolyte and electrolyte additives, and this important issue has 

generally been treated in a very difficult, time-consuming and imperfect manner in the past. At ARL, we believe that future 

progress requires an in-depth understanding of the mechanism of the SEI formation. Accordingly, the SEI has been investigated 

using surface analysis methods including X-ray photoelectron spectroscopy, Nuclear Magnetic Resonance spectroscopy (NMR), 

and Fourier Transform Infrared spectroscopy (FTIR). Such studies and our results are described in a featured article by Xu et al. 

(J. Mater. Chem., 2011, 21, 9849). Jow et al. (J. Electrochem. Soc., 2012, 59(5), A604) studied the Li+ charge transfer kinetics 

at graphite anode/electrolyte and LiFePO4 cathode/electrolyte or LiNi1-y-zCoyAlzO2 cathode/electrolyte interfaces at the same 

time in full cells. They found that the activation energy of the Li+ charge transfer kinetics varied with the electrode materials. 

This clearly suggests that the different electrodes create different interphases in the same electrolyte and result in different Li+

charge transfer kinetics. The understanding of how the electrode surface reacts with the electrolyte is critical in advancing the 

Li-ion energy storage technology.
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Emerging Lithium Battery Systems and Novel Solid Electrolytes

The energy density of Li-ion batteries based on intercalation reactions has its limitations. What is beyond Li-ion? At ARL and 

other laboratories, researchers are once again exploring the use of pure Li as a negative electrode, pairing it with an air or 

sulfur cathode that has a higher theoretical energy density than Li-ion. The fact that dissolved sulfur eliminates Li dendrites 

may overcome the usual safety problem. The use of a ceramic electrolyte in a Li/air battery may similarly overcome the lithium 

safety problem. 

A key issue in the development of a Li-S cell has been the excessive reaction of soluble polysulfide (formed during cell discharge) 

with the Li anode. Most prior work in this area has involved the attempt to reduce polysulfide solubility. Zhang et al. (J. Power 

Sources, 2012, 200, 77) were successful in pursuing a completely different approach to this problem. Instead of resolving the 

solubility issue of the polysulfide, they treated the soluble product as a soluble cathode and discovered an additive that could 

protect the Li anode from reacting with the dissolved sulfur compounds. Further research is required to refine this approach. In 

the area of Li-air batteries, there are many challenges to making the air electrode operate properly. Ren et al. (J. Mater. Chem., 

2011, 21, 10118) demonstrated that the catalyst causes a shift in the discharge reaction pathway, making the discharge 

capacity more effectively utilized.

To utilize pure Li metal as an anode for Li-S and Li-air batteries, there is a critical need for a solid electrolyte membrane that is 

compatible with Li, has high Li+ conductivity (>10-4 S/cm), and is thin and mechanically strong. The Li+ solid electrolyte based on 

the garnet atomic structure (Li7La3Zr2O12) [LLZO] can potentially meet the needs. However, the challenge is to stabilize the LLZO 

cubic structure, to provide higher Li+ conductivity, at room temperature. Presently, a lower conductivity tetragonal phase is more 

stable at room temperature. Allen and Wolfenstine et al. (J. Power Sources, 2012, 206, 315–319) succeeded in achieving the 

stabilization of LLZO in the cubic phase and achieved the highest known conductivity for a Li+ solid conductor by Ta substitution. 

The value achieved, 8.9 x 10-4 S/cm, is close to values for liquid electrolytes. They also showed that Ta substitution in the Li 

sublattice is preferred over the previously reported Al, which substitutes for Li. Through hot pressing, a near-100% dense material 

can also be achieved. Thus, Ta doping is very promising. Future work will focus on chemical stability, mechanical properties and 

making test cells of the Ta-doped material.

Computational Exploration of Electrolytes and Electrodes

Despite the progress ARL has made in materials development, there are still many materials issues that need to be understood 

to assist in the advancement of higher energy density and better performance energy storage devices. In particular, a basic 

theoretical understanding of factors controlling stability of electrolytes and their reactivity pathway with electrode materials for 

guiding future materials development is urgently needed. Borodin as a primary author and his coauthors (Xing et al, , Density 

Functional Theory Study of the Role of Anions on the Oxidative Decomposition Reaction of Propylene Carbonate, J. Phys. Chem. 

A, 2011, 115, 13896–13905) demonstrated that the presence of anions such as PF6- or ClO4- not only significantly reduces the 

oxidative stability of the carbonate electrolyte solvents but also stabilizes the solvent-anion oxidation decomposition products 

and changes the order of the oxidation decomposition paths. The primary oxidative decomposition products of PC-PF6- and PCClO4- were CO2 and acetone radical, which is in good agreement with available experimental data. Density Function theory (DFT) 

calculations also explained why toxic fluoro-organics were experimentally observed at elevated temperatures in a carbonateLiPF6 electrolyte in the presence of a cathode material and were not observed during electrolyte oxidation at room temperature. 

The ARL-initiated Collaborative Research Alliance (CRA) on Multiscale Multidisciplinary Electronic (MSME) Materials will further 

enhance and expand our efforts for enhancing the basic understanding of factors controlling stability of electrolytes and 

electrode materials and providing materials predictive and design capability for achieving over 300 Wh/kg energy density goals 

for electrochemical energy storage devices.

ENERGETICS 

ARL has an overarching objective to develop a science-based understanding of novel ways to store and release chemical 

and structural energy for use in explosive and propellant applications that may lead to leap-ahead weapons capabilities. 

The fundamental research initiative is capitalizing on recent breakthroughs in advanced experimentation and computational 

chemistry and physics to identify, understand and characterize processes and mechanisms that control energy storage and 

release. ARL, as a recognized leader in theoretical chemistry research, is directing its resources and expertise into a rigorous and 

innovative program that will lead to a new generation of energetics whose development is computationally guided. One example 

of ARL’s paradigm-breaking research is the quest to store large quantities of structural energy in simple molecular systems, 
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which are inert under ambient conditions but can be compelled to form extended solids by compressing them under extreme 

pressure. Of interest are high pressure forms of polymerized nitrogen, theorized for decades through quantum mechanical 

prediction but only recently synthesized by compressing it in a diamond anvil cell at temperatures and pressures exceeding 2000 

K and 110 GPa, respectively. The energy content of this polymerized nitrogen material is expected to significantly exceed any 

existing conventional explosive, thus ARL has initiated efforts to overcome the substantial scientific and technical challenges to 

create this material under less extreme conditions and stabilize it to an ambient state. New agile synthetic methods are being 

designed and developed to fabricate recoverable highly energetic metastable states of matter using combinations of extreme 

pressure, mechanochemistry, mechanophysics and laser photochemical processes. 

The exciting potential of the nanoscale also forms a large part of ARL’s novel energetics thrust. Based on recent advances 

in nanoscale materials and nanofabrication, ARL is investigating ways of engineering energetic materials with fundamentally 

new and technologically useful characteristics. Over the past decade considerable evidence has accumulated showing that 

energetic nanoparticulates may react to produce enhanced power compared to their more conventional counterparts of the 

same chemical composition. 

Extended Solids: A new class of energetic materials 

Extended solids are emerging as a novel new class of energetic materials. Large-scale quantum molecular dynamics (QMD) 

calculations are being used to identify and characterize candidate extended solid structures and determine what factors affect 

stability. Polymeric nitrogen is the current frontrunner due to its unusually large energy content (its energy release and rapid 

conversion to molecular form), which could be classed as “superexplosive.” Mattson and Balu (Physical Rev. B, 2011, 83, 174105) 

are using QMD simulations to study the shock response of the cubic gauche form of nitrogen, the only crystalline polymeric form 

of nitrogen synthesized to date. Their findings revealed that this cubic gauche nitrogen exhibits unusual mechanical properties, 

which produce an extremely complex shock behavior, far different from that seen in conventional energetic materials. Rather 

than initiating chemical reaction behind the compression front, the shock impact induced phase transitions and material defects 

that resulted in unusual energy dispersion mechanisms that slowed the shock wave. To foster the rapid release of the stored 

structural energy in this material, we must develop new methods to trigger high-rate chemical reactions. One method being 

pursued at ARL is the imposition of mechanical stresses through a process known to accelerate solid-state reaction rates. 

Another new approach is introducing shear into the fabrication process to lower the transition pressures needed to form novel 

high-pressure-phase materials when compressed. The goal is to enable the scale-up synthesis of these new materials for further 

test and evaluation. As such, ARL researchers Ciezak and Jenkins (Rev.of Scientific Instruments, 2011, 82, 073905) have 

designed a special rotational diamond anvil cell for generating high pressure and shear simultaneously. Preliminary studies 

using this cell indicate that both amorphous and crystalline polymeric nitrogen can be synthesized at room temperature and 

pressures of ~ 40 GPa when shear is introduced. 

ARL is also developing novel energetic materials for fuel cells, solar cells and actuators in microrobots. We have made great 

strides in these areas building upon the development of a new type of energetic material based on nanoporous silicon (PS). When 

silicon oxidizes it produces a significantly greater energy release than conventional energetic materials. Of course, the rate of 

energy release depends on intimate mixing of the oxidant with the silicon-based fuel. Capitalizing on the high reactivity (oxidation 

potential) of PS, Becker et al. (Galvanic Porous Silicon Composites for High-Velocity Nanoenergetics, Nano Letters, 2011, 11, 

803–807) have successfully engineered highly explosive composite materials composed of PS infused with an oxidizer that can 

produce very high combustion velocities (~3 km/s). Churaman et al. (J. Microelectromechanical Systems, February 2012, 21(1)) 

used this novel material for thrust actuation of an autonomous jumping microelectromechanical system (MEMS) microbot. The 

actuation pulse lasted several microseconds allowing the robot to reach a vertical height of approximately 8 cm. 

Multiscale Response of Energetic Materials 

A key challenge facing energetic material design and development is to understand the response of an energetic material within 

a munition to unexpected initiation. Given the multi-scale nature of the response of the energetic material in a munition (i.e., 

decomposition and energy release occur at the molecular scale where material response is manifested at the macroscale), 

considerable technology gaps exist in both the measuring and modeling processes. 

ARL is spearheading a research program to explore the multiscale response of energetic materials, with a goal to develop 

methods and models to predict response of energetic materials, with all models validated by advanced experimentation at scales 
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ranging from molecular to continuum. Although in its infancy, this program has yielded notable theoretical and experimental 

advances. As a demonstration of this multiscale approach, ARL has targeted a simple explosive formulation for quantum-based 

multiscale modeling. This formulation, composed of cyclotrimethylene trinitramine (RDX, a commonly used energetic material) 

and polyethylene, has been subjected to extensive quantum mechanical calculations and the results have been homogenized 

into atomistic models for molecular dynamics (MD) simulations. To allow the atomistic simulation results to subsequently be 

upscaled to continuum-level descriptions, Munday et al. (J. Phys. Chem. B, 2011, 115, 4378–4386) have performed a series 

of comprehensive MD simulations using this model to provide information about its atomic-level dynamic response, properties, 

localized heating and mechanical deformation, looking specifically at the various conditions this material will experience under 

extreme compression. Their results demonstrated that molecular deformation is coupled with material phase transitions and 

could activate new slip systems while blocking slip systems of other phases. Such information about deformation mechanisms is 

crucial for the emerging plasticity models for RDX that are being developed at ARL. Clayton and Becker (J. Applied Physics, 2012, 

111, 063512) recently described the elastic-plastic behavior of single crystals of RDX under spherical indentation using this new 

model. Their results will be used in the continuum simulations of macroscale response to mechanical stresses.

At present, system size and simulation constraints prevent the representation of microstructural features inherent in a 

composite energetic formulation in atomistic simulations. To overcome this challenge, ARL has developed a mesoscale modeling 

methodology, similar in spirit to a MD simulation, but that can bridge the system sizes and simulation times necessary to explore 

the dynamic response and evolution of microstructural features. This methodology, Dissipative Particle Dynamics (DPD), uses 

coarse-grain (CG) representations of the materials, and has been shown to work well in describing dynamic and rheological 

properties of soft matter such as liquids or polymers using very simple models of the interparticle interactions. ARL has made 

tremendous advances in the state of the art with the development of density-dependent CG models of a condensed phase 

explosive that properly depict structure and shock properties of RDX for use within the DPD method. Izvekov et al. (J. Chem. 

Phys., 2011, 135, 044112) have produced accurate particle-based CG models directly from atomistic-level interactions for 

RDX through force matching to quantum-based MD simulations. The development of this model is a notable step forward to 

bridging the scales by providing a model that allows for the dynamic response of a realistic composite energetic material at the 

microscale level. 

Direct experimental observation of the key phenomena at the various scales also presents a formidable challenge due to the 

numerous processes occurring at the same time, at extremely rapid rates, and under conditions of extreme temperatures and 

pressures. To tackle these obstacles, ARL has invested in advanced experimental tools and systems to unravel the details of the 

fundamental chemical and physical steps involved in the combustion or detonation of an energetic material. Gottfried’s (Applied 

Optics, 2012, 51(7), b13–b21) recent work using laser-induced breakdown spectroscopy to monitor the time-resolved chemical 

reactions of metallic nanoparticles with RDX at high temperatures is a prime example. In this study, Gottfried showed that the 

plasma chemistry of RDX in the presence of metallic nanoparticles is substantially different from that of the pure material, 

suggesting that introducing aluminum into explosive formulations not only produces extra heat due to aluminum oxidation, but 

also affects the chemical reactions that are occurring. 

In addition to interrogating the subscale chemistry in energetic material response, ARL is addressing the challenges inherent in 

characterizing macroscale material response within extreme dynamic environments by developing novel experimental capabilities, 

as typified in the study by Densmore et al (High-speed two-camera imaging pyrometer for mapping fireball temperatures, Applied 

Optics, 2011, 50(33), 6267–6271). In this work, ultrafast imaging methods were created to monitor temperature in explosive 

events using a specially designed high-speed imaging pyrometer that allows direct imaging and determination of temperature 

profiles within fireballs produced by explosions. Incorporation of this information into experimentally validated multiscale 

modeling and simulation schemes will provide a true predictive capability of energetic material response based on fundamental 

physics and chemistry parameters linked to engineering/continuum models. 

The examples highlighted in this monograph are but a small sample of the numerous cutting-edge research initiatives within ARL 

that have the potential to revolutionize the design, manufacture and implementation of novel, advanced energetic materials, all 

at substantially reduced risk, time and cost.
9

Research@

Improved Cycle Life of Fe-Substituted LiCoPO4

Jan L. Allen, T. Richard Jow, and Jeffrey Wolfenstine

U.S. Army Research Laboratory, Adelphi, MD

ABSTRACT

Fe-substituted LiCoPO4 exhibits greatly improved cycle life relative to LiCoPO4. 

Whereas, pure LiCoPO4 loses more than half of its discharge capacity at the 

10th cycle, the Fe-substituted LiCoPO4 retains about 100% of its discharge 

capacity at the 10th cycle and about 80% of its capacity at the 500th cycle. It 

is suggested that improved cycle life results from Fe3+ substitution on the Li 

and Co sites. The partial substitution of Li+ by Fe3+ and Co2+ by Fe2+ and Fe3+

was evidenced from Rietveld analysis of X-ray powder diffraction data, infrared 

spectroscopy, X-ray photoelectron spectroscopy and Mossbauer spectroscopy. 

The majority of the Fe3+ substitutes at the Co2+ site. The composition of Fesubstituted LiCoPO4 is Li0.92Co0.8Fe2+

0.12Fe3+0.08PO4 for a sample of starting 

composition LiCo0.8Fe0.2PO4.

Journal of Power Sources, Volume 196, Pages 8656-8661 (2011)
10 Journal of Power Sources 196 (2011) 8656–8661

Contents lists available at ScienceDirect

Journal of Power Sources

journal homepage: www.elsevier.com/locate/jpowsour

Improved cycle life of Fe-substituted LiCoPO4

J.L. Allen∗, T.R. Jow, J. Wolfenstine

U.S. Army Research Laboratory, 2800 Powder Mill Road, Adelphi, MD 20783-1197, USA

a r t i c l e i n f o

Article history:

Received 17 May 2011

Received in revised form 14 June 2011

Accepted 15 June 2011

Available online 22 June 2011

Keywords:

Li-ion battery

High voltage

XRD

Mössbauer

Cathode

LiCoPO4

a b s t r a c t

Fe-substituted LiCoPO4 exhibits greatly improved cycle life relative to LiCoPO4. Whereas, pure LiCoPO4

loses more than half of its discharge capacity at the 10th cycle, the Fe-substituted LiCoPO4 retains about

100% of its discharge capacity at the 10th cycle and about 80% of its capacity at the 500th cycle. It is

suggested that improved cycle life results from Fe3+ substitution on the Li and Co sites. The partial

substitution of Li+ by Fe3+ and Co2+ by Fe2+ and Fe3+ was evidenced from Rietveld analysis of X-ray

powder diffraction data, infrared spectroscopy, X-ray photoelectron spectroscopy and Mössbauer spectroscopy. The majority of the Fe3+ substitutes at the Co2+ site. The composition of Fe-substituted LiCoPO4

is Li0.92Co0.8Fe2+0.12Fe3+0.08PO4 for a sample of starting composition LiCo0.8Fe0.2PO4.

Published by Elsevier B.V.

1. Introduction

LiMPO4 compounds, where M = Fe, Mn, Co or Ni, have been

the focus of intense study both for scientific and practical reasons as Li-ion battery energy storage materials since the pioneering

work of Padhi et al. [1]. The voltage of the redox couple varies

with transition metal from 3.4V for Fe [1], 4.1V for Mn [1],

4.8V for Co [2] and 5.1V for Ni [3]. High voltage batteries are

desirable because the stored energy is proportional to the voltage and the power is proportional to the square of the voltage.

For example, LiFePO4 has potentially an energy storage capability of 578Wh kg−1 (3.4V × 170Ah kg−1) and LiCoPO4 about

802Wh kg−1 (4.8V × 167Ah kg−1). Hence, there is a keen interest

to move beyond already commercialized LiFePO4 to the other transition metals. As LiCoPO4 in particular has the potential to increase

energy ∼40% compared to LiFePO4, we have focused our work on

this material. Initial work on LiCoPO4 led to improved rate capability but capacity fade soon emerged as an impediment to further

progress [4–6]. The initial discharge capacity and rate capability

of LiCoPO4 were improved by varying the oxygen partial pressure

during synthesis [6], carbon coating [7] and substitution on the Co

site [8]. However, until now, LiCoPO4 has shown a severe loss of

discharge capacity upon charge–discharge cycling. For example,

Tadanga et al. [4] observed a 10th cycle discharge capacity of ∼52%

of the initial capacity, Bramnik et al. [5] reported a 10th cycle dis-

∗ Corresponding author. Tel.: +1 301 394 0291; fax: +1 301 394 0273.

E-mail addresses: jan.l.allen8.civ@mail.mil, jallen@arl.army.mil (J.L. Allen).

charge of ∼59% of the initial capacity and Wolfenstine et al. [6]

reported ∼53% capacity retention at the 10th cycle. This capacity

fade has been attributed to irreversible structural changes such as

amorphization [5,6] of the material and/or electrolyte degradation

[4]. In this paper, we will show that a substitution of Li+ by Fe3+ and

Co2+ by Fe3+ and Fe2+ improves not only rate capability but also dramatically reduces capacity fade. In addition, even further reduction

in capacity fade is observed when the Fe-substituted LiCoPO4 was

used with electrolyte containing tris(hexafluoroisopropyl) phosphate (HFiP) [9].

2. Experimental

LiCoPO4 samples were prepared via a citrate complexation

route. Co(OH)2, LiH2PO4, and citric acid, 1, 1.01, 1.02, molar ratio,

respectively, were mixed into deionized water until all solids were

dissolved. The resulting solution was evaporated to dryness via a

microwave oven. The resulting dried mass was removed, ground

lightly with mortar and pestle and heated in air at a rate of

10 ◦C min−1 to 600 ◦C and the reactant mixture was held at this

temperature for 12 h.

In order get Fe substitution on both the Li and Co sites,

Co(OH)2, LiH2PO4 and FeC2O4·2H2O with a nominal stoichiometry

of LiCo1−xFexPO4, x = 0.05, 0.1, 0.2 were weighed and then dissolved

in 1 M HNO3 (aq). The resulting nitrate solution was evaporated to

dryness via a microwave oven in a fume hood and then heated

under N2 at a rate of 10 ◦C min−1 to 600 ◦C and held at this temperature for 12 h. During the decomposition of the co-precipitated

nitrates, the decomposition of the nitrate ion provided an oxidizing

0378-7753/$ – see front matter. Published by Elsevier B.V.

doi:10.1016/j.jpowsour.2011.06.057
J.L. Allen et al. / Journal of Power Sources 196 (2011) 8656–8661 865711

component to the N2 atmosphere which transformed a portion of

the Fe2+ to Fe3+.

Carbon coating to improve electronic conductivity was done by

ball milling the samples of LiCoPO4 and Fe-substituted LiCoPO4 for

30 min with 5% by mass acetylene black followed by heating for 1 h

at 600 ◦C under N2. The improvements of the activity of LiCoPO4

after carbon coating [5] and after a short ball milling (<1 h) have

been previously reported [10].

Phase purity was evaluated using X-ray powder diffraction.

Data were collected using a Rigaku Ultima III diffractometer. Lattice constants were calculated from peak positions using Rietveld

refinement of the pattern collected in a parallel beam geometry

or with the use of a NIST certified silicon standard for collection in a Bragg–Brentano geometry using Riqas software (Materials

Data Inc.). Samples were further evaluated spectroscopically using

Attenuated Total Reflectance Fourier-Transform Infrared (ATRFTIR) Spectroscopy, X-ray Photoelectron Spectroscopy (XPS) to

evaluate site occupancy and oxidation states, respectively. Additional information about the oxidation state of Fe was obtained

from Mössbauer spectroscopy (collected at SEE Company, Edina,

Mn) and elemental analysis via inductively coupled plasma optical

emission spectroscopy (ICP-OES, data collected at Galbraith Laboratories, Inc.).

For electrochemical testing, a composite electrode was fabricated by a slurry coating method. Using N-methylpyrrolidone

(NMP) as solvent, a slurry was used to coat an Al foil substrate to produce a composite electrode of 80 wt.% active, 10 wt.%

polyvinylidene fluoride (PVDF) and 8 wt.% super-P carbon and

2 wt.% conductive carbon nanotube composite (CheapTubes.com).

The electrode film was cutinto small discs with an area of 0.97 cm2,

dried under an infrared lamp in air before use and thereafter in a

heated vacuum oven (∼100 ◦C). In a dry room (dew point < −80 ◦C),

Li/active coin cells (Hohsen Al-clad CR2032) were assembled using

Celgard® 3501 as the separator and a 1.0 molal LiPF6 solution in

a 3:7 (wt.%) mixture of ethylene carbonate (EC) and ethyl methyl

carbonate (EMC) electrolyte with and without 1 wt.% HFiP. Electrochemical testing was performed using a Maccor Series 4000 tester.

For calculation of C-rate, a capacity of ∼170 mAh g−1 was assumed.

3. Results and discussion

The partial substitution of Co2+ by Fe2+ was explored because is

known that substitution with Fe2+ can improve the rate capability

ofthe other olivines [1]. For example, Padhi et al.[1] showed that Fe

substitution for Mn2+ in LiMnPO4 enabled it to be cycled whereas

pure LiMnPO4 was found to be essentially electrochemically inert.

At the time, Padhi et al. proposed that the Fe3+–O–Mn2+ interactions destabilize the Mn2+ level and stabilize the Fe3+ level so as

to make the Mn3+/Mn2+ energy accessible. Thus, we report here

on the substitution of Co2+ by Fe2+ as a means to improve the rate

capability in an analogous fashion. However, there is no reported

improvement in capacity fade owing to the substitution of Co2+ by

Fe2+.

The partial substitution of Li+ and Co2+ by Fe3+ was explored

to address capacity fade through improved structural stability of

LiCoPO4/CoPO4 during cycling. This substitutional strategy is based

upon the speculation that decomposition of LiCoPO4 or CoPO4 may

result from a loss of oxygen during charging–discharging as shown

below.

For the case of CoPO4, the proposed decomposition reaction is:

2CoPO4 → Co2P2O7 + 1/2O2 (1)

This reaction results in the release of O2 during the discharge.

This potential mechanism is proposed based on the reported

Fig. 1. X-raypowderdiffractionpatternof LiCoPO4 preparedvia citrate aqueousprecursor route, bottom, and the X-ray powder diffraction pattern of Li0.92Co0.8Fe0.2PO4,

top. The peaks are labeled with the Miller indices of the phospho-olivine structure,

Pnma spacegroup.

decomposition of CoPO4 to Co2P2O7 and O2 during heating under

reducing conditions [11].

For the case of LiCoPO4, the proposed decomposition reaction

is:

2LiCoPO4 → Co2P2O7 + Li2O (2)

A similar mechanism leading to the electrochemical formation

of Li2O was reported by Armstrong et al. [12] during the electrochemical charge of Li2MnO3:

Li2MnO3 → Li2O + MnO2 (3)

We believe that Fe3+ substitution on the Li+ and Co2+ sites might

be a means to slow these types of reactions owing to the higher

affinity of Fe3+ to oxygen relative to Co2+ [13] and/or owing to

changes in the underlying electronic structure of Fe3+ substituted

LiCoPO4 (or substituted CoPO4) relative to pristine LiCoPO4 (or

CoPO4).

LiCoPO4 was prepared for comparison to substituted samples.

A typical X-ray diffraction pattern is shown as the lower pattern in Fig. 1. The pattern confirms that a single phase LiCoPO4

was prepared. A typical X-ray diffraction pattern is shown as

the upper pattern in Fig. 1 for a sample of starting composition

LiCo0.8Fe0.2PO4. As with LiCoPO4, there is no evidence of any impurity phases. ICP was used to calculate the amount of Fe3+ in the

sample of starting composition LiCo0.8Fe0.2PO4 from the assumption that the Fe3+ will be compensated by Li+ ion vacancies. This

atomic ratio was measured to be 0.92, which indicate 55% Fe2+

and 45% Fe3+. Li+ is volatilized during the synthesis in order to

accommodate the Fe3+. The product for the sample of starting composition LiCo0.8Fe0.2PO4 was thus Li0.92Co0.8Fe2+0.12 Fe3+0.08PO4 or

Li0.92Co0.8Fe0.2PO4, for short. Table 1 shows the lattice constants for

the series of compounds prepared for this study.

Fig. 2 demonstrates the effect of Fe substitution on the capacity fade and the importance of the HFiP electrolyte additive. The

Li0.92Co0.8Fe0.2PO4 composition was chosen to examine the cycle

life since it had the largest capacity atthe higher rate. The rate study

will be discussed later in the paper. The cells were cycled between

2.5 and 5.3V via a constant current method at C/5 rate except for

the first two cycles which used a C/10 rate. The time of charge was
128658 J.L. Allen et al. / Journal of Power Sources 196 (2011) 8656–8661

Table 1

Lattice constants for the series of Fe-substituted LiCoPO4 compounds. Numbers in

parentheses are the estimated standard deviation of the last significant digit.

Starting composition a (Å) b (Å) c (Å) Vol. (Å3)

LiCoPO4 10.1950(3) 5.9179(1) 4.6972(1) 283.39(1)

LiCo0.95Fe0.05PO4 10.1913(4) 5.9190(2) 4.6983(2) 283.42(2)

LiCo0.9Fe0.1PO4 10.1925(4) 5.9211(2) 4.6991(2) 283.60(2)

LiCo0.8Fe0.2PO4 10.1981(3) 5.9252(1) 4.6986(1) 283.92(1)

also limited to 10 h for C/10 rate and 5 h for C/5 rate so that during

the first few cycles the discharge capacity increased after the solid

electrolyte interphase (SEI) was formed on the cathode [9].

From Fig. 2, several points can be made.We first, used a standard

Li-ion electrolyte (1 m LiPF6 in 3:7 EC:EMC) to compare LiCoPO4

(blue open triangles) to the nominal Li0.92Co0.8Fe0.2PO4 composition (blue open squares). (For interpretation of the references

to color in text, the reader is referred to the web version of the

article.) For this case, it is clear that the Fe-substituted sample

demonstrates considerable improvementinreducing capacity fade.

However, capacity fade is still evident. Second, using a high voltage electrolyte (1 m LiPF6 in 3:7 EC:EMC + 1% HFiP additive) we

compare the capacity fade of the Li0.92Co0.8Fe0.2PO4 composition

(orange solid squares) to the same composition with the standard electrolyte (blue open squares). For this comparison, there

is additional decrease of the capacity fade with this change in electrolyte. Third, in order to discriminate fully between the effect of

the high voltage electrolyte and the substitutional effects, we compare the LiCoPO4 with standard electrolyte (blue open triangles)

to LiCoPO4 with the high voltage electrolyte (orange solid triangles). In this comparison, there is little discernible difference in the

fading. Both samples evidence rapid capacity fade. The electrolyte

has little effect. Thus, it is clear that structural decomposition of

LiCoPO4 or CoPO4 is primarily responsible for the discharge capacity fade of the LiCoPO4 electrode. In quantitative terms, about a

33% drop in capacity is observed between LiCoPO4 (blue open

triangles) and nominal Li0.92Co0.8Fe0.2PO4 (blue open circles) at

the 10th cycle using a standard electrolyte. The drop in capacity

between LiCo0.8Fe0.2PO4 with high voltage electrolyte (orange solid

squares) and Li0.92Co0.8Fe0.2PO4 with standard electrolyte (blue

open squares) is 12%. Thus, the capacity fade is mainly a result of

Fig. 2. Effect of LiCoPO4 modification and HFiP electrolyte additive. The “standard

(std.) electrolyte” is 1 m LiPF6 in EC:EMC, 3:7. The HFiP electrolyte is 1 m LiPF6 in

EC:EMC, 3:7 with 1% HFiP electrolyte additive.

Fig. 3. Long term cycling of composition Li0.92Co0.8Fe0.2PO4 using 1 m LiPF6 in

EC:EMC, 3:7 with 1% HFiP electrolyte additive.

LiCoPO4/CoPO4 structural decomposition and, to a lesser degree, a

result of electrolyte decomposition.

Fig. 3 shows the cycling performance of the Li0.92Co0.8Fe0.2PO4

composition over 500 cycles in a coin cell with Li metal as the

anode. The coulombic efficiency is about 97%. Although, a noticeable fade is still evidentitis a vastimprovement overprior literature

reports [4–6]. To reiterate, Tadanga et al. [4] observed a 10th cycle

discharge capacity of ∼52% of the initial capacity, Bramnik et al.

[5] reported a 10th cycle discharge of ∼59% of the initial capacity and Wolfenstine et al. [6] reported ∼53% capacity retention at

the 10th cycle. We observe approximately, 100% capacity retention

at the 10th cycle and about 80% capacity retention at the 500th

cycle. Fig. 4 shows the X-ray diffraction pattern of the cycled nominal Li0.92Co0.8Fe0.2PO4 cathode composite (Li0.92Co0.8Fe0.2PO4 with

carbon and PVDF) onAlfoil.All peaks can be assigned to the LiCoPO4

olivine structure, indicating structural integrity after cycling of

Fig. 4. X-ray diffraction of Li0.92Co0.8Fe0.2PO4 after electrochemical cycling. The

peaks are labeled with the Miller indices of the Pnma, phospho-olivine structure.
J.L. Allen et al. / Journal of Power Sources 196 (2011) 8656–8661 865913

Fig. 5. Unit cell volume as a function of the nominal substitution of Co by Fe in

LiCoPO4.

the Li0.92Co0.8Fe0.2PO4 strikingly different from the reported amorphization of LiCoPO4 during cycling [5,6].

Thus, we have shown the improved cycle life of Fe-substituted

LiCoPO4 and we assertthatit mainly results from the substitution of

Fe3+ on the Li and Co sites. In order to more fully characterize this

substitution we turn to X-ray diffraction, XPS, FTIR spectroscopy

and Mössbauer spectroscopy. First, from Rietveld analysis of the

X-ray diffraction data, the unit cell volumes of LiCoPO4 and Fesubstituted LiCoPO4 were determined. Fig. 5 shows the effect on

unit cell volume by the nominal substitution of Fe for Co2+ in

LiCoPO4. The observed linear increase in unit cell volume is consistent with the larger unit cell volume of LiFePO4 relative to LiCoPO4.

However, the line extrapolated to zero does not intercept at the

unit cell volume of Fe free LiCoPO4 as would be expected for

LiCo1−dFedPO4 where only Fe2+ substitution for Co2+ is observed.

The “extrapolated volume” is 283.26A ̊ 3 and the measure volume

is 283.39A ̊ 3. This smaller unit cell volume results from the partial

substitution of smaller Fe3+ for Li+ and Co2+.

Rietveld refinements were done to look at the anti-site defects,

e.g., Fe3+ or Co2+ on the Li site. The results are shown in Fig. 6. Hence,

this confirms that a small amount of Fe3+ is substituting at the Li

site (∼1.8 ± 0.5% for the nominal LiCo0.8Fe0.2PO4 composition). The

difference between the Li anti-site defect disorder concentration as

a function of Fe contentfalls within the measurement error. There is

on average for all samples about a 1.5 ± 0.5% concentration of antisite defects on the Li site. Pujana et al. [14] previously reported the

site preference of Fe3+ for the Li site in Li1−3xFexCoPO4. They did

not report Fe3+ substitution on the Co site. This difference from our

materials is a result of different starting compositions and different

reaction conditions.

Pujana et al. [14] also reported that the infrared (IR) spectra

of Fe3+ substituted samples have broadened peaks owing to an

increase of disorder resulting from the substitution of a portion

of the lithium ions by Fe3+ and the creation of 2 vacancies per Fe3+

atom. The IR spectra of Fe-substituted LiCoPO4 and LiCoPO4 are

shown in Fig. 7. We observed a small broadening upon substitution

of Li by Fe3+, in agreement with Pujana et al. [14] Furthermore, as

additional confirmation of Fe3+ in our samples, XPS revealed the

presence of both Fe2+ and Fe3+ in a sample of nominal composition LiCo0.8Fe0.2PO4. Two peaks were observed in the Fe2p spectra

which can be assigned to Fe3+ and Fe2+.

Mössbauer spectra and analysis were obtained from the SEE Co.

(Edina, MN) to corroborate the ICP data with respect to the relative

amounts of Fe3+ and Fe2+ and to learn about the coordination enviFig. 6. Percent anti-site disorder on Li site from Rietveld refinement of X-ray diffraction data.

ronment of the Fe2+ and Fe3+. The room temperature spectrum is

shown in Fig. 8. First, we identify the peaks as following: the doublet with the larger splitting (3.0 mm s−1) is typical high spin (S = 2)

Fe2+. The doublet with the smaller splitting (0.8 mm s−1) and shift

(0.44 mm s−1) is typical high spin (S = 5/2) Fe3+ [14]. Second, the

sharpness of the peaks gives information about the local environment. The Fe2+ lines are very sharp indicating that Fe2+ exclusively

sits at one site, the Co site of LiCoPO4 and locally all Fe2+ experiences a similar environment [15]. On the other hand, the lines of

the Fe3+ doublet are broad. The broadening of the peak results from

the differences in next nearest neighbors. The M2 (Co) octahedron

shares edges with two M1 (Li) octahedra and corners with four M2

octahedra. Thus, the local environment of the M2 site is mainly

controlled by the differences in the occupancy of the edge-sharing

M1 sites. An Fe3+ sited on the M2 will be adjacent to an M1 site

containing in order of likelihood either Li+, a vacancy, or Fe3+. Furthermore, vacant M1 sites will cluster around the Fe3+ therefore

Fig. 7. Comparison of the infrared spectra of samples of nominal composition

Li0.92Co0.8Fe0.2PO4 and LiCoPO4.
148660 J.L. Allen et al. / Journal of Power Sources 196 (2011) 8656–8661

Fig. 8. Mössbauer spectrum of composition Li0.92Co0.8Fe0.2PO4.

creating a more heterogenous environment for Fe3+ than for Fe2+.

Li and Shinno [15] described this “next nearest neighbor effect” in

detail for the mineral ferrisicklerite, Li1−xMn2+1−xFe3+xPO4, a mineral isotypic with the phospho-olivines. Third, from the area of the

peaks we can quantify the relative ratio of Fe2+/Fe3+. The Mössbauer

spectrum yields 60% Fe2+ and 40% Fe3+ in excellent agreement with

the ICP results. The analysis of the Fe2+/Fe3+ ratio by 2 independent

methods and 2 different laboratories is summarized in Table 2. All

are in excellent agreement.

Finally, having shown the improvements in capacity fade

because of the Fe substitution into LiCoPO4, we now show the

favorable effects of Fe substitution on capacity as a function of

rate in Fig. 9. Typical discharge curves for an electrode of nominal

starting composition LiCoPO4 (blue), Li1−ıCo0.9Fe0.1PO4 (orange)

where ı is the Li vacancy concentration and Li0.92Co0.8Fe0.2PO4

(black) are shown. Solid lines represent a C/10 discharge rate and

the dashed lines correspond to a 2C discharge rate. Several observations can be noted from Fig. 9. First and foremost, it is clear

that Fe-addition improves the capacity at both C/10 and 2C compared to pure LiCoPO4. Second, at the C/10 rate we observe that the

nominal Li1−ıCo0.9Fe0.1PO4 has the longest plateau at 4.8V. The

discharge capacity for the nominal Li0.92Co0.8Fe0.2PO4 is equal, but

a portion of the capacity is observed at around 3.5V, which corresponds to the Fe2+/Fe3+ couple. The LiCoPO4 discharge shows

evidence of polarization (lower average discharge voltage) compared to the Fe-substituted LiCoPO4. Turning our attention to the 2C

discharge curves, we observe that the nominal Li0.92Co0.8Fe0.2PO4

(black dashed curve) has the highest average discharge voltage and

the highest capacity. The nominal Li1−ıCo0.9Fe0.1PO4 and LiCoPO4

have similar discharge voltages though Li1−ıCo0.9Fe0.1PO4 has a

larger capacity. At 2C, the Fe3+/Fe2+ plateau is not observed for

either Fe containing composition.

Fig. 9. Discharge curves at different rates for starting compositions, LiCoPO4,

Li1−ıCo0.9Fe0.1PO4, and Li0.92Co0.8Fe0.2PO4.

The high rate performance may seem surprising owing to the

presence of anti-site disorder on the Li site. As discussed previously,

the nominal LiCo0.8Fe0.2PO4 composition has about 1.8 ± 0.5% disorder on the Li site. Intuitively, one would expect that the presence

of Fe on the Li site would impede the rate performance by blocking

the Li ion channels. However, we postulate that anti-site Fe does

not impede the rate for the following reasons. First, the anti-site

defect concentration in the Fe substituted samples is not much

greater than that reported for typical LiCoPO4 [16]. Second, the

anti-site defects will most likely cluster in some fashion. If the

defects preferentially occupy certain Li+ conduction channels they

may effectively leave other channels clear to enable Li+ conduction

leading to higher conductivity than if anti-site defects are present

in all the channels. The aggregation of defects is a well-known

phenomenon. For example, Unger and Perlman [17] showed that

divalent impurities in KCl and NaCl form impurity-vacancy dipoles

that aggregate in order to form dipolar pairs. The vacancy is a center

of excess negative charge and the divalent impurity has an excess

positive charge. This creates an electric dipole which will prefer

to interact with another dipole in order to collectively lower the

energy of the two defects. Similarly, in Fe-substituted LiCoPO4 the

Fe• •

Li is a center of excess positive and its associated vacancies, V�

Li,

are centers of excess negative charge and a lower energy state may

be achieved by clustering of these defect-vacancy complexes into

certainchannels leaving other channels free to conduct Li andhence

higher conductivity than anti-site defects in all channels. Finally,

the rate capability is also a function of the electronic structure and

so we further postulate thatthe rate capability ofthe Fe-substituted

LiCoPO4 is enhanced relative to pristine LiCoPO4 because of the

substitution of Fe for Co.

Table 2

Analysis of the Fe3+ content for sample of starting composition LiFe0.2Co0.8PO4.

Analysis method Atom % Fe2+ of total Fe Atom % Fe3+ of total Fe Laboratory

Mössbauer 60 40 SEE Co.

ICP-OES (from Li/M) 55 45 Galbraith

Laboratories, Inc.

Average ∼58 ∼42
15 J.L. Allen et al. / Journal of Power Sources 196 (2011) 8656–8661 8661

4. Conclusion

A synthesis method for Fe-substituted LiCoPO4 electrode material has been reported. The improved rate and dramatically reduced

capacity fade is striking. The capacity fade of LiCoPO4 results mainly

from structural decomposition of LiCoPO4/CoPO4 and to a lesser

degree results from electrolyte decomposition. Fe3+ on the Li and

Co sites appears to stabilize the structure. X-ray diffraction, XPS,

Mössbauer Spectroscopy, ICP-OES and FTIR spectroscopy confirm

the presence of Fe3+ on the Li and Fe3+/Fe2+ on the Co sites. Use of

the HFiP containing electrolyte further reduced capacity fade during charge–discharge cycling of the Fe-substituted-LiCoPO4. This

Fe-substitutionincombination withtheHFiP containing electrolyte

shows dramatically improved discharge capacity retention as well

as improved rate capability relative to un-substituted LiCoPO4 with

“standard” Li-ion electrolyte.

Acknowledgements

We thank ARL colleagues Kang Xu and Arthur von Cresce for

providing electrolyte, Unchul Lee for XPS measurements and Bruce

Poese for obtaining FT-IR spectra. We thank Thomas Kent (SEE Co.)

for obtaining and assisting with the analysis ofthe Mössbauer spectrum.

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[10] M.E. Rabanal, M.C. Gutierrez, F. Garcia-Alvarado, E.C. Gonzalo, M.E. Arroyo-de

Dompablo, J. Power Sources 160 (2006) 523.

[11] N.N. Bramnik, K. Nikolowski, D.M. Trots, H.E. Bramnik, Electrochem. Solid-State

Lett. 11 (2008) A89.

[12] A.R. Armstrong, M. Holzapfel, P. Novak, C.S. Johnson, S.H. Kang, M.M. Thackeray,

P.G. Bruce, J. Am. Chem. Soc. 128 (2006) 8694.

[13] R.H. Perry, D.W. Green, J.O. Maloney (Eds.), Perry’s Chemical Engineers’ Handbook, sixth ed., McGraw-Hill Book Company, New York, 1984, p. 3-149.

[14] A. Pujana, J.L. Pizzarro, A.A. Goni, T. Rojo, M.I. Arriortua, Anal. Quim. Int. Ed. 94

(1998) 383.

[15] Z. Li, I. Shinno, Miner. J. 19 (1997) 99.

[16] S. Adams, J. Solid State Electrochem. 14 (2010) 1787.

[17] S. Unger, M.M. Perlman, Phys. Rev. B 10 (1974) 3692.

17

Research@

Electrolyte Additive in Support of 

5 V Li Ion Chemistry

Arthur von Cresce and Kang Xu

U.S. Army Research Laboratory, Adelphi, MD

ABSTRACT

An electrolyte additive based on highly fluorinated phosphate ester structure 

was identified as being able to stabilize carbonate-based electrolytes on 5 V 

class cathode surfaces. The synthesis and structural analysis of the additive 

are described, and preliminary yet encouraging results from electrochemical 

characterization showed that this additive participates in forming a protective 

interphasial chemistry not only on transition metal oxide cathode at high voltage 

(~5 V vs Li) but also on graphitic graphite at low voltage (~0 V vs Li), making 

it possible to formulate an electrolyte supporting reversible Li+-intercalation 

chemistry in the coveted 5 V region.

Journal of The Electrochemical Society, Volume 158, Number 3, 

Pages A337-A342 (2011)
18

Electrolyte Additive in Support of 5 V Li Ion Chemistry

Arthur von Cresce* and Kang Xu*,z

Electrochemistry Branch, Power and Energy Division Sensor and Electron Devices Directorate,

U.S. Army Research Laboratory, Adelphi, Maryland 20783, USA

An electrolyte additive based on highly fluorinated phosphate ester structure was identified as being able to stabilize carbonatebased electrolytes on 5 V class cathode surfaces. The synthesis and structural analysis of the additive are described, and preliminary yet encouraging results from electrochemical characterization showed that this additive participates in forming a protective

interphasial chemistry not only on transition metal oxide cathode at high voltage 5 V vs Li but also on graphitic graphite at

low voltage 0 V vs Li, making it possible to formulate an electrolyte supporting reversible Li+-intercalation chemistry in the

coveted 5 V region.

© 2011 The Electrochemical Society. DOI: 10.1149/1.3532047 All rights reserved.

Manuscript submitted October 27, 2010; revised manuscript received December 1, 2010. Published January 21, 2011. This was

Paper 581 presented at the Las Vegas, Nevada, Meeting of the Society, October 10–15, 2010.

Li-ion chemistry is established upon reversible intercalation/

deintercalation of Li+ into/from host compounds.1 The voltage of

such an electrochemical device is dictated by the chemical natures

of anode and cathode, where Li+ is accommodated or released at

low potentials in the former and high potentials in the latter. In the

state-of-the-art Li-ion batteries, the operating voltage ranges between 3.5 and 4.2 V, as determined by the potentials of delithiated

olivine ironphosphates or transition metal oxides against that of

Li+-intercalated graphite, providing energy density between 110 and

175 Wh/kg, respectively.2 In pursuit of both a higher energy density

and a higher potential platform to deliver the energy, there have

been persistent attempts in quest of a new cathode chemistry that

can operate in the vicinity of 5 V, and various candidates have demonstrated this novelty, including LiCuxMn2−xO4   4.9 V,

3

LiNi0.5Mn1.5O4   4.7 V,

4 LiNixCo1−xPO4 4.8–5.1 V,5

Li2FCoPO4   5.1 V,

6 etc. However, a practical application of

any of these 5 V class chemistries is possible only with the removal

of a severe hurdle set by the electrolyte: the anodic stability limit of

the nonaqueous components, especially solvents, on highly oxidative surfaces of these charged cathodes. In today’s market, the majority of Li-ion batteries use organic carbonate esters as electrolyte

solvents, whose sustaining decomposition above 4.5 V vs Li sets an

intrinsic upper limit to the candidate cathode chemistry.7,8

Efforts were made in the past to seek nonaqueous solvent alternatives to carbonate esters, and anodic stability higher than 5 V was

reported for various systems, including asymmetric members from

sulfone family,9 and ionic liquids based on ammonium, imidazolium, and pyrrolidinium cations and bistrifluromethanesulfonyl

imide anion.10 However, the intrinsic weaknesses of sulfone-based

electrolytes, i.e., high viscosity and inability to form a protective

solid electrolyte interphase SEI on graphitic anode surfaces, have

either led to the formulation using low viscosity solvents, such as

linear carboxylate ester at the expense of anodic stability,11 or limited their use to applications where no graphitic carbon is used at the

expense of both operating voltage and energy density12; meanwhile,

ionic liquids introduced additional issues such as cathodic stability

of these onium-based cations on graphitic carbon, high viscosity,

and limited low temperature performance due to the supercooling of

the resultant electrolytes. Especially, because most ionic liquids still

rely on the use of fluorinated imide anions to ensure their low freezing temperature, there is a frequent concern about corrosion of the

cathode substrate at high potentials despite various efforts to achieve

passivation, while there is a constant concern about the cost.

In order to formulate an electrolyte composition that supports

5.0 V Li-ion intercalation cathode in combination with the low potential  0.20 V graphitic anode, we have explored various approaches, including design and synthesis of both new solvents and

additives. In the latter, we tried to minimize the disturbance to the

bulk composition of the state-of-the-art electrolyte systems by using

a trace amount of additives, which were already proven successful

on graphitic anode,13 with the hope to tailor an interphasial chemistry on the cathode surface that can protect the bulk electrolyte components from oxidative decomposition. In this short communication,

we report preliminary yet encouraging results obtained from such an

additive based on a phosphate ester with highly fluorinated alkyl

moieties.

Experimental

Synthesis of trishexafluoro-iso-propylphosphate

C3HF6O3PO or HFiP was conducted in a glassware setup

equipped with water-chilled condenser, addition funnels, and outlets

filled with silica drying agents. Thus, 0.45 mol 1,1,1,3,3,3-

hexafluoro-2-propanol C3H2F6O, Fluka was dissolved in 50 mL

diethyl ether Et2O at 0–5°C under continuous stirring, followed

by the addition of 0.45 mol lithium hydride LiH, Aldrich through a

solid-addition funnel. Once the generation of hydrogen stopped and

the reaction suspension became clear, 0.15 mol phosphorus oxychloride POCl3, Aldrich was introduced dropwise through an addition funnel. Precipitation of solid lithium chloride LiCl was observed immediately and persisted until the end of the reaction. Heat

generation at the initial addition of POCl3 was intense enough to

keep the Et2O refluxing; but near the late stage, the reaction mixture

had to be further refluxed by gently heating for 5 h to ensure the

completion of phosphorylation of alkoxides. Ether was removed by

distillation, and HFiP was purified through repeated distillations.

Purified product appeared as a clear needlelike crystal with melting

and boiling points of approximately 27°C and 50–55°C/3.0 mm

Hg, respectively. For a higher purity, HFiP can also be sublimed

under vacuum at 30°C. Electrospray ionization mass spectrum

EISMS and NMR spectra of 1

H-, 13C-, and 13F-nuclei were used

to confirm its chemical structure.

All procedures involving electrolyte solutions and cell assembly

were carried out in dry room, whose dew point is approximately

−90°C, or glove box, where moisture level is below 5 ppm. Ethylene carbonate EC, propylene carbonate PC, ethyl methyl carbonate EMC, and methylpropyl carbonate MPC were purchased

from Ferro, Corp., dried over molecular sieves, and stored in an

argon-filled glove box. The moisture level in the dried solvents was

below 10 ppm according to Karl–Fischer titration. Lithium

hexafluorophosphate LiPF6 from Stella Chemifa was used as received; 1.0 M LiPF6 in EC/EMC 30:70 by weight was selected as

baseline electrolyte, while a high voltage electrolyte was formulated

by dissolving 1% weight HFiP in the baseline.

CR2032 coin cells with aluminum-clad cap were used as testing

hardware, and 0.8 mm lithium foil cut to 12.7 mm diameter * Electrochemical Society Active Member.

z E-mail: cxu@arl.army.mil

Journal of The Electrochemical Society, 158 3 A337-A342 2011

0013-4651/2011/1583/A337/6/$28.00 © The Electrochemical Society

A337
19

1.27 cm2 was used as the counter electrode/anode. LiNi0.5Mn1.5O4

cathode material synthesized as reported was coated onto Al foil and

punched into 11.1 mm diameter 0.97 cm2 circular electrodes. Celgard 3501 with a hydrophilized surface was used as separator. All

test cells were assembled with a Hohsen automatic crimper. Galvanostatic cycling of the assembled cells was conducted on a Maccor Series 4000 Battery Tester, wherein a constant charge/discharge

current of 0.5 mA from 3.5 to 4.95 V with no rest in between

charge and discharge cycles. All cycling tests were performed at

room temperature 21°C.

Results and Discussion

Synthesis and structural identification of HFiP.— Because the

bulky fluorinated groups in hexafluoro-iso-propanol constitute spatial hindrance to the access of its hydroxyl, conventional approach

using trialkyl amine as an organic base catalyst usually lead to a low

yield  30% of the title compound HFiP,14 leaving behind a reaction mixture of corresponding mono- and biesters of phosphoric

acid. To ensure that POCl3 be completely alkylated at maximum

yield, we activated the nucleophilicity of the alcohol by converting

it into lithium alkoxide, which reacts exothermically with POCl3 at

much improved yields Scheme 1. The reaction yield of the alkoxide with POCl3 ranged averagely above 50%, as assessed from the

crude products.

HFiP can be purified by repeated distillation or recrystallization

in appropriate solvents; if desired, highly purified HFiP can also be

obtained by sublimation under high vacuum in a warm water bath.

HFiP exists as a crystalline solid at room temperature; above its

melting temperature, it acts as a viscous and weakly polar solvent

that shows negligible solubility toward most of the lithium salts,

likely due to the presence of high fluorine density in the molecule.

However, in most cases, it is miscible with typical nonaqueous electrolytes based on organic carbonate esters, especially when linear

carbonate such as EMC or DMC is the dominating cosolvent. A

phase separation occurs only when high HFiP concentration

10% is used in a solution where cyclic carbonate such as EC is

above 30%.

Purified HFiP was subjected to structural analysis, and Figs.

1a-1c showed the NMR spectra collected on this compound in deuteriated acetonitrile CD3CN. In addition to signals arising from the

solvent and internal reference 0.00 ppm, singlet, from tetramethylsilane; 1.95 ppm, multiplet, from CD3CN; and 2.15 ppm, singlet,

from trace moisture in CD3CN, there is only one NMR-anisotropic

proton as detected in 1

H spectra, which is represented by a collection of multiplets ranging between 5.6 and 5.9 ppm with complicated structures. Upon closer examination Fig. 1a, left side inset,

above, these multiplets seem to consist of two sets of heptets that

partially overlap with each other. We speculated that this split pattern correspond to the substructure as shown in Scheme 2a, where

the proton is simultaneously split by six neighboring 19F nuclei and

one 31P-nucleus, giving rise to total of 7  2 = 14 peaks. An estimation based on the heptet pattern yielded a coupling constant of

5.19 Hz, typical of a spin-spin coupling between 1

H and

19F-nuclei in vicinal positions.15

In order to simplify the spectra and confirm the source of these

spin-spin interactions, we decoupled the interaction between 1

Hand 19F-nuclei by exercising irradiation at the characteristic Larmor

frequency of the latter, and the above multiplets turned into a doublet Fig. 1a, left side inset, below, with a coupling constant of

11.99 Hz, typically observed between 1

H and 31P-nuclei.15 On the

other hand, 19F-spectra collected detected only one NMRanisotropic 19F-nucleus that is split into a doublet Fig. 1a, right side

inset, above, with a coupling constant of 5.26 Hz, which is in a

good agreement with the reciprocal coupling constant of 5.19 Hz as

estimated from the 1

H-spectra. Furthermore, the above doublet

merged into a singlet when 1

H-decoupling is conducted Fig. 1a,

Scheme 1. Synthesis of HFiP.

(a) (b)

(c) (d)

Figure 1. Color online a 1H-spectra of

HFiP in CD3CN without 19F-decoupling;

JH–F = 5.19 Hz; left inset, above: two sets

of 1

H-heptet due to simultaneous spinspin couplings with both 19F- and 31P-nuclei; left inset, below: 1

H-spectra

with 19F-decoupling; JH–P = 11.99 Hz;

right inset, above: 19F-spectra without 1

H-decoupling; JF–H = 5.26 Hz; right inset, below: 19F-spectra with 1

H-decoupling. These spin-spin couplings

along with individual coupling constants

confirm the substructure shown in Scheme

2a. b 31P-Spectra of HFiP with 1

H-decoupling in CD3CN using H3PO4 as

external reference; inset: 31P–1

H spin-spin

coupling with JP–H = 11.94 Hz confirming

substructure in Scheme 2b. c 13C-Spectra of HFiP with 1

H-decoupling

in CD3CN; inset: simultaneous 31P- and 19F coupling generate two sets of heptets

with JCa–F = 3.02 Hz, JCb–F

= 150.902 Hz, and JCa–P = 36.20 Hz, respectively, confirming substructure in

Scheme 2c.

A338 Journal of The Electrochemical Society, 158 3 A337-A342 2011
20

right side inset, below. This confirms that 1 there is only one

NMR-isotropic fluorine in the molecule, and 2 they are all located

at a vicinal position with a single NMR-isotropic 1

H-nucleus. Because of the remote locations of these two atoms in HFiP molecule,

there is no spin-spin coupling found between 19F- and 31P-nuclei,

which is consistent with the structure of HFiP as illustrated in

Scheme 1.

The 31P-spectra Fig. 1b also detected only one NMRanisotropic 31P-nucleus, which was found between 2.7 and 2.9 ppm

as a quartet and turned into a singlet at 2.862 ppm upon

1

H-decoupling irradiation. This 31P–1H spin-spin interaction has a

coupling constant of 11.94 Hz, again in a good agreement with

11.99 Hz of the 1

H–31P interaction. More importantly, the quartet

pattern with 1

H-coupling, which can only arise from a substructure

where 31P is surrounded by three NMR-isotropic 1

H-nuclei Scheme

2b, confirmed that the compound has to be a triester of phosphoric

acid, instead of partial alkylation products of phosphoric acid such

as mono- or biesters. 31C-spectra were shown in Fig. 1c, where, in addition to the

solvent molecules 1.33 ppm, singlet, and 118.20 ppm, multiplet for

CD3CN, two NMR-anisotropic 31C-nuclei were detected, which are

represented by a multiplet at 72.50–74.0 ppm for Ca and a quartet

between 123.86–118.41 ppm for Cb as shown in the substructure of

Scheme 2c. Note that the rightmost peak of the quartet overlapped

with the solvent signal at 118.41 ppm. The coupling constant between Cb and F is 150.902 Hz, consistent with the germinal spinspin interaction between a 13C- and 19F-nuclei. Of special interest is

the 13Ca

-nucleus at 72.50–74.0 ppm, which is obviously composed

of two sets of heptets inset, Fig. 1c. Because these twin heptets

were obtained under the condition of 1

H-decoupling, the pattern has

to arise from simultaneous splitting by six 19F- and one 31P-nuclei,

whose coupling constants are JC–F = 3.02 Hz and JC–P = 36.21 Hz,

respectively. This split pattern acts as the structural key that links the

F-nuclei in alkyl moiety and P-nucleus in phosphate moiety, thus

confirming the substructure represented in Scheme 2c.

Mass spectra were used to further characterize HFiP. As shown

in Fig. 1d, initial spectra failed to show the molecular peak m/e

548, which may indicate the reactivity of its radical cation with a

possible connection to its decomposition at the positively charged

electrode surfaces. The major peak at m/e 396.9 is a fragment arising from HFiP losing one of the fluoroalkyl side arms, probably

through a McLafferty-like rearrangement mechanism with the generation of a very stable radical Scheme 3.

On the other hand, there are multiple signals representing larger

molecular weight species m/e 660.8, 794.8, 1192, 1628, etc.,

which are apparently the dimeric, trimeric, and even polymeric

forms of HFiP or clusters of its fragments with new bonds formed.

The exact bonding configurations in these oligomeric species are

unclear at the moment and would require analytic tools that can

provide more detailed structural connectivity information. However,

in a general sense, this propensity of HFiP to polymerize may hint at

the preference of the molecule to form a network under oxidative

condition and even might be the source of its electrochemical behavior on a 5 V-class cathode surfaces.

Eventually, by employing a mild soft-ionization technique to prevent fragmentation, which is usually used to analyze fragile molecules, we successfully obtained the protonated and hydrated forms

of the molecule in high abundance M+H+, m/e 549 and M

+ H3O+, m/e 567, which are shown as inset of Fig. 1d.

Based on the structural information summarized above, we have

firmly established the chemical identity of the synthesized compound HFiP.

Rationale for tailoring cathode interphasial chemistry.— Since

the birth of Li ion chemistry, the central role of the electrode/

electrolyte interphase in this device has been recognized. However,

the majority of the investigations have hitherto focused on the anode

side for various reasons, the most important of which being the need

to stabilize the “fragile” graphene structure against a possible solvent cointercalation.16 The understanding of the oxidation mechanism of electrolyte components, especially the carbonate solvents,

on the surfaces of cathode materials has been scarce,7,17 and there is

still an on-going debate regarding whether an interphase such as the

one observed on graphitic anode surface also exists on the cathode,

despite spectroscopic detection of similar chemical deposits in the

latter. However, there is a consensus among researchers that the

commonly used nonaqueous electrolyte solvents based on organic

carbonate esters are not stable at potentials beyond 4.5 V vs Li,7,8,10

thus setting technical challenge to the application of a high voltage

Li ion chemistry.

With increasing interests in developing new Li ion chemistries of

higher energy density, researchers have resumed in recent years the

quest for 5 V class cathode materials, and several approaches have

been taken from solid chemistry perspective to mitigate the strongly

oxidative nature of these 5 V class cathode surfaces, such as latticedoping or surface coating,18 with various degree of success; on the

other hand, we believe that the most economical approach based on

the consideration of processing cost should be from an electrolyte

perspective, whose effectiveness has already been proven with the

various electrolyte additives nowadays widely employed by the Li

ion battery industry to modify the interphasial chemistry on graphitic anode. Similar to anode, the cathode in a newly assembled Li

ion cell is in “discharged state,” whose potential is only gradually

brought up to the “charged state” during the initial activation stage

of the cell; therefore, we envision that this stepwise potential increase would provide us the opportunity of tailoring the interphasial

chemistry on the cathode surface by selecting appropriate chemical

Scheme 2. Substructures derived from the spin-spin couplings as observed

in Fig. 1: a 1

H-nucleus is split simultaneously by six vicinal 19F-nuclei and

one vicinal 31P-nucleus; b 31P is split by three vicinal 1

H-nuclei; c 13Ca is

split simultaneously by six vicinal 19F-nuclei and one 31P-nucleus, while 13Cb by six germinal 19F-nuclei.

Scheme 3. Proposed fragmentation routes of HFiP under ionization state.

Journal of The Electrochemical Society, 158 3 A337-A342 2011 A339
21

species, whose sacrificial oxidation occurs before oxidation of the

bulk electrolyte components and whose oxidation products passivate

the catalytic sites of the cathode surface. Leveraging knowledge in

interphasial chemistry on graphitic anode as well as corrosionprotection science of metal surfaces, we have designed and synthesized several families of chemicals that have the potential to act as

protective passivates on transition metal oxide surfaces. In testing

the effectiveness of these chemicals on a host of a 5 V class cathode

materials, a member of the phosphate ester family, HFiP, stands out

conspicuously, reminiscent of the phosphate-treatment techniques

adopted for corrosion-prevention in steel industry,19 the corrosioninhibitor used to passivate aluminum surfaces in the high voltage

electrolytic capacitor,20 and similar modification of transition metal

oxide surfaces to modulate catalysis.21 Extensive tests on diversified

5 V cathode materials have confirmed the versatility of HFiP under

a high voltage constraint, which will be described in future publications. Here we selected to highlight the interphasial novelties of

HFiP on the surfaces of a candidate “5 V” cathode LiNi0.5Mn1.5O4

and a graphitic anode.

Electrochemical characterization on a 5 V cathode surface.—

Spinel structured LiNi0.5Mn1.5O4 with a redox potential between 4.5

and 4.7 V was selected as the primary testing platform for the high

voltage electrolyte formulated in this study. Figure 2a showed

its voltage profile in a cathode half cell constructed with

baseline electrolyte and Li as counter electrode and cycled galvanostatically between 3.5 and 4.95 V. Although reversible

Li+-intercalation/deintercalation was observed in the expected potential range, both the rapidly decreasing capacity utilization and the

increasing ohmic drop as characterized by the voltage hysteresis

between charging and discharging profiles indicated that there is a

severe and sustaining electrolyte decomposition on the cathode surface. The ac impedance spectra measured before and after cycling

have confirmed the cell resistance increase with cycling not shown

here, while differential capacity plots made for the 5th and the

100th cycle best represented the worsening irreversibility of the cell

chemistry with the widening separation between the lithiation and

delithiation peaks Fig. 2b. As comparison, with only 1% HFiP

present in the baseline electrolyte, both capacity fading and the

ohmic drop between charging and discharging profiles were minimized as shown in Figs. 3a and 3b. Figure 4 compared the cycle life

of LiNi0.5Mn1.5O4 in both baseline and high voltage electrolytes in

the span of 200 cycles, which is close to the limit set by the lithium

dendrite growth in a half cell. Although the capacity fading still

exists with HFiP-presence, there is already a significant improvement as compared with the baseline, and we are expecting that ad3

3.5

4

4.5

5

0 0.5 1 1.5

Capacity/mAh

80th 50th 1st

1.0 m LiPF6

in EC/EMC (30:70)

on LiNi0.5Mn1.5O4

3.4 3.6 3.8 4 4.2 4.4 4.6 4.8 5

Potential vs. Li/V

5 mAhg-1V-1

5th cycle

100th cycle

(b)

(a)

Figure 2. Color online a Voltage profiles of a 5 V class cathode

LiNi0.5Mn1.5O4 in baseline electrolyte LiPF6/EC/EMC 30:70. Only selected

cycles were shown for graphic clarity. b Comparison of differential capacity of LiNi0.5Mn1.5O4 in baseline electrolyte for the 5th and the 100th cycles.

3.4 3.6 3.8 4 4.2 4.4 4.6 4.8 5

Potential vs. Li/V

5 mAhg-1V-1

100th cycle

5th cycle

(b)

(a)

Figure 3. Color online a Voltage profiles of a 5 V class cathode

LiNi0.5Mn1.5O4 in baseline electrolyte with 1% HFiP as additive. Only selected cycles were shown for graphic clarity. b Comparison of differential

capacity of LiNi0.5Mn1.5O4 in electrolyte with HFiP as additive for the 5th

and the 100th cycles.

A340 Journal of The Electrochemical Society, 158 3 A337-A342 2011
22

ditional stabilization is possible with the combination of the current

electrolyte approach and the lattice stabilizations/surface modifications from the solid chemistry perspective.18

To further confirm that it was the presence of HFiP that stabilized

the interphase between the 5 V cathode and electrolyte, we took the

test of our electrolyte to an extreme condition, where only a linear

carbonate MPC was used as the neat solvent and no EC was present.

It was well known that most of those linear carbonate are not anodically stable beyond 4.0 V vs Li,8,9b and EC presence is required

not only for interphasial chemistry on graphitic anode but also for

anodic stability on the 4.0 V class cathode materials currently used

in Li ion industry. As shown in Fig. 5, LiPF6/MPC electrolyte did

not survive initial charging to 4.95 V and was never able to support

any Li+-intercalation/deintercalation chemistry between 4.5 and

4.7 V; instead, the voltage profile behave much like that of a

double-layer capacitance due to the high resistance built on the cathode surface and the lack of Faradaic processes only the first cycle

was shown. On the other hand, LiPF6/MPC with only 1% HFiP not

only provided the Li+-intercalation/deintercalation chemistry in the

range of 4.4–4.90 V but also, surprisingly, survived the prolonged

cycling with even a slight improvement in the polarization. While

the detailed mechanism of the decreasing resistance is still unclear,

it is unequivocal that HFiP must have participated in some interphasial process on spinel LiNi0.5Mn1.5O4. Based on its test on other

5.0 V cathode platforms to be published, this participation of HFiP

at a high potential might be universal.

Effect on graphitic anode.— In order for a new electrolyte additive to be used in Li ion system, its effect on every cell component

that it would be in contact with must be taken into consideration. In

the current case, the importance of graphitic anode is only second to

the cathode because it is known that the stacked graphene layer

structure is very sensitive to variation in electrolyte.22 We selected

mesocarbon microbead MCMB as the representative of the various

graphitic anode materials used in commercial Li ion batteries and

evaluated the potential impact of HFiP. Figure 6a showed the voltage profile of such an MCMB anode in both baseline and the high

voltage electrolytes formulated with 1% HFiP. As we have expected,

due to the small presence of additive, the difference between the

baseline and the high voltage electrolyte is negligible; if there is any

difference, both the capacity utilization and Coulombic efficiency in

the first cycle of HFiP-containing electrolyte are slightly better than

those of baseline. Upon closer examination, there seems to be an

early electrochemical reduction event in HFiP-containing electrolyte

at potentials around 1.2 V inset, Fig. 6a, which we tentatively

attribute to the initial reductive decomposition of HFiP; but at the

moment, we must caution that a conclusion on the role of this 1.2 V

event for a better SEI formed cannot yet be drawn with certainty, as

we have learnt in the past from the 1.7 V event of bisoxalatoborate

BOB anion.23 Further studies are needed to understand the nature

of this reductive process, and its impact on the performance of graphitic anode materials in electrolytes containing HFiP and related

additive materials.

To further reveal the cathodic property of HFiP on graphitic anode, we took it to another extreme condition, where PC was used as

the neat solvent. It was well known that graphitic structure is sensitive to solvent cointercalation, and PC is the most typical solvent

that fails to form a stable anode SEI to stop extensive grapheneexfoliation due to its cointercalation.8,22 The behavior of graphite

anode in a PC-dominant electrolyte is thus typically characterized

by a flat voltage profile at 0.7–0.8 V, where indefinite solvent decomposition proceeds without any Li+-intercalation chemistry involved, as shown in Fig. 6b. However, with only 1% HFiP presence

in PC, the above cointercalation/exfoliation process was quickly

suppressed by a new interphasial chemistry, and reversible

Li+-intercalation chemistry was eventually realized after a brief initial decomposition bump that corresponds to PC cointercalation. The

subsequent lithiation/delithiation for more than 30 cycles only the

initial four cycles were shown for clarity proved that this new interphasial chemistry, apparently with chemical signature of HFiP,

serves as a protective SEI as one usually formed in EC-containing

electrolyte. This unexpected merit of HFiP certainly adds on its

potential to be a part of the 5 V electrolytes.

It should be noted here that the above scenarios of employing

MPC and PC as neat electrolyte solvents were not meant to demonstrate either one as a practical system; rather, they created such

Figure 4. Color online Galvanostatic cycling performance of

LiNi0.5Mn1.5O4 in baseline and additive-containing electrolytes. Capacity retention was used for convenience of comparison, while the actual utilization

of LiNi0.5Mn1.5O4 is 98 mAh/h with baseline electrolyte and 108 mAh/g

with HFiP-containing electrolyte.

Figure 5. Color online Confirming HFiP-induced interphasial chemistry on

cathode surface under extreme condition: voltage profiles of LiNi0.5Mn1.5O4

in electrolytes based on a single solvent electrolyte using a linear carbonate

MPC with or without 1% HFiP. Only the first cycle was shown for

LiPF6/MPC because its subsequent cycles essentially showed no difference

from the first.

Journal of The Electrochemical Society, 158 3 A337-A342 2011 A341
23

conditions under which the effect of our target HFiP could be “exaggerated” to stand out. Apparently, both Figs. 5 and 6b served this

purpose perfectly.

Conclusion

Through electrolyte approach to stabilize nonaqueous electrolytes on a 5.0 V class cathode materials, we synthesized and identified an additive that is based on a highly fluorinated phosphate triester and carried electrochemical characterizations in a Li ion

environment. The preliminary yet encouraging results indicate that

the new additive not only significantly improved the anodic stability

of the electrolyte on a “5.0 V class” spinel LiNi0.5Mn1.5O4 but also

provided unexpected protective SEI chemistry on the graphitic anode. Further experiments are underway to understand how this additive participates in interphasial chemistries chemically and mechanistically and to evaluate its practical application in the 5.0 V full Li

ion batteries.

Acknowledgment

This project was partially funded by Department of Energy Advanced Battery Research DOE-ABR Program. The authors thank

Professor Amatucci of Rutgers University for providing cathode materials samples, and Dr. Yiufai Lam and Dr. Yue Li for providing

NMR services and Mass Spectrum Services, respectively.

U.S. Army Research Laboratory assisted in meeting the publication costs

of this article.

References

1. M. Winter and R. J. Brodd, Chem. Rev., 104, 4245 2004.

2. M. S. Whittingham, Chem. Rev., 104, 4271 2004; J. B. Goodenough and Y. Kim,

Chem. Mater., 22, 587 2010.

3. Y. Ein-Eli and M. F. Howard, Jr., J. Electrochem. Soc., 144, L205 1997; Y.

Ein-Eli, W. Howard, S. Lu, S. Mukerjee, J. McBreen, J. Vaughey, and M. Thakeray, J. Electrochem. Soc., 145, 1238 1998.

4. Y. Gao, K. Myrtle, M. Zhang, J. N. Reimers, and J. R. Dahn, Phys. Rev. B, 54,

3878 1996; T. Ohzuku, S. Takeda, and M. Iwanaga, J. Power Sources, 81, 90

1999; S. Mukerjee, X. Yang, X. Sun, S. Lee, J. McBreen, and Y. Ein-Eli, Electrochim. Acta, 49, 3373 2004.

5. J. Wolfenstine and J. Allen, J. Power Sources, 136, 150 2004; J. Wolfenstine, J.

Power Sources, 142, 389 2005.

6. D. Wang, J. Xiao, W. Xu, Z. Nie, C. Wang, G. Graff, and J.-G. Zhang, J. Power

Sources, In press.  10.1016/j.jpowsour.2010.10.021.

7. L. Yang, B. Ravdel, and B. Lucht, Electrochem. Solid-State Lett., 13, A95 2010.

8. K. Xu, Chem. Rev., 104, 4303 2004.

9. K. Xu and C. A. Angell, J. Electrochem. Soc., 145, L70 1998; K. Xu and C. A.

Angell, J. Electrochem. Soc., 149, A920 2002.

10. S. Martha, E. Markevich, V. Burgel, G. Salitra, E. Zinigrad, B. Markovsky H.

Sclar, Z. Pramovich, O. Heik, and D. Aurbach, J. Power Sources, 189, 288 2009;

V. Borgel, E. Markevich, D. Aurbach, G. Semraub, and M. Schmidt, J. Power

Sources, 189, 331 2009.

11. Y. Watanabe, S.-I. Kinoshita, S. Wada, K. Hoshino, H. Morimoto, and S.-I. Tobishima, J. Power Sources, 179, 770 2008.

12. A. Abouimrane, I. Belharouak, and K. Amine, Electrochem. Commun., 11, 1073

2009.

13. S. Zhang, J. Power Sources, 162, 1379 2006.

14. J. March, Advanced Organic Chemistry, 3rd ed., p. 357, John Wiley & Sons, New

York 1985.

15. R. Freeman, A Handbook of Nuclear Magnetic Resonance, 2nd ed., Longmans,

New York 1997.

16. M. Winter, Z. Phys. Chem., 223, 1395 2009.

17. K. Kanamura, S. Toriyama, S. Shiraishi, and Z. Takehara, J. Electrochem. Soc.,

142, 1383 1995.

18. J. Liu and A. Manthiram, Chem. Mater., 21, 1695 2009; J. Liu and A. Manthiram, J. Phys. Chem. C, 13, 15073 2009.

19. M. J. Franklin, D. C. White, and H. S. Isaacs, Corros. Sci., 33, 251 1992; G. Li,

L. Niu, J. Lian, and Z. Jiang, Surf. Coat. Technol., 176, 215 2004.

20. J. J. Randall, T. H. Nguyen, and C. E. Hutchins, U.S. Pat. 4,470,885 1984; S. D.

Ross, U.S. Pat. 4,541,037 1985; S. D. Ross and M. Finkelstein, U.S. Pat.

4,580,194 1986.

21. G. Ramis, P. P. F. Rossi, G. G. Busca, V. V. Lorenzelli, A. A. La Ginestra, and P.

Patrono, Langmuir, 5, 917 1989; M. Francisco, W. Cardozo, and Y. Gushikem,

Langmuir, 20, 8707 2004.

22. M. R. Wagner, J. H. Albering, K. C. Moeller, J. O. Besenhard, and M. Winter,

Electrochem. Commun., 7, 947 2005; K. Xu, J. Electrochem. Soc., 156, A751

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A117 2003.

Figure 6. Color online a Exploring HFiP-effect on interphasial chemistry

of graphitic anode: comparison of voltage profiles of graphitic anode in a Li

half cell with baseline electrolytes containing none or 1% HFiP. b Confirming HFiP effect on interphasial chemistry of graphitic anode under extreme

condition: comparison of voltage profiles of graphitic anode in a Li half cell

with electrolytes based on PC containing none or 1% HFiP.

A342 Journal of The Electrochemical Society, 158 3 A337-A342 2011

25

Research@

Interfacing Electrolytes with Electrodes 

in Li Ion Batteries

Kang Xu and Arthur von Cresce

U.S. Army Research Laboratory, Adelphi, MD

ABSTRACT

Since its birth in early 1990s, Li ion battery technology has been powering the 

rapid digitization of our daily life and finally made its debut in 2010 into the 

large format application for electrified vehicles such as the Nissan Leaf and GM 

Chevrolet Volt; however, much of the chemistry and processes underneath this 

amazing energy storage device still remain to be understood, among which is 

the interphase between electrolyte and electrodes. Interphases are formed in 

situ on electrode surfaces from sacrificial decomposition of electrolytes. Their 

ad hoc chemistry supports the reversible Li+-intercalation in both anode and 

cathode materials at extreme potentials, while preventing parasitic reductions/

oxidations on the reactive surfaces of those electrodes; but their existence 

places restrictions on energy and power densities of the device by impeding 

Li+-transport and setting operating voltage limits, respectively. It has been the 

dream of battery engineers to maximize the former and minimize the latter. 

This review summarizes the most recent knowledge about the chemistry and 

formation mechanism of this elusive battery component on both anode and 

cathode surfaces. The attempts to tailor a desired interphasial chemistry via 

diversified means were also discussed.

Journal of Materials Chemistry, Volume 21, Pages 9849-9864 (2011)
Interfacing electrolytes with electrodes in Li ion batteries

Kang Xu* and Arthur von Cresce

Received 9th December 2010, Accepted 7th February 2011

DOI: 10.1039/c0jm04309e

Aqui.

Onde a terra termina

e o mar ja na   ̃o comec ̧a

—A poetic definition of a binary liquid/solid interface by Luıs de Camo ̃es (1572).

Since its birth in early 1990s, Li ion battery technology has been powering the rapid digitization of our

daily life and finally made its debut in 2010 into the large format application for electrified vehicles such

as the Nissan Leaf and GM Chevrolet Volt; however, much of the chemistry and processes underneath

this amazing energy storage device still remain to be understood, among which is the interphase

between electrolyte and electrodes. Interphases are formed in situ on electrode surfaces from sacrificial

decomposition of electrolytes. Their ad hoc chemistry supports the reversible Li+

-intercalation in both

anode and cathode materials at extreme potentials, while preventing parasitic reductions/oxidations on

the reactive surfaces of those electrodes; but their existence places restrictions on energy and power

densities of the device by impeding Li+

-transport and setting operating voltage limits, respectively. It

has been the dream of battery engineers to maximize the former and minimize the latter. This review

summarizes the most recent knowledge about the chemistry and formation mechanism of this elusive

battery component on both anode and cathode surfaces. The attempts to tailor a desired interphasial

chemistry via diversified means were also discussed.

1. Introduction: Interface in electrochemical devices

To material scientists whose major interest lies with the bulk

properties, an interface is where one material ends and the other

begins; to electrochemists, however, the interface is the essential

component that separates electrochemistry from conventional

chemical reactions.1 It was because of these interfaces between

electrode and electrolyte materials that the otherwise random

and chaotic electron exchanges in a conventional redox reaction

are now forced to proceed in an orderly manner, producing

orientation movement by electrons (current) in an external

circuit, and a simultaneous movement by ions within the cell.

Thus, in any electrochemical device, these electrode/electrolyte

interfaces are the only ‘‘legitimate’’ locations where cell reactions

should happen; otherwise, the reactions are deemed ‘‘parasitic’’

and are usually irreversible and detrimental.

A cell reaction can be either spontaneous, as in a fuel cell,

primary battery or rechargeable battery during discharge, or

non-spontaneous, as in an electrolytic or electroplating cell, or

rechargeable battery during charge.2 In the latter case an external

electron source is needed to drive the energy-consuming reactions. During the electrochemical cell reactions, the electrode

materials are the only chemical reactants consumed or produced,

while electrolyte, at least in theory, must remain chemically inert

when providing ionic flow to balance the change in electroneutrality induced by the oxidation or reduction in the electrodes. In reality, however, the role of an electrolyte is by no

means so ‘‘inert’’ in cell reaction. More often than not, its

chemical composition is significantly interwoven with the

formation chemistry and the subsequent properties of the interfaces, which directly or indirectly dictate how well an electrochemical device functions. In some scenarios the electrolyte

formulation even determines whether a desired cell reaction can

occur electrochemically or not, and Li-based batteries present the

best examples of such scenarios.3

Since electron exchanges can only happen at electrode/electrolyte interfaces, in all the electrochemical scenarios these

interfaces always dictate the cell reaction kinetics, which, for a

battery, represents the rate at which energy can be released

(or captured if a rechargeable system is considered). For

a rechargeable battery, an interface also participates in determining whether the cell reactions can be reversed, and if they do,

with what Coulombic and energy efficiencies, which represents

the reversibility and the cycle life of the rechargeable device.

Furthermore, the interface could even indirectly affect the energy

density by setting limits of voltage window, within which the

devices can operate safely and reversibly. This bottleneck effect

of electrolyte becomes especially conspicuous in the efforts of

developing a coveted ‘‘5 V’’ class Li ion chemistry.

Electrochemistry Branch, Power and Energy Division, Sensor and Electron

Devices Directorate, U. S. Army Research Laboratory, 2800 Powder Mill

Road, Adelphi, MD, 20783

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In modern electrodics, the electrode/electrolyte interface

has been treated as a 2-D layer for mathematic convenience,

which is an approximation close to reality in most aqueous

electrochemical devices. In those devices the low operating voltage

ensures that most of the electrolyte components are thermodynamically stable on the surfaces of both cathode and anode

materials, therefore the interfaces mainly consist of both electrolyte solvents and solutes adsorbed on the surface of the electrode,

also known as the inner-Helmhotz layer. The thickness of interfaces is estimated to be less than 2 nm, so that trans-interface

electron-tunneling is possible, and kinetics of such electron

transfer process can be described by the Butler-Volmer equation.1

In the most common configuration of electrochemical devices,

both electrodes are solid and stationary, while the electrolytes

sandwiched between them are either liquid or solid, thus the

interface locations are usually stationary and well-defined.

However, there are certain electrochemical scenarios where

complications in defining electrolyte/electrode interfaces arise

due to the use of liquid or mobile electrode materials. One such

example is set by the flow battery that employs cell reactants

suspended in electrolytes and in constant movement relative to

current collectors. In these non-classical cases, the interfaces exist

only instantaneously at the surfaces of suspended particles that

are colliding with the current collector.4 Despite the complication, the essence of an interface remains unchanged nevertheless:

it is the demarcation between an electronic conductor (electrode)

and an ionic conductor (electrolyte).

2. Interphase in Li ion batteries

The attraction of lithium as the ‘‘Holy Grail’’ anode originates

from the fortuitous combination of three factors that can never

be duplicated by any elements in the periodic table: (1) its lowest

atomic weight as a metal providing the largest theoretical specific

capacity of 3860 mAh/g; (2) its smallest density (5.4  102 Kg

m3

) as a metal providing the highest possible gravimetric energy

density; and (3) its most negative electrochemical potential

(3.10 V vs. SHE) proving the highest possible cell voltage when

coupled with given cathode materials. To circumvent the

morphological catastrophe arising from lithium dendrite growth,

graphite intercalation compounds (GIC) of Li+ were used instead

of lithium metal, resulting in a loss of ca. 0.2 V cell voltage and

more than 90% original capacity.5 Nevertheless, when coupled

with a 4.0 V class cathode, which is also an intercalation host for

Li+, the Li-GIC still produce the rechargeable chemistry of the

highest energy density. The state-of-the-art Li ion battery is thus

based on the reversible intercalation/de-intercalation chemistries

of Li+ into/from the interstitial spaces in graphite anode and

transition metal oxide or phosphate cathodes, generating 100–

200 Wh/Kg energy at 3.5–4.0 V.6

Because of the reactivity of anode and cathode materials

employed in Li-based batteries, there is almost no known electrolyte that can remain thermodynamically stable on their

surface at charged states; alternatively, upon the initial activation

of a Li ion battery, there is always trace amount of electrolyte

sacrificially decompose, whose subsequent products deposit on

these active electrode surfaces, and passivate the catalytic sites,

preventing sustained decomposition of the bulk electrolyte.

Since the subsequent cell reaction, which requires Li+

-

intercalation/de-intercalation, can still proceed with the presence

of this passivating layer, its main characteristic is to insulate any

electron-transfer or tunneling between the bulks of electrode and

electrolyte but allow Li+

-migration.7,8 In other words, it acts like

a thin-layer electrolyte. Obviously this passivating layer,

different from both electrode and electrolyte both chemically and

morphologically, is no longer a simple 2-D existence as in the

classic interface; rather, it is an independent phase. Peled named

it ‘‘sold electrolyte interphase (SEI)’’ after its electrolyte nature.8

The crucial role of SEI in supporting the reversible Li ion

intercalation chemistry is now well recognized.

Perhaps the best example highlighting the importance of

interphase is the so-called ‘‘EC-PC disparity’’ and its impact on

the history of Li ion technology.9 GIC with Li+ as guest intercalant was first discovered in early 1950’s by reacting graphite

with either molten lithium metal or lithium vapor.10 Almost

immediately thereafter, efforts were made to prepare this GIC

electrochemically, with obvious hope to explore the possibility of

a new anode material. However in the following four decades,

despite the discovery of various intercalation hosts as cathode

candidates6 and the maturing understanding of solid intercalation chemistry,11 there was never a single success in preparing LiGIC through electrochemical means. The failures are almost

exclusively due to the lack of understanding in interphasial

chemistry on graphitic surface, and to most part, should be

attributed to one non-aqueous polar solvent used popularly at

the time, propylene carbonate (PC), whose methyl substituent

happens to render the resultant interphasial deposition nonprotective.9,12 Highly graphitic anodes thus exfoliate in electrolytes based on PC, which reductively decomposes at 0.70 V and

prevents any Li+

-intercalation chemistry from occurring. At

certain point the frustration even led to the belief that it was

beyond any possibility to prepare Li+-GIC electrochemically.

Not until 1990s did people realize that the above ‘‘insurmountable’’ obstacle can be easily resolved by simply using a close

cousin of PC, ethylene carbonate (EC).7 Eventually the interphasial chemistry based on EC catalyzed the birth of Li ion

battery in today’s configuration. In retrospect, it is without much

exaggeration to state that the difference between interphasial

chemistries, induced by a single methyl group, delayed the

emergence of Li ion technology by nearly four decades!

Like a classic interface, SEI in Li ion devices still significantly

affects – and dictates in some cases – the kinetics of cell reaction

(or power density of the device); but more importantly, SEI

determines the reversibility of the cell chemistry and hence cycle

life of the device. In a Li ion battery that employs intercalation

hosts as cathode and anode, Li+ is a limited source, and any loss

of Li+ is at the expense of capacity and energy density of the

device. The formation chemistry of SEI leads to such consumption of Li+

, which are converted into lithium salts insoluble in

electrolyte solutions and hence ‘‘electrochemically inactive’’.

Thus, the effectiveness of SEI is often measured by the amount of

Li+ consumed in the initial forming cycles as well as the average

loss in following operation cycles. These two parameters, known

as ‘‘irreversible capacity’’ and ‘‘capacity fading (rate)’’, respectively, determines how reversible the cell reactions proceed and

how long the service life of the device can last. In the past two

decades major efforts on developing new electrolytes and additives have been focused on improving these two parameters.

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28

An excellent kinetic analysis on SEI was conducted by Dahn and

coworkers, who developed a high precision coulometry technique

to investigate the irreversible loss associated with the reversible

lithiation and delithiation chemistries of various electrodes.13,14,15

They attributed this lost capacity at each cycle to the continuous

growth of SEI, based on the assumption that Li+ turns electrochemically inactive once becoming part of the interphase. It was

found that, at the given temperature, the thickness of SEI increases

with time1/2, while its growth rate is inversely proportionate to the

thickness. In this sense, decomposition of electrolyte to form the

interphase mathematically never ends with cycling, and these

reactions are actually dependent on howmuch time the electrolyteis

exposed to the electrode at theselithiation or delithiation potentials,

but independent from whether the cells are cycled or not. This study

revealed the ‘‘leaky’’ feature of the interphases, which were often

misinterpreted as a simple ionic conductor and electronic insulator.

In reality, electron-tunneling or transfer can still occur across these

interphases under certain conditions, one of which being sufficient

polarization potential, as evidenced by the use of redox shuttle

chemicals for the purpose of over-charge protection.

On the other hand, although these interphases are known to be

pervious to Li+ transport, Li+ conductivity across them was never

directly measured. A very elegant diagnosis done by Harris et al.

recently convinced us that Li+ ‘‘trapped’’ in the SEI are actually

rather mobile, although with lower diffusivity compared with

their counterparts in bulk electrolyte or electrode interior.16 They

formed SEI on a Cu substrate by using isotopic 6

Li-based electrolyte, which was then soaked in a regular 7

Li-based electrolyte

for various time intervals. Subsequent analysis by TOF-SIMS

technique found that most of the 6

Li+ ‘‘trapped’’ in the SEI can be

displaced by 7

Li+ in the electrolyte solution on a time scale of

hours. Thus, the compositions of interphases are ‘‘dynamic’’ in

this sense rather than ‘‘static’’. In an approximation these interphases can be visualized as a cation (Li+

) exchange layer, whose

framework is based on organic anions and which allows Li+-

conduction to proceed in a Grotthuss-like mechanism.

Apparently, SEI functions as a protective layer that prevents the

sustained decomposition of electrolyte components and safeguards

the reversible Li+

-intercalation chemistry; at the same time, this

new phase also presents a rather resistive component that is serial

in the complete circuit with the bulk electrolyte and electrodes.17 Its

bottleneck effect on mass-transfer becomes more pronounced at

low temperatures or high drain rates. Therefore, it is of great

significance to the Li ion battery industry to reduce the impedances

that is associated with the interphases while maintaining its

protective function, especially on the graphitic anode side.

The current state-of-the-art electrolytes used in Li ion batteries

are almost exclusively based on organic carbonate esters as

solvents, and lithium salts with inorganic fluorinated anions as

solutes, the results of compromise when a collection of requirements

that an ideal electrolyte should meet are applied.3 SEI formed in

such electrolytes can passivate both GIC anode down to the

potential lithium metal electrode (0 V vs. Li) and various oxide or

phosphate cathodes up to >4.0 V vs. Li. However, these interphases

cannot stabilize the carbonate-based electrolytes beyond 4.5 V limit,

constituting a severe challenge to the development of new Li ion

cathode chemistry that aims at delivering energy at higher voltages

in the neighborhood of 5.0 V.18 Unfortunately, the understanding

about formation chemistry of interphase on oxidative surfaces has

been rather limited as compared with the efforts made on interphases of reductive surfaces, especially graphitic carbons. The

recent thrust driven by U. S. Department of Energy (DOE) to target

new Li ion chemistry of 5 V is expected to divert more resources into

this under-investigated area.

2.1 On graphitic carbon anode

Despite the diversified cathode chemistries used by various

manufacturers, the anode material in today’s Li ion batteries stays

a constant, which is highly graphitic carbon.5 Li+ intercalates into

its inter-layer spaces at voltage ca. 0.20  0.05 V with high

reversibility, generating a capacity of 372 mAh/g. Compared with

cathode materials, in which the lattice is built upon strong Coulomic forces or covalent bonds, the layer structure of graphite is

held together only by weak van der Waals interactions between the

sp2

-hybridized graphene slabs, and is therefore rather susceptible

to co-intercalation of solvent molecules. It was for this reason that,

although Li+

-GIC was discovered in 1950s, its electrochemical

preparation came much later than the cathode counterparts

(which was discovered in late 1970s). It was also because of this

solvent-sensitivity of graphitic anode that the investigations on

SEI in the past two decades were mostly placed on anode side,

leaving the cathode surface to oblivion until recently. Thanks to

the monumental contributions from Aurbach et al.19–21 and

Besenhard and Winter et al.,

22,23 we now have gained substantial

understanding about the chemistry as well as the formation

mechanism of these interphases on graphitic anodes.

2.1.1 Chemistry. During the earlier unsuccessful attempts

to prepare Li+

-GIC in PC-based electrolytes, people identified

Li2CO3 as the major deposit on the exfoliated graphite surface,

and a two-electron reduction mechanism was proposed:12

in which the organic moiety is completely lost in the form of

gaseous product, leaving behind an inorganic interphase. However,

this mechanism was soon challenged by Aurbach et al.,

19 who

proposed that, considering the poor availability of electrons on the

graphite surface, an incomplete reduction should occur instead,

consisting of a radical anion as intermediate and subsequent intermolecular electron-transfer between the two radical anions:

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29

Averagely there is only one electron involved per PC molecule,

hence this mechanism is called ‘‘single-electron’’ reduction

pathway. Obviously, in this single-electron mechanism only half

(50%) of the organic moiety is lost in the form of gas, while the

rest is maintained in the final anionic product, known as either

alkylcarbonate or semi-carbonate, and becomes part of the

interphasial layer. Based on the extensive characterization work

carried out in Aurbach’s group and others, the spectroscopic

signatures of this class of compound were established, which

include the asymmetric carbonyl stretching at 1650 cm1 in

FTIR20,21 and the C 1s electron binding energy at 289 eV in

XPS.24 Since lithium salts of these alkylcarbonates are extremely

moisture-sensitive, their instantaneous conversion to Li2CO3

upon exposure to even trace amount of moisture explained the

origin of the two-electron mechanism proposed earlier, which

was most likely the consequence of poor moisture-exclusion

practice.25,26 The moisture-interference issue is of critical importance to all ex-situ experimental tools employed to investigate

the interphase in Li ion devices.

Having established the above single-electron reduction mechanism of carbonate solvents on graphitic anode surface, Aurbach

went further to suggest that the key ingredient of SEI on anode

side was actually such an alkylcarbonate from EC, dilithium

ethylene dicarbonate (LEDC), based on its overwhelming presence in spectroscopy:

This in part accounts for the fact that EC is an indispensable

electrolyte co-solvent in nearly all Li ion batteries on market

despite unique (and proprietary) electrolyte formulations used by

each individual manufacturer.

In more systematic effort to establish a baseline knowledge

about the interphase formed through reductive pathways, this

group at Army Research Lab (ARL) teamed up with scientists at

Lawrence Berkeley National Lab (LBNL) and synthesized these

alkylcarbonate compounds in vitro with high purity, and

collected the spectroscopy of these standard references to form

a reliable database.27,28 These alkylcarbonates consist of an entire

series of homologous lithium alkylcarbonates, lithium methylcarbonate (LMC), lithium ethylcarbonate (LEC), dilithium

propylene dicarbonate (LPDC) and LEDC, which are supposed

to form from the corresponding single-electron reduction pathways of all commonly-used electrolyte solvents in commercial

Li ion batteries, as shown in Fig. 1.

The selected NMR and FTIR spectra of these alkylcarbonate

compounds are shown in Fig. 2, from which rich structural and

morphological information can be extracted. For example, in

interpreting the presence of FTIR absorptions, ab intio computation on the chemical structure indicated that a better fitting to

the experimental data could be achieved if LEDC takes a dimer

or even multi-mer conformation,27 which was not observed for

LPDC although the latter is also a dilithium salt. It is very likely

that LEDC is more prone to form a network instead of

remaining a monomer in the SEI as well.

The real importance of the spectroscopic database lies in their

service as reliable references of comparison when the cycled

electrode surfaces are studied. Often, detailed chemical information can be extracted to either support or deny certain

chemical species that were suggested earlier as possible SEI

building blocks. For example, based on their dynamic function

theory (DFT) computation, Wang et al. once proposed that

alkylcarbonate of higher order (such as dibutylenecarbonate)

should be more thermodynamically favored through the singleelectron pathway and radical anion recombination.29 However,

the NMR spectra unequivocally excluded this possibility, as the

unique alkyl substructure should be easily detected. Instead, in

all the scenarios, simple alkylcarbonate were the overwhelming

Fig. 1 Model Compounds LMC, LEC, LEDC, and LPDC, as postulated reduction products from DMC, EMC, DEC, EC and PC, respectively, via the

single-electron reduction process on the graphitic anode surface (reproduced with permission by the Electrochemical Society from Ref. 26).

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30

products as supported by the reference spectra thus established.26–28

The value of the standard spectra can be best exemplified by

the comparison made between synthesized LEDC and the nickel

electrode surface polarized to lithiation potential in electrolyte

LiPF6/EC/DMC with 30 : 70.27 The chemical species in the latter

is almost an exact replica of the standard LEDC, confirming that

at the lithiation potential the single-electron reduction indeed is

the major process (Fig. 3). In addition to the below finger-print

match between spectra, what surprised us more at the time was

that LEDC was the only species detected on the electrode

surface, despite the major population of linear carbonate DMC

in the electrolyte composition. In closer examination, there is

obvious a ‘‘shift’’ between the bulk electrolyte composition and

the interphasial chemical species that originated from the various

electrolyte solvents. NMR spectra collected from the graphitic

anode surface in varying electrolyte compositions verified this

disparity. This finding led to the latter investigation of relative

competiveness of cyclic and linear carbonate for Li+

-solvation

shell and its implication in the eventual interphasial chemistry.

The reduction process on graphitic anode is more complicated

than what was detected on nickel surface. In addition to LEDC

that remains the major product in SEI, other species including

oxalate and alkoxide were also identified.30,31 The former (oxalate)

obviously resulted from CO2, but its source is still unclear. Under

the reductive environment at graphitic anode CO2 is possibly

reduced immediately and the resultant radical anion possesses lifetime long enough for a bi-molecular recombination reaction:

The latter (alkoxide) is possibly formed in a new two-electron

process that involves acyl-oxygen cleavage instead of the alkyloxygen cleavage as proposed by Aurbach et al.:

In a more recent effort to monitor the gas products generated

during the initial forming stage of Li ion batteries, Onuki et al.

proved the existence of this new process with the detection of CO.32

It is still unclear why such complications arose on graphitic

anode as compared with a simple electrode surface such as nickel.

We speculate that the rich functionalities on the graphene edge

sites (hydroxyls, alkoxides or quinones, etc) might have participated as mediators in the reactions, resulting in diversified

reduction pathways. Further investigation on this topic is

deemed necessary for a more thorough understanding of the

reduction chemistry on graphitic anode materials.

2.1.2 Formation mechanism. Unlike the chemistry of SEI, the

formation mechanism has been a subject under debate. The

initial models have treated the surface reduction like a simple 2DFig. 2 The 1

H- and 13C-NMR of the synthesized lithium 1,2-ethylene dicarbonate (LEDC, left), the proposed reduction product of EC, and

1,2-propylene dicarbonate (LPDC), the proposed reduction product of PC through single electron mechanisms (Reproduced with permission by the

Electrochemical Society from Ref. 26).

Fig. 3 The FTIR spectra collected from (a) synthesized standard reference LEDC, (b) surface species on a Ni-electrode that has been negatively

polarized to lithiation potential, and (c) Li(EC)x solvates (Reproduced

with permission by the Electrochemical Society from Ref. 27).

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surface process, while Besenhard et al. proposed a mechanism

that involved the unique structure of graphite anodes.22 Based on

their earlier studies of intercalation phenomenon of graphite,

they know that the electrolyte solvents, especially PC, tend to cointercalate and form the so-called ternary GIC. Thus it is

reasonable to speculate that the formation of SEI must also

experience a similar step, i.e., before the potential of graphite

anode becomes reductive enough to induce decomposition of

electrolyte solvent molecules, these molecules can co-intercalate

and exist in the graphene interlayer. In other words, the interphase thus formed, as schematically shown in Fig. 4, should

concentrate near the edge sites of the graphite particles and

partially penetrate the graphite interior. This formation mechanism is thus known as ‘‘3-D model of SEI’’.

During the past decade, with more studies focused on the SEI

of graphitic anode, the above mechanism gained more experimental evidence, while the most convincing proof came from the

in situ XRD performed on a graphite anode in PC-based electrolyte at fast scan down to 0.50 V vs. Li.33 The new series of

diffraction peaks appearing at 2q 1⁄4 24� and 27.5� obviously

indicate the reversible formation of ternary GIC with interlayer

distance as wide as 1.59 nm, which can accommodate solvated

Li+ with 3–4 solvent molecules in the primary solvation sheath.

Based on these experimental evidences, one can reasonably infer

that similar ternary GIC was also formed in commercial Li ion

batteries during the formation stage, where the main electrolyte

solvents transiently co-intercalate into the graphite interior.

A more recent refinement of the above 3-D model was made

by Xu et al.,

34,35 who discovered from the impedance studies

that the activation energy barrier (EAC) for Li+

-transfer

measured at electrolyte/graphite interface depends on electrolyte

composition, or the ratio between cyclic component EC and Li+,

in a rather particular manner. Using typical electrolyte solvent

mixtures of cyclic (EC) and linear carbonates (DMC or EMC),

they identified an almost universal dependence for these systems

with EAC reaching a near constant when EC/Li ratio is above 4,

which happens to be the average solvation number of Li+ in nonaqueous media.

The above correlation leads to the immediate suspicion that Li+

-

solvation sheath is the underlying cause. Cyclic carbonate molecules such as PC or EC, whose dielectric constants are 65 and 89,

respectively, should be the preferred members within the solvation

sheath of Li+ as compared with the linear carbonate molecules

such as DMC, DEC and EMC, whose dielectric constants remain

below 4. Considering that ternary GIC is the necessary ‘‘intermediate state’’ before an SEI is finally formed, and these inner

members of Li+ solvation sheath are most likely the preferred

precursor of interphasial chemistry, Xu et al. argued that in the

above 3-D model, the first wave of solvent molecules that cointercalate with Li+ are predominately cyclic carbonates as long as

there are enough in the bulk electrolyte solution to be recruited by

Li+

, and consequently the resultant SEI will preferably consist of

reduction products of these cyclic carbonates through the singleelectron reductive pathway. Assuming an average solvation

number of 4 for Li+

, the threshold solvent ratio for linear

carbonate molecule to be present in Li+

-solvation sheath (and also

subsequently in the interphase) is 30 : 70 by weight. In other

words, in majority of the commercial Li ion batteries, SEI would

consist of single-electron reduction product of EC because its

content was usually above that threshold ratio.

The above arguments concerning the ‘‘unsymmetrical’’ contribution from cyclic and linear carbonate molecules to interphasial chemistry have been receiving increasing support from

interphasial analysis efforts. For example, a NMR analysis on

Fig. 4 (a) Schematic drawing of a 3D SEI formed at graphite edge site via an intermediate of ternary GIC; (b) In situ XRD evidence of the ternary GIC

composed of Li(PC)4-Graphite (Reproduced with permission by Elsevier from Ref. 33).

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interphasial species as collected from cycled graphitic anode

corroborates that the alkylcarbonate originating from EMC does

not appear until the EC/EMC ratio is lower than 20 : 80;35 whereas

in a separate effort, Onuki et al. were able to trace the origin of

gaseous products during the forming stage of a Li ion battery by

labeling the cyclic and linear carbonates with isotopic 13C respectively, and found that the majority of alkene or CO produced were

from EC instead of linear carbonates, even though the electrolyte

used in the study is still dominated by linear carbonate DEC:32

All these analyses point to a formation mechanism that is

coordinated by the solvation sheath structure of Li+ in these

electrolytes.

Based on the above ‘‘solvation sheath’’ argument, the modified

formation mechanism of SEI on graphite surface was

schematically represented in Fig. 7, in which an electrolyte containing more than 30% EC will lead to an EC-dominated Li+

-

solvation sheath, and eventually an interphase that bears heavy

–– if not exclusive –– signature of EC-reduction products.

The Li+

-solvation factor might also play an important role in

the reduction kinetics of solvents molecules. According to

computational approach based on DFT,29 the reduction of a neat

solvent molecule is thermodynamically forbidden with a high

energy barrier for the initial ring-opening process; however,

coordination by Li+ significantly enhances the reduction by

stabilizing the high energy intermediates. It is very likely that the

bonds of the solvent molecules within the supermolecular

structure of a Li+

-solvation sheath are relaxed due to the effect of

a Li+ in their vicinity.

2.1.3 Efforts to tailor a desired chemistry. The interphase on

graphitic anode has been recognized as the main contributor of

cell impedance; therefore, in the past two decades, the Li ion

battery industry has been favoring a thinner interphase in order

to accelerate the Li+

-transport across this resistive component.

The most common approach to this thinner SEI is through

incorporating selected additives in electrolytes, whose sacrificial

decomposition before any other electrolyte components would

form an interphase of different chemistry.36 Since SEI thickness

on composite electrodes cannot be directly measured (or even

reasonably defined in some cases), the irreversible capacity in the

initial cycle was often used to quantify this ‘‘thickness’’ as

a measure for the effectiveness of those additives.

The additive approach was made possible by one characteristic

of Li ion battery, i.e., the electrode potentials before the initial

cell activation are at the discharged state, and open circuit

voltage between the lithiated cathode and non-lithiated graphite

usually lies below 1.5 V. Only during the initial forming was the

potential of graphitic anode gradually lowered to that of Li+

-

intercalation, accompanied by corresponding potential rise and

delithiation of cathode. This stepwise change of electrode

potential offered an opportunity to reduce only certain ‘‘chosen’’

species before the rest of the electrolyte solution are involved.

This selectivity is not possible with a metallic lithium electrode on

which reductions occur instantaneously and which is indiscriminate to any species it contacts.3

Due to insufficient understanding of the formation chemistry,

selection of these additives has been conducted on a semiempirical basis. Peled once suggested a kinetic model, assuming

that if a species can reductively decompose at faster rate on

graphite surface, it will most likely form an SEI of less thickness.37 Using reduction rate constants of chemicals in aqueous

media as a relative quantifying tool, he successfully established

such correlation for a series of chemicals. Later on an improved

model resorted to the thermodynamics instead of kinetics to

evaluate the additive candidates, arguing that the electrochemical potential at which the candidate decomposes is more

decisive than kinetic rate.38,39 Since the reduction/oxidation

potentials of a given species are more easily accessible via

computation than kinetic data, this latter model is now widely

adopted by researchers to screen and evaluate potential SEI

additives. Nevertheless, those two models share the common

belief that a more reactive species might be a better candidate for

interphase.

Fig. 5 ‘‘Universal’’ dependence of charge-transfer activation energy at

graphite surface on the cyclic/acyclic carbonate ratio in several common

electrolyte systems (Reproduced with permission by the Electrochemical

Society and the American Chemical Society from Ref. 34 and 35).

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In practice, computation approaches were usually taken to

estimate the reduction or oxidation potentials of candidate

species, which are related to the energy levels of lowest unoccupied molecular orbit (LUMO) and highest occupied molecular

orbital (HOMO), respectively. During the past two decades the

main attention were placed on how to improve the SEI on the

graphitic anode surfaces,36 therefore the additive species of

a lower LUMO are favored because it indicates a higher potential of reduction during the initial lithiation of graphite, hopefully before any component of the electrolyte is reduced. Such

examples include vinyl carbonate (VC),40 SO2,

41,42 ethylene sulfite

(ES),43 butyl sulfite,44 butyl sultone (BS),45,46 vinyl ethylene

carbonate (VEC),47–49 methylpropargyl sulfonate (MPS),50,51 etc.

In particular, VC, whose reduction potential starts at  2.0 V vs.

Li, can form a very thin interphase and has been widely used in

many commercial electrolyte formulations.52

However, despite its theoretical appearance, this HOMO/

LUMO approach is still empirical in nature, because there is no

necessary connection between the reduction potential and the

effectiveness of the resultant SEI, and the physicochemical

properties of the decomposition compounds, which cannot be

predicted by mere computation, were not taken into consideration. An ideal SEI additive is required not only to be reduced or

oxidized before any electrolyte component does, but also to be

able to decompose into such compounds that have minimum

solubility in the bulk electrolyte solvent, maximum Li+

-conductivity and thermal and chemical stability against all components

present, etc. Nevertheless, the HOMO/LUMO calculation did

provide an easy tool for the preliminary screening of additive

candidates, and numerous reports have described its successful

application in identifying useful additives.36,38,39

Most of the interphases in Li ion devices are formed in situ, i.e.,

naturally during the slow process of bringing potentials of

graphitic anode or cathode down to that of Li+ intercalation. On

the other hand, with increasing understanding about the chemical composition of interphases, there were also efforts to form

these interphases artificially with known chemicals that are

structurally similar to alkylcarbonates, either during the electrode processing stage or after it as a surface coating via various

deposition techniques. An obvious advantage of the artificial

approach is the minimization or even complete elimination of

Li+-consumption by irreversible processes during the otherwise

in situ SEI formation.

Since authentic ingredients of SEI on graphitic anode, lithium

alkylcarbonates, are sensitive to moisture and almost insoluble in

any non-aqueous solvents, early researchers used stable chemicals as surrogate interphasial building-blocks. Thus various

metal oxides or phosphates including Al2O3, ZrO2 and AlPO4,

etc, have been thus applied directly on graphite particles through

sol–gel processes,53–58 with improvements in either reduced

Fig. 6 13C-NMR spectra of interphasial species collected from the graphitic anodes cycled in electrolytes with varying ratios of cyclic/acyclic carbonates

(Reproduced with permission by the American Chemical Society from Ref. 35).

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irreversible capacity in the initial cycle, prolonged cycle life,

durability against cycling at elevated temperatures, or enhanced

safety behavior. Amorphous carbon layer was also applied in

a similar manner onto highly graphitic carbon surfaces.59

In a similar but much simpler approach, graphite particles

were soaked in Li2CO3 aqueous solution, which forms a thin

coating over the composite electrode upon drying.60 When tested

in a PC-rich electrolyte that is often a challenge to most highly

graphitic structures, it was found that the artificial interphase of

Li2CO3 not only reduces the irreversible capacity in the initial

forming cycle, but also significantly suppresses the self-delithiation rate of a charged graphite anode during long term storage at

elevated temperatures. A similar work, inspired by the discovery

that lithium bis(oxalato)borate (LiBOB) can form a unique

interphasial chemistry capable of withstanding high temperature,

employed aqueous solutions of mixed boric and oxalic acids.61

The interphase in this case is supposed to consist of an oxalatoborate network and also functions efficiently to protect the

graphitic anode from a PC-rich electrolyte.

Vacuum deposition techniques were also employed to

place desired chemicals onto graphitic anodes. Marrassi and

coworkers deposited various metal layers of sub-nano thickness

over partially-oxidized graphite via vacuum evaporation, and

observed improvement of electrochemical performances in all

cases but especially in the case of Cu and Sn.62,63 This increase in

metallic nature of graphite surface helps to reduce the resistance

to Li+

-transport, but more interestingly, their measurement of

the activation energy barrier seems to suggest that this

improvement in Li+-transport may come not only from the mere

increase of conductivity at the interphase, but more likely from

an accelerated Li+-desolvation process.63 To minimize the

thickness of the artificial interphase, Jung et al. used Atomic

Layer Deposition (ALD) technique to coat Al2O3 over the

particles of both anode and cathode materials, and demonstrated

that natural graphite thus treated can be cycled in PC-based

electrolyte without exfoliation.64 They claimed that ALD does

not block the inter-particle electric pathways as sol–gel may do,

thus keep the rate capability from being compromised at the

expense of protection.

It must be pointed out that, while the interphasial chemistry

and physical thickness in these deposition approaches are wellunderstood and can even be readily manipulated, there are

certain disadvantages that prevent their actual applications, the

most conspicuous of which is scalability and cost. Electrified

vehicle applications, in particular, require large format Li ion

batteries with quantities of materials which cannot be supplied

with any deposition methods carried out under vacuum. In

comparison, electrolyte additives provide a more economical

approach.

The pursuit of a merely thinner SEI has also been challenged

by the recent revelation of Li+-desolvation on Li+-transport.65

The possible role of Li+-solvation in interphasial processes have

been ignored until very recently, and the credit for bringing this

factor under light should go to Ogumi’s group,66–70 whose series

of pioneering work identified the desolvation process of Li+ as

a highly energy consuming process and initiated a sequence of

efforts which led to the discovery that Li+

-solvation sheath

structure plays an important role not only in SEI formation

mechanism but also in subsequent Li+

-transport across the

interphase.

Given its small ionic radius, a naked Li+ is known to exert

strong Coulombic forces on its surrounding solvent molecules,

and usually the innermost three to five molecules were considered

as the members of its primary solvation sheath, which stays with

the core Li+ without being displaced, at least statistically, while

Li+ migrates throughout the bulk electrolyte solution.1 Such

solvation stabilizes Li+ in electrolyte solution by compensating

the increase of Gibbs free energy when lithium salt is dissociated

from its original lattice. However, when such a solvated Li+

migrates to the edge sites of graphitic anode, the formed interphase would only allow naked Li+ through, therefore the solvation sheath must be stripped off before a naked Li+ can be

intercalated into the graphite interior (Fig. 7). Thus, during the

Li+-transport across the formed interphase, there exists a high

energy transition state, which corresponds to the generation of

a naked Li+ (Li+-desolvation) and migration of this naked Li+

subsequently through the interphase (Fig. 9, above). Since such

a desolvation process consumes energy to complete and would

Fig. 7 Refined model of the ‘‘3-D’’ formation mechanism in which Li+-

solvation sheath plays the central role of forming the intermediate ternary

GIC and dictates the initial ingredients of resultant SEI (Reproduced

with permission by the American Chemical Society from Ref. 35).

Fig. 8 Role of Li+-solvation on kinetics of EC reduction (Reproduced

with permission by the American Chemical Society from Ref. 29).

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naturally induce additional interphasial resistance to the Li+-

transport.

Unfortunately, the above two sub-processes are closely interwoven, and the so-called ‘‘interphasial resistance’’ in the earlier

literature has actually referred to their compounded effects

inclusively without differentiation. Not until Ogumi’s work was

it realized that the Li+

-transport consist of two sub-processes:

Li+-desolvation and the subsequent Li+-migration. The differentiation of these sub-processes was eventually made possible by

employing an intercalation anode, Li4Ti5O12,

65 whose surface is

free of any interphase because the reversible intercalation of Li+

occurs at 1.5 V vs. Li, a potential where no electrolyte

component is either reduced or oxidized. It was found that the

contribution from Li+

-desolvation to the activation energy

barrier at electrolyte/graphite junction is 50 kJ mol1

. By

comparing with the compounded energy barrier of 6070 kJ

mol1 as derived from a graphitic anode (Fig. 9), it was

concluded that Li+

-desolvation is the rate-determining step at the

electrolyte/graphite anode interphase.

In this context, the traditional approach of pursuing thinner

SEI certainly encounters its own limitation, because no matter

how thin an interphase is, as long as it acts as a layer of electrolyte nature at anode surface, i.e., forbidding any electrontunneling, the desolvation should always exist and its contribution to the interphasial activation energy barrier would

remain a constant. Considering the dominant role of Li+-desolvation to the overall cell resistance, we believe that any

further improvement in interphasial kinetics than what allowed

by current electrolyte additive approach would require a new

chemistry that assists in breaking the solvation sheath of Li+.

This group at ARL has been working on this concept and is

making progresses.

2.2 On Si and other alloy surfaces

The above 3D-formation mechanism of interphase is graphitespecific. When the electrode does not possess a structure for

intercalation, a simple 2D surface reaction is likely to be responsible for the formation of an interphase. Such examples

include metallic lithium or other metals as alloy host for lithium.

Although the resultant interphase no longer has the characteristic features as it does on graphite surface, i.e., primary location

at graphitic edge versus basal sites and partial penetration into

the bulk of interstitial spacing, the interphase thus formed is still

a 3D and independent entity.

Alloy hosts such as Si and Sn have attracted tremendous

interest as the next generation anode materials to replace

graphite, due to their much higher capability to accommodate

lithium; however, the accompanying huge volume expansion

(200%300%) associated with the lithiation/de-lithiation cycle,

as compared with <10% in the case of graphite, has been the

major hurdle preventing their practical application. The severe

stress during the volume change brings mechanical disintegrity

and results in loss of active materials (Fig. 10).71,72 Meanwhile,

both formation and chemistry of SEI is also perpetually overshadowed by the above volume change. Considering the potentials at which those hosts alloy with lithium (<200 mV vs. Li),

electrolyte components, especially carbonate esters, will inevitably be reduced into species that are structurally similar to

what identified on graphitic anodes. However, since fresh metal

surfaces are created repeatedly in each cycle due to the local

disintegrity, additional reduction of electrolyte has to take place

to cover those newly exposed surfaces, similar to what happens

when a metallic lithium electrode is being cycled in non-aqueous

electrolytes. This continuous process builds up interphasial

resistance, consumes the limited source Li+, and would eventually disable the cell, if it has not been disabled by the sheer loss of

active species. Thus, one of the keys to the successful application

of Si or other alloy hosts as Li ion battery anode is that the SEI

formed thereon must withstand a higher mechanical stress.

To circumvent the problem caused by volume change, Si

electrodes with nano-structure or in thin-film configuration

(#200 nm) are often employed in place of bulk configuration, as

illustrated in Fig. 10.73–81 With the absence of binder and carbon

additives that are often present in graphite composite electrodes,

the spectroscopy of these electrodes can usually provide

compositional information free of interference. It seems that

performance and possibly interphasial chemistry of these anodes

Fig. 9 Differentiating contributions to the interphasial resistance from

two sub-processes: (a) generation of naked Li+ (Li+-desolvation) and (b)

Migration of naked Li+ across the interphase (Reproduced with

permission by the American Chemical Society from Ref. 65).

Fig. 10 Schematic illustration of (a) bulk and (b) nano-structured Si

during the electrochemical lithiation/delithiation (Reproduced with

permission by the Nature Publishing Corp. from Ref. 74).

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are sensitive to the synthesis conditions;82–84 however, a constant

interphasial composition is always detected on Si: the fluorinated

derivatives SiFx, which is obviously arising from the reaction of

PF6

-based Li salt and whose formation is thermodynamically

favored, driven by the high bonding energy of Si–F (565 kJ

mol1

).85,86 Given the potential as well as the electron-poor

characteristic of Si that are similar to graphite, single-electron

reduction products from electrolyte solvents were also identified,

which include inorganic species such as Li2CO3, LiF, and Li2O,

and organic species such as alkylcarbonates, carboxylates, and

polyethylene oxide.71–86 What surprised the researchers is that

CFx was also detected by XPS, which was not present on

graphitic anode. It has been speculated that HF, produced due to

the hydrolysis of PF6

-anion by trace moisture, attack on both Si

and SiO2, and SiF6

2 resulting from HF also catalyzes formation

of CFx species. Interestingly, both SiO2 on older surfaces and Si

on freshly exposed surfaces should compete for HF, but it is

difficult to tell from spectroscopy which process or combination

results in the SiFx seen in XPS.86 For Si–C alloys and carboncoated silicon, it is a reasonable assumption that interphases

thereon are more similar to those formed on graphite with a lack

of SiFx species,85–88 because the surface functionalities mediating

the decomposition of the electrolyte are graphite or graphite-like

rather than Si lattice.

Significant presence of siloxane was also found in the inner

layer of the interphase, which was probably formed by the

reaction between solubilized Si with carboxyl radicals in solution

and re-deposits on the carbon layer during initial SEI formation.86 The presence of siloxane in the SEI has been shown to

benefit the reversibility of Si electrodes,77 making this a potentially useful improvement to the integrity of the silicon SEI layer.

In general, the presence of –Si–C– and –Si–O–C– terminated

components bonded with the anode surface would imply some

degree of SEI stability, at least at the interface, due to the

increased stability of –Si–C– and –Si–O– (318 and 452 kJ mol1

)

compared with their carbon counterparts –C–C– and –C–O–

(346 and 352 kJ mol1

).

Since the presence of moisture (and hence HF) is inevitable in

non-aqueous electrolyte that employs fluorinated Li salt anion,

the SEI on a Si anode has the additional responsibility of

shielding the surface underneath from HF attack. Both Si and its

native oxide film are vulnerable to HF, which convert Si into

electrochemically-inactive SiF6

2 and results in capacity loss.89

Compared with graphite surface, the interphase on Si-containing

alloy hosts have much higher presence of P–F moieties for this

reason, which might have originated from PF6

 anion reduction,

HF formation, and resulting incorporation of PF6

 products into

the network. Si might have directly participated in pulling F from

PF6

 anion.

Analysis on a thin-film composite Cu2Si alloy host revealed

evidence of alkyl esters and carbonates with only a weak signal

from semi-carbonate compounds like LEDC, very distinct

from what is found on graphite and Si anodes. This difference

is most likely related to the increased availability of electrons

on the Cu-rich anode surface and may result in more twoelectron reduction products in the interphase layer. More

importantly, this result points to the significance of the surface

composition itself in controlling the composition of its interphase layer.78–80

A number of researchers have also attempted to manipulate

interphasial chemistry on Si by using electrolyte additives.

Lithium bis(oxalatoborate) (LiBOB) was used as a minority salt

because of its well-studied network forming abilities,90,91

and was able to maintain the spectroscopic presence of semicarbonate species in the SEI through reductive decomposition

of the bis(oxalate) borate anion.75 Monofluoroethylene

carbonate (FEC), on the other hand, presumably provides

a more fluorinated SEI layer built on the fluoro-LEDC derivatives.76 Both efforts were reasonably successful in maintaining

the electrochemical reversibility of a thin (<200 nm) thick Si

anode, though only reported for less than 100 cycles, leaving

questions about the long-term stability of these SEI layers.75

Similar work also claimed that addition of 5 wt% trimethoxymethyl silane extended the cycling life of a thin-film Si anode,

though the mechanism is unclear.77

For bulk Si and other alloy anodes, as opposed to their nanostructured or thin-film configurations, interphase is still at the

mercy of the cyclic volume change, and no SEI is known to

inhibit this process. Further studies are needed to understand the

failure modes of interphase on these alloy host electrodes, and

interphases with unique chemistry and morphology dedicated to

these alloy anodes should be tailored to accommodate the new

requirements.

2.3 On cathode surfaces

Since the birth of Li ion chemistry in early 1990, most research

resources were invested in the pursuit of diversified new cathode

chemistries; however, a complete reversal exists in the resource

distribution of interphasial studies, and the chemistry on cathode

surfaces remains a little-understood regime as of today.

Similar to the scenario with graphitic anode, the potentials of

most cathode materials used in Li ion batteries are too high

(>3.5 V vs. Li) for almost any electrolyte component to remain

thermodynamically stable. Thus the reversible Li+

-intercalation/

de-intercalation must be preceded by the formation of certain

interphasial existence, although it is expected that the varying

composition, lattice and morphology of cathode materials would

render the chemistry and formation mechanism very distinct

from (and much more complicated than) the process on their

graphitic counterpart.

While the interphase on graphitic anode is usually more

resistive than that on cathode, the latter tends to increase in

magnitude at quicker pace than the former during long term

cycling, and would eventually take over the role as the most

resistant component at certain point. Cycling or aging at

elevated temperatures accelerates this process. The interphase

formed on the oxidizing surfaces is under-investigated for

historical reasons. While the overwhelming emphasis on

graphitic anode can be attributed to graphite sensitivity toward

solvent molecule co-intercalation, i.e., driven by ‘‘necessity’’,

several factors are responsible for the scarcity of knowledge on

cathode interphase, including the interference of a native

surface film of Li2CO3 that exists on most of the transition

metal oxides, the reaction between this native layer with acidic

electrolyte, and the entanglement of cathode species decomposition with the oxidation of organic electrolyte components.

Recent studies showed that even olivine phosphate, which

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operates at lower potential (3.5 V) and often considered

chemically ‘‘inert’’ than oxides, can react with both ambient

moisture and electrolytes based on acidic LiPF6.

92

Proving the existence of a SEI on cathode surface is not as

straightforward, and surface spectroscopy tools that have been

effective for graphite surface yield ambiguous results more often

than not. Till today there are still occasional arguments that an

SEI in traditional sense does not exist on cathode side. However,

more and more attention is drawn toward the cathode side,

especially after a DOE-wide effort was started to pursue 5 V-class

new Li ion chemistry, which requires in-depth understanding

about the oxidation chemistries of electrolyte components.

2.3.1 Oxidative processes on cathode. Unlike the interphase

formation process on graphitic anode, Li+

-solvation structure

plays only negligible role on cathode surface during the initial

activation of a Li ion cell, primary because this charging process

involves the removal of Li+ from the cathode lattice, and corresponds to Li+

-solvation rather than desolvation. Thus, the

oxidation no longer has the preference for those solvent molecules in the solvation sheath, but rather offers ‘‘equal opportunity’’ to all electrolyte components that are adsorbed on the

cathode surface.

Similar to the reductive decomposition, it is speculated that

a cation radical is involved as an intermediate, and the only

experimental evidence to the best of our knowledge is the ESR

observed on the electrolyte in contact with a charged LiCoO2:

93

The fate of the radical cation was not clear, and there was

suggestion that interphase bears chemical similarities to its

counterpart on graphite anode,94,95 but polycarbonate has been

suggested as a possible oxidation product based on surfaceenhanced FTIR studies:96

Based on an earlier study of polymerization of EC, there is

suspicion about the chemical stability of polycarbonate due to

the high density of carbonyls in the monomeric units;97 and more

likely, the polymer would lose part of the carbonyls in the form

of CO2 and result in a copolymer of carbonate and ethylene

oxide units. The detection of CO2 during cathode aging at

elevated temperatures seems to support this possibility.98,99

Nevertheless, a recent study on the EC oxidation on ‘‘5 V’’

cathode LiNi0.5Mn1.5O4 using both XPS and FTIR identified

polycarbonate species as the main product.18

Whatever the chemical species were, it is certain that their

stabilization of the electrolyte/cathode junction is valid only

below 4.5 V, above which sustained electrolyte oxidation would

occur.18 This upper limit set by the interphasial chemistry originating from organic carbonate solvents has been a major roadblock to the efforts to pursue Li ion chemistries of higher voltage

quality (5.0 V).

The thickness of SEI on cathode has also been a subject of

controversy. In general the interphases were considered as much

thinner when compared with its counterpart on graphitic anode.

Reports have placed a wide range from a few to 102 nm.100–102

However, with continuous cycling and especially high temperature aging, the interphase increases in thickness, while its resistance takes over as the kinetic bottleneck. Dupre et al. carried out

a rather interesting study on the interphase formed over

LiNi0.5Mn0.5O2 by using 7

Li-magic angle spinning NMR technique.102 By soaking the cathode materials with electrolyte, they

monitored the evolution of 7

Li-signals and discovered that the

reaction between active species and electrolyte solvents happens

in the timescale of 30 s, accompanied by the global disappearance

of the native Li2CO3 film. The eventual species identified include

decomposition products from PF6

 anion, polymeric or organic

species, and also alkylcarbonate salt that is a popular species

often found on graphitic anode surface. Combined SEM/TEM

analysis estimated an average of 2–20 nm thickness for the film

after the reaction, but since the cathode surface is not covered by

a homogenous layer, they thus argued that this reaction is not

a ‘‘real passivation’’.102

2.3.2 Solvents and additives for cathode passivation. While the

anions of lithium salts have been recognized to participate in the

passivation chemistry on surfaces of cathode materials and their

substrates,3 the hope of improving anodic stability of electrolytes

has always been with the solvents. This is mainly due to the

higher difficulty associated with developing a new lithium salt,

whose qualification must be subject to many more criteria

beyond oxidation chemistry. Instead, modern chemical reservoirs offer a rich source of solvents and especially additives with

diversified structures and redox properties for considerations of

tailoring an interphase with tolerance of higher potential.

With the appearance of 4 V class cathode chemistry based on

transition metal oxides, the selection of solvents has been

increasingly focused on organic esters, especially dialkyl esters of

carbonic acid. A few attempts to seek after alternative structures

have yielded varying results.

Sulfones were perhaps the first family of solvents shown to

have better anodic stability over the carbonates. By rendering the

molecules unsymmetrical so that their melting points are

depressed to the vicinity of room temperature, Xu et al. described

a series of sulfone compounds whose solution of lithium salts can

withstand potentials beyond 5.5 V vs. Li on the surface of

LixMn2O4, as shown in Fig. 11a.103 More interestingly, when the

sulfones were mixed in different ratios with acyclic carbonates

DMC and EMC in order to reduce the viscosity of the results

electrolyte, it was found that the presence of linear carbonates

has negligible effect on the anodic stability (Fig. 11b),104 which is

counter intuitive with the common knowledge that these

carbonate are not stable beyond 4.0 V and would compromise

the overall oxidation stability limit. This indicates that, even in

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the presence of linear carbonates at high ratios, the interphasial

chemistry on cathode surface is still dominated by the sulfone,

which is the more polar co-solvent in the system. In other words,

the more polar component seems to have more significant

contribution to the formation of interphase, in striking similarity

to what observed on graphite surface where the more polar EC is

favored. Considering the oppositely charged surfaces as reaction

locations in those cases, one would infer that the unsymmetrical

contribution to interphasial chemistry must be related to the

stronger interaction of the more polar solvent molecules with

either cation (Li+) or anion (PF6

). This phenomenon certainly

merits further investigations to understand the interphase

formation mechanism.

The composite electrode based on LixMn2O4 used in the above

work serves as a more convincing working electrode in demonstrating anodic stability limit than does non-porous surfaces such

as glassy carbon or platinum; however, it is not a real high (5 V)

voltage cathode, and practicality of sulfones needs further

confirmation. In more recent efforts, these sulfone solvents were

revisited with either carboxylate or carbonic esters act as cosolvents, and cycling results on a 5 V class cathode based on

LiNi0.5Mn1.5O4 led to the conclusion that anodic stability superior to neat carbonates was obtained.105,106 Especially, Abouimrane et al. demonstrated impressive stability in the prolonged

cycling of a ‘‘5 V’’ class cathode cell using a sulfone-linear

carbonate mixture (Fig. 12), confirming that the presence of

acyclic carbonate EMC co-solvent does not compromise the

oxidation stability at those high potentials.106

Organic nitriles represent another family of solvents that were

investigated as alternative to carbonates.107–111 Nitrile functionality is one of the few organic compounds that are polar enough

to dissolve various alkali salts (including most lithium salts

whose lattice energy ranks top) but still remain aprotic, and there

have been numerous reports on their use in electrochemical

capacitors, the most popular representative being acetonitrile.

There have been conflicting data regarding its electrochemical

stability, and the source of ambiguity seems to be the manner in

which such stability is determined.107 However, several high

order derivatives of nitrile have been tested on different cathode

surfaces and seem to remain stable at potentials near 4.0 V, such

examples including acrylic acid nitrile on LixMn2O4,

108 and 3-

ethoxypropionitrile, 3-(2,2,2-trifluoro)ethoxypropionitrile,109

and adiponitrile110 on LiCoO2. The only dinitrile molecule that

was tested on a true ‘‘5 V’’ class cathode is sebaconitrile, which

was reported to be able to support Li+ deintercalation above

5.0 V in Li2NiPO4F, although no cycle life data was presented.111

Despite the electrochemical stability window of 6.0 V as

measured on non-porous electrodes and positive effect on cycling

behavior with these cathodes, whether they can truly withstand

oxidation on surfaces of 5 V cathode species still remain to be

investigated.

Compared with the efforts of developing alternative bulk

solvents to replace the carbonate solvents, the additive approach

proves more convenient and economical and therefore more

likely to be acceptable to the industry, because their presence in

the electrolyte system is often too small to cause any undesirable

impact on the already-established systems in terms of both

properties and costs. Additives to improve the interphases on

graphitic anode surfaces have been widely used in Li ion battery

industry, while the additives designed to tailor cathode surface

has been very rare. The few such additives found in the reports

are the natural extension from the work on anode side. For

example, vinylethylene carbonate (VEC), which has been

known as additive to form polymeric species on graphite surface,

was also tested as an additive on cathode surface, and reported

to stabilize the electrochemical performance of LiNi0.8-

Co0.2O2;

112,113 while furan- and lactone-derivatives were also

described as being able to inhibit electrolyte oxidation on

Li1.17Mn0.58Ni0.25O2 surface at potentials up to 4.9 V.114

Recently, spinel LiNi0.5Mn1.5O4 was more and more targeted

as a reliable ‘‘5 V’’ class cathode candidate, and efforts aiming

at stabilizing conventional carbonate-based electrolytes on this

cathode began to take momentum. Cresce et al. identified an

additive, tris(hexafluoro-iso-propyl)phosphate (or HFiP), which

is based on phosphate ester structure with highly fluorinated

alkyl arms, whose presence in a carbonate-based electrolyte at

merely 1% can effectively stabilize both interphasial impedance

and capacity retention of spinel LiNi0.5Mn1.5O4 cathode

Fig. 11 (a) Anodic stability of various sulfone-based electrolytes on

spinel LiMn2O4 cathode surface obtained by linear sweep (Reproduced

with permission by the American Chemical Society from Ref. 104). (b)

Effect of linear carbonates on anodic stability of sulfone-based electrolytes on spinel LiMn2O4 cathode surface (Reproduced with permission by

the American Chemical Society from Ref. 104).

Fig. 12 Cycling of a ‘‘5 V’’ cathode LiNi0.5Mn1.5O4 in an electrolyte

based on the mixture of tetramethylene sulfone (TMS) and EMC

(Reproduced with permission by the Elsevier from Ref. 106).

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(Fig. 13a and b);115 meanwhile, Abouimrane et al. reported

a thiophene-based additive, which can form polymeric film on

the same cathode surface and yielded similar improvement.116

While the electrolyte additive remains a more economical

approach for the industry to accept, the interphasial chemistry

involved remains unclear and merits further investigation.

Different routes were also explored in order to directly coat the

cathode with an artificial interphase. Similar to what was done

on graphitic anode, Al2O3 and AlPO4 was used for LiCoO2

through either sol–gel or atomic layer deposition approach, and

better cycling stability was claimed when more than half of the

stoichoimetric Li+ can be removed from lattice of LiCoO2, which

was otherwise inaccessible due to both electrolyte decomposition

and lattice instability at potentials above 4.2 V.43,117–120 The

concern on cost and scalability remains.

Among these coating approaches, perhaps the most promising

and feasible was the electrostatic self-assembly technique that

Liu et al. developed.121–124 By adjusting pH value of the solution

so that the surface of the cathode particles remain negatively

charged, they successfully covered it with a series of oxide or

phosphate adsorbents, which bear positive charges in the

suspension and, after high temperature processing, turns into a

nm-thickness protective layer as shown in Fig. 14a. They applied

this technique onto several high voltage cathode materials and

obtained rather reversible operations in the region near 5.0 V

(Fig. 14b).

It must be noted here that these nano-coatings, as well as any

such coatings that were directly applied to the electrode material

particles in the bulk before they are attached to a current

collector, are not interphases themselves in the traditional sense,

otherwise the active mass would be electronically insulated from

the circuit; rather, during the cell operation, the electrons must be

allow to tunnel through them in order for the electrochemistry to

occur. In other words, an additional SEI should be formed from

the electrolyte components on top of these nano-coatings once

the Li ion cell is charged for the first time. The role of those nanocoatings might be more than just a physical barrier that regulates

the flow of Li+

; instead their specific chemistry and morphology

could be interfering or even directing the decomposition chemistry of the electrolyte components. The eventual ‘‘interphase’’ in

this case could be a composite film consisting of the inorganic

inner and an organic outer layer.

3. Concluding remarks

As the past two decades witnessed the glory and frustration of Li

ion batteries, it is well recognized that the interphases between

electrolyte and electrodes in Li ion batteries played important

roles in determining the performance of the device. Thanks

to advancements from fundamental research, we have learned

much about the chemistry and formation mechanism of this

elusive component, although it remains the least understood

Fig. 13 (a) Voltage profiles of ‘‘5 V’’ class cathode LiNi0.5Mn1.5O4 in standard electrolyte LiPF6/EC/EMC (30 : 70) without additive (a) and with 1%

HFiP as additive (b) (Reproduced with permission by the Electrochemical Society from Ref. 115).

Fig. 14 TEM picture of artificial interphases consisting of Al2O3, ZnO, Bi2O3 and AlPO4 on LiMn0.42Mn1.42Co0.16O4 through electrostatic selfassembly (a) and the cycling performance (b) (Reproduced with permission by the American Chemical Society from Ref. 124).

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within the device and merits our further and intensified investigation. It is predicted that the interphases on cathode surfaces,

especially on those of the high voltage (‘‘5 V’’) materials, will

attract most of the attention in the next 5–10 years.

Acknowledgements

The work was partially funded by U. S. Department of Energy

Advanced Battery Research (DOE-ABR) Program. The authors

want to thank Dr Peter Faguy of DOE and Prof. Jeff Dahn of

Dalhousie University for useful discussions.

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9864 | J. Mater. Chem., 2011, 21, 9849–9864 This journal is a The Royal Society of Chemistry 2011

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43

Research@

Distinguishing Li+ Charge Transfer Kinetics 

at NCA/Electrolyte and Graphite/Electrolyte 

Interfaces, and NCA/Electrolyte and LFP/

Electrolyte Interfaces in Li-Ion Cells

T. Richard Jow, Michelle B. Marx, and Jan L. Allen

U.S. Army Research Laboratory, Adelphi, MD

ABSTRACT

In examining the Li+ charge transfer kinetics at the graphite anode and the 

lithium nickel cobalt aluminum oxide, LiNi0.80Co0.15Al0.05O2 (NCA), cathode in 

a full cell, we found that the activation energy, Ea

, for the charge transfer at 

the graphite/electrolyte interface is about 68 kJ/mol, which is consistent with 

recently reported values. However, the Ea

 for the charge transfer at the NCA/

electrolyte interface is about 50 kJ/mol, which is lower than at the graphite 

anode. With desolvation as the predominate step for limiting the kinetics and 

both electrodes subjected to the same electrolyte, the difference in Ea

 suggests 

that it is greatly influenced with respect to the nature of the electrode materials 

and their associated SEIs. This is further confirmed by the examination of Li+

charge transfer at the LiFePO4 (LFP)/electrolyte and the graphite/electrolyte 

interfaces using a LFP/graphite full cell.

Journal of The Electrochemical Society, Volume 159, Number 5, 

Pages A604-A612 (2012)
44A604 Journal of The Electrochemical Society, 159 (5) A604-A612 (2012)

0013-4651/2012/159(5)/A604/9/$28.00 © The Electrochemical Society

Distinguishing Li+ Charge Transfer Kinetics at NCA/Electrolyte

and Graphite/Electrolyte Interfaces, and NCA/Electrolyte and

LFP/Electrolyte Interfaces in Li-Ion Cells

T. Richard Jow, ∗,z Michelle B. Marx,∗ and Jan L. Allen∗

U.S. Army Research Laboratory, Adelphi, Maryland 20783, USA

In examining the Li+ charge transfer kinetics at the graphite anode and the lithium nickel cobalt aluminum oxide,

LiNi0.80Co0.15Al0.05O2 (NCA), cathode in a full cell, we found that the activation energy, Ea, for the charge transfer at the

graphite/electrolyte interface is about 68 kJ/mol, which is consistent with recently reported values. However, the Ea for the charge

transfer at the NCA/electrolyte interface is about 50 kJ/mol, which is lower than at the graphite anode. With desolvation as the

predominate step for limiting the kinetics and both electrodes subjected to the same electrolyte, the difference in Ea suggests that

it is greatly influenced with respect to the nature of the electrode materials and their associated SEIs. This is further confirmed by

the examination of Li+ charge transfer at the LiFePO4 (LFP)/electrolyte and the graphite/electrolyte interfaces using a LFP/graphite

full cell.

© 2012 The Electrochemical Society. [DOI: 10.1149/2.079205jes] All rights reserved.

Manuscript submitted November 22, 2011; revised manuscript received February 9, 2012. Published March 2, 2012. This was Paper

488 presented at the Las Vegas, Nevada, Meeting of the Society, October 10–15, 2010.

In developing high power Li-ion batteries, reducing resistance that

limits the charge and discharge rates is important to improve the power

capability of Li-ion cells. Various resistances existing in cells such as

contact resistance between the current collector of the electrodes and

the cell container and electrolyte resistance can be reduced through

engineering and the use of a more conductive electrolyte, respectively.

For the same electrode material, the use of thinner electrodes will result

in lower resistance cells as the length of the electrodes is increased

when packaging them in the same size of cells such as 18650. This

is simply due to the fact that the resistance is proportional to the

thickness of the electrodes and inversely proportional to the area of

the electrodes. The reduction of the particle size of the active materials

can increase the number of electrochemical reaction sites for Li+

and reduce the time to utilize the active materials. However, the Li+

charge transfer resistance, Rct, the resistance that Li+ encounters when

moving from a solvated ionic state in the electrolyte solution crossing

the electrode-electrolyte interface and inserting into the electrodes is

one critical source of resistance that requires further understanding

and reduction.

During the charge process, for a state-of-the-art Li-ion battery, an

electron leaves the lithium metal oxide cathode via an external circuit and moves to the graphite anode. To retain the charge neutrality

of the cathode, Li+ is released from the cathode moving across the

cathode SEI and entering the electrolyte. The Li+, which is solvated

by the solvent molecules in electrolyte, needs to be desolvated before

moving across the graphite anode SEI, inserts into the graphite providing charge neutrality by compensating or accepting the electron

coming from the cathode through the external circuit. This process is

reversed during discharge. The Li+ charge transfer process, described

in this paper, involves the desolvation of the solvated Li+ in the liquid electrolyte, crossing of Li+ through the SEI layer formed at the

electrolyte/electrode interface and the acceptance of an electron from

the external circuit while inserting into the intercalation type of electrode materials. The resistance resulting from this process is Rct. If the

charge transfer across the interface is a thermally activated process,

Rct follows the relationship,1

1

Rct

= Ao e−Ea /RT , [1]

where Ao, Ea, R and T are a constant, the activation energy, the gas

constant and the temperature in Kelvin, respectively. The Ea of the

Li+ charge transfer process or kinetics represents the barrier that the

∗ Electrochemical Society Active Member.

z E-mail: t.r.jow.civ@mail.mil

Li+ needs to overcome to cross the interface between the electrolyte

and the electrode. The value of Ea can be obtained from the slope of

a log(1/Rct) versus the inverse of temperature (1/T) plot.

In a practical cell, however, the cell resistance is consisted of other

resistances in addition to Rct as shown below.

Rcell = Re + Relectrolyte + RSEI + Rct, [2]

where Re is the contact resistance, Relectrolyte is the resistance of the

electrolyte and RSEI is the resistance of the SEI layers on the electrodes.

The Rct value is usually determined using an AC impedance technique that potentially can separate the various resistances by measuring the impedance in wide frequency ranges. At higher frequencies,

RS is often used to represent the combined Re and Relectrolytye resistances. RSEI is usually determined using the impedance spectrum at

the medium frequencies and Rct at lower frequencies. A DC pulse

impedance technique can also be used to estimate the value of Rct

when Rct is the dominating resistance in the cell, which is often the

case at low temperatures. Using this technique, the DC resistance is

estimated from the voltage change by applying a pulse current to the

cell. This downside of this technique is that it is unable to separate the

various resistances in a cell.

Li+ charge transfer kinetics across various electrode/electrolyte

interfaces was extensively studied by the group led by Ogumi and

Abe2–8 using an AC impedance technique. The study of HOPG (highly

oriented pyrolytic graphite)/liquid electrolyte [1 m LiClO4 in EC

(ethylene carbonate):DMC (dimethyl carbonate) (1:1)] interface as

reported by Abe et al.2 and Yamada et al.3 found the value of Ea for

Li+ charge transfer across this interface to be 58 kJ/mol in a temperature range of 25◦C to around 50◦C. Abe et al.4 also reported the Li+

charge transfer kinetics at the interface between the solid-state Li+

conductor, Li0.35La0.55TiO3 (LLT), and the Li+ liquid electrolyte. At

this interface, the Li+ moves between a solvated state in the liquid

to a de-solvated state in the solid. The Li+ charge transfer, in this

case, involves only the Li+ de-solvation and solvation step without

the electron transfer step associated with the electrode (anode or cathode)/electrolyte interface, in which the electron is involved. It is also

found by the same authors that the Ea value varies with solvent selected. Yamada et al.5 reported an Ea of 51 kJ/mol for the interface of

LLT/1 m LiClO4 in EC:DMC (1:1) electrolyte. Considering this value

is similar or slightly lower than that at the HOPG/electrolyte interface,

this seems to affirm that de-solvation is the predominant step for the

Li+ charge transfer across the graphite/electrolyte interface.

Li+ charge transfer kinetics at the interface between thin

film lithium transition metal oxide electrodes such as Li4Ti5O12,

6

LiMn2O4

7 and LiCoO28 and the electrolyte of 1 m LiClO4 in EC:DEC

(diethyl carbonate) or in PC (propylene carbonate) solvent have

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Journal of The Electrochemical Society, 159 (5) A604-A612 (2012) A605 45

recently been reported by Doi et al.6 and Yamada et al.7, 8 The

Li4Ti5O12 thin film electrode was prepared using sol-gel process while

LiMn2O4 and LiCoO2 were prepared using pulse laser deposition technique. The values of Ea obtained for Li+ charge transfer across the

Li4Ti5O12/electrolyte, LiMn2O4/electrolyte and LiCoO2/electrolyte

interfaces are about 44–48 kJ/mol, 50 kJ/mol and 46 kJ/mol,6–8 respectively, in a temperature range of 10◦C to 40–45◦C. These values are

lower than that for Li+ charge transfer across the HOPG/electrolyte

interface with varying amounts. However, the Ea values are still quite

large for Li+ to move across the lithium transition metal oxide and

electrolyte interface. It seems reasonable to attribute the large Ea values to the de-solvation step as identified by Ogumi and Abe et al.2–8

Li+ charge transfer resistance at the interface between lithium iron

phosphate (LiFePO4 or LFP) electrode and electrolyte of LiPF6 in

EC with linear carbonate mixtures was recently reported by X. Liao

et al.9 and L. Liao et al.10 In these studies, the electrodes were made of

a mixture of LFP active material with carbon diluents and binder. As

with previous studies,2–8 the cells with Li as a counter electrode and

a reference electrode were used in these studies. L. Liao et al.10 also

reported activation energy values of 33.8 kJ/mol for LFP electrode at

the fully charged state (or fully de-lithiated state). Li et al.11 recently

also reported activation energy of electrode reaction of 58 kJ/mol

and 66 kJ/mol at LFP and graphite electrodes using LFP/Li and

graphite/Li half cells in LiPF6 in EC:PC:EMC mixtures, respectively,

in a configuration of two electrodes button cells. Liaw et al.12 studied

the power and capacity fades of NCA/graphite 18650 cells containing

LiPF6 in EC:EMC electrolyte correlating with the cell impedance

through the evaluation of activation energies of these parameters.

Activation energy values of 50–55 kJ/mol for the initial stage of

degradation of various degrees of aged cells were reported. The activation energy of the static heat generation rate of un-aged cells about

55 kJ/mol was also reported. The reported cell impedance of the cells

aged in different conditions in medium frequencies seems to correlate

with the power and capacity degradation and static heat generation.

As the Li ions are brought from the electrolyte to Li in the electrode

by crossing the electrode/electrolyte interfaces for completing a charge

or discharge process, it is reasonable to suggest that the solvated Li+

in the electrolyte goes through the same solvation and desolvation

steps during charge/discharge at either the graphite/electrolyte or the

cathode/electrolyte interface. As suggested by previous reports, the

desolvation of solvent molecules from Li+ is the rate limiting step.

This suggests that the Ea for Li+ charge transfer at either electrode

would be similar. However, at different interfaces, the different SEIs,

or different interphases formed on electrode surfaces and different

nature of the electrodes could have played a role in affecting the Li+

charge transfer kinetics and their activation energy. Such distinction

has not been well studied before.

The Rct was recently examined at the NCA/electrolyte interface

and at the graphite/electrolyte interface at the same time in a 3-

electrode pouch cell using a DC pulse current measuring resistance

changes at both electrodes at the same time at temperatures from 25◦C

to −40◦C.13 It was found that the Ea value at the graphite/electrolyte

interface is 61 kJ/mol, which is consistent with the values reported by

Jow et al.14 and Xu.15 However, the Ea value for the NCA/electrolyte is

41 kJ/mol, which is substantially lower than at the graphite/electrolyte

interface. This result suggests that the nature of the electrode and the

associated SEI or interphases formed at the interface could have played

significant roles in influencing the Li+ charge transfer kinetics.

The Rct values measured using a DC pulse current technique includes other resistance such as electrolyte resistance and contact resistance, which could compromise the accuracy of the Rct values although

insignificant at temperatures at or below −20◦C. While believing the

conclusion is qualitatively correct, it was decided to further confirm

these results by measuring Rct using an AC impedance method on

NCA/Graphite. LFP/graphite systems were also added, in which the

LFP electrode is of different nature from other lithium transition metal

oxides, for closer examination of the Li+ charge transfer kinetics at

both the graphite/electrolyte and LFP/electrolyte interfaces. The importance of the nature of the electrode and how the electrode kinetics

could be affected are what were set out to be understood through such

examinations.

Experimental

Electrode materials, electrodes and cell configurations.— Two

different Li-ion cell chemistries were examined in this study,

NCA/graphite and LFP/graphite. The 3-electrode NCA/graphite

pouch cells include the same cathode and anode as those used in

the Saft VL6A cells plus a third electrode made of pure Li as a

reference electrode. Both the NCA and the graphite electrodes are

porous electrodes. The dimension of the NCA cathode was 6.35 cm

× 3.81 cm, the dimension of the graphite anode was 7.62 cm

× 5.08 cm, and the reference electrode was roughly 2 cm × 1 cm.

The electrolyte is made of 1M LiPF6 in EC:DMC:MB (methyl butyrate) + VC (vinylene carbonate) mixtures with MB content greater

than 50 volume percent and VC content smaller than 3 but larger

than 0.5 weight percent. In each cell, 1.5 mL of electrolyte was used.

This electrolyte could stay as liquid and perform at temperatures as

low as −40◦C.16 The nominal capacity at C rate at 4.1 V/2.5 V at

25◦C is 16 mAh. The 3-electrode LFP/graphite pouch cells were similarly constructed with the dimensions of the cathode, anode, and

reference electrode being the same as for the NCA cell. The electrolyte and the amount of electrolyte used was also the same. The

nominal capacity of LFP/graphite cells at C rate at 3.6 V/2.5 V at

25◦C is approximately 15 mAh. For NCA/electrolyte cells, the AC

impedance measurements were carried out at 3.91 V and 0.22 V vs.

Li/Li+ for the NCA cathode and the graphite anode, respectively. For

LFP/graphite cells, the impedance measurements were carried out at

3.43 V (at a 50% SOC) and 0.14 V vs. Li/Li+ for the LFP cathode and

the graphite anode, respectively. The potentials of graphite anode vs.

Li/Li+ in NCA/graphite and LFP/graphite cells are both in the highly

lithiated state. At this state, the graphite electrode is conductive and

of similar resistance. The potential of the graphite electrode changed

only little from the initial chosen potential under our experimental

conditions.

DC pulse current resistance and AC impedance measurements.—

The DC pulse current resistance of the cell, cathode, and anode

were determined by the following steps. A pulse current, Ipulse, was

applied to the cell, which was equilibrated for one hour at a voltage,

Vo, before applying a current, and measuring the cell voltage, Vt, as

a function of time for fifteen seconds. The DC pulse resistance of

the cell during charge, Rc, or during discharge, Rd, was obtained by

dividing the voltage change, V = Vt−Vo, by Ipulse depending on

the directions of the current. The experiment was carried out at 25, 0,

−20, −30, and −40◦C. For LFP/graphite cells, the cell was charged

and then discharged for one hour at a rate of C/2 to assure the cell was

at the 50% of the state-of-charge for each measurement. The pulse

current values, for the NCA/graphite cell, from 25◦C to −40◦C, were

80 mA (5C), 16 mA (C), 8 mA (C/2), 2.67 mA (C/6), and 1 mA

(C/16) respectively. The pulse current values for the LFP/graphite

cell, from 25◦C to −40◦C, were 30 mA (2C), 15 mA (C), 7.5 mA

(C/2), 1.875 mA (C/8), and 0.75 mA (C/20) respectively. The reason

different Ipulse values were used for the cells was based on the resistance

values at 25◦C and adjusted so that the change in voltage would be

noticeable but not too extensive. A Solartron ModuLab Potentiostat

with a Frequency Response (FRA) module was used to measure the

DC pulse resistance.

The AC impedance of the cell, cathode, and anode were determined using a galvanostatic AC impedance method. An AC current

was applied across the graphite anode and the cathode terminals of the

cell, in which both the graphite anode and the cathode were subjected

to the same current perturbation. The AC voltage responses of the cell

and each electrode monitored using the reference Li electrode were

then recorded and the AC impedances calculated. The AC impedance

of the cell and electrodes were measured from 100 kHz to 10 mHz.

The experiments were carried out at the same temperatures as those

used for the DC pulse methods. The amplitudes for the NCA/graphite

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46A606 Journal of The Electrochemical Society, 159 (5) A604-A612 (2012)

cell, from 25◦C to −40◦C, were 5 mA (C/3.2), 2 mA (C/8), 1.5 mA

(C/10.67), and 0.5 mA (C/32) respectively. The amplitudes for

the LFP/graphite cell, from 25◦C to −40◦C, were 7.5 mA (C/2),

3.75 mA (C/4), 2.5 mA (C/6), 1.5 mA (C/10), and 1.07 mA (C/14).

The reason different amplitudes were used for the cells was based on

the impedance values at 25◦C and adjusted so that the change in voltage would be noticeable but not too extensive. A Solarton ModuLab

Potentiostat with a Frequency Response Analyzer (FRA) module was

used to measure the AC impedance.

Results and Discussion

The Li+ charge transfer kinetics at graphite/electrolyte,

NCA/electrolyte and LFP/electrolyte interfaces were examined using two full NCA/graphite and LFP/graphite pouch cells with Li as a

reference electrode with both cell chemistries in the same electrolyte.

A DC pulse resistance method and an AC impedance method were

employed to measure the resistance associated with the Li+ charge

transfer kinetics. The resistance, R. was measured at temperatures

from 25◦C down to −40◦C using each method. The R value as shown

in Eq. 1 is determined by the pre-factor Ao and the activation energy,

Ea, for the thermally activated process. Ao could be varied with the

accessibility of intercalation sites or the availability of lithium ions

on the electrode surfaces. Ao could be different for different electrode

materials and electrode morphologies in the same electrolyte. For this

study, Ao is assumed to be a constant and independent of temperature.

DC pulse resistance of NCA/graphite and LFP/graphite cells.—

The DC pulse resistance of the NCA/graphite cell, NCA cathode and

graphite anode at various temperatures were measured and reported

by Jow et al.13 At 25◦C, the cell resistance is largely due to the NCA

cathode with less than 10% contribution coming from the anode resistance at 15 s charge or discharge. With decreasing temperature, the

contribution of the graphite anode resistance increases. As a result, the

activation energy, Ea, for Li+ charge transfer at the graphite/electrolyte

interface is 61 kJ/mol, which is higher than the 41 kJ/mol that was

found at the NCA/electrolyte interface. A similar trend was found in

the DC pulse resistance of LFP/graphite cells as shown in Fig. 1.

The slopes of the plot of log(1/R) vs. 1/T for both the graphite

anode and LFP cathode, as shown in Fig. 2, were used to calculate

the Eas of both electrodes. The Ea values for the Li+ charge transfer

across the LFP/electrolyte and the graphite/electrolyte interfaces were

31 kJ/mol and 58 kJ/mol, respectively. The Ea value for the graphite

anode in the LFP/graphite cell is consistent with that for the graphite

anode in the NCA/graphite cell. However, a lower Ea value was found

at the LFP/electrolyte interface than at the NCA/electrolyte interface.

AC impedance of NCA/graphite cells.— A galvanostatic AC

impedance method was used to measure the impedance of 3-electrode

pouch cells. The 3-electrode cell configuration allowed for impedance

measurements of the full cell and the individual electrodes simultaneously. The AC impedance of the graphite anode, the NCA cathode

and the full cell plotted as Nyquist plots at temperatures 25◦C, 0◦C,

−20◦C, −30◦C and −40◦C are shown in Fig. 3. The potentials of

the NCA cathode at 3.91 V and the graphite cathode at 0.22 V were

selected for impedance measurement. The results show that there are

two depressed and overlapped semicircles with a straight line at low

frequencies indicating resistance associated with diffusion at the NCA

cathode. At −30◦C or below, the diffusional behavior has disappeared.

At the graphite anode, only one semicircle was observed for the most

part. The distinct differences in impedance spectra between the two

electrodes reflect a clear distinction in Li+ charge transfer kinetics between the graphite anode and the NCA cathode. The graphite anode

behaviors at 25 and 0◦C are hard to explain. The semicircle turned

into another semicircle below the x-axis and then went up as almost

a straight line at lower resistance values. This behavior could be due

to the substantially low resistance at the graphite anode such that the

inductive component of the measuring system became dominant at

these two temperatures. The diameter of the semicircle above the xFigure 1. The cell impedance, Rc and Rd, the graphite anode impedance,

RA,c and RA,d, and the LFP cathode impedance, RC,c and RC,d, during charge

and discharge, respectively, at 25◦C and −30◦C measured at LFP: 3.43 V and

graphite: 0.15 V.

axis was assigned as the real resistance of the anode. The measured

impedance was so small that the impact of the inductive loop on the

determination of Ea is insignificant.

Interpretation of impedance data of NCA/graphite cell.—The electrodes used in this study are industrial Li-ion electrodes, which are

Figure 2. The log (1/R) versus 1/T plots for the LFP cathode and the graphite

anode during charge and discharge of the 3-electrode pouch cell at 15 s.

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Journal of The Electrochemical Society, 159 (5) A604-A612 (2012) A60747

Figure 3. Nyquist plots of galvanostatic AC impedance

of a NCA/graphite cell at (a) 25◦C, (b) 0◦C, (c) −20◦C,

(d) −30◦C and (e) −40◦C, where the graphite anode is shown

in blue, the NCA cathode is green and the full cell is red.

porous and may contain active materials of different particle sizes. The

Nyquist plots shown in Fig. 3 could not be fitted using the equivalent

circuit (EC) made of simple RCs and Warburg impedance as shown

below.

Rs C1

R1

C2

R2

W1

[EC1]

One simple way to look at the data is to use apparent resistance (or

the total length of resistance on the x-axis) as the total resistance of

the Li+ charge transfer process. The apparent resistance is defined as

resistance from the intersection of the semicircle on the x-axis at the

high frequency end to the intersection of the semicircle at the lower

frequency, which for some plots needed to be extended in order to

intersect the axis. The Rct values obtained by this method were used

to estimate the activation energy for Li+ charge transfer kinetics at

both electrodes as discussed in the following sub-section.

Since the Li+ charge transfer resistance that was measured reflects the nature of heterogeneity of the SEI and possibly distributed

time-constant for charge transfer reactions, a constant phase element

(CPE)17 approach was adopted. The equivalent circuit made of RS,

two circuits in series, each is made of R and CPE in parallel, and a

Warburg impedance as shown in EC2 was used to fit the impedance

of the NCA electrode at 25◦C, 0◦C, −20◦C and −30◦C.

Rs CPE1

R1

CPE2

R2

W1

[EC2]

At −40◦C, the equivalent circuit made of RS, two circuits, each is

made of R and CPE in parallel without a Warburg impedance as

shown in EC3 was used to fit the impedance of NCA electrode.

Rs CPE1

R1

CPE2

R2

[EC3]

The impedance of the graphite anode in the NCA/grphite cell was

fitted using the equivalent circuit made of one single R-CPE in parallel

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48A608 Journal of The Electrochemical Society, 159 (5) A604-A612 (2012)

Figure 4. The curves that fit the Nyquist plots

shown in Fig. 3 using the equivalent circuits EC2,

EC3 or EC4 are plotted. Only fitting curves at 0◦C,

−30◦C and −40◦C are shown. R1 is the charge

transfer resistance due to the faster process (at higher

frequencies) and R2 is the charge transfer resistance

due to the slower process (at lower frequencies).

(a) Graphite electrode at 0◦C (fitted with EC4),

(b) NCA electrode at 0◦C (fitted with EC 2),

(c) Graphite electrode at −30◦C (fitted with EC4),

(d) NCA electrode at −30◦C (fitted with EC2),

(e) Graphite electrode at −40◦C (fitted with EC4),

(f) NCA electrode at −40◦C (fitted with EC3).

as shown in EC4 at all temperatures.

Rs CPE1

R1

[EC4]

Using these equivalent circuits, a reasonably well fit was obtained.

Data fitting using the equivalent circuits, EC2 to EC4 at 0◦C, −30◦C

and −40◦C for the graphite anode and NCA cathode are shown in

Fig. 4. Through the above impedance data fitting, R values obtained

from R1 and R2 processes were used to estimate Eas of these processes

as discussed in the following sub-sections.

Activation energy for Li+ charge transfer at NCA and at graphite

electrodes.—The R values obtained in the previous section were used

to estimate the Ea values of the Li+ charge transfer kinetics at both

the graphite and NCA electrode. These values are calculated from the

slopes of the log(1/R) vs. 1/T plots. Fig. 5 shows the log(1/R) vs.

1/T for the total apparent resistance estimated from Fig. 3. The Ea

values for the NCA cathode and the graphite anode are 49 kJ/mol and

68 kJ/mol, respectively, as calculated from the slopes and are also

shown in the figure.

R values could also be obtained by fitting the data in Nyquist

plots using the equivalent circuits containing CPE such as EC2 to

EC4 as shown in Fig. 4. The activation energy of the NCA cathode

and graphite anode could also be calculated from the slopes of the

log(1/R) vs. 1/T plots. For the NCA cathode, the impedance spectra

suggest two processes represented by R1 and R2. R1 could be due to

surface films or the interphases formed on the NCA cathode and R2

could be the charge transfer resistance. For the graphite anode, the

impedance spectra suggest only one major process R2 representing

charge transfer. The R1 representing SEI layer at the graphite anode

was too small relatively to R2 to be observed. Fig. 6 shows the log(1/R)

vs. 1/T for R1 and R2 for the NCA cathode and R2 for the graphite

anode. The Ea values for the NCA cathode are 55 kJ/mol (R1) and

51 kJ/mol (R2) and the graphite anode 68 kJ/mol as calculated from the

slopes and are also shown in Fig. 6. The Ea value for the graphite anode

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Journal of The Electrochemical Society, 159 (5) A604-A612 (2012) A609 49

Figure 5. Log(1/R) vs. 1/T plot for R values obtained using the total apparent

charge transfer resistance for the NCA/graphite cell as seen from the x-axis of

the Nyquist plots shown in Fig. 3.

in the NCA/graphite cell is consistent with that reported by Jow et al.14

and Xu.15 The activation energy for the two processes at the NCA

cathode is similar. Liaw et al.12 reported 50–55 kJ/mole activation

energy values for the initial stage of degradation of NCA/graphite

cell and the static heat generation of un-aged 18650 NCA/graphite

full cells in a temperature range of 25 to 65◦C. These numbers are

close to the average of the activation energies for NCA/electrolyte and

graphite/electrolyte obtained here. This also suggests that there is a

link between the charge transfer resistance, cell heat generation and

cell degradation.

AC impedance of LFP/graphite cells.— The AC impedance of the

graphite anode and the LFP cathode plotted as Nyquist plots are shown

in Figs. 7 and 8, respectively. The plots at a smaller scale are shown

as an insert in the figures in order to see the results obtained at 25◦C

and 0◦C. The potentials of the LFP cathode at 3.45 V and the graphite

anode at 0.15 V were selected for impedance measurement.

The impedance spectra of the LFP electrode also exhibited more

depressed semicircles than the graphite electrode in the LFP/graphite

cell similar to that of the NCA electrode in NCA/graphite cell. Similar

to the NCA/graphite system, the Nyquist plots shown in Figs. 7 and

Figure 6. Log(1/R) vs. 1/T plot for R values obtained using the equivalent

circuit containing CPE for the NCA/graphite cell shown in the Nyquist plots

in Fig. 4.

Figure 7. Nyquist plots of galvanotat AC impedance of the graphite anode

at 25, 0, −20, −30 and −40◦C in a LFP/graphite cell. The insert shows the

impedance curves at a smaller scale.

8 could not be fitted using the equivalent circuit made of simple RCs

and Warburg impedance. With the use of the constant phase element

(CPE),17 the Nyquist impedance plots could be fitted nicely. One

example is shown in Fig. 9 for the LFP and the graphite impedance

fitting in a LFP/graphite cell at −20◦C. The values of R1 and R2

for the LFP and graphite anode at various temperatures are shown in

Fig. 10a. R1 represents the resistance resulting from SEI or surface

films on either graphite anode or LFP cathode. R2 represents the

charge transfer resistance, Rct, at either interface. The figure clearly

indicates that R values from the R2 process for the graphite anode

increases faster than those for the LFP electrode. The Ea values for

the LFP cathode are 2.3 kJ/mol (R1) and 33 kJ/mol (R2) and the

graphite anode are 7 kJ/mol (R1) and 67 kJ/mol (R2) as calculated

from the slopes of the log(1/R) versus 1/T plots as shown in Fig. 10b.

Figure 8. Nyquist plots of galvanostatic AC impedance of the LFP cathode

at 25◦C, 0◦C, −20◦C, −30◦C and −40◦C in a LFP/graphite cell. The insert

shows the impedance curves at a smaller scale so that the results from 25◦C

and 0◦C can be observed.

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50A610 Journal of The Electrochemical Society, 159 (5) A604-A612 (2012)

Figure 9. (a) The Nyquist plots of the impedance

of the LFP/graphite cell, LFP cathode and graphite

anode. (b) The fitting curve of the LFP cathode

using EC2. (c) The fitting curve of the graphite

anode using EC3.

The Ea value for the graphite anode in the LFP/graphite cell is also

consistent with that reported by Jow et al.14 and Xu.15 The obtained

Ea for charge transfer process is very close to that obtained by L. Liao

et al.10 at the fully charged state and 32 kJ/mol Ea values calculated

based on Eq. 1 using Rct values reported by X. Liao et al.9 However,

these values are lower than that obtained by Li et al.11 using LFP film

electrode and two electrodes button cells.

Distinguishing Li+ charge transfer kinetics at different electrode/electrolyte interfaces.— The impedance spectra for graphite anode, NCA cathode and LFP cathode obtained from the AC impedance

measurements using NCA/graphite and LFP/graphite full cells as

shown in Figs. 3, 7, and 8 provide an insightful look at how electrode

materials affect the impedance behavior and therefore the kinetics at

each electrode. For the graphite anode, the semicircle as shown in the

impedance spectra indicates that the charge transfer process dominates

with little or no observable RSEI contribution. For the NCA cathode,

two depressed semi-circles represent two processes, one at medium

frequencies represents process due to surface film or interphases layer

formed at the NCA/electrolyte interface and one at lower frequencies

represents charge transfer process. Both processes contribute almost

equally to the impedance of the NCA electrode. For the LFP cathode,

two depressed semicircles were also observed while only one at the

lower frequencies dominates. The faster one, R1, is small relative to

the charge transfer process.

By analyzing the impedance spectra using the equivalent circuits

as shown in EC2 to EC4, we determined the charge transfer resistance, Rct, or R2 process for each interface. The activation energy,

Ea, values for Li+ charge transfer at the different interfaces obtained

using the AC method along with the DC method are summarized in

Table I. The Ea values obtained using the DC method are lower than

that using the AC method as R values obtained using DC methods

include more than just the charge transfer resistance, Rct, which is

less dominant at temperatures above −20◦C. Qualitatively, the results

are in good agreement that the Ea values at the graphite/electrolyte

interface are significantly larger than at either the NCA/electrolyte

or the LFP/electrolyte interface. The Ea values of 68 kJ/mole or

67 kJ/mol obtained from AC impedance method using equivalent

circuit fitting for graphite/electrolyte interface in either NCA/graphite

or LFP/graphite cells were consistent with earlier reports.13–15 Ea values at the NCA/electrolyte interface appear to be consistent with those

at other lithium transition metal oxide/electrolyte interfaces.6–8

The Ea values of 52 kJ/mol and 33 kJ/mol for the NCA/electrolyte

interface in the NCA/graphite cell and LFP/electrolyte interface

in LFP/graphite cell, respectively, strongly suggest that the electrode materials have a significant impact on the activation energy

for the Li+ charge transfer kinetics across the electrode/electrolyte

interface.

How the nature of the electrode material is impacting the activation energy of the Li+ charge transfer across the electrode/electrolyte

interfaces is not clear. However, the fact that the nature of the SEI

formed on the electrode surface is impacted by the electrode material

is clear. El Ouatani et al.18, 19 reported that the polymerization of VC,

which was used as an additive in the electrolyte, was detected only on

the surfaces of certain electrode materials even the electrochemical

potential of these electrodes relative to Li/Li+ were similar. For example, the VC polymerization is observed at the surface of the graphite

electrode and LiCoO2 electrode in LiCoO2/graphite cells.18 However,

the VC polymerization is not observed at the surface of the LFP electrode when VC is present in the LFP/graphite cells.19 The difference in

surface films, which could be SEI or other interphases, on the cathode

is consistent with the fact that Co4+, which is present in the charged

Li1−xCoO2 electrode, will be more oxidizing than the Fe3+ found in

FePO4, the charged composition of the iron phosphate cathode. These

results appear to support the observation that the activation energy

for the Li+ charge transfer kinetics at the LFP/electrolyte interface is

smaller than that at the NCA/electrolyte interface.

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Journal of The Electrochemical Society, 159 (5) A604-A612 (2012) A61151

Table I. Activation energy, Ea, values for Li+ charge transfer at the NCA and the graphite electrodes in the NCA/graphite cell and at the LFP

and the graphite electrodes in the LFP/graphite cell obtained using both DC and AC methods, where R1 and R2 are obtained by fitting the data

using the equivalent circuits.

Ea, kJ/mole

At Cathode/electrolyte Interface At Graphite/electrolyte interface

Measuring Method DC Pulse AC Impedance DC Pulse AC Impedance

Resistance R R, Total Apparent R1 R2 R R, Total Apparent R1 R2

NCA/Graphite cell 41 49 55 52 61 68 − 68

LFP/Graphite cell 31 − 2.3 33 58 − 7 67

At the graphite/electrolyte interface, a rather well defined SEI,

solid electrolyte interface (or interphase), is formed.20 Desolvation

is identified as the limiting factor for the Li+ charge transfer at the

graphite/electrolyte interface.2–5 This is supported by the measurements that the activation energy for solvated Li+ to de-solvate and

get into the solid electrolyte is similar to that for solvated Li+ to desolavate and get into SEI and into graphite.2–5 At the NCA/electrolyte

or LFP/electrolyte interface, no well defined SEI layers like that on

graphite/electrolyte have been reported. It is believed that the surface

films formed on the cathode resulting from oxidative reactions could

be thin and varied with the nature of the electrode materials as inFigure 10. (a) The R1 and R2 impedance values of the LFP cathode and the

graphite anode from the LFP/graphite cell. (b) The log (1/R) vs. 1/T plots for

the LFP cathode and the graphite anode. The Ea values calculated from the

slopes of each linear fitting are also shown.

dicated by El Ouatani et al.18, 19 This may also attributed to the low

activation energy for Li+ charge transfer at the cathode/electrolyte

interface compared to that at the graphite/electrolyte interface.

From a practical point of view, at temperatures above 0◦C, the

larger Rct values at the cathodes, NCA or LFP, rather than at the

graphite anode dictates the cell rate performance. At temperatures

below −20◦C, the cell kinetics is increasingly dominated by the higher

Rct at the graphite/electrolyte interface because of higher activation

energy at the graphite/electrolyte interface.

Conclusions

The Li+ charge transfer kinetics at the graphite/electrolyte interface, the NCA/electrolyte interface, and the LFP/electrolyte interface were studied in the NCA/graphite and the LFP/electrolyte full

cells with a Li reference electrode using DC pulse resistance and AC

impedance methods.. With such cells, the impedance behavior and the

Li+ charge transfer kinetics at the graphite anode and NCA cathode or

LFP cathode could be studied at the same time in the same electrolyte.

Impedance behaviors at different interfaces were distinctively diffferent. A much less depressed semicircle observed on graphite electrode

than those observed on NCA or LFP cathode suggests a better defined

SEI existed at the graphite anode. The charge transfer resistance, Rct,

values derived from the DC resistance and AC impedance measurements were used to calculate the activation energies at various interfaces. The activation energy values were consistent between these

two methods. The activation energies at the graphite/electrolyte interface, the NCA/electrolyte interface and the LFP/electrolyte interface

obtained from the AC impedance method are 67 kJ/mol, 52 kJ/mol

and 33 kJ/mol, respectively. The differences in activation energies at

different interfaces link to the different nature of the electrode materials and their surface films, which could be SEI or other interphase

materials, formed on the electrodes. The higher activation energy at

the graphite/electrolyte interface was reasonably attributed to welldefined SEI on graphite with desolvation as the limiting factor. The

lower activation energy at the NCA/electrolyte or the LFP/electrolyte

interface could be attributed to less well defined thin surface films of

different nature. In brief, the difference in Ea suggests that Li+ charge

transfer kinetics is greatly influenced with respect to the nature of the

electrode materials and their associated SEIs.

Acknowledgments

The authors thank K. Nechev, B. Deveney and S. Rickman at Saft

America, Inc. for providing electrodes and helpful discussion and

the DOE Office of Renewable Energy and Vehicle Technologies, Applied Battery Research for Transportation Program for partial financial

support.

References

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53

Research@

A New Direction for the Performance 

Improvement of Rechargeable 

Lithium/Sulfur Batteries

Sheng S. Zhang and Jeffrey A. Read

U.S. Army Research Laboratory, Adelphi, MD

ABSTRACT

In this work we introduce a new direction for the performance improvement of 

rechargeable lithium/sulfur batteries by employing an electrolyte that promotes 

Li anode passivation in lithium polysulfide solutions. To examine our concept, 

we assemble and characterize Li/Li2S9 liquid cells by using a porous carbon 

electrode as the current collector and a 0.25 m Li2S9 solution as the catholyte. 

Results show that Li/Li2S9 liquid cells are superior to conventional Li/S cells in 

specific capacity and capacity retention. We also find that use of LiNO3 as a cosalt in the Li2S9 catholyte significantly increases the cell’s Coulombic efficiency. 

More importantly, the cells with LiNO3 have a ~2.5 V voltage plateau before the 

end of charging and demonstrate a steep voltage rise at the end of charging. 

The former is indicative of the formation of elemental sulfur from soluble 

lithium polysulfides on the carbon electrode, and the latter provides a distinct 

signal for full charging. Electrochemical analyses on Li plating and stripping 

in Li2S9 catholyte solutions indicate that LiNO3 participates in the formation of 

a highly protective passivation film on the Li metal surface, which effectively 

prevents the Li anode from chemical reaction with polysulfide anions in the 

electrolyte and meanwhile prevents polysulfide anions from electrochemical 

reduction on the Li surface.

Journal of Power Sources, Volume 200, Pages 77-82 (2012)
54 Journal of Power Sources 200 (2012) 77–82

Contents lists available at SciVerse ScienceDirect

Journal of Power Sources

j ou rnal homepage: www.elsevier.com/locate/jpowsour

A new direction for the performance improvement of rechargeable lithium/sulfur

batteries

Sheng S. Zhang∗, Jeffrey A. Read

U.S. Army Research Laboratory, RDRL-SED-C, Adelphi, MD 20783-1197, USA

a r t i c l e i n f o

Article history:

Received 26 September 2011

Received in revised form 17 October 2011

Accepted 19 October 2011

Available online 25 October 2011

Keywords:

Lithium/sulfur battery

Sulfur

Polysulfide

Catholyte

Lithium plating

a b s t r a c t

In this work we introduce a new direction for the performance improvement of rechargeable

lithium/sulfur batteries by employing an electrolyte that promotes Li anode passivation in lithium polysulfide solutions. To examine our concept, we assemble and characterize Li/Li2S9 liquid cells by using

a porous carbon electrode as the current collector and a 0.25 m Li2S9 solution as the catholyte. Results

show that Li/Li2S9 liquid cells are superior to conventional Li/S cells in specific capacity and capacity

retention. We also find that use of LiNO3 as a co-salt in the Li2S9 catholyte significantly increases the

cell’s Coulombic efficiency. More importantly, the cells with LiNO3 have a ∼2.5V voltage plateau before

the end of charging and demonstrate a steep voltage rise at the end of charging. The former is indicative

of the formation of elemental sulfur from soluble lithium polysulfides on the carbon electrode, and the

latter provides a distinct signal for full charging. Electrochemical analyses on Li plating and stripping in

Li2S9 catholyte solutions indicate that LiNO3 participates in the formation of a highly protective passivation film on the Li metal surface, which effectively prevents the Li anode from chemical reaction with

polysulfide anions in the electrolyte and meanwhile prevents polysulfide anions from electrochemical

reduction on the Li surface.

Published by Elsevier B.V.

1. Introduction

Lithium/sulfur (Li/S) batteries have attracted increasing interest in developing high density energy storage devices due to their

high theoretical capacity. Based on the complete reduction of elemental sulfur to lithium sulfide (Li2S), Li/S batteries can deliver a

specific capacity as high as 1675 mAh g−1 sulfur. However, the specific capacity of a practical cell is lower than the theoretical value

and the cell suffers low charging efficiency, high self-discharge and

short cycle life [1,2]. All these problems are known to be related to

the high solubility oflithium polysulfides,the series of sulfur reduction intermediates, in organic electrolyte solutions. Dissolution of

lithium polysulfides not only results in the loss of sulfur active

materials from the cathode, but also causes serious “redox shuttle” reactions between polysulfide anions in the electrolyte and the

Li metal anode. Recently, a number of publications have reported

a reduction in the dissolution of lithium polysulfides by making sulfur-carbon composite materials [3–14]. Based on physical

adsorption, these composites in different contexts reduce the dissolution of lithium polysulfides from the cathode. However, these

approaches are fundamentally ineffective since polysulfide anions

carrynegative charges,indischarging the electricfieldbetweentwo

∗ Corresponding author. Tel.: +1 301 394 0981; fax: +1 301 394 0273.

E-mail address: shengshui.zhang@us.army.mil (S.S. Zhang).

electrodes will drive polysulfide anions migrating toward Li anode.

Furthermore, the incorporation of electrochemically inert carbons

reduces the gravimetric energy density of Li/S batteries. In addition,

we have noticed that most of decent capacities reported previously

were obtained through two lows,thatis, low sulfur contentin composition and low sulfur loading in cathode. In many cases, the total

sulfur content in the cathode is not more than 65% by weight and

the sulfur loading is not higher than 2 mg sulfur per cm2 of cathode

[4–7,14–18].

Since dissolution of lithium polysulfides (Li2Sx, x > 2) in organic

electrolytes is inevitable, in this work we propose a different

approach for the performance improvement of rechargeable Li/S

batteries by employing a liquid electrolyte that is able to promote

the formation of a highly protective passivation film on lithium

surface in lithium polysuifide solutions. We expect that the resulting passivation film not only protects lithium metal from chemical

reaction with the polysulfide anions but also prevents polysulfide

anions from electrochemical reduction on the Li anode. Our effort

will be focused on increasing Li cycling efficiency in highly concentrated lithium polysulfide solutions. To examine our idea, we

selected Li/Li2S9 liquid cell [19,20], instead of the conventional

Li/S cell, as the testing vehicle by employing a porous carbon electrode as the cathode current collector and a Li2S9 solution as the

catholyte. Due to the known ability of LiNO3 in facilitating the formation of a better passivation film on Li metal surface [21,22], in

this work we study the effect of LiNO3 on cycling performance

0378-7753/$ – see front matter. Published by Elsevier B.V.

doi:10.1016/j.jpowsour.2011.10.076
78 S.S. Zhang, J.A. Read / Journal of Power Sources 200 (2012) 77–82 55

of Li/Li2S9 liquid cells and on cycling efficiency of Li metal in

Li2S9 catholyte solutions by adding LiNO3 as a co-salt of the Li2S9

catholyte.

2. Experimental

Elemental sulfur (S8, 99.5%), lithium sulfide (Li2S, 99%), and

lithium nitrate (LiNO3, 99.99%) were purchased from Aldrich

and used as received. Lithium bis(trifluoromethane sulfone)imide

(LiN(SO2CF3)2, LiTFSI, 3M Company) was dried at 110 ◦C under

vacuum for 10 h and triethylene glycol dimethyl ether (TG3, 99%,

Aldrich) was dried over 4A ̊ molecular sieves for a week. For conventional Li/S cells, a liquid electrolyte was prepared by dissolving

0.2 m (molality) LiTFSI in TG3 in Ar-filled glove-box and a sulfur

cathode with a composition of 77% S, 20% Super-P carbon and

3% binder by weight was coated onto a carbon-coated aluminum

foil by using poly(acrylonitrile-methyl methacrylate) (ANMMA,

AN/MMA= 94:6, MW= 100 000, Polysciences Inc.) as binder and Nmethyl pyrrolidinone (NMP) as solvent. Resulting cathode, dried at

80 ◦C under vacuum for 2 h, had an average loading of 2.4 mg sulfur per cm2. For Li/Li2S9 liquid cells, a carbon electrode having a

composition of 90% Super-P carbon and 10% binder by weight was

prepared using the same materials and procedure as was used in

preparation ofthe sulfur cathode and dried at 110 ◦C under vacuum

for 10 h. The carbon electrode was measured to have an average

loading of 0.78 mg carbon per cm2.

Following Rauh et al’s procedure [19,23], lithium polysulfide (Li2Sx, x = 9) solution was prepared by adding stoichiometric

amounts of elemental sulfur (S8) and Li2S into the electrolyte solution, followed by heating at 80 ◦C with magnetic stirring for 6 h. In

this work, we prepared two catholyte solutions with composition

below:

Catholyte-A: 0.25 m Li2S9–0.2 m LiTFSI in TG3

Catholyte-B: 0.25 m Li2S9–0.1 m LiTFSI–0.2 m LiNO3 in TG3

Both solutions were red-dark color and had moderate viscosity.

Li/S coin cells with an electrode area of 1.27 cm2 were assembled

using sulfur cathode and filled with 10 L electrolyte. Li/Li2S9 liquid cells were assembled by using a 1.27 cm2 carbon electrode

as the cathode current collector (also serving as catalyst for the

reduction of elemental sulfur and lithium polysulfides) and 10 L

catholyte solution as the electrolyte and cathode. It should be noted

that the assembly of Li/Li2S9 liquid cells should be conducted in an

oxygen-free environment so as to avoid the oxidization of polysulfide anions by oxygen. For easy comparison with conventional Li/S

cells, the specific capacity of Li/Li2S9 liquid cells was normalized as

“mAh g−1 sulfur”. Thus, the theoretical capacity of Li2S9 was calculated to be 1489 mAh g−1 sulfur (i.e., 1675 × 8/9 = 1489), and each

Li/Li2S9 liquid cell contained ∼1.65 mg sulfur as calculated based

on the amount of Li2S9 in the catholyte.

Both Li/S and Li/Li2S9 cells were cycled at 0.2 mAcm−2 on a Maccor Series 4000 cycler with a 1.5V discharge cutoff voltage. The

charging process was terminated either by a 3.0V cutoff voltage

or by a capacity equaling to 150% of the last discharge capacity,

whichever came first. To measure Li plating and stripping efficiency, a Li/Ni cell with a 1.27 cm2 electrode area was assembled

and cycled by discharging (Li plating) at 0.2 mAcm−2 for 1 h and

then charging (Li stripping) until the cell’s voltage reached 1.0V.

Coulombic efficiency of Li cycling was defined as the percentage of

charging time over discharging time. In addition, three-electrode

coin cells were assembled for electrochemical measurements by

using a 0.97 cm−2 Ni foil as working electrode, two Li foils as the

counter electrode and reference electrode, respectively. Detailed

descriptions about cell structure and assembly procedure are

referred to our previous works [24,25]. The impedance and cyclic

voltammetry measurements were run on a Solartron SI 1287

Electrochemical Interface and a SI 1260 Impedance/Gain-Phase

Analyzer. Impedance was measured at open-circuit potential(OCP)

in the frequency range from 0.01 to 100 kHz with an ac oscillation

of 10 mV amplitude. Before each test, the cell rested for 1 h after Li

plating or Li stripping to get the same conditions.

3. Results and discussion

3.1. Li/S cell vs. Li/Li2S9 liquid cell

Fig. 1 shows voltage profiles of the first and fifth cycles of a conventional Li/S cell and a Li/Li2S9 liquid cell, respectively. For the

conventional Li/S cell (Fig. 1a), the initial discharging consists of

three voltage regions:(1) a short plateau at 2.3V as indicated by the

arrow, (2) a linear sloping decline, and (3) a long plateau at ∼2.0V

until the end of discharge. Combining the conclusions of previous

publications [16,26–29], we ascribe these three discharging voltage

regions to the following three reactions:

S8(insoluble) + 2Li → Li2S8(soluble) (1)

Li2S8(soluble) + 6Li → 4Li2S2(insoluble) (2)

Li2S2(insoluble) + 2Li → 2Li2S(insoluble) (3)

In the following charge step, the cell voltage responds to the

reversible processes of Eqs. (2) and (3) until 2.5V, at which the

voltage stays constantly, indicating that Eq. (1) is irreversible. In

the fifth cycle, the voltage plateau observed at 2.3V in the first

cycle no longer appears, and the charge voltage does not exceed

2.5V. For the Li/Li2S9 liquid cell (Fig. 1b), the first discharge does

not show voltage plateau near 2.3V since no elemental sulfur exists

in the system. However, in the following charge, the cell voltage

reaches 3.0V, achieving a cycling efficiency of 91%. Discharging

and charging voltage profiles of the fifth cycle are very similar as

those observed in the conventional Li/S cell. We tested many other

Li/Li2S9 cells and the similar results were repeatedly observed. It

should be mentioned that except for the initial few cycles, both the

Li/S and Li/Li2S9 cells cannot be charged to higher than 2.5V, i.e.,

the reversal process of Eq. (1) does not occur. The reason is because

soluble lithium polysulfides (Li2Sx, x > 2), especially those having

long S–S chain, dissolve and diffuse to the surface of the Li anode,

where polysulfide anions not only are electrochemically reduced as

described by Eq. (4) but also are chemically reacted with Li metal

as described by Eq. (5).

(x−y)Li2Sx + 2yLi+ + 2ye− →xLi2Sx−y (4)

(x−y)Li2Sx + 2yLi →xLi2Sx−y (5)

Thus, the soluble polysulfide anions act as a redox shuttle

between the cathode and Li anode, resulting in low charging efficiency.

Fig. 2 compares capacity retention of the conventional Li/S cell

and Li/Li2S9 liquid cell. It is shown that these two cells have the

similar initial capacities (about 610 mAh g−1 S) in spite of the fact

that Li2S9 has a relatively lower theoretical capacity than elemental sulfur. However, the Li/Li2S9 cell exhibits more stable capacity

retention. By the end of testing (50 cycles), the Li/Li2S9 cell remains

452 mAh g−1 capacity (equaling to 72% ofthe initial capacity), while

the Li/S cell retains only 196 mAh g−1 S (32% of its initial capacity).

These results reveal that Li/Li2S9 liquid cells are superior to the

conventional Li/S cells in capacity retention.
56 S.S. Zhang, J.A. Read / Journal of Power Sources 200 (2012) 77–82 79

Fig. 1. Discharging voltage profiles of the first and fifth cycles of (a) a conventional Li/S cell and (b) a Li/Li2S9 liquid cell.

Fig. 2. Comparison of the capacity retention of a conventional Li/S cell and a Li/Li2S9

liquid cell.

3.2. Effect of LiNO3 on cycling performance

Reactions (4) and (5) are believed to be the main cause for

the low charging efficiency and high self-discharge rate of Li/S

batteries. Furthermore, their reaction products will permanently

deposit on the surface of the Li anode if insoluble Li2Sx (x ≤ 2) is

formed. When this happens, the performance of Li/S cells will be

dramatically affected. For this reason, in this work we focused our

effort on suppressing Reactions (4) and (5) by adding LiNO3 as a cosaltin the Li2S9 catholyte. Fig. 3 compares discharging and charging

voltage profiles of two Li/Li2S9 cells with and without LiNO3 as cosalt. As indicated in Fig. 3a, the most significant difference between

these two cells is that Cell-2 with LiNO3-containing Catholyte-B can

repeatedly be charged to 3V while Cell-1 with Catholyte-A free of

LiNO3 can be charged above 2.5V only in the first cycle. Another difference is that Cell-2 exhibits a pair of additional voltage plateaus

Fig. 4. Voltage profiles of the first and fifth cycles of a Li/Li2S9 liquid cell with

Catholyte-B.

at ∼2.3V in discharge and charge processes, as indicated by two

arrows in Fig. 3b. This pair of voltage plateaus can be attributed to

the reversible reaction of Eq. (1).

To further verify this, we plot discharge and charge curves for

the first and fifth cycles of Cell-2 in Fig. 4. It can be seen thatthe first

discharge does not show voltage plateau near 2.3V although Li2S9

has a longer S–S chain than elemental sulfur (S9 vs. S8), instead, the

fifth discharge shows a distinct plateau at 2.3V. In charging, both

the first and fifth cycles distinctly show a 2.3V plateau, followed

by a steep voltage rise to the cutoff voltage (3.0V). These results

indicate that the 2.3V plateau is not related to the length of polysulfide anions (i.e., the x value in Li2Sx), instead to the two-phase

Reaction-1 occurring between Li2Sx in solution and elemental sulfur in cathode.

Fig. 3. Voltage profiles of the discharging and charging cycles of two Li/Li2S9 liquid cells with Catholyte-A and Catholyte-B, respectively: (a) overall view of the voltage

profiles and (b) cell voltages of a typical cycle.
80 S.S. Zhang, J.A. Read / Journal of Power Sources 200 (2012) 77–82 57

Fig. 5. Effect of LiNO3 co-salt on cycling performance of Li/Li2S9 liquid cells: (a) capacity retention and (b) Coulombic efficiency of cycling.

The effect of LiNO3 on cycling efficiency and capacity retention is presented in Fig. 5a and b, respectively, where Cell-1 uses

Catholyte-A free of LiNO3 and Cell-2 employs Catholyte-B containing 0.2 m LiNO3.As compared with Cell-1, Cell-2 not only has higher

capacity (Fig. 5a), but also has significantly higher Coulombic efficiencies (Fig. 5b). The latter would be the most important benefit

added by the use of LiNO3 co-salt. For Cell-1, except for the first

cycle the charging voltages can never reach the cutoff voltage (3.0V,

see Fig. 3), all charging processes are ended by the pre-set charging

time. Since we setup 150% of the last discharge capacity as one of

the charge limits, Coulombic efficiencies for all cycles are 66.7% (i.e.,

the reciprocal of 150%). For Cell-2, charge voltages repeatedly reach

the pre-set cutoff voltage and the charge processes are terminated

by the cutoff voltage, resulting in higher Coulombic efficiencies.

3.3. Understanding the role of LiNO3 in Li/S cell

Assuming that Reactions (4) and (5) are the main reason for low

Coulombic efficiency of Li/S cells and that both reactions occur on

the Li anode surface, we believe that the important role of LiNO3 in

Li/S cells can be understood from the viewpoint of the Li anode.

Therefore, we place our attention on the plating and stripping

behavior of Li metal in highly concentrated Li2Sx (x > 2) solutions.

Fig. 6 exhibits the potential curves of the first plating and stripping

of Li metal on a fresh Ni surface in Catholyte-A and Catholyte-B,

respectively. In Catholyte-A, the potential of Ni working electrode

shows multiple plateaus between OCP and 1.3V, followed by a slow

and smooth decline. Even atthe end ofthe plating test,the potential

of Ni still remains at +0.15V vs. Li/Li+. This is because the Li metal

plated on the Ni surface is not dense, measured is a mixed potenFig. 6. Potential profiles of the first Li plating and stripping on a fresh Ni surface,

which were recorded at 0.2 mAcm−2 for 3 h on a 3-electrode cell: (1) Catholyte-A

and (2) Catholyte-B.

tial ofthe Li metal and polysulfide solution. Subsequent Li stripping

test gives a 26% Li cycling efficiency. In Catholyte-B, the potential

remains a plateau for 82 min, and then steeply drops to a constant

negative value (−50 mV vs. Li/Li+), at which Li plates progressively,

giving a 50% cycling efficiency as shown by the subsequent Li stripping test. It is clearly shown that in both catholyte solutions, Li

metal cannot be plated until a protective passivation film is formed

on the Ni surface, and that such a film must be formed by the irreversible electrochemical reduction of polysulfides, as expressed by

Eq. (4).

Coulombic efficiencies of Li plating and stripping in Catholyte-A

and Catholyte-B are compared in Fig. 7. Obviously, the efficiencies in Catholyte-B are significantly higher than those obtained

in Catholyte-A, verifying that LiNO3 promotes the formation of a

more protective (normally denser) passivation film. Itis interesting

to note that in both Cathode-A and Cathode-B, Li cycling efficiencies are increased slowly with cycle number, probably because the

passivation film becomes denser and denser, resulting in better

protection of Li metal from Reactions (4) and (5).

On the other hand,the essential role of LiNO3 in the formation of

a passivation film on the Li surface can be examined by impedance

analyses. Fig. 8 shows the impedance spectrums of the Ni surface

after Li plating and after subsequent Li stripping. For the Li-plated

Ni surfaces indicated by Curves 1 and 2 in Fig. 8, the impedance

spectrum is composed of two flatted semicircles. In general, the

semicircle in higher frequency range corresponds to a passivation film and the one in lower frequency range corresponds to

the charge-transfer process occurring on the electrolyte–electrode

interface [30,31]. For Li metal electrode,the charge-transfer usually

reflects the following reversible electrochemical process:

Li − e− ⇔ Li+ (6)

Fig. 7. Coulombic efficiency of Li plating and stripping on Ni surface in Catholyte-A

and Catholyte-B, respectively.
58 S.S. Zhang, J.A. Read / Journal of Power Sources 200 (2012) 77–82 81

Fig. 8. Impedance spectroscopes of a Ni electrode in Catholyte-A and CatholyteB, respectively, which were recorded at open-circuit potential on a 3-electrode

cell: (1) after Li plating at 0.2 mAcm−2 for 3 h in Catholyte-A, (2) after Li plating

at 0.2 mAcm−2 for 3 h in Catholyte-B, (3) after Li stripping at 0.2 mAcm−2 to 1.0V

following (1), and (4) after Li stripping at 0.2 mAcm−2 to 1.0V following (2).

Comparison of Curves 1 and 2 shows that the Li metal plated

in Catholyte-B (Curve-2) has a higher passivation resistance and

a much lower charge-transfer resistance. The higher passivation resistance is a good indication that the passivation film

formed in the presence of LiNO3 co-salt is denser and hence more

protective. After subsequent Li stripping, the charge-transfer process as expressed by Eq.(6) no longer takes place due to the absence

of Li metal, and therefore the semicircle in the lower frequency

range disappears as indicated by Curves 3 and 4, respectively. By

comparing Curve-1 with Curve-3 and Curve-2 with Curve-4, one

sees that the passivation film still remains even after Li stripping,

indicating that the passivation film is permanently present on the

electrode surface once it has formed.

The benefit of LiNO3 also can be observed from cyclic voltammetric results as shown in Fig. 9. First, the cyclic voltammogram

(CV)in Catholyte-B has much higher peak currents. This result coincides with the impedance results (i.e., much lower charge-transfer

resistance in Catholyte-B). Second, the CV in Catholyte-B gives

much higher Li cycling efficiency (i.e., 65% vs. 29% in CatholyteA, calculated from the cyclic voltammograms). This result agrees

with those obtained from galvanostatic plating and stripping tests

Fig. 9. Cyclic voltammograms of Ni electrode in Catholyte-A and Catholyte-B, which

were recorded by scanning potential at 2 mVs−1 between 0.1V and 0.5V vs. Li/Li+.

as shown in Fig. 7. Third, Li platting in Catholyte-B has no polarization while that in Catholyte-A shows polarization, as indicated by

the arrow in Fig. 9.

Based on the discussion above, the charging voltage profiles

of Cell-2 in Figs. 3a, b and 4 can be explained as follows: as Eq.

(5) shows, Li metal and Li2Sx chemically react with each other.

While protecting Li metal, the passivation film meanwhile protects Li2Sx from attack by the highly reductive Li metal. Moreover,

the passivation film serves as a solid electrolyte interphase (SEI) to

prevent soluble Li2Sx (x > 2), especially those species having long

S–S chain near the full charge state, from being electrochemically

reduced, as expressed by Eq. (4). Instead, the soluble Li2Sx species

are electrochemically oxidized to elemental sulfur on the cathode,

resulting in an additional voltage plateau at 2.3V, higher charging efficiency, and accordingly a steep voltage rise upon the full

charge.

4. Conclusions

This work demonstrates an alternative approach for the performance improvement of rechargeable Li/S batteries. While

dissolution of lithium polysulfides (Li2Sx, x > 2) in organic electrolytes is inevitable, research efforts focusing on the protection

of the lithium anode to increase Li cycling efficiency in highly

concentrated polysulfide solutions may be more feasible. LiNO3 is

excellent in promoting the formation of a denser and more protective passivation film on the Li surface. The film formed not only

increases Li cycling efficiency, but also protects soluble polysulfide anions from chemical and electrochemical reductions on the

Li anode. Use of LiNO3 as a co-salt makes the two-phase reaction

of “S8 (insoluble) + 2Li→Li2S8 (soluble)” reversible, resulting in a

2.3V plateau, higher specific capacity and higher charging efficiency. Moreover, Li/S cells with LiNO3 co-salt can be repeatedly

charged to cutoff voltage (>2.5V) and indicate a steep voltage rise

as the signal of full charge. On the contrary, the conventional Li/S

cells can only be charged to 2.5V at which the voltage stays constant untilthe charging process is manually terminated, resulting in

low specific capacity and low charging efficiency, due to the chemical and electrochemical reductions of polysulfide anions on the Li

anode.

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61

Research@

Oxygen Reduction Reaction Catalyst on 

Lithium/Air Battery Discharge Performance

Xiaoming Ren, Sheng S. Zhang, Dat T. Tran, and Jeffrey A. Read

U.S. Army Research Laboratory, Adelphi, MD

ABSTRACT

Lithium/air batteries have the potential to substantially outperform the best 

battery system nowadays on the market. Oxygen reduction reaction (ORR) 

at the cathode in an aprotic organic lithium electrolyte is well-known to limit 

the discharge rate and capacity of the lithium/air batteries. In this study, the 

discharge characteristics of Li/air cells with cathodes made of different carbon 

materials were examined. The results showed that the ORR kinetics in the 

lithium/air batteries can be drastically improved by using an effective catalyst, 

achieving higher discharge voltage and rate. The discharge capacity of the 

lithium/air battery was found to be correlated to the cathode pore volume, to 

which the mesopore volume of the carbon material has a large contribution. An 

ORR mechanistic model involving a reaction product deactivating the catalytic 

sites on the carbon surface is proposed to explain the experimental results.

Journal of Materials Chemistry, Volume 21, Pages 10118-10125 (2011)
Oxygen reduction reaction catalyst on lithium/air battery discharge

performance

Xiaoming Ren,* Sheng S. Zhang, Dat T. Tran and Jeffrey Read

Received 30th November 2010, Accepted 25th January 2011

DOI: 10.1039/c0jm04170j

Lithium/air batteries have the potential to substantially outperform the best battery system nowadays

on the market. Oxygen reduction reaction (ORR) at the cathode in an aprotic organic lithium

electrolyte is well-known to limit the discharge rate and capacity of the lithium/air batteries. In this

study, the discharge characteristics of Li/air cells with cathodes made of different carbon materials were

examined. The results showed that the ORR kinetics in the lithium/air batteries can be drastically

improved by using an effective catalyst, achieving higher discharge voltage and rate. The discharge

capacity of the lithium/air battery was found to be correlated to the cathode pore volume, to which the

mesopore volume of the carbon material has a large contribution. An ORR mechanistic model

involving a reaction product deactivating the catalytic sites on the carbon surface is proposed to explain

the experimental results.

1. Introduction

Lithium/air batteries have the potential of achieving a higher

energy density and are safer than the leading commercial lithium

batteries, primary or secondary, by utilizing ambient air as the

oxidant, thus removing the need to store any oxidant within

enclosed cells. A Li/air battery is expected to have a theoretical

charge capacity 5–10 times higher than that of a lithium battery,1

which is limited by the amount of lithium that can be reversibly

inserted in the intercalation cathode material such as LixCoO2,

0.5 < x < 1. Since the first introduction of a rechargeable Li/air

battery by Abraham and Jiang in 1996,2 there has been much

progress being made in enhancing and understanding its

performance. For practical applications, electrolyte solutions

made with non-hydrolytic lithium salts and aprotic organic

solvents with low volatility and high electrochemical stability

were used.3 Other electrolyte systems were also reported, such as

Deng et al.4 for the use of a hydrophobic ionic liquid–silica–

PVdF–HFP polymer composite electrolyte membrane to minimize moisture exposure to the lithium anode and Wang and

Zhou5 for an aqueous electrolyte at the air cathode and aprotic

organic electrolyte at the lithium anode with the two electrolyte

solutions separated by a super-ionic glass conductor. In aprotic

organic electrolyte, the cell electrode reactions involved in the

discharge process are:

Anode: Li / Li+ + e�

Cathode: 2Li+ + O2 + 2e� / Li2O2 (solid)

and the complete cell reaction:

2Li + O2 / Li2O2 (solid), Vcell0 1⁄4 2.96 V.

The lithium anode has shown very little polarization at the

discharge current density of interest, at less than a few mA cm�2

,

6

and the loss in cell voltage from the reversible value (Vcell0) occurs

largely at the cathode. Many studies have been devoted to

address this cathode voltage loss by exploring oxygen reduction

reaction (ORR) catalysts for the cathode and by studying the

ORR mechanism in aprotic organic electrolyte solutions.

Cathode catalysts, such as carbon supported MnOx,

7 a-MnO2

nanowires,8 carbon supported nanosized g-MnOOH (manganite),9 carbon supported Fe2O3, Fe3O4, CuO, CoFe2O4 and

Co3O4,

10 carbon supported Pt and Au,11 and carbon supported

pyrolyzed Co macrocyles,2 have not shown much improvement

to the Li/O2 (air) battery discharge voltage over what has been

obtained with using carbon only as the ORR catalyst in the

cathode. From such observations, it has been assumed that the

ORR in a Li–air cathode is not a catalytically sensitive process or

that the carbon itself can provide better catalytic activity than

those catalysts of interest which themselves are supported on

carbon.11 Furthermore, study on the ORR mechanism in aprotic

organic lithium electrolyte solution on glassy carbon electrode by

Laoire et al.12 using cyclic voltammetry and rotating disc electrode technique showed the initial formation of lithium superoxide (LiO2), which can further be converted to lithium peroxide

U.S. Army Research Laboratory, RDRL-SED-C, Adelphi, MD,

20783-1197, USA. E-mail: xiaoming.ren@us.army.mil; Fax: +1 (301)

394-0273; Tel: +1 (301) 394-0379

10118 | J. Mater. Chem., 2011, 21, 10118–10125 This journal is a The Royal Society of Chemistry 2011

Journal of Dynamic Article LinksC<

Materials Chemistry

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63

either by disproportionation reaction (eqn (4)) or by further

reduction reaction:

2LiO2 / Li2O2 (solid) + O2

LiO2 + Li+ + e� / Li2O2 (solid)

Computational studies13,14 indicate that LiO2 is likely very

unstable at room temperature at less than 1 atm. O2 pressure,

favoring the disproportionation to Li2O2, which has been identified as the major reaction product by ex situ examination of the

cathode products from discharged Li/O2 cells using Raman

spectroscopy2 and by oxygen consumption stoichiometry in the

discharge process.3 At a low discharge current density, Zhang

et al.6 found part of the initially deposited Li2O2 at the cathode

can be further converted to Li2O at a highly polarized voltage

below 2 V.

The precipitation of Li2O2 as a solid product on the carbon

surface in the cathode further complicates the ORR process in

discharging a Li/air cell by either blocking the oxygen diffusion

pathways in the pores or by occupying and deactivating the

catalytic sites. Zhang et al.6 found the impedance of the air

electrode is progressively increased with polarization cycles,

indicating the carbon surfaces are gradually covered by the

insoluble products, which prevents oxygen from diffusing to the

reaction sites on carbon. Mirzaeian and Hall15 found the

discharge capacity and discharge voltage of Li/air cell depend on

the morphology of carbon, which exerts a combined effect of

pore volume, pore size and surface area of carbon on the storage

capacity, with carbon with a larger pore volume and a wider pore

size preferred. Williford and Zhang16 analyzed several

approaches in designing the air electrode by considering the

electrode porosity and catalyst reactivity distributions to minimize diffusion limitations and maximize air electrode material

utilization. Tran et al.17 proposed a model mechanism based on

gas diffusion electrode passivation by the reaction products in

blocking small pores and thus preventing them from further

utilization, again emphasizing the use of carbon materials possessing high surface area and large pore diameter in the cathode.

On the contrary, Read et al.18 found the BET surface area of the

carbon in the air electrode is not a significant factor in determining the discharge capacity, which is correlated to oxygen

transport in organic electrolyte. Dramatic decrease in the

discharge capacity at a high current density was linked to the

rapid decrease in cell voltage and uneven distribution of Li2O2

deposition concentrating at the air interface.6 The uneven Li2O2

deposition may have been one of the factors in explaining the

vast difference in the specific discharge capacities reported in

literature, up to a high value of 5800 mA h per g carbon,19 when

normalized to a low carbon mass in the cathode. Xu et al.20

explored using tris(pentafluorophenyl)borane as a functional

additive and co-solvent in electrolytes to dissolve part of Li2O

and Li2O2 in order to achieve higher Li/air battery discharge

capacity, although other factors introduced by the same additive

adversely affect the discharge capacity.

It is evident that the poor ORR cathode performance severely

limits the discharge rate and capacity of a Li/air cell. However,

no clear and systematic results have been obtained in finding an

effective catalyst to improve the ORR kinetics. In addition, it is

still not clear on the air cathode performance loss mechanism as

whether it is from Li2O2 deposits blocking the narrow pores of

the carbon material or Li2O2 covering and deactivating the ORR

catalytic sites on the carbon surface. To address these issues, we

examined in this study the discharge characteristics of Li/air cells

with cathodes made of three carbon materials: Super P carbon

(SP-carbon), Ketjen carbon (K-carbon) and pyrolized CuFe

macrocycle compounds on Ketjen carbon (CuFe catalyzed

K-carbon). These three carbon materials differ in the type and

number of the catalytic sites for the ORR, and in the carbon pore

volume and pore distribution. Comparison of the Li/air cell

discharge behaviors under identical test conditions for the cells

made with these cathode materials could provide the opportunity

to identify the key contributing factors to the cell performance in

terms of cell discharge voltage, rate and capacity, and thus shed

light on the complex Li/air discharge process. Based on experimental results, an ORR mechanistic model involving a reaction

product deactivating the catalytic sites on the carbon surface is

proposed.

2. Experimental

Super P carbon (SP-carbon) was obtained from TIMCAL

Graphite Carbon, and Ketjenblack EC-600 JD carbon black (Kcarbon) from Akzo Nobel. The CuFe-catalyzed Ketjenblack EC600 JD (CuFe-catalyzed K-carbon) was a non-precious metal–

oxygen reduction catalyst produced on a commercial scale and

provided by Acta SpA, Italy. This carbon supported catalyst was

made first by absorbing a mixture of iron and copper complexes

with phthalocyanine-based ligands onto the carbon support, and

then heat-treated at between 800 and 900 �C in Ar atmosphere.

The CuFe-catalyzed K-carbon has a Cu content at 1.7 wt% and

a Fe content at 1.5 wt%, corresponding to an atomic ratio of Cu

to Fe at 1 to 1.

X-Ray powder diffraction (XRPD) patterns of the carbon

materials were collected on a Rigaku Ultima III diffractometer

using a Bragg–Brentano geometry and Cu Ka radiation

(l 1⁄4 1.5418 A ̊ ) over a 2q scan range of 10–90� with a resolution of

0.02� at a scan rate of 0.2� per min. The surface areas of the

carbon materials were measured with a Micromeritics ASAP

2010 system using N2 gas as adsorptive. The pore size distributions of the carbon materials were obtained from Barrett–Joyner–Halenda (BJH) desorption pore distribution using ASAP

2010 V1.00 software.

Lithium triflate (LiSO3CF3, 96%, Aldrich) was dried at 100 �C

under vacuum for 8 h. Electrolyte grade propylene carbonate

(PC, Ferro) was used as received. Tris(2,2,2-trifluoroethyl)

phosphate (TFP) was synthesized by reacting sodium trifluoroethoxide with phosphorus oxychloride and purified by

repeated fractionation under reduced pressure. Detailed

descriptions on the synthesis and characterization of TFP are

referred to in ref. 21 and 22.

Electrolyte solution with a fixed solvent composition of

PC : TFP at a 7 : 3 weight ratio containing 0.2 M lithium triflate

was prepared in a glove-box. Physical properties of this electrolyte solution, such as kinematic viscosity, ionic conductivity,

oxygen solubility, and boiling point, are referred to in ref. 23.

This journal is a The Royal Society of Chemistry 2011 J. Mater. Chem., 2011, 21, 10118–10125 | 10119

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64

Cathodes with a composition of 90 wt% carbon materials and

10 wt% polytetrafluoroethylene (PTFE) were prepared by mixing

calculated amounts of carbon materials and PTFE emulsion

(Teflon�, solid content 1⁄4 61.5%, DuPont Co.), and then rolling

the resulting paste mixture into a freestanding cathode sheet,

which was punched into disks with an area of 0.97 cm2 and dried

at 100 �C under vacuum for over 8 h. Li/air cells with an air

access window of 0.97 cm2 were assembled in a dry-room having

a dew point below �90 �C by stacking sequentially a Li foil,

a Celgard� 3500 membrane, a cathode, a Ni mesh (as cathode

current collector), and an air-window frame into a coin cell cap.

To activate the cell, a volume of 200 mL liquid electrolyte solution

was added through the air-window, followed by applying

vacuum for 20 s to ensure complete filling of the electrolyte into

the cell internal space. Excessive electrolyte solution over the

cathode outer surface was removed by gently swiping with a filter

paper over the Ni mesh. The electrolyte-activated cell was rested

for 2 h before commencing discharging tests, which were carried

out at room temperature (22 �C) in the dry-room on an Arbin

BT-2000 tester galvanostatically from OCV until reaching the

cutoff voltage at 1.5 V. Specific capacity of the cell was

normalized by the mass of the carbon materials in the air

cathode, which was in the range of 6–7 mg cm�2

.

3. Results and discussions

3.1. Characterization of carbon materials and cathodes

The XRD patterns of the carbon materials shown in Fig. 1 reveal

the relative degree of graphitization of the three carbon materials. SP-carbon has the highest degree of graphitization of the

three carbon materials as shown by the sharp graphitic basal

plane (002) peak at a d-spacing of 3.573 A ̊ . Relative to K-carbon,

the heat-treatment at around 800–900 �C in producing CuFe

catalyzed K-carbon increased carbon graphitization degree, and

shrank basal plane d-spacing (from 3.722 to 3.620 A ̊ ) towards

that of an ideal graphite (3.354 A ̊ ). The sizes of the crystallites

parallel to the graphite basal plane for the carbon materials listed

in Table 1 are calculated by using the equation L 1⁄4 1.84l/(Bcos

q), where l is the wavelength of the X-ray beam, B is the angular

width of the basal plane diffraction (002) peak at the halfmaximum intensity, and q is the Bragg angle. The Brunauer–

Emmett–Teller (BET) surface areas obtained using N2 gas

adsorption at 77 K for these three carbon materials are listed in

Table 1, and the carbon pore distribution and accumulated pore

volume as a function of pore diameter obtained using BJH

method is plotted in Fig. 2. SP-carbon possesses little internal

pore volume and a rather low surface area, as compared to those

of K-carbon and CuFe catalyzed K-carbon. The loss of carbon

pore volume by the heat treatment used in making CuFe catalyzed K-carbon is consistent with the increase in graphitization

degree of the carbon material as revealed by the XRD pattern.

Also, the pore volume distribution of CuFe catalyzed K-carbon

remains the same as that of the original K-carbon, which

suggests a uniform structural collapse occurred during the heat

treatment. This evidence does not support the suggestion of

CuFe filling up or blocking some pores of K-carbon, especially at

a metal loading at less than 3% of the carbon support.

The porosities of the cathodes made of the three carbon

materials were measured using two separate methods: solvent

filling and wet thickness. With the solvent filling method, the

cathodes were vacuum filled with PC solvent, followed by

removing excessive solvent on the electrode surfaces by gently

dabbling on both sides with filter papers until reaching a stable

mass. The carbon electrode porosity obtained using the solvent

method is calculated according to the following equation:

r 1⁄4



Wwet � Wdrydc 

Wwet � Wdrydc þ Wdrydsol

100%

where Wdry is the dry mass of the electrode, Wwet is the electrode

mass wetted with solvent, dc is the density of carbon (2.0 g cm�3

),

and dsol is the density of solvent (1.206 g cm�3 for PC). With the

thickness method, the porosities were calculated from the dry

mass and the wet thickness when filled with PC solvent. The

porosity obtained using the thickness method is calculated

according to the following equation:

r 1⁄4



1 � Wdry

d Adc



100%

where A is the electrode area and d is the thickness of electrode

wetted with solvent. Both these methods yielded consistent

results for the electrode porosity measurements listed in Table 1.

A trend observed is that a cathode made of a carbon material

possessing a higher internal volume of mesopores (and consequently a higher surface area) has a higher electrode pore

volume.

3.2. Li/air cell discharge curves

In Fig. 3, cell voltage curves during discharge at a constant

current density of 0.2 mA cm�2 for the Li/air cells with cathodes

Fig. 1 XRD patterns for SP-carbon, K-carbon and CuFe catalyzed Kcarbon.

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made of three carbon materials are compared. Without catalyst,

a cell with a K-carbon cathode out-performed a cell with a SPcarbon cathode in terms of cell discharge voltage and capacity. A

cell with a CuFe catalyzed K-carbon cathode exhibited two

distinguishable voltage plateaus in its discharge voltage curve,

with the first discharge plateau at a cell voltage above 2.5 V

demonstrating a substantial improvement in cell discharge

voltage of over 200 mV higher than that of a cell with a K-carbon

cathode, and over 500 mV higher than that of a cell with a SPcarbon cathode; while the second discharge plateau at a voltage

window between 1.7 to 1.5 V has a substantially higher discharge

capacity than that of a cell with a K-carbon cathode. The

discharge curve of the cell with a SP-carbon cathode did not have

the second discharge plateau in the voltage window studied.

Results of further efforts carried out to identify the origin of

the second discharge plateau observed for the Li/air cells with Kcarbon and CuFe catalyzed K-carbon cathodes, especially from

the charge contribution by the possible cathode process of Li2O2

to Li2O conversion, are presented in Fig. 4. As will be shown by

the experimental results, there is negligible charge contribution

by the Li2O2 to Li2O conversion for the Li/air cells. Shown in

Fig. 4 (top), the Li/air cells were first discharged at a constant

current density of 0.2 mA cm2 from point A to point B to reach

a state of discharge at a capacity of 300 mA h g1

. After that, the

two cells were reassembled in sealed coil cells to close off the air

access to the cells, and then discharged as Li/C cells at a current

density of 0.05 mA cm2 from point B until reaching the cutoff

voltage at 1.5 V. For the cell with a K-carbon cathode, the Li/C

cell discharge step reached a discharge capacity close to 80 mA h

g1

; while for the cell with a CuFe catalyzed K-carbon cathode,

a discharge capacity over 480 mA h g1

, which is by far higher

than what would be expected if the charge were originated from

further reduction of Li2O2 produced at a charge capacity of 300

mA h g1 during the initial Li/air cell discharge step (from point

A to point B in Fig. 4a). To further verify if there is any charge

contribution by the possible cathode process of Li2O2 to Li2O

conversion, hermetic Li/C cells with fresh cathodes were built

and discharged in the absence of oxygen access. As shown in

Fig. 4b, the charges obtained with the fresh cathodes almost

Table 1 Physical properties of carbon materials, cathode porosity and specific discharge capacity of Li/air cells with cathodes made of the carbon

materials

SP-carbon K-carbon CuFe catalyzed K-carbon

Graphitic basal plane d-spacing/A ̊ 3.573 3.722 3.620

Crystallite size parallel to basal plane/A ̊ 39.6 20.8 30.2

BET surface/m2 g1 69.3 1413 751

Total pore vol. @ >20 A ̊ dia./cm3 g1 0.14 2.06 1.23

Electrode porosity by solvent methoda 77.3  1.8% 90.8  0.5% 86.8  0.7%

Electrode pore vol.,b

/cm3 g1

carbon 1.89 5.46 3.64

Electrode porosity by thickness methoda 75.3  2.2% 90.9  1.1% 87.9  1.5%

Li/air cell discharge capacity, mA h g1

carbon

@ 0.05/mA cm2 531 1286 1339

@ 0.20/mA cm2 356 761 817

@ 0.50/mA cm2 205 430 597

@ 1.00/mA cm2 — 165 390

a Average and standard deviation from measurements on 6 samples. b Calculated from electrode porosity by solvent method.

Fig. 2 Incremental pore volume and accumulated pore volumes as

a function of pore diameters obtained using BJH method for SP-carbon,

K-carbon and CuFe catalyzed K-carbon.

Fig. 3 Li/air cell discharge voltage curves at 0.2 mA cm2 for cells with

cathodes made of (1) SP-carbon, (2) K-carbon and (3) CuFe catalyzed

K-carbon.

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66

account for all of the charges obtained from the corresponding

two cells shown in Fig. 4a in the discharge step starting from

point B. The two discharge curves shown in Fig. 4b and that in

Fig. 4a starting from point B for cells with CuFe catalyzed Kcarbon cathodes are re-plotted in Fig. 4c, where the two curves

overlap each other completely, verifying the negligible charge

contribution by the Li2O2 to Li2O conversion for the Li/air cells

tested. This conclusion could be rationalized by the non-electronic conductivity and low solubility of Li2O2 in the chosen

electrolyte solution.

The second discharge plateau observed for the cell with a Kcarbon cathode (shown in Fig. 3, 4a from point B, and 4b) could

be attributed to the surface adsorbed oxygen on the carbon

material, and for the cell with a CuFe catalyzed K-carbon

cathode to both of the surface adsorbed oxygen on the carbon

material and solvent reduction. The lack of surface adsorbed

oxygen on SP-carbon could be attributed to its much lower

surface area and a more ordered surface structure. Solvent

reduction often occurs on the carbonaceous anode in the

formation cycling of Li-ion for Li rechargeable batteries, albeit at

a lower voltage plateau around 0.9 V without using catalyzed

carbon materials at the electrode. To understand the reaction

mechanism of the CuFe catalyzed carbon in catalyzing the

solvent reduction further investigation is needed. With

purposeful selection and optimization of the solvent and catalyst

pairs, the solvent redox process catalyzed by the cathode material

could be utilized as new cell chemistry in designing rechargeable

Li batteries.

3.3. Li/air cell discharge rate and capacity

The discharge voltages obtained for cells with K-carbon and

CuFe catalyzed K-carbon cathodes at various discharge current

densities are shown in Fig. 5, and these results are summarized in

Table 1 and in Fig. 6 in a plot of cell energy density as a function

of discharge current density. The discharge voltage of the cell

with a CuFe catalyzed K-carbon cathode is over 200 mV higher

than that of the cell with a K-carbon cathode at a low discharge

current density, and over 1 V higher at a high discharge current

density. The Li/air cell with a CuFe catalyzed K-carbon cathode

discharged at a relatively high current density of 1 mA cm�2

demonstrated an energy density over 0.8 Wh g�1, which is more

than twice of that of the cell with a K-carbon cathode.

The discharge capacities of the Li/air cells, shown in Fig. 3 and

5a, are plotted as a function of the cathode pore volume in Fig. 7.

There is a good correlation between the cell discharge capacity

and electrode pore volume for these three carbon materials,

especially at a low discharge current density of 0.05 mA cm�2

,

where it is estimated 10% of the cathode pore volume is filled by

the electrode reaction product of Li2O2 at the end of the

discharge process. At a higher discharge current density, there is

a rapid decrease in the discharge capacity. At the fast discharge

rate of 1 mA cm�2

, no visible depositions of the reaction products

Fig. 4 a) Li/air cells with cathode made of (1) K-carbon and (2) CuFe

catalyzed K-carbon discharged from point A to point B at 0.2 mA cm�2

for 300 mA h g�1, then the cells containing Li2O2 produced were re-assembled in enclosed coin cells to block off the air access, and then discharged as Li/C cells without oxygen access from point B at 0.05 mA cm�2

to 1.5 V. (b) Li/C cells discharge voltage curves obtained at 0.05 mA cm�2

for cells with cathodes made of (10) K-carbon and (20) CuFe catalyzed Kcarbon. Initial capacity at voltage greater than 2 V is due to the oxygen

absorbed in electrolyte solution. (c) Discharge voltage curves for cells

with CuFe catalyzed K-carbon, where curve (2) is obtained in (a) from

point B discharge process and curve (20

) in (b).

Fig. 5 Li/air cell discharge voltage curves obtained at various discharge

current densities for cells with cathodes made of (1) K-carbon and (2)

CuFe catalyzed K-carbon.

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67

at the cathode out-surface were observed, likely due to the low

capacity obtained. For the three types of carbon cathodes, the

difference in the obtained discharge capacity is significant: the

capacity of SP-carbon is at near zero, and that of CuFe catalyzed

K-carbon is about twice that of the original K-carbon. These

results would argue against the clogging pore at the cathode

surface being the limiting factor for the discharging capacitance

at the fast discharging rate. Furthermore, if it were clogging

pores at the carbon surface in limiting the discharge capacity, the

original K-carbon would have a higher capacitance than the

CuFe-catalyzed K-carbon due to its higher carbon pore volume.

The rapid build-up of lithium peroxide covering the active sites

for ORR on the carbon surface at a higher discharging current

density may have caused the fast passivation of the cathode, as

shown by the slant discharge voltage plateau, attributing to the

early termination of the cell discharge process.

3.4. Micrographs of cathodes after discharge

Because of the low solubility of the Li2O2 produced in discharging a Li/air battery, several factors are expected to affect its

production and precipitation during the discharging process, and

subsequently, its final location within the cathode, which can be

examined with a microscope after the discharge step is

completed. It has been observed that as the discharge current

density increases, the Li2O2 particles tend to be found within the

cathode at a location close to the air interface rather than the

electrolyte interface. Even at the same discharge current density,

more Li2O2 deposits were observed on the surface of the cathode

facing air for a cell with CuFe catalyzed K-carbon than those

with K-carbon, as shown in Fig. 8, where large crystals of Li2O2

were found on the cathode surface and in the cracks close to the

air surface with CuFe catalyzed K-carbon. The formation of

large Li2O2 crystals on the surface and within the cracks of the

electrode implies a meaningful solubility and mobility of Li2O2 in

the electrolyte solution needed for the crystal growth from the

Li2O2 initially produced on the ORR catalytic sites during the

discharging process. The fact that Li2O2 was found at a location

close to the air surface of the cathode also indicates the rather

slow diffusion of dissolved oxygen in the non-aqueous electrolyte

solution. At a high discharging current density or with a cathode

material possessing a higher density of active sites as offered by

the CuFe catalyzed K-carbon cathode material, the ORR reaction zone moves to the air surface of the cathode. The severe

Fig. 6 Plot of energy density as a function of discharge current density

from data shown in Fig. 5 for Li/air cells with cathodes made of (1)

K-carbon and (2) CuFe catalyzed K-carbon.

Fig. 7 Plots of Li/air cell specific discharge capacities as a function of

cathode pore volume.

Fig. 8 Optical microscope pictures of cathodes facing air for cathodes

made of (1) K-carbon and (2) CuFe catalyzed K-carbon in Li/air cells

discharged at 1.0 mA cm2 to 1.5 V, rested for 5 h, and then discharged at

0.05 mA cm2 to a total capacity of 850 mA h g1

. Deposits from cathode

reaction of discharging Li/air cells were seen as the white spots on the

black carbon surface and as colorless crystals formed (shown by arrows)

within the cracks of the cathode.

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decrease in cell discharge capacity at high current density is

attributed to both increased cathode passivation and uneven

distribution of Li2O2 deposits within the pore volume of the

cathode.

3.5. Li/air cell polarization

The dynamic Li/air cell discharging polarization curves, shown in

Fig. 9 dashed lines, were obtained at a current density scan rate

of 0.02 mA cm2 s1 for cells with K-carbon cathode and CuFe

catalyzed K-carbon cathode after both cells were discharged to

a charge capacity of 300 mA h g1 and then followed by a rest

period of two hours. Previous study has shown that the lithium

anode has negligible polarization, and the cell voltage polarization is largely attributed to the cell cathode polarization and cell

IR drop. Under the dynamic scan conditions, the polarization

curves for both cells are similar, with the cell with a CuFe catalyzed K-carbon cathode showing a slightly better ORR kinetics

at the low current density range, and the cell with a K-carbon

cathode a slightly better mass transport at the high current

density range. However, when the cell polarization curves representing the steady state discharge as obtained from data shown

in Fig. 5a–d are compared, there are large differences between the

dynamic and steady-state conditions for the two cells. For the

cell with a CuFe catalyzed carbon cathode, the polarization

curve at steady-state shifts downward by about 80 mV from that

at dynamic scan; while for the cell with a K-carbon cathode, the

corresponding downward shift in polarization curve is over

300 mV, and increasing with the current density to over 500 mV

at 1 mA cm2

. It was expected that the ORR polarization from

the increased mass transportation limitation from the dynamic

state to steady-state is similar for the two cells. The much larger

voltage downward shift, especially at a higher current density, for

a cell with a K-carbon cathode than that for a cell with a CuFe

catalyzed K-carbon cathode indicates that there is added sluggishness in the ORR kinetics under the steady-state conditions.

Under the steady-state conditions, there is a continuous formation of Li2O2 on the ORR catalytic sites, and its subsequent

removal by dissolution and crystal growth keeps the number of

active sites constant, and the number of available free active sites

decreases with the increase in discharge current density. The

higher ORR catalytic activity provided by a higher level of free

site density found in CuFe catalyzed K-carbon material, which

contains catalytic sites consisting of CuFe and those on Kcarbon, decreases the ORR polarization under the steady-state

conditions.

3.6. ORR electrode process

It becomes evident from above test results and observations that

the performance of a Li/air cell is limited by the ORR electrode

process, which could involve the following key steps:

O2 (g) / O2 (sol) (1)

O2 (sol) + *–CS / O2–CS (2)

2Li+ + O2–CS + 2e / Li2O2–CS (3)

Li2O2–CS / *–CS + Li2O2 (sol) (4)

Li2O2 (sol) / Li2O2 (solid)Y (5)

In this cathode reaction mechanism, the oxygen gas molecule

(O2(g)) from air first dissolves in the electrolyte solution forming

dissolved oxygen molecules (O2(sol)), which then adsorbs onto

a free ORR catalytic site (*–CS) on the carbon surface. The

subsequent ORR reaction produces one Li2O2 molecule attached

to the catalytic site (Li2O2–CS) resulting in deactivating the

catalytic site for further ORR. The deactivated catalytic site by

attached Li2O2 is only freed by removal of the Li2O2 molecule

into the surrounding electrolyte solution. The driving forces for

such a removal process come from the Li2O2 diffusion into the

electrolyte solution within the electrode pore volume, and from

the aging-growth process of some Li2O2 crystals at nearby

locations. The sustainable Li/air cell discharge rate would thus

depend on the initial number of free catalytic sites on the carbon

surface and the balancing act of Li2O2 production and its

removal in maintaining a sufficient number of free catalytic sites

to sustain the ORR. At an excessively high discharge current

density, the free catalytic sites diminish rapidly because of the

relatively slow rate of Li2O2 removal from the deactivated

catalytic sites, resulting in fast increase in ORR polarization. As

shown in Fig. 5d, the Li/air cell with a CuFe catalyzed K-carbon

cathode has two distinguishable cell voltage plateaus at above

2 V when discharged at a relatively high current density of

1 mA cm2

. The two voltage plateaus reflect the ORR process

carried out at two different types of catalytic sites in the cathode

material. The first discharge plateau arises from ORR catalyzed

by the more active CuFe catalytic sites. After these CuFe sites are

deactivated at the end of the first discharge plateau, the ORR

proceeds at the less active catalytic sites on the supporting carbon

surface, forming a second voltage plateau, which is at a similar

voltage to what was observed for a Li/air cell with a K-carbon

cathode where only the active sites provided by the K-carbon

exist. At the fast discharge rate, all active sites are deactivated

and cause rapid termination of the discharge process. For the

CuFe-catalyzed K-carbon, the highly active CuFe catalytic sites

are deactivated first at a high cathode voltage, followed by the

Fig. 9 Li/air cell polarizations at steady-state (solid line) and at dynamic

scan at 0.02 mA cm2 s1 (dashed line), measured after discharged to

300 mA h g1 for cells with cathodes made of K-carbon (-) and CuFe

catalyzed K-carbon (:).

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69

deactivation of the less active sites of the carbon surface at

a lower cathode voltage. The charge contribution from absorbed

oxygen only occurs at the start of the discharging process, and

since oxygen supply is not the limiting factor in an operating

Li/air cell, the overall impact of absorbed oxygen on the

discharge capacity of a Li/air cell at a cell voltage >2.0 V is small

at less than 35 mA h g1

. The charge contribution from solvent

reduction occurs at a cell voltage bellow 1.8 V.

As for the Li/air battery discharge capacity obtained at a low

discharging current density, available electrode pore volume is

a determining factor. Carbon material possessing a high mesopore volume at a pore diameter greater than 20 A ̊ provides

additional electrode pore volume accessible for the ORR and for

accommodating the Li2O2 deposit. The Li/air cell discharge

capacity is also impacted by the cell discharging current density,

which affects the rate of cathode catalytic sites deactivation and

the distribution of Li2O2 within the cathode. The ability to

redistribute the Li2O2 reaction product more evenly throughout

the cathode pore volume in the thickness direction could provide

substantial improvement in the Li/air cell discharge capacity,

from currently 10% of the potential capacity estimated based on

the full occupancy of the electrode pore volume by the Li2O2

deposit. There are several possible approaches to address this

issue, and these measures include: (1) increasing the solubility of

Li2O2 in the electrolyte solution with a better selection of solvents

and electrolytes, (2) a further decrease in the lithium electrolyte

concentration to move the reaction zone from the air interface

towards the electrolyte interface of the cathode and (3) modifying

electrode structures to provide less tortuous diffusion paths for

the oxygen and Li2O2.

4. Conclusions

Based on the observations and test results of the Li/air batteries

with cathodes made of different carbon materials, a mechanism

for the ORR electrode process is proposed where the newly

formed Li2O2 molecules deactivate the catalytic sites, and the

density and activity of free catalytic sites on the surface of carbon

and the removal rate of attached Li2O2 from the deactivated

catalytic sites determine the ORR polarization, and thus the

discharge cell voltage.

It has been demonstrated that by increasing the catalytic site

density and activity on the carbon surface for the ORR, as shown

from the graphite-like SP-carbon, to the defect-rich amorphous

K-carbon, and to the CuFe catalyzed K-carbon, the ORR

polarization experienced during the Li/air discharging process

can be significantly decreased, thus considerably increase the

discharge cell voltage and rate for the Li/air batteries. Cells with

a CuFe catalyzed K-carbon cathode demonstrated a higher cell

discharging voltage of over 200 mV than that with K-carbon,

and of over 500 mV than that with SP-carbon.

It was observed that there is a negligible amount of Li2O2

being converted to Li2O during the Li/air cell discharge process.

A cell with a CuFe catalyzed K-carbon cathode has a distinguishable second discharge plateau in the voltage window from

1.7 to 1.5 V attributed to the solvent reduction catalyzed by the

CuFe catalyst.

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71

Research@

Effect of Substitution (Ta, Al, Ga) on the 

Conductivity of Li7La3Zr2O12

Jan L. Allen1, Jeffrey Wolfenstine1, Ezhiylmurugan Rangasamy2, and 

Jeff Sakamoto2

 1 U.S. Army Research Laboratory, Adelphi, MD 2 Department of Chemical Engineering and Materials Science, 

 Michigan State University, East Lansing, MI

ABSTRACT

Cubic garnets of composition Li6.75La3Zr1.75Ta0.25O12, Li6.15La3Zr1.75Ta0.25Al0.2O12, 

and Li6.15La3Zr1.75Ta0.25Ga0.2O12 were prepared from a co-precipitated precursor 

and consolidated by hotpressing to a relative density of ~96–98%. The total 

Li-ion conductivities at 298 K and activation energies (in parentheses) of 

Li6.75La3Zr1.75Ta0.25O12, Li6.15La3Zr1.75Ta0.25Al0.2O12 and Li6.15La3Zr1.75Ta0.25Ga0.2O12

were 0.87 mS cm−1 (0.22 eV), 0.37 mS cm−1 (0.30 eV) and 0.41 mS cm−1

(0.27 eV), respectively. The above results suggest that cubic stabilizing 

substitutions outside of the Li-ion sub-lattice are preferable to obtain faster 

Li-ion conductivity.

Journal of Power Sources, Volume 206, Pages 315-319 (2012)
72 Journal of Power Sources 206 (2012) 315–319

Contents lists available at SciVerse ScienceDirect

Journal of Power Sources

j ou rnal homepage: www.elsevier.com/locate/jpowsour

Short communication

Effect of substitution (Ta, Al, Ga) on the conductivity of Li7La3Zr2O12

J.L. Allena,∗, J. Wolfenstinea, E. Rangasamy b, J. Sakamoto b

a US Army Research Laboratory, RDRL-SED-C, 2800 Powder Mill Road, Adelphi, MD 20783, United States b Department of Chemical Engineering and Materials Science, Michigan State University, East Lansing, MI 48824, United States

a r t i c l e i n f o

Article history:

Received 27 December 2011

Received in revised form 18 January 2012

Accepted 21 January 2012

Available online 30 January 2012

Keywords:

Garnet

Li

Ta

Cubic

Ionic conductivity

Hot-pressing

a b s t r a c t

Cubic garnets of composition Li6.75La3Zr1.75Ta0.25O12, Li6.15La3Zr1.75Ta0.25Al0.2O12, and

Li6.15La3Zr1.75Ta0.25Ga0.2O12 were prepared from a co-precipitated precursor and consolidated by hotpressing to a relative density of ∼96–98%. The total Li-ion conductivities at 298K and activation energies

(in parentheses) of Li6.75La3Zr1.75Ta0.25O12, Li6.15La3Zr1.75Ta0.25Al0.2O12 and Li6.15La3Zr1.75Ta0.25Ga0.2O12

were 0.87 mS cm−1 (0.22 eV), 0.37 mS cm−1 (0.30 eV) and 0.41 mS cm−1 (0.27 eV), respectively. The above

results suggest that cubic stabilizing substitutions outside of the Li-ion sub-lattice are preferable to

obtain faster Li-ion conductivity.

Published by Elsevier B.V.

1. Introduction

Li-ion solid electrolytes have been long sought after in order

to obtain a safe, reliable, long-lived battery system without use

of flammable, volatile and relatively unstable organic liquid-based

electrolytes. The advantages are numerous including long shelf

life, ability to operate at high voltage or high temperature and

ease of manufacturing. However, solid electrolytes have not been

able to provide sufficient conductivity and/or have been unstable

to Li, air or moisture. The Li-stuffed Al-substituted cubic garnet

Li7La3Zr2O12 (LLZO) is a promising fast Li-ion conducting ceramic

electrolyte owing to its fast ionic conductivity (�total ∼ 0.2 mS cm−1

at 298K) and Li and moisture stability [1,2]. High voltage Liion cathodes [3] that increase the demands on flammable, liquid

organic electrolytes for high voltage stability [4] and safety considerations make a solid state electrolyte an attractive option.

Furthermore, Li-air or Li–S batteries [5] can be enabled by the discovery and development of faster Li-ion conducting solids.

The challenge has been to stabilize the cubic form since LLZO

can also exist with a tetragonal garnet structure [6,7] of lower Liion conductivity. The tetragonal structure results from ordering of

the Li-ion sub-lattice. The first approach to stabilization ofthe cubic

form was the serendipitous discovery that the inclusion of Al as an

impurity from contamination during processing [2,7] stabilized the

higher conducting cubic phase. Later Al was intentionally added

∗ Corresponding author. Tel.: +1 301 394 0291; fax: +1 301 394 0273.

E-mail addresses: jan.l.allen8.civ@mail.mil, jallen@arl.army.mil (J.L. Allen).

[8–11] and a dense (relative density > 90%) material with a cubic

structure was obtained. Geiger et al. [7] have suggested that Al substitutes for Li and thereby stabilizes the cubic phase relative to the

tetragonal phase. An Al substitution for Li will lead to a reduction

in Li content in order to maintain electro-neutrality (1Al3+ = 3Li+).

Others have speculated that Al addition yields a denser material

by acting as a sintering aid [9,10]. A similar stabilization of cubic

phase is observed with Ga substitution [12]. Thus it is desirable to

look at other substitutions which like the Geiger et al. approach

may stabilize the cubic garnet structure through a reduction in Li

content and/or increased Li vacancy concentration. Ta substitution

for Zr follows this approach and it is desirable for a couple reasons. First, Ta is stable relative to Li [13]. Second, Ta substitution

on the Zr site will not hinder Li-ion motion whereas Al substitution

on Li sites will (as discussed later). Li et al. [13] recently reported

Li6La3ZrTa012 cubic garnet with relatively high total Li-ion conductivity (0.18 mS cm−1 at 298K). The sample was prepared in an

alumina crucible and contained 1.3 wt% Al. Also, Logéat et al. [14]

recently reported that Al-free, Ta substitution for Zr could stabilize the cubic garnet (Li7−xLa3Zr2−xTax012). Conductivity data was

not included in their report. Similarly to Ta, Nb substitution for

Zr should also lower the Li content of the LLZO and a high Li ion

conductivity (0.8 mS cm−1 at 298K) cubic garnet sample has been

reported with Nb substitution for Zr by Ohta et al.[15]. However, Ta

is preferable to Nb since Nb is not chemically stable with Li [16–18]

and would change LLZO from an ionic conductor to a mixed conductor. In this paper, we have chosen a level of Ta substitution equal to

the level of Nb substitution reported by Ohta et al. [15] to have the

highest conductivity and lowest activation energy. It is the purpose

0378-7753/$ – see front matter. Published by Elsevier B.V.

doi:10.1016/j.jpowsour.2012.01.131
316 J.L. Allen et al. / Journal of Power Sources 206 (2012) 315–319 73

ofthis paper to report on the effects of Ta, Ta + Ga and Ta +Al substitutions in order to further the understanding of Li-ion conductivity

in LLZO.

2. Experimental

Li6.75La3Zr1.75Ta0.25O12, Li6.15La3Zr1.75Ta0.25Al0.2O12 and

Li6.15La3Zr1.75Ta0.25Ga0.2O12 were prepared for study from a

co-precipitated precursor. Li2CO3, La(OH)3, “[ZrO2]2·CO2·xH2O”

(zirconium carbonate, basic hydrate; equivalent ZrO2 content

determined from thermogravimetric analysis) Al(OH)3 and Ga2O3

weighed in the desired stoichiometry were dissolved in ∼1.4 M

HNO3 (aq). A 3 wt% excess of Li was used to compensate for Li

volatilization during synthesis. The desired amount of TaCl5 was

first dissolved in anhydrous ethanol and then added to the acid

solution. The resulting clear solution was evaporated to dryness

in a microwave oven contained inside a fume hood. Evolution of

NOx was observed during this step. The dried co-precipitate was

lightly ground with a mortar and pestle and pressed into a pellet

using a Carver laboratory die and press. The pellet was placed on a

ZrO2 plate and heated in air at 923K for 15 h and subsequently at

1273K for 3–4 h. The furnace was turned off and the sample was

removed.

Li6.75La3Zr1.75Ta0.25O12, Li6.15La3Zr1.75Ta0.25Al0.2O12 and

Li6.15La3Zr1.75Ta0.25Ga0.2O12 dense discs were prepared by hotpressing. The powders were hot-pressed at 1323K at 40 MPa

pressure for 1 h under air. From the hot-pressed discs rectangular

parallelepipeds were cut using a low-speed diamond saw for

density, microstructural and electrical property measurements.

X-ray diffraction (Cu K radiation) was used to characterize the

phase purity of the powders and the material after hot-pressing.

Lattice constants were calculated from the peak positions of the Xray diffraction pattern using Rietveld refinement of data collected

in parallel beam geometry. The bulk density of the hot-pressed

samples was determined from the weight and physical dimensions.

The relative density values were determined by dividing the bulk

density by the theoretical density. The microstructure of the hotpressed samples was examined on fracture surfaces using scanning

electron microscopy (SEM).

AC measurements were undertaken to determine ionic conductivity as a functionoftemperature inthe range of 298–373K.Au was

sputter coated on to the top and bottom surface of the specimens.

AC impedance was measured using a Solatron 1260 Impedance

Analyzer in the frequency range of 1–106 Hz.

3. Results and discussion

Conventional garnets are described by the formula A3B3O12

where A, B, and C have 8, 4, and 6 oxygen coordination, respectively. In a cubic stuffed garnet, La occupies A sites, Zr occupies C

sites and Li occupies the B sites and additional distorted octahedral

sites that are unoccupied in the conventional garnet [19]. Awaka

et al. [20] label the tetrahedral B site Li1 and the distorted octahedral site Li2 and they describe the basic unit of the Li sub-lattice

to be loop of Li1 and Li2 sites. The Li1 site is the only connection

between loops and these junctions of the basic loop unit form the

complete 3-D Li ion sub-lattice and enable fast Li-ion conduction.

A very high conductivity [2] was reported for cubic “Li7La3Zr2O12”

that was later shown to be a result of Al substitution for Li [7]. We

believe that Al substitution lowers the Li content and stabilizes a

cubic structure as hypothesized by Geiger et al. [7]. Ga can stabilize cubic garnet in a fashion similar to Al [12]. However based on

the Awaka et al. structural model of Li-stuffed cubic garnet [20],

the presence of Al or Ga on the Li site will be detrimental to Liion conduction and we therefore focused on substitutions that are

Fig. 1. The X-ray diffraction patterns of Li6.75La3Zr1.75Ta0.25O12,

Li6.15La3Zr1.75Ta0.25Al0.2O12 and Li6.15La3Zr1.75Ta0.25Ga0.2O12.

not part of the Li-ion conduction pathway. Nuclear magnetic resonance (NMR) data has shown Al occupying both the Li1 and the Li2

sites with preference for the Li1 site [7]. We speculated that Al on

these sites, and in particular Al on the Li1 site, might hinder ionic

conductivity since this site forms the junction between loops.

Thus, we wanted to control the Li content in order to stabilize the cubic stuffed garnet structure without placing ions

such as Al or Ga on the Li sub-lattice. With this in mind, we

prepared Li6.75La3Zr1.75Ta0.25O12 where La occupies A sites, Zr

and Ta occupy C sites and Li occupies B (Li1) sites as well as

distorted octahedral sites (Li2) that are unoccupied in conventional garnet. We also prepared Li6.15La3Zr1.75Ta0.25Al0.2O12 and

Li6.15La3Zr1.75Ta0.25Ga0.2O12 in order to determine how substitutions of Li by Ga or Al would affect the Li-ion conductivity and to

shed light on the role of Ga or Al as either a substituting atom on

the Li sub-lattice which reduces the Li content or as a sintering aid.

Based on crystal chemistry, we expect Al to have a site preference

for the tetrahedral site (B, Li1) and Ga to prefer the distorted octahedral site (Li2) [21]. We therefore expected that there may be some

benefit to using Ga over Al since it may not preferentially occupy

the critical Li1 junction.

The X-ray diffraction patterns for Li6.75La3Zr1.75Ta0.25O12,

Li6.15La3Zr1.75Ta0.25Al0.2O12 and Li6.15La3Zr1.75Ta0.25Ga0.2O12 after

hot-pressing are shown in Fig. 1, confirming a predominantly single

phase of cubic symmetry with space group Ia3 ̄ d. From Fig. 1, it is

observed that Ta does indeed stabilize the cubic LLZO phase at room

temperature. The lattice parameters are shown in Table 1. By comparison, Logéat et al. [14] reports a lattice parameter of 12.96A ̊ for

Li6.75La3Zr1.75Ta0.25O12 which is in exact agreement with our result.

Li et al.[13] reported a lattice parameter of 12.89A ̊ for the compositionLi6La3ZrTaO12. The latticeparameter change withincreasing Ta

contentis consistent withthe relative ionic radii of Ta5+ (64 pm) and

Zr4+ (72 pm)[22]. Turning back to our samples, Ga orAl substitution

on the Li sub-lattice has little effect on the lattice constant.

A representative micrograph of the fracture surface of the

hot-pressed Li6.75La3Zr1.75Ta0.25O12 sample is shown in Fig. 2.

From SEM analysis, a couple of important points are noted. First,

the material is very dense in agreement with the relative density ∼96%, determined from the physical dimensions, weight and
74 J.L. Allen et al. / Journal of Power Sources 206 (2012) 315–319 317

Table 1

Density, conductivity and lattice constants of samples prepared for this study.

Composition Lattice constant (Å) Relative density (%) Total Li-ion  at 298K (mS cm−1) Activation energy (eV)

Li6.75La3Zr1.75Ta0.25O12 12.96 96 0.87 0.22

Li6.15La3Zr1.75Ta0.25Al0.2O12 12.95 98 0.37 0.30

Li6.15La3Zr1.75Ta0.25Ga0.2O12 12.95 98 0.41 0.27

the theoretical density. Almost no porosity is observable. A high

relative density is extremely importantfor device applications. Second, no second phases were observed at grain boundaries. Third,

the fracture surface is very flat indicating transgranular fracture,

revealing high grain boundary strength which should lead to low

inter-granular ionic resistance. The Li6.15La3Zr1.75Ta0.25Al0.2O12

and Li6.15La3Zr1.75Ta0.25Ga0.2O12 samples also have near theoretical density (see Table 1) and exhibit microstructures similar to

that shown in Fig. 2 in that very little porosity is observed and the

fracture surface is very flat. In contrast the microstructure of the

only other Ta-substituted LLZO material, Li6La3ZrTa012 [13], consolidated by conventional sintering revealed a higher percentage

of intergranular fracture, which reveals that the grain boundaries

are not as strong as for the present study. This should lead to a

higher inter-granular resistance compared to the hot-pressed samples used in the present study. In addition, the microstructure

of cubic LLZO without Ta consolidated by conventional sintering

[9,10,23,24] tends to reveal a higher percentage of intergranular fracture when compared to the present hot-pressed materials

which should also result in a higher inter-granular resistance.

In a device, the total Li-ion conductivity rather than bulk Liion conductivity is the important measure for practical application

and so we will focus on total Li-ion conductivity. Furthermore,

our samples because of their desirable, high density and negligible

grain-boundary impedance have a total ionic conductivity nearly

equal to their lattice conductivity. This point is illustrated by the

room temperature Nyquist plots of all 3 samples (Fig. 3) which

revealed only a nearly straight line at a slope of ∼45◦ without any

clear semi-circles. Since we used Au Li-ion blocking electrodes, the

shape of the curve represents a material which is a pure Li-ion conductor. The lack of clear semi-circles at the highest frequencies is

Fig. 2. Representative SEM image of a fracture surface of hot-pressed

Li6.75La3Zr1.75Ta0.25O12.

characteristic of a material with low grain boundary impedance.

This is in agreement with the results of Kotobuki et al. [9] who suggested that the lack of clear semicircles at the highest frequencies

in their impedance plots implied that grain boundary impedance

is negligible compared to lattice impedance and Kumazaki et al.

[8] who suggested that the lack of a clear semicircle in the highest conductivity sample reveals almost complete removal of the

sample grain boundary resistance.

Fig. 4 shows the Arrhenius plot of the total Li-ion conductivity of Li6.75La3Zr1.75Ta0.25O12, Li6.15La3Zr1.75Ta0.25Al0.2O12 and

Li6.15La3Zr1.75Ta0.25Ga0.2O12 as a function of temperature. The

activation energy was estimated from the slope of the line

in the temperature range of 298–373K. From Fig. 4, several

important points can be made. First, the total Li-ion conductivity of Li6.75La3Zr1.75Ta0.25O12 is highest at all temperatures.

Second, the activation energy (0.22 eV) for the total Li-ion conductivity of Li6.75La3Zr1.75Ta0.25O12 is considerably lower than

the other two samples (0.30 eV for Li6.15La3Zr1.75Ta0.25Al0.2O12

and 0.27 eV for Li6.15La3Zr1.75Ta0.25Ga0.2O12). It is also lower than

the other reported Ta-substituted LLZO garnet fast lithium-ion

conductors Li6La3ZrTaO12 (0.42 eV) [13] and lower than that

reported for (Al-substituted) Li7−xLa3Zr2O12 (0.30 eV) [2] and

Li6.75La3Zr1.75Nb0.25O12 (0.31 eV) [15]. Third, the conductivity of

the Ga, Ta substituted sample is roughly the same as the Al, Ta

substituted sample. Fourth, the activation energy of the Ga, Ta substituted sample is slightly lower than that of the Al, Ta substituted

sample. Understanding the slight difference in activation energy

between Ga, Ta-substituted LLZO and Al, Ta substituted LLZO is difficult owing to the small difference in activation energy. First, the

two materials have very similar or equal densities, lattice constants,

grain size, charge of the substituting atoms (Ta5+ and Al3+ or Ga3+)

Fig. 3. Nyquist plots of the ac impedance spectra for hot-pressed samples of

Li6.75La3Zr1.75Ta0.25O12, Li6.15La3Zr1.75Ta0.25Al0.2O12 and Li6.15La3Zr1.75Ta0.25Ga0.2O12

with Li-blocking Au electrodes at 298K in air.
318 J.L. Allen et al. / Journal of Power Sources 206 (2012) 315–319 75

Fig. 4. Conductivity of Li6.75La3Zr1.75Ta0.25O12, Li6.15La3Zr1.75Ta0.25Al0.2O12 and

Li6.15La3Zr1.75Ta0.25Ga0.2O12 as a function of temperature.

and Li content. The only difference is the size of Al3+ relative to Ga3+

which in turn leads to different site preferences in garnets [21]. The

trend is in agreement with the supposition that Ga would prefer to

occupy the distorted octahedral Li2 site and Al the tetrahedral Li1

site, as discussed earlier. However, the small differences in conductivities and activation energies imply that most likely Al and Ga

occupy both sites, though perhaps with a higher concentration of

Ga on Li2 relative to Ga on Li1 and a higher concentration of Al on

Li1 relative to Al on Li2.

Table 1 summarizes the data for all samples. Since our samples are of similar, high density (∼96–98%) a good relative

comparison of the effect of substitutions is possible. A couple

important points can be made. First, as discussed for Fig. 4, the Ta

substitution for Zr without Al or Ga leads to the highest conductivity (0.87 mS cm−1 at 298K) and a very low activation energy

(0.22 eV). Second, the results in Table 1 provide confirmation

of the hypothesis that substitution within the Li sub-lattice by

Al or Ga would be detrimental owing to its hindrance of the

Li-ion conduction path. As discussed by Ohta et al. [15], the Li-ion

conductivity is a function of the Li-ion concentration and the Li-ion

mobility. The calculation of the Li-ion concentration difference

between Li6.75La3Zr1.75Al0.2O12 and Li6.15La3Zr1.75Ta0.25Al0.2O12

and Li6.15La3Zr1.75Ta0.25Ga0.2O12 analyzed via the method

of Ohta et al. [15] shows that the Li concentrations of

Li6.15La3Zr1.75Ta0.25Al0.2O12 and Li6.15La3Zr1.75Ta0.25Ga0.2O12

are 90% of Li6.75La3Zr1.75Al0.2O12. This 10% difference in Li-ion

concentration cannot explain the >2-fold increase in conductivity.

Therefore, the Li-ion mobility difference is the predominant

factor which controls the conductivity difference. This is further

supported by the lower activation energy of Li6.75La3Zr1.75Al0.2O12

(0.22 eV) relative to Li6.15La3Zr1.75Ta0.25Al0.2O12 (0.30 eV) and

Li6.15La3Zr1.75Ta0.25Ga0.2O12 (0.27 eV). Al or Ga on the Li sublattice hinders the Li-ion mobility. Third, at least for hot pressing,

Al or Ga substitution is not needed in order to obtain a sample

with high density and fast Li-ion conductivity.

By way of comparison, Li et al. [13] reported a total Li-ion

conductivity at 298K of 0.18 mS cm−1 for the Ta-substituted LLZO

garnet of composition Li6La3ZrTaO12. Al was also found in the sample at a weight percent of 1.3 which is about 0.4 M or about twice as

much as our Al-containing sample. This is the only other reported

conductivity for Ta-substituted LLZO however the comparison is

notideal since we do not know the density ofthe Li et al. sample and

the Ta content and Al contentis different. In contrast, our Ta, Al substituted sample (Li6.15La3Zr1.75Ta0.25Al0.2O12) had a conductivity of

∼0.37 mS cm−1. Qualitatively, we can say that first, the lower conductivity of the Li et al. [13] sample may result from lower sample

density and the higher Al (reduced Li-ion mobility) and Ta content

(reduced Li-ion concentration). Second,the activation energy ofthe

Li et al. sample is considerably higher (0.42 eV) which is probably a

result of higher grain boundary resistance as well as some blockage

of the Li sub-lattice by the higher content of Al.

Also, as comparison, the total Li-ion conductivity and activation energy of our sample of composition Li6.75La3Zr1.75Ta0.25O12

(0.87 mS cm−1 and 0.22 eV) can be compared to the Nb-substituted

LLZO of composition, Li6.75La3Zr1.75Nb0.25O12 (0.8 mS cm−1 and

0.31 eV) [15]. The conductivity values are in excellent agreement. The activation energy of our sample is lower as a result of

the higher relative density (96%) obtained via hot-pressing relative to sintering (89–92%). These values approach the calculated

(via molecular dynamics simulation) bulk lattice conductivity of

1 mS cm−1 reported by Adams et al. [25].

Turning once again to our samples, these results support the

supposition that high ionic conductivity can be achieved by Ta substitution for Zr cubic garnet stabilizing substitutions that do not

place the substituting atom (Al or Ga) in the conductive pathway

of Li. Both Nb [15] and Ta (this work) substituted LLZO exhibited

higher conductivity than Al-substituted LLZO [2]. Ta is preferred

owing to its higher stability against Li reduction.

4. Conclusion

Cubic LLZO can be stabilized at room temperature by

the addition of Ta. Hot-pressed Li6.75La3Zr1.75Ta0.25O12,

Li6.15La3Zr1.75Ta0.25Al0.2O12 and Li6.15La3Zr1.75Ta0.25Ga0.2O12

have a relative density of ∼96–98%. The room temperature AC

conductivity results for the hot-pressed Li6.75La3Zr1.75Ta0.25O12

sample yields a total ionic conductivity value ∼0.87 mS cm−1 and

activation energy of 0.22 eV. Addition of Al or Ga to Ta-substituted

LLZO reduces the conductivity and increased the activation energy

suggesting that Al or Ga substitutions in the Li-ion sub-lattice tend

to lower the Li-ion conductivity.

Acknowledgments

J.L.A. and J.W. would like to acknowledge support of the US

Army Research Laboratory (ARL). E.R. and J.S. authors would like

to acknowledge the support of the US Army Research Office (ARO).

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77

Research@

Shock-Induced Behavior of Cubic Gauche 

Polymeric Nitrogen

William D. Mattson and Radhakrishnan Balu

U.S. Army Research Laboratory, Aberdeen Proving Ground, MD

ABSTRACT

Quantum-mechanical calculations based on density functional theory are 

used to study the shock response of the polymeric cubic gauche phase of 

nitrogen (cg-N), proposed as an alternative energetic ingredient to those 

used in conventional explosive formulations. The shocked polymeric nitrogen 

undergoes multiple complex phase transformations and spontaneously forms 

defects. The occurrence of these dynamic phenomena absorbs the shock 

energy which subsequently slows the compression wave. Additionally, no 

reaction occurs immediately behind the shock front; rather reactions result 

from the unraveling of the material at the free edge of the filament opposite to 

shock propagation. As the material unravels, numerous polyatomic transients 

are formed, including five-membered rings and polymeric chains, which 

subsequently undergo secondary reactions to form the final diatomic products. 

The speed at which these reactions propagate through the material is much 

slower than the sound speed, and combined with the slowing compression 

wave, indicates that the material may not detonate under these conditions.

Physical Review B, Volume 83, Page 174105 (2011)
78

PHYSICAL REVIEW B 83, 174105 (2011)

Shock-induced behavior of cubic gauche polymeric nitrogen

William D. Mattson and Radhakrishnan Balu

U.S. Army Research Laboratory, RDRL-WML-B (Bldg. 4600), Aberdeen Proving Ground, Maryland 21005-5069, USA

(Received 26 January 2011; revised manuscript received 8 April 2011; published 11 May 2011)

Quantum-mechanical calculations based on density functional theory are used to study the shock response of

the polymeric cubic gauche phase of nitrogen (cg-N), proposed as an alternative energetic ingredient to those

used in conventional explosive formulations. The shocked polymeric nitrogen undergoes multiple complex phase

transformations and spontaneously forms defects. The occurrence of these dynamic phenomena absorbs the

shock energy which subsequently slows the compression wave. Additionally, no reaction occurs immediately

behind the shock front; rather reactions result from the unraveling of the material at the free edge of the filament

opposite to shock propagation. As the material unravels, numerous polyatomic transients are formed, including

five-membered rings and polymeric chains, which subsequently undergo secondary reactions to form the final

diatomic products. The speed at which these reactions propagate through the material is much slower than the

sound speed, and combined with the slowing compression wave, indicates that the material may not detonate

under these conditions.

DOI: 10.1103/PhysRevB.83.174105 PACS number(s): 62.50.Ef, 82.40.Fp, 61.50.Ah, 61.72.Cc

I. INTRODUCTION

Pressure-induced transitions of molecular solids to polymeric forms have been the subject of intense activity for

several decades. One of the more widely studied systems both

theoretically1–24 and experimentally25–40 is nitrogen, which

in its ambient state is a homonuclear diatomic molecule

with one of the strongest bonds in nature, making it an

ideal candidate for energy storage if a stable high-density

polymeric form can be realized. Theoretical predictions have

indicated a variety of possible molecular and polymeric phases

over a range of pressures and temperatures; however, until

recently, only the diatomic molecular high-pressure phases

were observed experimentally within the predicted pressure

and temperature regimes. Phase transitions to the more exotic

polymeric31,33,34 and nondiatomic molecular phases36 have

been observed well above the predicted transition pressure

range for the diatomic to cubic gauche (cg-N) transition.5 This

cg-N form of nitrogen is a bulk three-dimensional covalent

network, similar to diamond, in which every atom has three

neighbors. This structure was observed in diamond-anvil-cell

experiments39 in which nitrogen is compressed to 180 GPa

and heated to 2200 K. The sample was only partially cg-N

and included amorphous polymeric nitrogen. The material

was not recoverable at ambient conditions, although phonon

calculations indicate that the cg-N structure should be stable

at ambient pressures.41 However, all theoretical calculations

have assumed an ideal crystal structure, which is not attainable under the extreme experimental conditions used in

its synthesis, and thus the presence of these other phases

most likely introduces instabilities preventing the recovery

of the cg-N at ambient conditions. Continued theoretical and

experimental explorations to understand the principle physical

and chemical laws which govern bonding, phase behavior, and

thermodynamic stability of high-pressure polymeric phases

are necessary for developing the technologies required to

recover polymeric forms of nitrogen at the ambient state.

These new types of materials are not only of academic

interest, as the large energy densities of some of the various

polymeric forms make them promising high-performing alternatives to currently fielded conventional energetic materials

(EM). The calculated energy of complete conversion of cg-N

to the ground state diatomic nitrogen ranges from 0.75 to

1.55 eV per atom,12 several times that of conventional

explosives.42 However, before cg-N can be considered for

use in a practical EM application, its stability and dynamic

response under a variety of conditions must be well understood. This study is the first in a series in which such factors

are explored, where we present quantum molecular-dynamics

(QMD) results of shocked cg-N initially at low temperature

and ambient pressure.

Finally, one of the primary goals of this paper is to determine

whether shock impact will initiate a self-sustained reaction

wave within the defect-free cg-N material. In addition to

presenting results demonstrating the dynamic response of

shocked cg-N, this paper will briefly describe the theoretical

methods, the computational details, software explored to

enable adequate treatment of the calculations, and results of

the calculations to date.

II. DETAILS OF THE CALCUATION

The first-principles calculations reported herein use the

Kohn-Sham density functional theory (DFT) with the generalized gradient approximation (GGA) exchange-correlation

functional of Perdew, Burke, and Ernzerhof (PBE).43 QMD

simulations were performed using the Gaussian and planewave-method code QUICKSTEP as implemented in CP2K

(Ref. 44) and the double-zeta with valence polarization basis.

Total energies and forces were converged to 2.6 × 10−5 and

1.6×10−4 a.u., respectively, for all calculations. The equations

of motion are from the Born-Oppenheimer approximation with

a time step of 1 fs for all simulations.

The direct shock-wave simulations here are implemented

with a long thin filament45 composed of blocks of eight-atom

simple cubic cells. Each block is constructed using two

instances of the primitive cell of cg-N (space group I2135)

with one centered at the origin (0,0,0) and the other at the

body-centered site (1/2,1/2,1/2). The blocks are used to create

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WILLIAM D. MATTSON AND RADHAKRISHNAN BALU PHYSICAL REVIEW B 83, 174105 (2011)

FIG. 1. A schematic of the filament used in the QMD simulations

of shock-impacted cg-N. The filament can be described as a series of

blocks composed of (1 × 3 × 3) unit cells of cg-N. The block at the

far right edge of the filament (outlined in black) denotes the material

that is fixed during the simulation; to its left is the “test” block, used

to determine shock-wave approach (see text). The eight blocks at the

far left edge of filament contain the flyer plate atoms.

a “slab” which is defined to be a (1 × 3 × 3) supercell of

blocks with dimensions of (3.8,11.3,11.3) A. The filament  ̊

is then composed of 32 slabs of cg-N. The simulation cell,

shown in Fig. 1, contains 2304 atoms. In order to generate

an initial configuration for use in the shock simulation,

the 0 K configuration was first obtained through geometry

optimization of the filament. Next, the optimized simulation

cell was first subjected to an isothermal isochoric (NVT) MD

simulation at either T = 200 or 300 K for 5 ps. The final

configuration of the NVT-MD simulation was then used in an

isothermal-isobaric (NPT) MD simulation at either T = 200 or

300 K, P = 1 bar for 5 ps to produce the filaments that would

be used in the shock simulations. In all of the aforementioned

simulations, periodic boundary conditions were imposed in all

directions.

For the shock-wave simulations, computational schemes for

reducing the costs of MD simulations of shocked crystals were

employed; specifically, that outlined by Rice.46 In this scheme,

a shock wave is initiated in the filament composed of quiescent

material [denoted as cg-N(V) and shown in Fig. 1] through

flyer plate impact. Periodic boundary conditions are imposed

in the two directions perpendicular to the direction of the shock

impact and propagation (in this case, the x direction). For

these simulations, the shock will be introduced at the far-left

side of the filament by a “flyer plate” composed of cg-N. These

correspond to the eight leftmost slabs shown in Fig. 1. The edge

of the simulation cell opposite to shock impact is bounded by a

slab of constrained atoms (located at the right edge of Fig. 1).

Within this scheme, as the shock wave approaches the rigid

layer, additional undisturbed material is inserted before the

rigid layer to allow continuance of the shock wave. The shock

wave can be initiated through instantaneously increasing the x

component of the velocity of each atom within the flyer plate by

an additional 10 km/s. Following Rice et al., a (1 × 3 × 3) test

slab was monitored for shock-wave approach and is located

one slab length to the left of the rigid slab of atoms in the

filament (see Fig. 1). The center-of-mass displacement within

the “test slab” is evaluated at each integration step and when

this value exceeded that of thermally equilibrated bulk cg-N by

0.1 A, material composed of thermally equilibrated bulk cg-N  ̊

is inserted between the rightmost slab of moving atoms and

the rigid layer. The filament has a “free edge” at the leftmost

region (the flyer plate) with vacuum to the left. We confirmed

that the material with the free edge was stable at both T =

200 or 300 K using NVT-MD simulations integrated for 15 ps.

In these simulations, the initial configuration of the free-edge

filament was that of the final configuration of the heretofore

described NPT-MD simulation in which periodic boundary

conditions were imposed in all three dimensions. During these

NVT-MD simulations the atoms at the free edge exhibited

larger oscillations about their lattice sites than those in the

interior of the filament; however, no reaction was observed.

Also, relaxation of the atoms near the free surface occurred,

resulting in a slight decrease in lattice spacing relative to that

within the interior of the filament.

A potential source of error is the small cross-sectional

area (11.3 × 11.3 A) normal to the direction of the shock  ̊

front. However, previous quantum molecular- dynamics shock

simulations have successfully used cells with similarly sized

cross-sectional areas.47,48 These calculations were performed

in conjunction with the multiscale shock technique (MSST)

for systems of either 64 water molecules or 360 carbon

atoms. In both studies, exploration of the effect of system

size showed that the simulation cells were sufficiently large to

reach converged results. Also, QMD calculations of the shock

Hugoniot of nitrogen over a large range of temperatures and

densities using systems composed of 32, 54, and 64 atoms

showed that of the measured properties, only the Gruneisen

parameter showed a dependence on size, and only for the

region of the shock Hugoniot where the smaller system

predicted a negative value.49 Since our simulations have

substantially larger cross-sectional areas than in Ref. 49, we

do not expect the results to be affected by our choice of system

size.

Another potential source of error is that these simulations

were performed using Born-Oppenheimer QMD and therefore

the electrons in the system are treated adiabatically. Shock

simulations on deuterium50 and water47 has shown that treating

the electrons nonadiabatically within Car-Parrinello QMD

(CPMD) had little effect on simulations with shock velocities

less than 15 and 12 km/s, respectively. Since the electronion coupling constant for nitrogen is smaller than that for

deuterium and hydrogen, it is expected that shock simulations

of nitrogen will be insensitive to the nonadiabatic deviations

of CPMD at the shock velocity investigated in this study

(10 km/s). While the nonadiabatic excitation reflected by

CPMD might not adequately represent that in a shock event, the

insensitivity of the simulations to CPMD excitations for shock

velocities of 10 km/s implies that the Born-Oppenheimer

approximation is reasonable for the simulations presented

here. The shock velocity for the simulation was chosen to

impart sufficient energy that might lead to reaction, but was

not likely to induce nonadiabatic effects.

To illustrate the evolution of the material under shock, temporal profiles of various properties within the filament will be

described; this is accomplished by taking a weighted moving

average of the properties in local regions of the material along

the direction of the shock. For each plane normal to the x

xis on intervals of 0.5 A, the weighted moving average is  ̊

calculated by summing the weighted value of the property

of each atom within 4 A of the plane under consideration.  ̊

The weight assigned for each atom is proportional to its

distance from the plane. For example, at 4 A the weight is  ̊

0, and at 0 A from the plane the weight is 0.25. This scheme  ̊

produces temporally averaged values that make the relevant

aspects of the simulation clear, and maintains correspondence

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of the property depicted in the temporal profiles to that of the

material property. Finally, in this work, chemical bonds are

depicted between atom pairs with interatomic separations that

are less than 1.6 A. This value was selected after observing the  ̊

temporal behavior of such bonds in the shocked lattice and in

various vibrationally hot product species.

III. RESULTS

The shock-wave simulations for filaments at 200 and

300 K were performed within the microcanonical ensemble

(NVE-MD) using the final configurations and atomic velocities

from the NVT-MD simulations of the free-edge filament (with

flyer plate atoms given the appropriate impact velocity as

described heretofore). As the shock wave propagated through

the material, its effects reached the test slab after 480 fs, at

which time the simulations were stopped. The features of the

shocked material were similar for both 200 and 300 K filaments

up to this time, including phase transitions and spontaneous

formation of defects, as will be discussed hereafter. Thus,

further investigation is limited to the 300 K filament. For the

300 K filament simulation, an additional 25 slabs (containing

1800 atoms) were inserted immediately to the left of the

rightmost rigid slab to allow the continuation of the shock

simulation, increasing the total number of atoms in the filament

to 4104. As shown in earlier classical molecular-dynamics

studies of shocked filaments, such a scheme allows for

simulation of only the area of immediate interest (i.e., the

shocked region), with the idea that simulating infinitely large

material well ahead of the shock discontinuity provides no

extra information on the shock phenomena.46 All results

discussed hereafter refer to the simulation using the 300 K

filament.

A. Complex phase transitions and defects

A series of snapshots of atomic positions at various times

during the simulation (Fig. 2) illustrates the very complex

behavior of shocked cg-N, with the formation of numerous

phases as well as interesting reaction intermediates displayed.

Figure 2(a) shows that the original honeycomb structure of

the cg-N collapses to a herringbonelike lattice as the flyer

plate drives the material forward. This phase [denoted HB(l)]

continues to grow as the simulation progresses to a maximum

width of ∼50 A and propagates through the material in the  ̊

direction of the initial shock at a gradually diminishing velocity

as will be discussed in greater detail in Sec. III C. In this

phase transition, no bonds are broken or formed. At 400 fs

[Fig. 2(c)] additional complexity in the form of a region of

compressed cg-N behind the leading phonon but ahead of the

HB(l) phase emerges. This region will be denoted hereafter

as the leading compression zone (LCZ). The LCZ continues

to develop behind the shock front for the remainder of the

trajectory. Also in Fig. 2(c), the appearance of a spontaneously

formed stacking fault is apparent between the HB(l) phase

(denoted as D1 and enclosed by a box) and the LCZ. In

Fig. 2(d), the material immediately to the right of D1 deforms

to a HB lattice configuration that corresponds to a shifting of

the HB(l) crystal along the vertical axis by half of the original

lattice vector [denoted hereafter as HB(r)]. The formation

of this stacking fault was unexpected and has significant

consequences, as will be discussed below. Finally, as the HB(l)

phase relaxes and expands on the left, the material reforms to

the cg-N structure. Note that the reformed cg-N phase has a

different crystalline orientation than the initial structure. The

undisturbed crystal ahead of the leading phonon has the long

axis of the large honeycomblike rings oriented vertically, but

the reformed region in Fig. 2(f) has the long axis of the rings

oriented horizontally. Therefore, the material in the reformed

cg-N phase will be denoted as cg-N(H) hereafter, whereas the

material ahead of the shock front will be denoted as cg-N(V).

The nearest neighbors of the nitrogen atoms in the cg-N(H)

region are the same as that of cg-N(V), as is the handedness

of the crystal.

More complex structures evolve in this shocked material as

the trajectory continues. Specifically, at 690 fs, a new “defect”

FIG. 2. (Color online) Snapshots of the

cg-N system at (a) t = 90 s, (b) 160 fs, (c)

400 fs, (d) 690 fs, (e) 830 fs, and (f)1040 fs.

For clarity, the block of stationary atoms at

the far-right edge of the simulation cell are

not shown; also, quiescent crystal is added

to the snapshots in (a)–(c) to have the same

x dimension as those in (d)–(f). See text for

explanation of the larger simulation cells used

in (d)–(f).

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appears (denoted D2), which is bounded by the cg-N(H)

structure on the left and a reformed cg-N(V) on the right. To the

right of these are the HB phases (which also have in their midst

D1, discussed above) and the LCZ. In D2, bonds are broken

and formed, unlike the material associated with D1. At 830 fs

[Fig. 2(e)], D1 has widened and by 1040 fs has transitioned

into the cg-N(H) phase [see the region enclosed by the box

in Fig. 2(f)]. The width of the HB(l) phase decreases during

the transition of D1 to the cg-N(H) phase, a result both from

back transition to cg-N(V) at its left boundary and from back

transition of D1 to cg-N(H) at its right boundary. The HB(r)

phase to the right of D1 is maintained as the LCZ transitions

to HB(r); the LCZ material to its right continues to widen.

B. Chemical reaction

Reactions in the shocked material are equally complex

and interesting. Initially, as seen in the snapshot at 400 fs

[Fig. 2(c)], most of the products ejected from the free cg-N

surface at the left are diatomic nitrogen, with a few short-lived

N3 molecules. As soon as two N3 molecules approach each

other, they rapidly react to form three N2 molecules. By 690 fs

[Fig. 2(d)], chains with one end connected to the leftmost

free surface start to form, break free at the surface, and

subsequently decompose. As chains continue to form, their

lengths increase and disintegration occurs within the chains

rather than detachment at the surface. The lengthening of these

chains indicates that surface disintegration of the cg-N lattice

occurs more quickly than formation of the more stable N2

product. By 1040 fs [Fig. 2(f)], large fragments of crystal have

disintegrated into chains. This process creates a reaction zone

with little to no N2 forming at the cg-N surface, but rather N2

resulting from reactions of the detached chains. Additionally,

other more complicated structures are formed including N5

rings [see, for example, the ring in Figs. 2(d)–2(f) bounded by

the red circle]. Some of the N5 rings also have chains attached

to one of the atoms in the ring. Figure 3 shows a time history

of the formation and progression of the N5 ring highlighted

in the snapshots in Fig. 2. This five-membered ring within the

red circle has a lifetime of in excess of 450 fs as evidenced by

the temporal history of the bonds labeled on the moiety shown

in Fig. 4. Up to 770 fs, the pentazole ring has a six-membered

nitrogen chain attached to it at which point the nitrogen

chain subsequently decomposes into three nitrogen molecules.

The chain decomposes first by simultaneously forming two

nitrogen molecules from the four terminal nitrogen atoms

of the chain, followed by the dissociation of the remaining

two nitrogen atoms from the pentazole ring to form the final

nitrogen molecule. By 820 fs, the three nitrogen molecules

oscillate about the equilibrium gas phase value (1.1 A),  ̊

while the bonds corresponding to the pentazole ring oscillate

about larger values (∼1.3–1.4 A) for the remainder of the  ̊

trajectory integration. Pentazole structures have been shown to

be stable51 and similar structures have recently been observed

in simulations calculating the high-temperature phase diagram

of nitrogen.52,53 Interestingly, both Boates et al.

52 and Donadio

et al.53 have observed a liquid-liquid phase transition where

the liquid molecular nitrogen transforms into nitrogen chains

and five-membered rings, similar to the reaction intermediates

shown in the simulations herein.

FIG. 3. Snapshots of one of the longer-lived polyatomic transients

formed in the reaction zone of the material. Times during the

simulation are denoted in the figure.

C. Spatial-temporal profiles of material properties

Spatial-temporal profiles of the velocity of particles of

the material in the supercell along the direction of shock

propagation are given in Fig. 5. In the figure the length of the

supercell in the direction of shock propagation corresponds

to the x axis and the simulation time corresponds to the y

axis. The red triangular region in the lower left-hand area

of the figure corresponds to the initial flyer plate. Partitions

corresponding to a portion of the vacuum region to the left

of the initial flyer plate position and corresponding to the

fixed slab (far right of the simulation cell) are not shown.

The phonons emanating in front of the shock discontinuity are

clearly apparent in Fig. 5, with an initial phonon generated by

the shock impact (∼150 fs) and progressing ahead of a series

of evolving shock fronts. We define a shock front as a nearly

discontinuous change in material state during the simulation.

Additionally, density profiles along the filament at later times

indicate a series of subsequent phonons emanating ahead of

the leading shock front. By the end of the simulation, five

very distinct phonons have appeared and are traveling at a

constant speed of 16.5 ±0.5 km/s. The compressed region

behind the phonons, however, is clearly slowing down. At the

initial stages of its formation, the velocity of the leading edge

of the HB(l) phase aligns with the initial phonon. Interestingly,

at 40 A,  ̊ ∼150 fs, a second phonon emanates, with an almost

discontinuous decrease in velocity of the shock front. A similar

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FIG. 4. (Color online) Time trace of the bond lengths of the bonds

shown in Fig. 3. The bonds are numbered in the inset moiety which

is a copy of the 680 fs snapshot in Fig. 3.

manifestation of this behavior is apparent at 90 A, 550 fs, in  ̊

which a third phonon appears with a corresponding decrease

in velocity of the shock front. Beyond that, two additional

phonons appear and the shock front continues to slow.

Figure 6 shows the instantaneous temperature profile of

the material in the supercell along the direction of shock

propagation. The local temperatures within the material were

calculated by first subtracting the weighted center-of-mass

velocity of all atoms in the local region around each plane from

the individual atomic velocities in the local region, and then

used to calculate an instantaneous local temperature. There is

significant heterogeneity within the material phases as shown

in Fig. 6. Of most importance is the region where the initial

spontaneously formed stacking fault (D1) resides. It is notable

that the temperature of the area around the fault remains lower

than the material on either side of it. The stacking fault, which

first appears at ∼400 fs, corresponds to the point in Fig. 6 at

FIG. 5. Temporal velocity profiles. Phonons emanate from the

shock front and propagate at a velocity of 16.5 km/s away from the

shock front. The shock front (HB transition front) itself slows down.

Legends for the color range are given in the figure for the range of 0

to 10.1 km/s.

FIG. 6. Temporal instantaneous temperature profiles. The temperature of the defect is lower than the surrounding compressed

region. Legends for the color range are given in the figure for the

range of 0 to 9000 K.

which the velocities of the second phonon, and the following

compressed region, separate. The region corresponding to D1

remains lower in temperature than the material ahead or behind

it all throughout the simulation, indicating that the stacking

fault acts as an energy sink.

IV. DISCUSSION

It was expected that cg-N, like all other forms of polymeric

nitrogen, would be an advanced, new type of energetic material

due to its stored structural energy. The transition to gaseous

nitrogen would release 1.55 eV/atom,12 a substantially higher

energy release than that of conventional energetic materials.

Unfortunately, the results of these simulations indicate that the

mechanical properties of cg-N in its purest, defect-free, form

enable it to disperse shock energy through extremely complex

phase transitions such that initiation to a self-sustained

exothermic shock-driven reaction is improbable. Most likely,

the reaction within the rarefaction region could continue until

the material is consumed but at a rate that would preclude cg-N

from consideration as an advanced high explosive.

It is worth noting at this point that Chen et al.

19 performed

first-principles calculations of mechanical properties and

failure modes for cg-N, and concluded that this material

has “exceptional mechanical properties.” They evaluated the

failure modes for cg-N using the method of Roundy et al.

54

in which the lattice vectors are incrementally deformed in the

direction of the applied stress and the atomic coordinates are

fully relaxed between increments. In this study they found that

the dominant mode of lattice failure is shear, a condition not

imposed in the shock simulations contained in this effort. The

size of the simulation cell used in this study is too small to allow

the formation of shock-induced shear bands and therefore, the

role of shear on dynamic response cannot be evaluated using

these results.

An equally important consideration is the role of defects on

the shock response of cg-N. Synthesized cg-N will include defects, which could influence the shock behavior, as evidenced

by the role spontaneously formed defects play in the dynamic

response of the pure material discussed above. Although it is

well beyond the scope of the current study, a thorough QMD

investigation of a shocked defected material could identify

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WILLIAM D. MATTSON AND RADHAKRISHNAN BALU PHYSICAL REVIEW B 83, 174105 (2011)

fundamental mechanisms that contribute to the physical

response of the material, including plasticity, fragmentation,

defect-nucleated reactions, and melting. In light of the possible

roles of shear and/or defects in the shocked material on reactivity and energy release, it would be premature to conclude

that cg-N would not be useful as an energetic material.

V. SUMMARY

Quantum MD simulations have been used to provide an

atomic-level description of the material response of the defectfree cg-N polymeric form of solid nitrogen to shock. The

cg-N solid showed extremely complex shock-wave structure

involving several structural phase changes and spontaneous

formation of material defects behind the shock front that

absorbed and dispersed the shock energy. Reaction did not

occur directly behind the shock front, but rather at the free

edge of the filament opposite to the direction of shock propagation, unraveling into a variety of polyatomic and polymeric

transients, including chains and five-member rings. These

transients subsequently decomposed or underwent collision

with other transients to form the final diatomic product gas. The

heat released in the reactions was localized in this dense-gas

region behind the disintegrating filament and not transferred

to the remaining intact filament. Although the simulations

suggest that the pure form of cg-N might not be a good

candidate for use as an explosive, its unique energy dispersion

properties due to the complex phase transitions initiated by

shock could be exploited for other applications in which such

a response would be advantageous.

ACKNOWLEDGMENTS

Calculations were performed using the DOD Supercomputing Resource Centers (DSRCs) located at the U.S. Army

Research Laboratory and the Air Force Research Laboratory

under the Computing Challenge Project C2L awarded by the

DOD High Performance Computing Modernization Program.

R.B. acknowledges support from the Research Participation

Program for the U.S. Army Research Laboratory administered

by the Oak Ridge Institute for Science and Education. The

authors wish to thank Betsy Rice for helpful comments and

discussions.

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85

Research@

Optical Cell for In Situ Vibrational 

Spectroscopic Measurements 

at High Pressures and Shear

Jennifer A. Ciezak1 and Timothy A. Jenkins2

1 U.S. Army Research Laboratory, Aberdeen Proving Ground, MD

2 National Institute for Standards and Technology, Gaithersburg, MD

ABSTRACT

An optical cell is described for performing simultaneous static high-pressure 

and shear experiments. This cell design is a modification of the previously 

designed megabar diamond anvil cell used by Mao and Bell that allows for 

controlled, remote shear. With this diamond anvil cell, it is possible to use 

a wide range of existing experimental techniques and pressure media. The 

cell was validated on a sample of calcite at 5 kbar. Raman measurements 

show the onset of the phase transformation from calcite to aragonite at 10◦

of rotation.

Review of Scientific Instruments 82, 073905 (2011)
86

REVIEW OF SCIENTIFIC INSTRUMENTS 82, 073905 (2011)

Optical cell for in situ vibrational spectroscopic measurements at high

pressures and shear

Jennifer A. Ciezak1,a) and Timothy A. Jenkins2

1US Army Research Laboratory, RDRL-WML-B (Building 390), Aberdeen Proving Grounds,

Maryland 21005, USA

2NIST Center for Neutron Research, 100 Bureau Drive, MS 6100, Gaithersburg, Maryland 20899, USA

(Received 19 May 2011; accepted 2 June 2011; published online 19 July 2011)

An optical cell is described for performing simultaneous static high-pressure and shear experiments.

This cell design is a modification of the previously designed megabar diamond anvil cell used by

Mao and Bell that allows for controlled, remote shear. With this diamond anvil cell, it is possible to

use a wide range of existing experimental techniques and pressure media. The cell was validated on

a sample of calcite at 5 kbar. Raman measurements show the onset of the phase transformation from

calcite to aragonite at 10◦ of rotation. [doi:10.1063/1.3606640]

I. INTRODUCTION

The diamond anvil cell, developed over 50 years ago1, 2

is a ubiquitous, versatile tool for the study of materials subjected to thermomechanical extremes of pressure, temperature, and stress. In recent years, a broad range of diamond

anvil cells have been designed and adapted for numerous

experimental techniques, such as laser heating, NMR, magnetic measurements, and synchrotron-based experiments, the

results of which have revealed new fundamental knowledge

within physics, chemistry, and materials science.3

While the phase behavior of materials subjected to quasihydrostatic pressure and temperature has been extensively

examined since the inception of the diamond anvil cell, it is

primarily only in the last 15 years that the effect of nonhydrostatic pressure and its resulting deviatoric stress on material

behavior have been examined. X-ray diffraction along the

radial direction of the diamond anvil cell has been previously

used to examine the elastic and rheological properties of

materials4, 5 as a function of pressure, but such techniques do

not allow for in situ control of the stresses within the cell.

Such capability is only possible with a rotational (or shear)

diamond anvil cell, which permits rotation of one diamond

anvil perpendicular to the axis of compression, while the

second anvil remains fixed. Originally designed by Bridgman

in the 1930s with tungsten carbide anvils,6 the rotation of

one anvil relative to the fixed anvil generates additional

shear stress and strain on the sample, resulting in severe

plastic deformation. The modification of the early Bridgman

design to substitute diamond anvils made exploration of

the material response to pressure and shear possible using

the numerous experimental probes available for diamond

anvil cells. Several measurements of utilizing this type of

diamond anvil cell have been reported during the past two

decades;7–21 these reports have studied phase behavior in

semiconductors,7, 8 C60 fullerene,9 sodium azide21 as well as

elements 10 and energetic materials.19

Despite the critical scientific insight provided by such experiments, rotational diamond anvil cells are not widely availa)Electronic mail: Jennifer.ciezak@us.army.mil.

able. Here we report on a rotational adaptation of the widely

used Mao–Bell megabar cell,22 which provides a robust experimental apparatus capable of generating simultaneous high

pressure and shear. The shear capabilities of the cell can be

used to pressures exceeding 50 GPa and, although not detailed

in this paper, to temperatures beyond 2500 K. Additionally,

unlike previous designs,7–21 the rotational capability can be

precisely controlled via an encoded motor. To illustrate the

room temperature capabilities of the cell, in situ Raman spectroscopic measurements of calcite at elevated pressures and

shear are presented.

II. CELL DESIGN AND USE

A cross section of the major components of the assembled rotational diamond anvil cell (RDAC) is shown in Fig. 1.

The main body of the RDAC, made of Vascomax T-250

steel,23 is an adaptation of the Mao-Bell megabar cell22 with

additional gearing for use in rotational shear experiments. The

additional gearing enables the rotation of the piston anvil relative to the fixed cylinder anvil. Using a set screw, the position

and orientation of the cylinder can be fixed to prevent concurrent rotation with the piston. A minimum clearance of 7 μm

between the piston and cylinder was necessary to maintain

anvil alignment while allowing unhindered rotation of the piston. Careful machining of the piston and cylinder is required

to allow for unhindered rotation. Polished tungsten carbide

disks with highly parallel flat faces were used as diamond

anvil mounts and fixed positionally with set screws to prevent rotation. The diamonds are then attached to the tungsten

carbide mounts with epoxy resin and aligned similar to a conventional diamond anvil cell. The preparation of the RDAC

for experimental measurements is similar to that of traditional

diamond anvil cells and the cell can be with any experimental technique adapted for use with the Mao–Bell megabar

cell.22

The lever arm and body assembly are constructed of

Vascomax T-250 steel (Ref. 23) and have dimensions comparable to previous designs. The lever arm is designed to

smoothly compress the piston and cylinder under a uniform

0034-6748/2011/82(7)/073905/4/$30.00 82, 073905-1

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87

073905-2 J. Ciezak and T. Jenkins Rev. Sci. Instrum. 82, 073905 (2011)

FIG. 1. Cutaway view of rotational DAC. A indicates the diamond anvil

cylinder, B is the diamond anvil piston, C is the cell body, D is the lever

arm, E is the thrust bearing, F is the diamonds and tungsten carbide mounts,

G is the gear assembly for the rotational cell, H is the Bellville washer and

bolt assembly, and I is threaded load block for loading the lever arms.

load. Rotation of the cylinder occurs via the additional gearing, as detailed in Fig. 2; a worm gear, with a worm gear to

cylinder ratio of 180:1, was chosen both for space constraints

and its anti-backlash capabilities. The worm gear shaft can

be easily interfaced with an encoded motor drive system to

allow for both precision control of the rotation and remote usage. Similar to the conventional Mao–Bell megabar cell,22 the

RDAC is pressurized by tightening a half inch hex bolt with

Belleville spring washers stacked in an alternating direction

with the lever arms yielding a force multiplier of 5.25. To facilitate the rotation of the piston, a thrust bearing is used in

conjunction with the thrust block. The thrust bearing is seated

in a centering mount on the thrust block and allows for removal of washers and screws that may be employed during

gas loading applications.24, 25

In the RDAC design, the maximum pressure is limited by

the thrust bearing, which, with a rotational speed of 50 rpm,

allows for upward of 50 GPa of pressure on 100 μm diameter diamond culets. As rotational speed increases, the available pressure decreases. Without the use of the thrust bearing

the cell operates similar to a conventional Mao–Bell megabar

cell,22 with the maximum pressure dictated by the culet size

of the diamond anvils.

III. IN SITU RAMAN SPECTROSCOPIC

MEASUREMENTS

To demonstrate the capabilities of the rotational diamond

anvil cell, Raman spectra were collected for polycrystalline

calcium carbonate (≥99%) (Calcite-I) using Type I diamond

anvils with 800 μm culets. A 600 μm sample well was drilled

via electron discharge machining into a rhenium gasket,

which was preindented to a thickness of 100 μm and then

mounted on the diamond culet. The in situ pressure was monitored during the experimental measurements using the ruby

fluorescence method, which correlates the Raman spectroscopic shifts of the R1 and R2 lines of ruby to a well-defined

equation of state.26 High-pressure Raman spectroscopic

experiments were performed in backscattering geometry

utilizing the 473 nm line of a cw diode pumped solid state

laser as the excitation source, with an output power of

∼50 mW. A 500 mm focal length f/6.5 imaging spectrograph

equipped with an 1800 grooves/mm grating and a resolution

of ±0.5 cm−1 was used. The intensities of all Raman spectra

were normalized and truncated at 1200 cm−1 as the first order

scattering of diamond precluded definitive identification of

the fundamental Raman active mode near 1430 cm−1.

The Raman spectrum of calcite at ambient conditions is

well known and characterized.27–29 At ambient pressure, calcite has five fundamental Raman active modes attributed to

two sets of doubly degenerate Eg modes near 155, 280, 710,

and 1430 cm−1 and an intense A1g mode near 1085 cm−1.

29

Representative Raman spectra of calcite collected upon only

isothermal compression at room temperature within the rotational diamond anvil cell to near 20 kbar are presented

in Fig. 3. At ambient pressure, four of the five fundamental Raman-active modes of calcite-I are observed A1g

(∼1085 cm−1) + 3 Eg (712 cm−1, 282 cm−1, and 151 cm−1)

with vibrational frequencies nearly identical to those previously reported.27–29 Upon compression, the bands shift

nearly linearly to pressures approaching 17 kbar (Fig. 4). At

17 kbar, several modifications in the spectral pattern were discerned which are associated with the well documented calcite

→ aragonite transition. The fundamental Eg Raman mode of

the CO3 bending mode splits into two resolvable maxima near

703 and 710 cm−1 (inset Fig. 3). At pressures above the transition, one new phonon band appears in the spectrum of aragonite near 190 cm−1 and the spectral intensity of the band

FIG. 2. Cutaway view of rotational diamond anvil design, showing the gearing in relation to the typical megabar cell.

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88

073905-3 J. Ciezak and T. Jenkins Rev. Sci. Instrum. 82, 073905 (2011)

FIG. 3. (Color online) In situ Raman spectra of calcium carbonate obtained

upon compression in the RDAC. During these experiments, the rotational

capability of the RDAC was not used. The inset shows the splitting of the Eg

bending mode that occurs upon the calcite – aragonite transition. Spectra are

normalized with respect to spectral intensity but offset for clarity.

∼215 cm−1 significantly increases. The band which appears

in the calcite spectrum near 215 cm−1 has not been previously

observed,27–29 so it is likely the sample used for the experiments detailed herein contained a small aragonite impurity.

Additionally, the appearance of two bands between 190 and

225 cm−1 suggests there may also be a concentration of vaterite (calcite-III). Upon decompression, no barrier to hysteresis was observed and the polymorphic phase transformation

was reversible.

FIG. 4. (Color online) Pressure dependence of the fundamental Raman active modes of calcite. Dashed lines indicate apparent phase boundaries. Linear fits are shown for the frequency shifts.

FIG. 5. (Color online) In situ Raman spectra of calcium carbonate held at

5 kbar and subjected to rotation (shear) in the RDAC. The degree of rotation

corresponds to the rotation of the piston anvil. Spectra are normalized with

respect to spectral intensity but offset for clarity.

Representative Raman spectra of calcite held at 5 kbar

and then subjected to incremental degrees of rotation, and thus

shear, within the rotational diamond anvil cell are presented

in Fig. 5. At the low pressures studied by these experiments,

rotation of the anvil did not change the in situ pressure

appreciably, concurrent with previous observations.20 A

comparison of the integrated spectral intensity of the vibrational feature near 215 cm−1 observed during rotation and

that present at ambient conditions indicates the onset of the

calcite → aragonite transition is near 10◦ at 5 kbar. When

the anvil was rotated 30◦ relative to its normal, the resulting

Raman spectrum was nominally identical to the spectrum

of aragonite obtained at 17 kbar and shown in Fig. 3. Upon

decompression, only the pressure was manually released and

interestingly the transition was found to be irreversible, suggesting frictional hysteresis may be associated with the shear

transition.

IV. CONCLUSION

A diamond anvil cell capable of generating simultaneous high pressure and shear for optical spectroscopies has

been developed. The cell can be pressurized to 50 GPa with

full 360◦ of rotation of the piston anvil relative to the fixed

cylinder anvil, which is locked into position with a set screw.

An encoded motor can be attached to the worm gear shaft

which controls the degree of rotation for precision measurements. Raman spectroscopic measurements of the shear induced phase transitions of calcite were used to demonstrate

the capabilities of the RDAC under combined pressure and

shear.

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91

Research@

The First Launch of an Autonomous 

Thrust-Driven Microrobot Using 

Nanoporous Energetic Silicon

Wayne A. Churaman1, Luke J. Currano1, Christopher J. Morris1, 

Jessica E. Rajkowski2, and Sarah Bergbreiter3

1 U.S. Army Research Laboratory, Adelphi, MD

2 Currently with The Mitre Corporation, McLean, VA

3 Department of Mechanical Engineering, University of Maryland, 

 College Park, MD

ABSTRACT

As the capability and complexity of robotic platforms continue to evolve from 

the macro to the micron scale, the challenge of achieving autonomy requires 

the development of robust, lightweight architectures. These architectures 

must provide a platform upon which actuators, control, sensing, power, and 

communication modules are integrated for optimal performance. In this paper, 

the first autonomous jumping microrobotic platform is demonstrated using a 

hybrid integration approach to assemble on-board control, sensing, power, 

and actuation directly onto a polymer chassis. For the purposes of this paper, 

jumping is defined as brief parabolic motion achieved via an actuation pulse 

at takeoff. In this paper, the actuation pulse comes from the rapid release 

of chemical energy to create propulsion. The actuation pulse lasts several 

microseconds and is achieved using a novel high-force/low-power thrust 

actuator, nanoporous energetic silicon, resulting in 250 μJ of kinetic energy 

delivered to the robot and a vertical height of approximately 8 cm.

Journal of Microelectromechanical Systems, Volume 21, Number 1, 

pages 198-205 (2012)
92198 JOURNAL OF MICROELECTROMECHANICAL SYSTEMS, VOL. 21, NO. 1, FEBRUARY 2012

The First Launch of an Autonomous Thrust-Driven

Microrobot Using Nanoporous Energetic Silicon

Wayne A. Churaman, Luke J. Currano, Member, IEEE, Christopher J. Morris, Member, IEEE,

Jessica E. Rajkowski, and Sarah Bergbreiter, Member, IEEE

Abstract—As the capability and complexity of robotic platforms continue to evolve from the macro to the micron scale,

the challenge of achieving autonomy requires the development of

robust, lightweight architectures. These architectures must provide a platform upon which actuators, control, sensing, power,

and communication modules are integrated for optimal performance. In this paper, the first autonomous jumping microrobotic

platform is demonstrated using a hybrid integration approach to

assemble on-board control, sensing, power, and actuation directly

onto a polymer chassis. For the purposes of this paper, jumping

is defined as brief parabolic motion achieved via an actuation

pulse at takeoff. In this paper, the actuation pulse comes from

the rapid release of chemical energy to create propulsion. The

actuation pulse lasts several microseconds and is achieved using a

novel high-force/low-power thrust actuator, nanoporous energetic

silicon, resulting in 250 μJ of kinetic energy delivered to the robot

and a vertical height of approximately 8 cm. [2011-0030]

Index Terms—Autonomy, microrobot, porous silicon.

I. INTRODUCTION

S MALL autonomous robotic platforms hold the potential to

help target the efforts of first responders, provide stealthy

surveillance, or add mobility to sensor networks. However,

existing microrobot platforms have limited autonomy [1]–[11].

Autonomous microrobot platforms require essential components including sensors, actuators, electronic circuits, and a

power supply. These components must further be integrated

to provide certain essential functions including conversion of

sensor data to actionable information, intelligence to make

decisions based on the information, and mobility to take action

on the decisions. Integration of sensing, control, power, and

actuation on a single chassis increases the utility of robots,

allowing them to be placed in environments where they can

sense, think, and act with limited or no human intervention

[1]–[5]. These enabling features of autonomy lead to robots that

Manuscript received January 28, 2011; revised September 23, 2011; accepted

September 28, 2011. Date of publication December 5, 2011; date of current

version February 3, 2012. This work was supported by the U.S. Army Research

Laboratory. This work was also supported in part by the National Science

Foundation under Award CNS0931878. Subject Editor C. Mastrangelo.

W. A. Churaman, L. J. Currano, and C. J. Morris are with the U.S. Army

Research Laboratory, Adelphi, MD 20783 USA (e-mail: wayne.a.churaman.

civ@mail.mil; luke.j.currano.civ@mail.mil; christopher.j.morris58.civ@

mail.mil).

J. E. Rajkowski was with the University of Maryland, College Park, MD

20783 USA. She is now with The Mitre Corporation, McLean, VA 22102 USA

(e-mail: jrajkowski@mitre.org).

S. Bergbreiter is with the Mechanical Engineering Department, University

of Maryland, College Park, MD 20783 USA (e-mail: sarahb@umd.edu).

Color versions of one or more of the figures in this paper are available online

at http://ieeexplore.ieee.org.

Digital Object Identifier 10.1109/JMEMS.2011.2174414

will eventually traverse environments unreachable by humans,

while providing critical data, and the ability to adapt to ever

changing surroundings.

As robot length scales decrease below the centimeter scale, it

becomes an increasing challenge to integrate systems to deliver

complex functionality on a single platform. Larger robots like

RoACH [6] or Alice [7] are assembled using commercial offthe-shelf (COTS) components including motors, sensors, and

electronics with little difficulty. Application-specific integrated

circuits (ASICs) provide a unique approach to developing control architectures for millimeter scale platforms by reducing the

chip area [5]. This paper will take a new hybrid integration

approach that includes both custom fabrication (similar to the

bottom-up design of ASICs) and COTS components (similar to

simple assembly available to larger robots).

Recent work in microrobotics has made progress toward

achieving autonomous functionality, which the authors define

as the integration of sensing, actuation, power, and control on a

single chassis. These characteristics are available to the robot

independent of the operating environment. A number of the

published microrobots can be characterized as “actuator only”

platforms because they rely on a controlled operating environment to function, cannot make decisions based on the environment, and cannot carry their power supply. The microrobot in

[8] maneuvers using an integrated scratch drive actuator and a

cantilever steering arm, allowing it to achieve speeds greater

than 200 μm/sec. The operation of the microrobot depends on

an underlying electrical grid, which provides power and control.

The “walking chip” in [9], designed to carry a maximum

external load of 2.5 g on its back, relies on external power in

the form of tethered wires that measure up to 10 cm in length

and lacks sensing and on-board control. The bio-inspired fly in

[10] demonstrates insect-like wing trajectories with integrated

actuation and mechanics, but does not yet possess integrated

on-board sensing, power, or control.

The complexity of the design and integration increases as

more functionality is incorporated directly onto the chassis.

The microrobot in [11] appears to come close to our definition

of autonomy using solar cells for power, electrostatic motors

for actuation, and a CMOS finite state machine for control.

These components are fabricated individually and integrated

using a hybrid approach involving wirebonding to electrically

connect each component. However, the robot does not have

integrated sensors to allow it to observe and react to changes

in the environment. The only robots that appear to include

all components necessary for individual autonomy are several

centimeters in size or larger [6], [12].

1057-7157/$26.00 © 2011 IEEE
CHURAMAN et al.: FIRST LAUNCH OF AN AUTONOMOUS THRUST-DRIVEN MICROROBOT USING NANOPOROUS ENERGETIC SILICON 93

Fig. 1. Computer model of jumping microrobot showing the polymer chassis

and control circuit. The energetic thrust actuator is underneath, on the “belly”

of the robot.

While these microrobots have demonstrated ingenious functionality, the goal remains to develop a platform that is truly

autonomous, where all required functionality is provided onboard the microrobotic platform. To overcome the challenges

associated with implementation and integration on a subcentimeter scale, the authors have developed a novel hybrid

integration approach to fabricate a microrobotic platform,

conceptually shown in Fig. 1. The microrobotic platform is

designed with a polymer chassis, upon which micron-scale

electrical traces are lithographically defined and patterned to integrate power (capacitors), sensing (a phototransistor), and the

most rudimentary form of decision making (a single transistor).

Nanoporous energetic silicon, a solid energetic formulation that

produces gas when ignited [13], [14], functions as a thruster to

provide mobility. The nanoporous energetic silicon is mounted

underneath the polymer chassis. The microrobotic platform can

detect a change in the ambient light and respond to the change

by triggering ignition of the energetic material. This produces

an upward thrust.

This work offers two key contributions toward greater autonomy in microrobotics along with a demonstration of this

autonomy. First, a novel energetic silicon-based actuator is

proposed to enhance mobility (Section II). Mobility is demonstrated in the form of vertical thrust as stored chemical energy

is rapidly converted to mechanical energy. While only a single

propulsion event is demonstrated, the nanoporous energetic

silicon can be arrayed on-chip to provide multiple propulsion

events in different directions. The second contribution develops

a hybrid integration approach used to mechanically and electrically integrate this actuator with sensing and control elements

(Section III). The integration of on-chip, chemical-based actuation with electronic controls results in sufficient autonomy for

applications ranging from mobile sensor networks to dynamTABLE I

PHYSICAL DIMENSIONS OF SYSTEM COMPONENTS

USED TO DESIGN THE HEXAPOD

ically controlled microthrusters. These two contributions are

combined to demonstrate the first robot at this size scale with

the ability to sense, think, and act (Section IV).

II. SYSTEM ARCHITECTURE

The goal for the microrobot system architecture was to

develop a robotic platform that could sense a change in light

intensity and respond by jumping. To simplify matters, the

microrobot would jump vertically with a takeoff angle of 90◦

rather than follow a particular direction. To accomplish this

goal, the system architecture required a sensor, control circuit,

actuator, and power, each of which needed to be integrated on a

polymer chassis. The electronics were designed and fabricated

with COTS parts to provide both on-board power and logic to

sense a change in light intensity and to provide an electrical

stimulus to trigger the energetic material. Table I summarizes

the size and weight of each of the components of the system,

and the following subsections describe in detail the function

of each part. The total mass of the assembled system was

approximately 300 mg. In our experiments, three hexapods

were assembled and tested. The overall mass of the hexapod

varied from 280 mg to 318 mg across the platforms that were

assembled. The largest contributor to this variation in mass was

an additional layer of polymer that was applied to the hexapod

once it was assembled to encapsulate the IC components and

increase the robustness of the robot.

A. Actuation

Instead of using a mechanical spring to store energy as seen

in [15], nanoporous energetic silicon was chosen to store and

release energy chemically. The primary benefit to this approach

is reduced complexity and lower part count due to eliminating the need for a separate electrically controlled actuator to

compress a spring. Nanoporous energetic silicon is formed

in an electrochemical etch process as outlined in [13]. It is

initially inert, and only becomes energetic when infused with an

oxidizer such as sodium perchlorate. The exothermic reaction

can be triggered with heat, friction, or focused light. Multiple

actuators can be integrated onto the chassis of a microrobot,

allowing it to achieve multiple jumps in different directions.

For this work, actuation was demonstrated using a single, 2 mm

diameter region of nanoporous energetic silicon.
94200 JOURNAL OF MICROELECTROMECHANICAL SYSTEMS, VOL. 21, NO. 1, FEBRUARY 2012

Fig. 2. Tethered jump of 100-mg hexapod platform, tethered by 12-mil wire.

The images are spaced 10 frames (10 ms) apart. In frames 1–3, a portion of the

silicon chip detached from the hexapod.

Propulsion of a hexapod platform was demonstrated using

the nanoporous energetic silicon. Fig. 2 shows an initial (tethered) experiment with a 2 mm diameter, 40 μm thick volume

of nanoporous energetic silicon incorporated on the polymer

hexapod chassis [16]. The nanoporous silicon was attached

to the underside of the hexapod, and connected via wires to

an external DC power supply. The components of the control

circuit were not attached to the hexapod in this experiment,

resulting in a significantly lower system mass. Including the

wire leads and the energetic silicon chip, the hexapod weighed

100 mg in this configuration. The hexapod travelled up out of

the ∼11 cm field of view. The takeoff velocity was measured

at 2.8 m/s from high-speed video frames captured at 1000

frames/s. Based on this, the projected jump height was 40 cm.

This experiment showed that nanoporous energetic silicon is a

viable actuator for a jumping robot, so the authors proceeded to

integrate sensing, control, and power functions onto the chassis.

B. Sensing and Control

The microrobotic platform is designed to jump in response

to “seeing” a change in light level, as an insect which may

jump in response to sensing the shadow of a predator. Sensing

is achieved using a light-sensitive phototransistor (SFH3710)

shown in Fig. 3. The space available to integrate electronic

control, sensing, and actuation is limited to the top and bottom

of the chassis, approximately 56 mm2 total area. The top of

the chassis is used to integrate a circuit composed of five

discrete surface-mount components (shown in Fig. 3), while the

underside of the chassis houses the energetic material used for

actuation. The control circuit itself consists of a 2N7002, single

n-channel enhancement MOSFET, and a 10 kΩ resistor used to

bias the gate of the MOSFET.

The spectral sensitivity of the NPN phototransistor chosen

is 350 nm to 950 nm with a maximum sensitivity at 570 nm.

When the lighting condition is varied from an illuminance

of ∼0.3 lux, which is comparable to a full moon on a clear

night [17], to approximately 26 lux, somewhat less than the

illuminance of a family room [18], the voltage at the gate of

the MOSFET rises to 5 V. This closes the MOSFET transistor.

Fig. 3. Circuit designed to provide sensing and actuation. The 3 Ω resistor

represents the hotwire used to ignite the energetic material.

A series connection of a 10 μF capacitor and 3 Ω resistor with

the MOSFET is shown in Fig. 3. The 3 Ω resistor is used to

model the electrical characteristics of an electrothermal initiator

used to ignite the energetic material. Details about the electrical

initiator design and performance are discussed in [13]. The

10 μF capacitor is precharged to 6 V for the experiments

discussed here. As the MOSFET closes, it provides a path for

current to flow through the electrical initiator, producing heat

to ignite the energetic material. The electrothermal initiator requires at least 150 mA to ignite the energetic material. However,

it should be noted that this current is only required for at most

100 μs. In this system, the capacitor discharges directly through

the electrothermal initiator wire, without voltage or current

regulation. This arrangement simplifies the circuit, reduces the

part count, and is sufficient to ignite the energetic silicon.

C. Power

The two capacitors shown in Fig. 3 provide power to the

phototransistor and the initiator on the nanoporous energetic

silicon. The 100 μF capacitor provides power to bias the

phototransistor. The 10 μF capacitor is used to provide power to

ignite the nanoporous energetic silicon once the phototransistor

detects the increase in light intensity. Both the 100 μF and

10 μF capacitors in Fig. 3 are precharged to 6 V. Both power

sources are isolated until an event is detected, triggering the

actuation response. A larger capacitor is chosen to power the

phototransistor to increase the time constant and supply a larger

amount of energy to accommodate leakage. While the longevity

of the capacitors as power sources is not investigated in detail,

the 100 μF capacitor is able to retain sufficient charge to power

the phototransistor for up to 8 min in a dimly lit room (∼0.3

lux). While the expected discharge time, based on the off-state

resistance, for the phototransistor and the FET is approximately

40 min, a significant amount of leakage current across the

phototransistor results from exposure to the low-level ambient

light. It should be noted that the 100 μF capacitor used to

bias the phototransistor is the heaviest single part of the robot

(Table I), with the much smaller capacitor only acting to initiate the energetic material. In the future, the requirement for

this larger capacitor may be removed through better sensors

or through the use of a low-volume, high-capacity thin film

battery. The energy associated with the capacitors and the

nanoporous energetic silicon is summarized in Table II.
CHURAMAN et al.: FIRST LAUNCH OF AN AUTONOMOUS THRUST-DRIVEN MICROROBOT USING NANOPOROUS ENERGETIC SILICON 95

TABLE II

SUMMARY OF ENERGY SUPPLIED BY THE CAPACITORS AND

ENERGY SUPPLIED TO AND GENERATED BY THE

NANOPOROUS ENERGETIC SILICON

III. HYBRID INTEGRATION

To electrically and mechanically integrate the sensing,

power, and control components with the actuator and robot

chassis with minimal added mass, a new hybrid integration

process was designed. This approach used a polymer chassis

for structural support to carry a payload as well as a substrate

for electrical traces for the integration of the control circuit.

A. Polymer Chassis

The hexapod style chassis was chosen to enable walking in

future iterations of the platform, and the polymer was chosen

for robustness, durability, light weight, and ease of fabrication.

The chassis provided the robot’s functional skeleton as well as

a point of assembly on which to integrate all other components

in this hybrid integration approach (Fig. 1). It is 4 mm by 7 mm

by 0.5 mm in dimensions and was fabricated using a rapid prototyping process [19]. A photodefinable polymer (Loctite3525

modified acrylic) was first spread evenly on a transparent plastic

sheet, and another transparent sheet was placed on top of the

polymer film. Glass slides were placed between the two sheets

to define the thickness. A photolithographic mask, with the

hexapod features defined as clear windows in an opaque field,

was then placed on top, and UV light was used to expose the

hexapod features through the mask. The exposed polymer was

hardened, and the unexposed polymer could then be removed

with a methanol rinse. The metal traces for assembly of the

control circuit were deposited and patterned (described in the

following section), then the legs were bent out of plane, and an

additional layer of polymer, [shown in Fig. 4(c)], was applied to

the inside corner between the legs and the body and cured. This

resulted in the 3-D hexapod structure similar to that in Fig. 4(d),

which allows for future integration of leg actuators for bimodal

(walking/jumping) mobility or to orient the robot for a jump in

a particular direction.

B. Metallization

A copper film was deposited on top of the polymer chassis

using metal evaporation at room temperature. Once deposited,

the electrical traces were lithographically patterned on the

copper film. With photoresist protecting the underlying copper

metal traces, the exposed copper was etched using a 2 : 1

mixture of water to nitric acid. To facilitate attaching the circuit

elements to a flexible substrate, the copper traces were then

coated with a low melting temperature solder (Indalloy 117

from Indium Corporation) using a dipping process discussed

Fig. 4. Process flow to fabricate 3-D hexapod structure. This process flow

shows metal patterned on the legs. For the purpose of our experiment, metal

was only patterned on the body.

Fig. 5. Metalized layer begins to crack when submerged in low-temperature

solder. Copper traces were subsequently widened to 472 μm to allow enough

solder to heal the cracks.

in [20]. Indalloy 117 melts at 47 ◦C, and was chosen because

the Loctite polymer softens above 100 ◦C and using a solder

iron to heat the solder would result in damage to the polymer

substrate. The circuit components can be attached by reflowing

the low-temperature solder without damaging the substrate.

For this process, the solder was kept at 65 ◦C to ensure the

robustness of the polymer chassis. A mixture of ethylene glycol

and hydrochloric acid was added to the solder dipping crucible

to remove any surface oxides on the copper traces and on the

solder. The polymer chassis was slowly submerged into the

molten alloy, and retracted after 10 s. About 50% of the time, a

visual inspection revealed incomplete coating, and the chassis

was submerged again to coat all copper traces. Initial results

indicated a tendency for the copper layer to crack when cooling

after removal from the solution, as shown in Fig. 5. Therefore,

the minimum copper feature size was increased from 133 μm

to 472 μm, providing greater surface area to capture enough

solder to “heal” any cracks that form.
96202 JOURNAL OF MICROELECTROMECHANICAL SYSTEMS, VOL. 21, NO. 1, FEBRUARY 2012

Fig. 6. Metal bond pads fixed to Si allow the energetic to be connected to the

control circuit.

C. Assembly

A serial process was used to assemble each of the components. The hexapod and circuit components were separately

coated with the low melting temperature alloy. A solder iron

was used to manually solder external wires to the capacitors

and the MOSFET before the components were assembled onto

the hexapod chassis. This allowed electrical leads to be easily

attached to charge the capacitor power sources after assembly. The polymer hexapod was then placed on a hotplate at

65 ◦C to melt the Indalloy, and the circuit components were

positioned using tweezers. Moderate pressure was applied manually between the component and the binding site to ensure

a robust connection. The hexapod was then removed from

the hotplate and allowed to cool. As the polymer substrate

cooled, the circuit components were fused to the binding

sites.

Attempts to solder directly to the energetic porous silicon

chip were unsuccessful because the thin-film initiator was not

robust enough to withstand the heat of a soldering iron. The low

melting temperature alloy was not attempted because of concern that the dipping process might contaminate the nanoporous

silicon. Although it was possible to wirebond to the silicon chip,

wirebonding to the polymer chassis was not successful because

the polymer was too soft. Therefore, gold-plated metal pads

shown in Fig. 6 were glued to the silicon chip, and wirebonds

were made between the integrated initiator and the metal pads.

Wires were then soldered between the control circuit and these

metal pads, and the nanoporous energetic silicon was attached

to the underside of the hexapod using a double-adhesive tape.

In future iterations, the pads could be eliminated by creating

vias in the polymer chassis and the nanoporous silicon substrate

and filling them with metal for direct connection using the lowtemperature solder.

Once the hexapod was assembled, the circuit and underlying

electrical traces were coated with an additional layer of Loctite

to prevent delamination of the rigid circuit elements from the

flexible polymer skeleton. This robust integration allowed the

hexapod legs to be flexed and potentially actuated out of plane,

without disturbing the functionality of the circuit. This unique

process therefore uses one material as the mechanical structure,

electrical substrate, and packaging material for the microrobotic

Fig. 7. Metal bond pads soldered to wires coming off of the capacitor

and MOSFET.

platform. Fig. 7 shows a circuit that has been fully assembled

onto the polymer chassis.

D. Packaging

The sodium perchlorate oxidizer used to make the

nanoporous silicon into an energetic material is strongly hygroscopic, and the absorption of water by the oxidizer degrades

the efficiency of the energetic reaction and produces stress

within the porous layer. The induced stress can also cause the

porous material to crack and the hotwire initiator to subsequently break. For this reason, the jumping experiments were

all conducted in a dry box plumbed with nitrogen. Even so,

the process of opening the door of the enclosure to charge the

capacitors prior to ignition allowed enough time for the oxidized sample to absorb water and degrade the efficiency of the

reaction.

Three techniques were employed to limit the rate at which

moisture was absorbed by the oxidizer. The first technique

involved applying one drop of oxidizer solution (3.2 M NaClO4

in methanol) using a syringe in a dry room controlled to < 1%

relative humidity and allowing the sample to dry for 20 min.

After the sample was dried, a thin layer of Loctite (∼500 μm

thick) was carefully applied on top of the porous silicon and UV

cured for 5 min. The second technique was to apply the oxidizer

and dry in a nitrogen dry box, then seal the porous silicon with

an evaporated parylene layer (∼12 μm thick). While loading

the samples into the parylene system, there was some exposure

to environmental moisture. An alternative approach which has

not been attempted yet would be to dry the sample under

the dry nitrogen inside the parylene coating chamber to avoid

transport.

The final (and simplest) humidity exposure minimization approach was to connect leads to the hexapod prior to application

of the oxidizer and run these leads outside of the controlled

environment chamber. The oxidizer was allowed to dry for

15 min. The door to the chamber was then opened very briefly

(less than 30 s) to disconnect the leads and then closed for the

experiment.

IV. RESULTS

Experiments were performed in a laboratory, and a moderately dark environment was simulated by turning off all
CHURAMAN et al.: FIRST LAUNCH OF AN AUTONOMOUS THRUST-DRIVEN MICROROBOT USING NANOPOROUS ENERGETIC SILICON 97

overhead fluorescent lights, and turning on a small 12-W

lamp in the corner of the room to allow some visibility

for positioning the device and connecting and disconnecting electrical leads. The resulting ambient light level at the

photodetector was ∼0.3 lux, as mentioned above. Oxidizer

was applied to the nanoporous energetic silicon using one of

the three different techniques discussed above to minimize

moisture exposure. The 100 μF and 10 μF capacitors were

charged to 6 V, and the hexapod was positioned approximately

12 inches under a high intensity 3 LED array light source. The

light source was turned on, triggering the phototransistor, which

closed the MOSFET switch, which in turn triggered ignition of

the energetic nanoporous silicon. Each jump was captured by a

high-speed camera recording at 1000 frames per second.

The first untethered experiment was performed with the

Loctite moisture seal encapsulating the porous silicon. The

experiment was done in ambient air; therefore an encapsulation

layer was needed to prevent moisture uptake and subsequent

degradation of the porous silicon and oxidizer mixture. This

hexapod had a total weight of 280 mg, and achieved a vertical

height of 1 cm when the light was turned on. The limited

height may have been due to packaging the sample with Loctite.

Deposition of the Loctite directly on the porous silicon resulted

in a thick layer (∼500 μm) of the Loctite remaining on the

surface. A portion of the output energy being converted into

thrust for movement would therefore be lost in breaking the

seal created by the Loctite. It was also possible that the Loctite

coating may have affected the oxidized porous silicon surface,

causing it to be less reactive. Similar results were achieved with

the parylene-coated samples. Further work must be done to

explore other alternatives to encapsulate the energetic without

degrading the performance.

The best jump heights were achieved with the simplest

humidity exposure minimization approach, which involved preconnecting the electrical leads to charge the capacitors before

applying and drying the oxidizer, and running the experiment

in a nitrogen-filled box. The box was only very briefly opened

(less than 30 s) to disconnect the leads. In this case, a 314-mg

hexapod jumped approximately 8 cm vertically (Fig. 8).

The authors analyzed approximately 200 frames of the jump

trajectory to determine if any significant acceleration sources

other than gravity affected the microrobotic platform during

flight. Several frames at the end of the trajectory were not

captured in the high-speed video, so the landing portion of the

trajectory was not shown. The center of the hexapod chassis was

chosen as the center of mass, and the height was measured as

the center of mass crossed each half centimeter marker on the

height scale. The resulting position versus time data is shown

in Fig. 9. Given the resolution of the camera, the accuracy

with which we were able to pinpoint the robot’s center of mass

in each frame, and parallax errors involved observation from

a fixed viewpoint, we estimated the uncertainly for each data

point as approximately 0.25 cm. Using this uncertainty, the

potential energy based on the maximum observed height was

between 246 μJ and 262 μJ.

The data points were then fitted to the parabolic curve

describing the motion of a projectile under constant free-fall

acceleration g and negligible air resistance. The parabolic curve

Fig. 8. Propelled hexapod when actuator triggered by light captured at 1000

frames per second. Frames shown above are 20 ms apart. The vertical scale bar

shown is spaced 1 cm per division.

Fig. 9. Trajectory of the hexapod jump. Experimentally determined height

shown as blue circles. The parabolic curve fit describes projectile motion under

acceleration due to gravity and negligible air resistance.

was then fitted to the data to determine the initial velocity vi, as shown in Fig. 9. This fit resulted in a calculated

take-off velocity of 1.263 m/s. The associated kinetic energy

is 250 μJ, which is within the potential energy uncertainty

bounds calculated based on measured maximum height. The

quality of the fit to the experimental data also demonstrates

that the microrobotic platform followed a drag-free parabolic

trajectory under acceleration due to gravity and negligible

air resistance.

Although the overall mass of the energetic chip as given in

Table I was approximately 25 mg, the mass of the reactive

region was estimated at only 0.21 mg. The 2 mm diameter

energetic region was able to propel the robot approximately

8 cm into the air, even though the robot was more than

1000 times the mass of the energetic material.

As a point of comparison, to achieve the same height with

a conventional actuator would require an actuator stiffness

of 2.4 kN/m and a force of at least 9.6 N (assuming the
98204 JOURNAL OF MICROELECTROMECHANICAL SYSTEMS, VOL. 21, NO. 1, FEBRUARY 2012

leg length of 4 mm is the actuator throw). This would be

very difficult if not impossible to achieve at this size scale.

To ignite the nanoporous energetic silicon, a maximum of

180 μJ of electrical energy was consumed (assuming that the

10 μF capacitor discharged entirely from 6 V to 0 V). Defining

an “effective actuation efficiency” as the ratio of mechanical

energy produced to electrical energy consumed, the 8 cm jump

here represented an effective efficiency of 137%. Based on

past measurements of the energy density of the nanoporous

energetic silicon reaction, approximately 1.9 J of chemical

energy was released [21]. Therefore, the chemical to kinetic

conversion efficiency was actually extremely low (0.0127%),

indicating that the vast majority of the available energy was

wasted.

There is consequently room for improvement, and the authors believe implementation of a nozzle in future designs will

increase this efficiency by directing gases in a more effective manner to produce maximum thrust. However, the fact

that a respectable height of over 11× the body length was

achieved despite the abysmal conversion efficiency underscores

the tremendous amount of energy available in the nanoporous

energetic silicon. The authors believe that this technology will

have many applications in future MEMS systems if the energy

can be more efficiently captured and put to use.

V. CONCLUSION

An autonomous jumping microrobot has been demonstrated,

achieving a vertical jump of 8 cm, which was over 11 times its

longest physical dimension. Autonomy was achieved through

on-board integration of a simple circuit constructed of discrete surface mount components. The resulting system has

optical sensing, control, power, and mobility. The rudimentary sensing and intelligence allow the hexapod to detect an

increase in light intensity and respond by jumping. In addition, the novel use of an energetic material for actuation

was shown, allowing significantly more mechanical energy to

be produced than the electrical energy needed to initiate the

reaction.

A hybrid integration approach has also been presented, which

uses a low-temperature solder and metal interconnects patterned directly on the robot’s polymer chassis to integrate various discrete electrical components. One of the unique aspects

to the integration approach is that a single UV-curable polymer

material was used for the robot chassis, as a substrate for the

assembly of the circuit components, and as a packaging material. The authors expect improvements in direction control and

development of an array of energetic actuators to allow multiple

jumps, leading to future technological leaps for autonomous

microrobots.

ACKNOWLEDGMENT

The authors would like to thank the Micro Robotics Laboratory at the University of Maryland for assistance in fabricating

the robotic polymer chassis.

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Wayne A. Churaman received the B.S. and M.S.

degrees in electrical engineering from the University

of Maryland, College Park, in 2005 and 2010, respectively, with a specific focus on microelectronics

and microrobotic platforms.

He currently works as a MEMS Engineer for

the U.S. Army Research Laboratory, Adelphi, MD.

His interests include experimental analysis of nanoenergetic materials, integration of nanoenergetic

materials with MEMS, and MEMS sensor design

and fabrication. He is currently developing novel

actuators using nanoenergetic porous silicon.
CHURAMAN et al.: FIRST LAUNCH OF AN AUTONOMOUS THRUST-DRIVEN MICROROBOT USING NANOPOROUS ENERGETIC SILICON 99

Luke J. Currano (M’09) received the B.S., M.S.,

and Ph.D. degrees in mechanical engineering from

the University of Maryland, College Park, in 2000,

2002, and 2010, respectively.

He has worked as a MEMS Engineer for the U.S.

Army Research Laboratory (ARL), Adelphi, MD,

since 2001. He is currently Team Leader of the

MEMS and Nanoenergetic Devices team at ARL.

His research interests include integration of energetic materials with MEMS, high force/displacement

MEMS actuators, acceleration switches, no-power

sensors, and biologically inspired MEMS devices.

Christopher J. Morris (S’04–M’07) received the

M.S. degree in mechanical engineering and the Ph.D.

degree in electrical engineering from the University

of Washington, Seattle, in 2000 and 2007, respectively.

Between his graduate studies, he was with Micronics, Inc., developing microfluidic devices. He is

currently with the U.S. Army Research Laboratory

in the Sensors and Electron Devices Directorate,

Adelphi, MD. His current research interests lie in the

integration of new materials and technologies with

microelectronics and MEMS, including on-chip energetic materials, electronics

packaging, and fluidic self-assembly.

Jessica E. Rajkowski received the B.S. and M.S.

degrees in mechanical engineering from the University of Maryland, College Park, in 2009 and 2010,

respectively, with a focus on microrobotics.

She is currently a Senior Systems Engineer at The

MITRE Corporation, McLean, VA, where she works

in the Nanosystems Group. Her research interests

include the design and fabrication of millimeter-scale

robotics and thermal micro-actuators.

Sarah Bergbreiter (S’03–M’07) received the B.S.E.

degree in electrical engineering from Princeton University, Princeton, NJ, in 1999, and the M.S. and

Ph.D. degrees from the University of California,

Berkeley, in 2004 and 2007, respectively.

She is currently an Assistant Professor of Mechanical Engineering, with a joint appointment in

the Institute for Systems Research at the University

of Maryland, College Park. Her research addresses

the challenges necessary to make robotic bugs

move, including effective locomotion, efficient actuation, micromechanism design and fabrication, power systems, sensing, and

integration.

Dr. Bergbreiter received a DARPA Young Faculty Award for her work toward

autonomous millimeter-scale robots in 2008.

101

Research@

Simulations of High-Pressure Phases in RDX

Lynn B. Munday1, Peter W. Chung1, Betsy M. Rice1, and Santiago D. Solares2

1 U.S. Army Research Laboratory, Aberdeen Proving Ground, MD

2 Department of Mechanical Engineering, University of Maryland, 

 College Park, MD

ABSTRACT

Using a fully flexible molecular potential in equilibrium molecular dynamics 

simulations, we study the a-and g-polymorphs of the energetic molecular 

crystal hexahydro-1,3,5-trinitro-s-triazine (RDX), their respective properties, and 

the conditions that contribute to the stress-induced g "a solid–solid phase 

transition mechanisms. We find the pressure-dependent atomic structure, 

mechanical properties, and transition behavior to be described reasonably well. 

Uniaxial deformation of a-RDX along the crystal axes is shown to result in three 

different crystal responses where compression of the c-axis results in the a " g

transition, compression of the b-axis causes a transition with resulting structure 

similar to stacking faults observed by Cawkwell et al. [J. Appl. Phys. 2010, 107, 

063512], and no transitions are observed for compression of the a-axis.

The Journal of Physical Chemistry B, Volume 115, Pages 4378-4386 (2011)
102

Published: March 24, 2011

r 2011 American Chemical Society 4378 dx.doi.org/10.1021/jp112042a |J. Phys. Chem. B 2011, 115, 4378–4386

ARTICLE

pubs.acs.org/JPCB

Simulations of High-Pressure Phases in RDX

Lynn B. Munday,†,‡ Peter W. Chung,*,† Betsy M. Rice,† and Santiago D. Solares‡

†

U.S. Army Research Laboratory, Aberdeen Proving Ground, Maryland 21005, United States

‡

Department of Mechanical Engineering, University of Maryland, College Park, Maryland 20742, United States

bS Supporting Information

1. INTRODUCTION

The events leading up to initiation of energetic materials is an

area of research that continues to draw much interest.1 The need

to assess the sensitivity of explosive materials for the sake of

safety and handling has fostered research that has led to better

understanding of the mechanisms that convert mechanical

stimuli to enable chemical events. Numerous efforts have described potentially relevant initiation mechanisms including uppumping,2 steric hindrance,3 dislocation avalanches,4 and electronic effects at extremely high compressions.5 Required in the

representation of these mechanisms is a presumption of the

polymorph at the point of initiation.

One of the more popular systems used to explore initiation

mechanisms is the energetic crystal RDX (C3H6N6O6 or hexahydro-1,3,5-trinitro-s-triazine), a cyclic nitramine widely used in

military applications. At the reference conditions, T = 300 K and

P = 1 atm, RDX is stable in the R-polymorphic form, which is an

orthorhombic unit cell occupied by eight molecules in space

group 61. The unit cell axes and origin for space group 61 can be

chosen in multiple ways and are represented by the symbols

Pbca6,7 or its alternative axes setting Pcab.

8 R-RDX material

properties are orthotropic in thermal expansion9 and elastic

tensor.8,10 At high pressure, it goes through a phase transition

to form the γ-polymorph, which also contains eight molecules

but belongs to the orthorhombic Pca21 space group 29.11 Other

observed phases include the β phase12 and the δ and ε phases, 13

which are summarized by Ciezak and Jenkins14 in a P�T phase

diagram. The experimentally determined crystal phases of

RDX6,11�14 are distinguishable by the so-called wag angle δ,

which is the angle between the N�N bond and the plane created

by the C�N�C triple (as shown in Figure 1a). Each RDX

molecule has three such wag angles. The wag angle normally falls

into three approximate ranges, less than �10, �10 to þ10, and

greater than þ10 in the positive sense of the angle shown in

Figure 1a. These segments of the molecule are in the so-called

equatorial (E), intermediate (I), or axial (A) configurations,

respectively. An R-RDX molecule has two N�N bonds in the

axial configuration and one bond in the equatorial configuration;

thus, the conformation of such a molecule is denoted as AAE.

The same conformation is shared by all eight molecules in the

unit cell of R-RDX. In contrast, among the eight molecules in the

γ-RDX unit cell, four have the AAI configuration and four have

the AAE configuration. The molecular conformations and unit

cells of R- and γ-RDX are presented in Figure S2 of the

Supporting Information. The β12 and ε-RDX13 phases both

belong to the Pca21 space group and contain molecules in the

AAA conformation.

Early experiments15,16 identified the existence of the γ-polymorph at high pressure. The accompanying discontinuous

decrease in volume across the R- to γ-transition and distinct

compressibility changes at high pressure11,14�16 make knowledge of this nonreactive behavior important for determining its

role in initiation. Hydrostatic compression experiments place the

critical transition pressure at 3.8 GPa, independent of temperature in the range of 150�375 K.17 Shock experiments on single

crystals have shown the phase transition at approximately 4.5

GPa to be independent of crystal orientation.18

Atomistic simulations have largely been qualitatively consistent with experiment. Nonreactive flexible molecule models of

RDX such as that of Boyd et al.19 or Smith and Bharadwaj20 (SB)

Received: December 19, 2010

Revised: March 2, 2011

ABSTRACT: Using a fully flexible molecular potential in

equilibrium molecular dynamics simulations, we study the Rand γ-polymorphs of the energetic molecular crystal hexahydro-1,3,5-trinitro-s-triazine (RDX), their respective properties,

and the conditions that contribute to the stress-induced γ f R

solid�solid phase transition mechanisms. We find the pressuredependent atomic structure, mechanical properties, and transition behavior to be described reasonably well. Uniaxial deformation of R-RDX along the crystal axes is shown to result in

three different crystal responses where compression of the c-axis

results in the R f γ transition, compression of the b-axis causes a transition with resulting structure similar to stacking faults

observed by Cawkwell et al. [J. Appl. Phys. 2010, 107, 063512], and no transitions are observed for compression of the a-axis.
103

4379 dx.doi.org/10.1021/jp112042a |J. Phys. Chem. B 2011, 115, 4378–4386

The Journal of Physical Chemistry B ARTICLE

have been shown to reproduce the various experimental conformations of RDX and mechanical properties (see Supporting

Information). The SB potential has furthermore been used to

simulate the observed dependence of shock strength and orientation on the nonreactive dynamic processes involved in the

elastic plastic response of RDX.21�24 On the basis of their

simulations, Cawkwell et al.21 and Ramos et al.22 proposed that

partial dislocation loops homogenously nucleated throughout

the material as stacking faults cause an anomalous hardening,

changing the elastic plastic response as observed in (111)- and

(021)-oriented shock experiments.22,25 They also determined

these stacking faults to be stabilized by stress applied normal to

the (001) plane and nucleated by shear stress in the [010]

direction.21 However, it was also noted that anomalous hardening may be observed if the stress was sufficient to induce the

transition into the γ-polymorph. Owing to the higher experimentally known transition pressure and the fairly isotropic

dependence on shock orientation,18 the involvement of the R

f γ transition was deemed unlikely.17,22

Bedrov et al.23 performed uniaxial constant stress Hugoniotstat simulations using the SB potential for [100] compression of

R-RDX, resulting in amorphous shear banding for pressures

above 9 GPa. Similar calculations of [001] compression of RRDX produced the R f γ transition for pressures above 2.0

GPa.23 Furthermore, Bedrov et al. found differing behavior

according to the crystal orientation. Whereas [100] compression

yielded amorphization and sudden changes in volume following

the initial compression, similar behavior did not occur for [001]

compression. Clearly, determining the dependence of this solid�solid phase transition upon compression, stress, and shear of

the material requires further exploration.

In this paper, we study the basic equilibrium properties of and

transitions to and from the R- and γ-polymorphs in crystalline

RDX based on the SB potential.20 Atomic-level effects in

preinitiation processes are difficult to discern from measurement,

but they likely play significant roles in the mechanisms that lead

to initiation. Thus, verified atomistic models may offer useful

insights to complement experiments. While it has been shown

that the γ-phase is observable using atomistic simulation approaches, to the best of our knowledge, a definitive lowest

possible pressure and properties of the stable γ-structure using

the SB potential have not been reported. It would therefore be

useful to determine the conditions that foster transition behavior

in simple models of RDX (such as that described by the SB

potential). To this end, we employ molecular dynamics (MD)

with the SB potential to model the quasi-static response of the

RDX crystal to hydrostatic and uniaxial compression. We also

compute the transitioned structure under uniform deformations

at room temperature to study the steric mechanisms. In addition,

we consider the roles of crystal anisotropy and load orientation

on these steric effects. The paper is organized as follows. In

section 2, the details of the simulation method are presented.

Results and discussions are presented in section 3 and, the

conclusions are given in section 4.

2. SIMULATION APPROACH

2.1. Molecular Potential. The nonreactive potential energy

function parametrized by Smith and Bharadwaj20 (SB) is used in

this work to describe RDX atomic interactions with fully flexible

bonded interactions and nonbonded electrostatic and dispersion

repulsion forces. The SB potential was parametrized to quantum

mechanical calculations for octahydro-1,3,5,7-tetranitro-1,3,5,7-

tetrazocine (HMX) but has been shown to be transferable to use

in atomistic simulations of RDX.21�24 The form of the SB

potential is

UTotal 1⁄4 ∑

N � 1

i1⁄41

∑

N

j>i

ke

qiqj

rij

þ Aij expð�BijrijÞ � Cij

r6

ij !

þ ∑

bonds

KS

2 ðr � r

o

Þ

2 þ ∑

angles

KB

2 ðθ � θo

Þ

2

þ ∑

proper dihedrals

KT

2 1⁄21 � cosðnφÞ� þ ∑

improper dihedrals

KD

2 φ2

ð1Þ

with the parameters published by Bedrov et al.26 The intramolecular bonded terms include harmonic bond stretching and

angle bending (KS and KB

), harmonic out-of-plane bending using

an improper dihedral (KD), and cosine series dihedrals (KT and

n). The improper dihedral angle definition used in this study is

shown in Figure 1b, where the black lines represent bonds and

the dashed red lines represent the improper dihedral. Nonbonded interactions are represented by exponential repulsion

(Aij and Bij), r

�6 dispersion (Cij), and electrostatic interactions

(partial charges qiand qj

), which are calculated between all atoms

separated by more than two bonds.

2.2. Simulation Methodology. Molecular dynamics simulations in both the isothermal�isochoric (NVT) and isothermal�isostress (NST) ensembles were performed. The MD

package DL-POLY 2.2027 was used to determine the thermalized

equilibrium structures of the RDX polymorphs at the reference

state, P = 1 atm and T = 300 K for R-RDX and P = 5.2 GPa and T

= 300 K simulations for γ-RDX. The supercell used in these

simulations is a block of 2 � 3 � 3 crystallographic unit cells of

RDX and contains 3024 atoms. The small supercell size will help

to isolate only those conditions that foster homogeneous phases

and are unable to account for seeding through a local defect.

Figure 1. (a) RDX wag angle δ and (b) definition of the improper

dihedral angle φijkl and vectors used to find wag angle δ used in this

study. The wag angle, δ, in (a) is related to the indices in (b) by i = N4, j

= C1, k = C3, and l = N7.
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The Journal of Physical Chemistry B ARTICLE

Crystallographic properties were calculated using NST-MD for

both the zero-pressure R-RDX phase characterized by Choi and

Prince6 and the high-pressure γ-RDX phase characterized by

Davidson et al.11 for P = 5.2 GPa. Both phases belong to

orthorhombic space groups (Pbca for R-RDX and Pca21 for γRDX) and contain eight RDX molecules per unit cell. In order to

provide a direct comparison of the γ-polymorph with R, γ-RDX

in this study is presented in the Pb21a space group, an alternative

setting to Pca21 in space group 29, presented in Figure S2c in the

Supporting Information. This effectively switches the b- and clattice vectors given by Davidson et al.11 to match those used by

Choi and Prince.6 Nonorthogonal parallelepiped periodic

boundary conditions were used to model the bulk material under

the least geometrically restrictive conditions. Additionally, thermalized equilibrium structures of these polymorphs were calculated for a variety of pressures and strains. Each new pressure or

strain increment used the final configuration from the prior

increment as its initial state.

The R-polymorph was subjected toNST-MD simulations over

the pressure range of 1 atm to 5.2 GPa using 0.5 GPa increments.

The γ-polymorph, beginning with the 5.2 GPa structure, was

subjected to NST-MD simulations over the same pressure range

and increment. The shock Hugoniot of R-RDX was also calculated using the method developed by Erpenbeck.28

The cutoff distance for nonbonded van der Waals and

electrostatic interactions is 10 Å. The electrostatic forces were

evaluated using Ewald sums. The velocity Verlet algorithms for

the NVT and NST ensembles as implemented in DL-POLY 2.20

were used to integrate the equations of motion with a 0.75 fs time

step. Both thermostat and barostat coupling times were set to 1.0

ps. Each system underwent a 7.5 ps equilibration trajectory, during

which, direct velocity scaling of the temperature was performed at

every 5 steps. After equilibration, trajectories were integrated for

75 ps, during which, ensemble averages were obtained.

2.3. Isochoric Uniaxial Deformation Simulations. Uniaxial

deformation of the R-RDX structure is simulated by manually

deforming the system and equilibrating it at constant temperature and volume using the NVT ensemble. The forces calculated

in these simulations can be used to determine the effective stress.

Then, assuming an orthotropic elastic form29 for the stress�strain

relationship, the effective elastic constants can be computed. In

contrast, NST calculations described in section 2.2 used the reverse

procedure; the resulting strain was determined based on an applied

set of forces corresponding to a prescribed stress. We performed

two sets of uniaxial calculations, the first using small increments of

0.1% up to a total of 1% and the second using larger increments of

1% up to a total of 10%. These calculations were repeated for each

of the three crystallographic axes. An analogous procedure was

performed for three unique shear planes. In each set, we started

with the thermalized structures of R-RDX at P = 0 GPa and of γRDX at P = 5.2 GPa. At each strain increment, the system was

thermalized for 75 ps. The computed elastic constants are given in

Table 1.

At each deformation increment, we map all atom coordinates

into a scaled volume using the 3 � 3 deformation gradient, which

is given in terms of the lattice vectors by29

F 1⁄4 hh�1

o ð2Þ

where ho and h are tensors containing columnwise lattice vectors.

The reference lattice vectors, ho, are from the NST R-RDX

simulation at P = 0 and T = 300 K. Then, specifying the deformed

lattice vectors h uniquely determines the nine terms in F in eq 2.

Because RDX is orthorhombic, the Cartesian coordinates, ei

,

align with the crystal axes with a = ae1, b = be2, and c = ce3. The

reference locations of the molecule centers of mass xo are

mapped to the deformed unit system by x = Fxo. This system

is then thermalized using the earlier described procedure. The

Lagrange strain, Eij, from Fij is given by29,30

E 1⁄4 1

2

ðFTF � IÞ 1⁄4 1

2

ðh�T

o hThh�1o � IÞ ð3Þ

where I is the identity.

Under small strains, RDX single-crystal elasticity has been

shown to be sensitive to the direction of loading.8,10 In refs 21, 22,

and 25, the material response below the phase transition pressure

is sensitive to the crystal direction in which the planar shock

travels. This is because the elastic material response is anisotropic

as a result of the orthorhombic unit cell and intermolecular

interactions. The material response to large deformation leads to

molecular and lattice distortions that are also dependent on the

crystal axis being loaded. By applying varying levels of uniaxial

deformation (a nonhydrostatic deformation and state of stress),

measurements of the anisotropy and its connection to underlying

crystal orientation can be determined.

We next simulate R-RDX under large uniaxial deformation. In

simulations of kinetic mechanisms, it is often appropriate to seed

the model with a means to disrupt homogeneity and therefore

instigate the kinetic response. A seeded lattice defect, for

instance, disrupts the otherwise homogeneous crystal by introducing nontrivial amounts of local shear. The resulting simulation is generally insensitive to the magnitude of the stress from

the seed, and it is therefore sufficient to perform simulations on a

slab geometry (external surfaces) or with finite void volume

(internal surfaces).31 Experiments such as those in refs 18, 21, 22,

and 25 in ostensibly single crystals can still promote the kinetic

mechanisms of interest due to significant surface effects or lattice

defects that introduce local shear and escape identification or

Table 1. RDX Elastic Constants and Third Order BMEOS

Parameters

R-RDX γ-RDX

this worka experiment this workb

Orthotropic Elastic Constants

C11 (GPa) 25.0 25.6c 80.3

C22 (GPa) 23.8 21.3c 67.0

C33 (GPa) 23.4 19.0c 57.9

C44 (GPa) 3.1 5.4c 11.9

C55 (GPa) 7.7 4.3c 16.3

C66 (GPa) 5.2 7.3c 13.4

C23 (GPa) 8.8 6.4c 43.9

C31 (GPa) 7.6 5.7c 37.0

C12 (GPa) 10.6 8.7c 37.8

Third-Order BMEOS

Bo (GPa) 13.0 12.1d

Bo

0 9.2 8.6d

Vo (Å3) 1634 1641d

a R-RDX constants at reference condition. b γ-RDX constants at P = 5.2

GPa and T = 300 K. c

Sun et al.10 d Olinger et al. P = 0�3.95 GPa and T =

293 K.15
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The Journal of Physical Chemistry B ARTICLE

characterization. The direct experimental investigation of shearinduced kinetic structural response of RDX can be accomplished

by exploiting such surface effects such as those through

indentation.32 Simulations have also investigated RDX in the

context of directional loading of the crystal.1421,33 This also

produces local crystal shear that can promote bulk kinetic effects

such as phase transitions. Common in these approaches are

imposed or generated nontrivial components to the deviatoric

stress tensor. Because the deviatoric stress tensor is defined as

s = σ1/3 Tr(σ),29 it is possible to load the crystal in a

nonhydrostatic but homogeneous sense and produce the simulated effect of a seeded model.

3. RESULTS AND DISCUSSIONS

3.1. NST Pressurization Simulations. We use NST-MD

simulations to study quasi-static pressurization and depressurization of the bulk crystal. Under these conditions, the Gibbs energy

is conserved and used to investigate the crystal phases before and

after the transition. The change in Gibbs energy relative to the

reference state, ΔEGibbs(P) = EGibbs(P)  EGibbs(P=0), versus

the mean strain is shown in Figure 2 for the R-RDX pressurization (filled circles) and γ-RDX depressurization (open diamonds). The mean strain is the average of the diagonal or volumetric

strain components in eq 3, with ho given by the undeformed Rlattice vectors at the reference state. A negative mean strain

indicates compression from the undeformed R-lattice. Under

quasi-static pressurization of R-RDX, the R-structure is maintained up to roughly 10 GPa, the highest pressure value that we

tested (see Figure S5 of the Supporting Information for RDF and

wag angle distributions reflecting the absence of phase transition). Re-equilibrating the crystal at a higher temperature (350

K) and pressures up to 10 GPa yielded no observable departure

from the R-structure. However, starting with the crystal in the γpolymorph at 5.2 GPa, the γ-phase equilibrates at 300 K and

remains stable. Upon release of pressure, the crystal converts to

the R-polymorph, as is evident from the Gibbs energy in Figure 2

and structural data (wag angles and RDF) presented in Figure S6

of the Supporting Information. The phase transition occurs at a

pressure value of 2.1 GPa and a corresponding mean strain of

em = 0.034 (V/Vo = 0.90). At high pressure, the γ-RDX crystal

is correctly shown to be lower in Gibb’s energy than highly

compressed R-RDX (Figure 2 inset).

The computed value of 2.1 GPa is in agreement with the

results of Bedrov et al.,23 where the γ-phase was obtained

through a Hugoniotstat approach using a uniaxially applied

pressure on the (001) plane of ∼2.0 GPa. Our results are

qualified by the use of a depressurization approach for evoking

the reverse γ f R phase transition. The transition pressure is also

in qualitative agreement with the results of Ramos et al.22 and

Cawkwell et al.,21 where the abrupt changes to the elastic plastic

response were computed to occur in the range of 1.02.25 GPa.

In the present calculations, the appearance of the R-phase from

the γ-phase equilibration at discrete pressures was sudden and

categorical across the entire simulation domain, though we have

not investigated the influence of the small simulation size.

However, the sudden change in lattice properties implies changes

in the slip plane thickness, dislocation mobility, and shear moduli

to which the plastic slip is sensitive. The phase transition that

evokes changes in basic lattice parameters is therefore likely to be

accompanied by changes in the plastic response of an RDX bulk

crystal, particularly one undergoing slip as the transition pressure

is reached.

These results indicate the following. First, no γ-structure can

be found under equilibration at pressures below a critical value

(2.1 GPa), which indicates that the R-phase is indeed the

reference state polymorph using the SB potential. Second, the

transition is not observable with increasing hydrostatic pressure

Figure 2. Change in Gibbs energy, ΔEGibbs = EGibbs(P)  EGibbs(P=0),

versus the mean strain relative to the R-RDX configuration at P = 0. The

inset shows the region near the transition with pressure values labeled in

GPa units.

Figure 3. Volume ratio (a) and lattice constants (b) versus pressure.

Solid lines in (a) are third-order BMEOS fit. Triangles denote experiments for R-RDX15 and γ-RDX,11 and circles and diamonds denote the

present results. Solid and open symbols denote R- and γ-RDX phases,

respectively.
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The Journal of Physical Chemistry B ARTICLE

of pristine R-RDX. Such idealized simulations may preclude the

transition to the γ-phase at higher hydrostatic pressures using the

SB potential. This is particularly evident in the fact that uniform

depressurization of the γ-phase indeed yields the R-phase. To

investigate the possibility that simulated hydrostatic compression

of an ideal crystal and the smallness of the simulation might

artificially preclude the appearance of the γ-polymorph, we

explored uniaxial deformation of the RDX model, as will be

detailed in a later section.

Basic equilibrium properties were computed using the NST

ensemble. The pressure-dependent lattice constants and volume

for the R- and γ-RDX crystals were found to be in very close

agreement with available experimental data,6,11 as shown in

Figure 3 and presented in Tables S1 and S2 in the Supporting

Information. The R-RDX PV data for simulation (circles) and

experiment (triangles) shown in Figure 3a are fitted to the thirdorder Birch�Murnaghan equation of state (BMEOS)34 by leastsquares fitting of the bulk modulus and its derivative at zero

pressure, Bo and Bo

0

, respectively. The third-order BMEOS

values for Vo, Bo, and Bo

0 fitted to R-RDX data points are given

in Table 1; the calculated Bo underpredicts the available experimental value by 7%. Table 1 also contains values from experiments of ref 10 and the extensive comparisons summarized

therein.

Experiments (triangles in Figure 3) have shown that under

quasi-static pressurization, the RDX crystal undergoes the R f γ

transition at P = 3.9 GPa, evidenced by the abrupt change in

volume and lattice constants.11,15 The simulation data (circles

and diamonds in Figure 3), under depressurization using the SB

potential, also show an abrupt change in volume and lattice

constants after the transition; however, this occurs at a lower

pressure, P = 2.1 GPa.

The R f γ transition involves changes in the molecular

conformation as well as abrupt lattice changes. To clarify the

potential impact of these conformational changes, it is helpful to

detail the variation in each component of the potential energy

described by the SB model given in eq 1 before and after the

transition; these are shown in Figure 4. Compression of R-RDX

leads to increases in the van der Waals and dihedral contributions

to the potential energy while the other energy components

decrease. At the R f γ transition, the angle energy increases,

and the dihedral energy decreases to an almost zero net change

from the uncompressed R-RDX crystal. Upon increasing compression of the γ-RDX structure, the angle and dihedral energies

remain constant. The Coulomb energy decreases for increasing

compression of R-RDX, and after a slight drop across the R f γ

transition, it begins increasing for increasing compression of γRDX. These observations show that the pressure-dependent

contributions of the SB potential terms in eq 1 are dependent

on the crystal phase.

As one of the key molecular structural changes between the

two crystalline polymorphs involves the wag angle δ defined in

Figure 1a, we have generated relative distributions (shown in

Figure 5 and Figures S4 and S6 in the Supporting Information)

for values of the three wag angles in crystalline RDX during the

depressurization of the γ-polymorph. Each wag angle in the

figure identifies the three associated nitro groups individually at

each pressure, (a) N4�N7, (b) N5�N8, and (c) N6�N9 as

labeled in Figure 1a. The distributions are taken from 30

configurations of the 144 molecules of the supercell recorded

during the NST simulations. In these figures, the red curves

depict distributions of AAI molecules, and the blue curves depict

the distributions of AAE molecules. The bimodal distributions in

Figure 5a and c for pressures greater than 2.0 GPa (reflecting a

mix of AAE and AAI molecules) collapse to single-peaked

distributions at lower pressure; these correspond to the AAE

conformation of the R-polymorph. The γ-RDX AAE (blue)

distribution widths and peak values change very little through the

γ f R transition, suggesting that the transition does not

significantly affect the intermolecular interactions between

AAE molecules and their neighbors. The AAI molecules (red)

exhibit changes in peak location and distribution width for both

Figure 4. Change in energy components of the SB potential energy

given in eq 1, relative to their respective values at em = 0 versus the mean

strain. The van der Waals term is shown on a different scale in the upper

frame of the figure.

Figure 5. RDX wag angle distributions versus pressure during depressurization of γ-RDX for wag angles (a) N4�N7, (b) N5�N8, and (c)

N6�N9 labeled on the molecule in Figure 1a. Red and blue lines

indicate those molecules that have AAI and AAE structure in the γphase, respectively.
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The Journal of Physical Chemistry B ARTICLE

the N4N7 group, switching from equatorial (E) to intermediate (I), and the axial (A) N6N9 group, becoming similar to the

N5N8 distribution (see also Figure S4 in the Supporting

Information for additional distributions). We find that the

symmetry and width of the wag angle distribution are strongly

influenced by the dependence on the wag angle of the SB

potential in eq 1. A large curvature of the energy leads to

narrower distributions. The N4N7 wag angle is therefore softer

and experiences larger vibrational amplitudes than the N6N9

wag angle on the AAI molecule in the γ-structure. The asymmetry of the distributions is caused mainly by a corresponding

asymmetry in the intermolecular dispersionrepulsion interactions. The bonded part of the energy is an order of magnitude

smaller and plays a minor role. The behavior of the distributions

in Figure 5 with respect to increasing pressure is consistent with

the γ f R transition at P = 2.1 GPa. The locations of the peaks of

the wag angle distributions from this work and experiment6,11 are

presented in Table S3 in the Supporting Information.

Wireframe representations of the eight molecules for the

simulated R-RDX unit cells at P = 0 GPa are shown in

Figure 6ac, and those of the γ-RDX unit cell at P = 5.2 GPa

are in Figure 7df. The molecules that are AAI in γ-RDX are

shown in red, and those that remain AAE are shown in blue;

oxygen and hydrogen atoms are not shown for clarity. The γRDX molecules in red that undergo the conformation change

from AAE to AAI also undergo a rotation about an axis oriented

normal to the ring. This rotation may be aided by the effective

overall size of the AAI molecule relative to the size of the AAE

molecule, which allows them greater freedom to rotate. The

molecular centers of mass change little between R- and γ-RDX in

Figure 6af, and only slight differences are observed in the radial

Figure 6. Wireframe depictions of the RDX unit cell for T = 300 K at the

prescribed axial stress and strain shown to the left of each row. Strain

values are found relative to R-RDX at P = 0. For ease of viewing, H and O

atoms are not shown. R-RDX at P = 0 and γ-RDX at P = 5.2 GPa from

section 3.1 are shown in rows 1 (ac) and 2 (df), respectively. R-RDX

uniaxially compressed along either the a-, b-, or c-axis as described in

section 3.3 are shown in rows 3 (gi), 4 (jl), and 5 (mo),

respectively. Each column represents a view along the indicated crystallographic axis (labeled at the bottom of each column). Molecules that are

red are AAI in γ-RDX. Molecules that remain AAE are indicated in blue.

Figure 7. Change in (a) enthalpy, (b) total SB potential energy, (b) van

der Waals, (c) angle, and (d) dihedral energy terms of the SB potential

for each compression direction and γ-RDX NST depressurization. Zero

mean strain signifies the R-RDX reference configuration at reference

conditions.
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4384 dx.doi.org/10.1021/jp112042a |J. Phys. Chem. B 2011, 115, 4378–4386

The Journal of Physical Chemistry B ARTICLE

distribution functions in Figures S3, S5, and S6 of the Supporting

Information. The main features of the R- to γ-RDX phase

transition are therefore the conformation changes and rotations

of the AAI molecules about the axis normal to the ring.

Figure 6a�f will be used as reference structures to allow

comparison of uniaxial compression results in Figure 6g�o

discussed in section 3.3.

3.2. Hugoniot Results. NST-MD simulations were used to

explore the R-RDX structure at increasing pressure and temperature for states near the static shock conditions given by the

Hugoniot equation, using a method similar to Erpenbeck’s.28 To

satisfy the conservation of mass, momentum, and energy at the

shock front described by the Hugoniot equation, an increase in

pressure requires an increase in internal energy (kinetic and

potential). This work uses the R-RDX crystal at P = 0 and T = 300

K as the reference point, and the Hugoniot equation is satisfied at

higher pressures by increasing the temperature. At P = 5 GPa, the

Hugoniot equation results in a Hugoniot temperature of THG =

350 K, raising the potential and kinetic energy by 11.8 kcal/mol.

Isothermal compression from the reference state to P = 5 GPa

(presented in section 3.1) raised only the potential energy by 6.3

kcal/mol. The Hugoniot simulations do not show the R- to γRDX phase transition, despite the increased temperature. Consequently, we note that the computed Hugoniot curve in P(V)

space, shown by the þ symbols in Figure 3a, is almost identical to

the isothermal curve for R-RDX at T = 300 K, where the slight

deviation is caused by the increased temperature of the Hugoniot

simulation. Table S5 in the Supporting Information presents

tabulated PVT data for the R-RDX Hugoinot curve.

3.3. NVT Uniaxial Deformation Results. A series of uniaxial

deformations were performed first to determine the full set of

orthotropic elastic stiffness constants at small strains below 1%

from the initial R-RDX equilibrated structure at P = 0 GPa and γRDX equilibrated structure at P = 5.2 GPa. Below 1% strain, the

stress�strain response remains linear elastic. The computed

coefficients, Cij, are listed in Table 1 and are in reasonable

agreement with experiment.6 The bulk modulus of R-RDX

calculated from Cij is B = 14.0 GPa, in close agreement with Bo

= 13.0 GPa obtained from the third-order BMEOS fit of

hydrostatic compression simulations from section 3.1, listed in

Table 1.

Using larger uniaxial deformations, we also investigated the

dependence of the loading direction on the appearance of new

phases down to an effective compressed volume ratio of V/Vo =

0.9. The uniaxially deformed unit cells along each of the three

crystal axes are depicted in Figure 6g�o. The strain-dependent

radial distribution functions and wag angle distributions are

presented in Figures S7�S9 in the Supporting Information.

These three cases of uniaxial deformation lead to three different

crystal responses where compression of the c-axis results in the Rto γ- transition, compression of the b-axis results in a phase

transition similar to the stacking faults observed in other simulation work,21,22 and no transitions are observed for compression

of the a-axis.

For compression along the c-axis, the deformed structures are

shown in Figure 6g�i for c/co = 0.9. They are visually indistinguishable from the γ-RDX structures at P = 5.2 GPa in

Figure 6d�f. Thus, unlike the hydrostatic pressurization calculations, uniaxial compression evokes the R- to γ-transition. This is

consistent with the findings by Bedrov et al.23 using the SB

potential and uniaxial pressure Hugoniotstat applied to the c-axis.

Figure 6j�l for b-axis compression at b/bo = 0.9 shows a large

shift in the location of the two center molecules, most easily

observed near the center of Figure 6l. The molecules also change

to the AEE conformation, confirmed by the wag angle distributions presented in Figure S8c�e in the Supporting Information.

Cawkwell et al.21 and Ramos et al.22 observed a similar molecule

conformation and stacking change for simulations of weak shock

loading at P = 1.3 GPa on the (111) and (021) planes. In their

simulations, this phase only appeared in partial dislocation loops,

which generated stacking faults on the (001) slip plane.21

For compression along the a-axis, Figure 6m�o shows the

configurations at a/ao = 0.9. There is no noticeable change in

molecular conformation or packing relative to the R-RDX

configuration at P = 0 (see Figure 6a�c). Uniaxial stress loading

along the a-axis or (100) plane of R-RDX is commonly studied in

the literature because it leads to no resolved shear stress on

experimentally known slip systems,24,32 making this direction

sensitive to shock loading according to the steric hindrance

model.2 The steric hindrance model predicts that sensitivity in

energetic molecular crystals is related to the absence of available

slip systems for dislocation-based deformation and that deformation in these directions leads to large molecular distortions and

the breaking of bonds. Experimentally, uniaxial shocks above P =

3.9 GPa applied to the a-axis result in the γ-RDX phase after a

100 ns so-called incubation period,18 but reproducing these

experimental conditions would be outside of the spatial and

temporal range of atomistic simulations. Simulated shocks on the

a-axis by Cawkwell et al.21 and Bedrov et al.,23 both using the SB

potential, showed no noticeable conformation changes for weak

shocks below 5.9 GPa. For stronger shocks above 9 GPa, they

both show the development of amorphous shear bands oriented

along planes of maximum shear stress irrespective of the

crystal axes.

We next examine the effects of crystal orientation and material

anisotropy. Earlier in this section and elsewhere,21�24 it was

found that RDX response is sensitive to the direction of loading

relative to its native crystal axes. We now attempt to clarify the

dependence of the transition on the direction of loading through

the analysis of steric effects under uniform and uniaxial strain.

Figure 7a�e shows the change in different energy components, ΔE, relative to their corresponding value for R-RDX at

reference conditions as functions of mean strain. Depicted are

the three uniaxial compression simulations, a (green), b (red),

and c(blue), and the NST depressurization simulation of γ-RDX

(purple). It is apparent that uniaxial deformation causes phase

transitions to occur at smaller effective volumetric strains and

therefore lower effective pressures. The effective mean strains at

which the transitions occur are indicated by the vertical colored

lines traversing all plots. The purple line represents the NST γ f

R transition at em = �0.034 (V/Vo = 0.90), the red line is the bcompression transition at em = �0.026 (V/Vo = 0.92), and the

blue line is the R f γ transition for c-compression at

em = �0.017 (V/Vo = 0.95). Once again, no transition is observed

for compression along the a-axis. The transition due to uniaxial

loading in the c-direction occurs earliest with c-axis stress �1.4

GPa, followed by the transition due to loading in the b-direction

at b-axis stress �2.58 GPa. No transition is observed for loading

in the a-direction.

The change in enthalpy, ΔH, gives a measure of the system

internal energy plus the work put into the system due to the

deformation. For a crystalline material at room temperature, the

Gibb’s energy is approximately equal to the enthalpy. In the NST

γ-RDX simulation, the potential energy contributes 69% to the
109

4385 dx.doi.org/10.1021/jp112042a |J. Phys. Chem. B 2011, 115, 4378–4386

The Journal of Physical Chemistry B ARTICLE

total system enthalpy at the maximum compression of P = 5.2

GPa but only 5% to the change in enthalpy from that state during

depressurization. The largest contribution to ΔH is external

work (PV). The abrupt change in ΔH for the NST simulation is

due to the abrupt volume change shown in Figure 3a and the

abrupt change in ΔH for NVT b-compression is due to an abrupt

change in stresses, shown in Figure 8b.

The principal noteworthy observation is the sensitivity of

molecular deformation, reflected by the angle and dihedral

energies in Figure 7d and e, to the crystal axis being compressed.

Compressions of the a- and c-axes of R-RDX result in decreasing

angle energy and increasing dihedral energy, as is seen in the NST

case. A transition is observed for c-compression strained beyond

em = 0.017, where the change in angle energy diminishes to

zero, as is observed for NST compression of γ-RDX. Compression of R-RDX along the b-axis follows opposite trends with an

increase in angle energy and decrease in dihedral energy.

The stress components that develop along each crystal axis

(σa,b,c along a,b,c-axes, respectively) are also the principal stresses

and are shown in Figure 8ac. The maximum shear stress, shown

by the dashed line, is equal to the minimum (σmin) minus the

maximum (σmax) principal stresses. The maximum shear stress is

an indicator of plastic material response when its magnitude

exceeds the known yield strength of the material and acts on a

plane oriented 45 to the directions of σmax and σmin.

Remarkably, the b- and c-compressed transitions occur consistently for (σmin  σmax) in the range of 1.0 to 1.25 GPa.

For b-compression, the principal stresses change abruptly after

the transition with σc (green line) decreasing and both σa (blue

line) and σb (red line) increasing. Prior to the phase change, the

maximum principal stress is oriented along the c-axis (σmax = σc),

which becomes oriented along the a-axis (σmax = σa) in the new

phase. The change in crystal phase leads to a reorientation of σmax

and in effect a reorientation of the plane of maximum shear stress.

Thus, new slip systems may become active, resulting in plastic

slip on new planes that block the R-RDX slip planes. Similar

effects were observed in the simulated plastic response of RDX

under oriented shock loading.21,22

4. SUMMARY AND CONCLUSIONS

Details at the atomic level provided by MD simulations are

useful in providing additional chemical and physical interpretations to experimental data. In this work, we studied the applicability of the SB potential to model R- and γ-phases of RDX

under various load conditions using simulations and comparing

to available experimental data. The simulations were used to

provide details of the roles of molecular conformation and

mechanical deformation on the phase transitions. The small

simulation cell and time scale limits the consideration to homogeneous phase transitions under hydrostatic and uniaxially

deformed conditions. The uniaxial conditions mimic the effect

of local shear that can occur in larger simulations. The nonreactive SB potential was appropriate for this work because

simulation conditions were limited to those processes preceding

chemical events.

The homogeneous crystal structures obtained were in agreement with experimental R- and γ-RDX structures, and the

orientation dependence of the phase transitions agree well with

other simulation work using the SB potential.2124 We showed

that under increasing hydrostatic loads, the R- to γ-transition

could not be observed for the time and space scales of the current

simulations. However, the reverse or hydrostatic depressurization transition was found at a critical computed pressure of

2.1 GPa.

Under uniaxial deformation, the transitions were shown to

depend on the crystal axis being strained. Two different phase

transitions were shown to occur at mean strains equal to

em = 0.017 and 0.026, both of which are less than the strain

em = 0.034 required to observe the transition under hydrostatic

stress. Compressive deformation along the c-axis reproduced the

γ-RDX configuration for σc = 1.5 GPa. Compression along the

b-axis did not result in the γ-RDX structure but instead led to a

structure similar to that identified in simulations by Cawkwell

et al.21 and Ramos et al.22 as a stacking fault. Deforming along the

different crystal axes had markedly different effects on the

molecule conformation, reflected in the bonded SB potential

terms where loading on the b-axis decreased the dihedral energy,

whereas it increased for loading along the c-axis. These results

suggest that the molecule conformation in RDX crystals plays a

significant role in the stability and accessibility of high-pressure

phases.

’ASSOCIATED CONTENT

bS Supporting Information. Pressure-dependent lattice

constants for the R- and γ-RDX simulations are given in Figure

S1 and Tables S1 and S2. Figure S2 contains the graphical

depictions of the contents of the R- and γ-RDX unit cells with the

asymmetric units overlaid on the corresponding space group

diagrams. Figure S3 presents a comparison of R- and γ-RDX

Figure 8. Stress versus mean strain for uniaxial compression along the

(a) a-axis, (b) b-axis, and (c) c-axis. Principal stress components are

labeled by the crystal axis (e.g., σa is the principal stress oriented along

the a-axis of R-RDX). Maximum shear stress (σmin  σmax) is shown by

the dashed line.
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4386 dx.doi.org/10.1021/jp112042a |J. Phys. Chem. B 2011, 115, 4378–4386

The Journal of Physical Chemistry B ARTICLE

radial distribution functions to the experimental data. Figure S4

presents wag angle distributions for each of the eight molecules in

the unit cells for the R- and γ-RDX simulations, and Table S3

gives locations of the peak. Figures S5S9 present RDF and wag

angle distributions for each series of NST pressurizations and

NVT uniaxial strain simulations. Figures S10 and S11 give

temperature-dependent lattice constants and strain used to

determine the coefficients of thermal expansion presented in

Table S4. Table S5 presents PVT points for the calculated

Hugoniot curve. This material is available free of charge via the

Internet at http://pubs.acs.org.

’ACKNOWLEDGMENT

Support is gratefully acknowledged from the U.S. Army

Research Laboratory STEP program and the DoD Higher

Performance Computing Modernization Office through the

Multiscale Reactive Modeling of Insensitive Munitions Software

Applications Institute. Computing support was provided by the

DoD Supercomputing Resource Center located at the Army

Research Laboratory.

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111

Research@

Elastic-Plastic Behavior of Cyclotrimethylene 

Trinitramine Single Crystals Under Spherical 

Indentation: Modeling and Simulation

J. D. Clayton and R. Becker

U.S. Army Research Laboratory, Aberdeen Proving Ground, MD

ABSTRACT

A nonlinear anisotropic elastic-plastic model is developed for single crystals 

of the energetic material cyclotrimethylene trinitramine (RDX). Numerical 

simulations of spherical indentation on oriented single crystals are performed; 

predictions are compared with experimental data and observations from 

the literature. Model predictions for elastic response and initial yield using 

elastic constants obtained from resonant ultrasound spectroscopy agree with 

experimental data; predicted forces using constants obtained from Brillouin 

scattering tend to exceed experimental data. Influences of elastic anisotropy 

and elastic nonlinearity are significant. Predicted slip system activity is in 

reasonable agreement with that deduced from experimental surface profiles 

when a uniform strength of G/20 is assigned to all six slip systems, with G 

an effective elastic shear modulus. Predicted indentation forces in the postyield regime exceed those observed in experiments, suggesting that surface 

and possibly subsurface fractures may contribute to a loss of stiffness in 

experiments at larger indentation depths

Journal of Applied Physics 111, 063512 (2012)
112

Elastic-plastic behavior of cyclotrimethylene trinitramine single crystals

under spherical indentation: Modeling and simulation

J. D. Claytona) and R. Becker

Impact Physics RDRL-WMP-C, U.S. Army Research Laboratory, Aberdeen Proving Ground,

Maryland 21005–5066, USA

(Received 5 January 2012; accepted 13 February 2012; published online 21 March 2012)

A nonlinear anisotropic elastic-plastic model is developed for single crystals of the energetic

material cyclotrimethylene trinitramine (RDX). Numerical simulations of spherical indentation

on oriented single crystals are performed; predictions are compared with experimental data and

observations from the literature. Model predictions for elastic response and initial yield using

elastic constants obtained from resonant ultrasound spectroscopy agree with experimental data;

predicted forces using constants obtained from Brillouin scattering tend to exceed experimental

data. Influences of elastic anisotropy and elastic nonlinearity are significant. Predicted slip

system activity is in reasonable agreement with that deduced from experimental surface profiles

when a uniform strength of G/20 is assigned to all six slip systems, with G an effective elastic

shear modulus. Predicted indentation forces in the post-yield regime exceed those observed in

experiments, suggesting that surface and possibly subsurface fractures may contribute to a loss

of stiffness in experiments at larger indentation depths. [http://dx.doi.org/10.1063/1.3695392]

I. INTRODUCTION

Defects in energetic materials are thought to affect their

initiation sensitivity. Stresses concentrate in the vicinity of

cracks, pores, or lattice defects, which in turn can affect initiation of reactions associated with burning or detonation. In

single crystals, availability of slip systems associated with

mobile dislocations may lower peak stresses and decrease

sensitivity to shock initiation.1

The focus of the present work is the mechanical behavior of the energetic material cyclotrimethylene trinitramine

(C3H6N6O6), referred to as RDX (Research Development

eXplosive). Single crystals of RDX belong to an orthorhombic space group with eight molecules per unit cell. Dislocations in RDX have been characterized using etch pit2 and x

ray topographic3,4 techniques. Likely slip systems in RDX

have been suggested from analysis of anisotropic hardness

profiles5 and indentation experiments.6,7 The latter experiments6,7 also provide an estimate of the critical resolved

shear stress associated with slip initiation, thought to be on

the order of the theoretical strength (i.e., � G/10-G/20, with

G a representative elastic shear modulus), which corresponds

to homogeneous dislocation nucleation. Inelastic behavior of

RDX crystals has also been probed using shock experiments1,8 and molecular dynamics simulations.8,9

Continuum crystal plasticity theory permits predictive

mesoscale modeling of materials’ behavior at length scales

larger than that feasible using molecular models, but with

greater resolution than that afforded by macroscopic elasticplastic models that omit anisotropy and slip system activity.

Grain interactions can be studied in direct numerical simulations via finite element models, wherein each crystal of a

polycrystal is resolved geometrically. Crystal plasticity

models have been implemented elsewhere to study shock

loading of energetic materials cyclotetramethylene tetranitramine (HMX)10 and pentaerythritol tetranitrate (PETN).11,12

One aim of the present work is development and implementation of a crystal plasticity model for RDX, a particular

model which, to the authors’ knowledge, has not been

published elsewhere.

The single crystal elastic-plastic model developed here

extends a previous model13 for cubic crystals loaded to possibly high pressures. Here, the model is applied to crystals

with orthorhombic symmetry characteristic of RDX. Anisotropic elastic constants and pressure-dependent compressibility are considered from experimental literature.14–16 Six slip

systems (from four different families of systems) are implemented following analysis of indentation loading profiles

and surface impressions.5,6 The model is applied to study indentation, with a spherical indenter of (001), (021), and

(210) faces of single crystals of RDX.

This paper is organized as follows: Constitutive theory

and material properties are described in Sec. II. Indentation

simulations are reported in Sec. III. Conclusions follow in

Sec. IV. Notation of continuum mechanics is used, e.g., boldface type for vectors and tensors all referred to fixed Cartesian

coordinates. Summation applies over repeated indices.

II. THEORY

A. Single crystal model

Let x 1⁄4 vð Þ X; t denote the motion of material points of

the body. The deformation gradient is

rv 1⁄4 F 1⁄4 FEFP; (1)

where r �ð Þ denotes the material gradient (i.e., FaA 1⁄4 rAva

1⁄4 @xa=@XAÞ; FE denotes thermoelastic deformation of the

a)Author to whom correspondence should be addressed. Electronic mail:

john.d.clayton1.civ@mail.mil.

0021-8979/2012/111(6)/063512/9/$30.00 111, 063512-1

JOURNAL OF APPLIED PHYSICS 111, 063512 (2012)
113

crystal lattice, and FP represents plastic deformation due to

slip. The spatial velocity gradient is

~_

FF

~�1 1⁄4 F_ EFE�1 þ FEFPFE�1

; (2)

where the plastic velocity gradient associated with slip rates,

c_

k, reference slip directions, sk

0, and reference slip plane normals, mk

0, on slip systems, k, is

LP 1⁄4 F_ PFP�1 1⁄4 X

k

c_

k

s

k

0 � mk0: (3)

Slip system geometry is pushed forward to the current configuration via sk 1⁄4 FEsk

0 and mk 1⁄4 FE�Tmk0, where ð Þ� T

denotes the transpose. From orthonormality of slip directions

and slip plane normals, plastic deformation is isochoric;

thus, J 1⁄4 detF 1⁄4 detFE

> 0. By the polar decomposition theorem, let FE 1⁄4 RE

UE. A logarithmic thermoelastic strain measure is defined as e 1⁄4 lnUE and is split into deviatoric and

volumetric parts as

e 1⁄4 e

0 þ ð Þ 1=3 �e 1; �e 1⁄4 tre 1⁄4 ln J; (4)

where 1 is the second-order unit tensor, and the trace of a

second-order tensor is tr(�). Let r denote the usual Cauchy

stress tensor; stress in an unrotated coordinate system is

S 1⁄4 RE–1

rRE.

Only the isothermal case is considered, an appropriate

assumption for applications of the model discussed in Sec.

III. Let the unrotated stress be split into deviatoric and

hydrostatic parts,

S 1⁄4 S0 þ S�1; S� 1⁄4 ð Þ 1=3 trS 1⁄4 ð Þ 1=3 trr 1⁄4 �p; (5)

with p the Cauchy pressure. The following operator extracts

the deviatoric part of a second-order tensor,

I0

1⁄4I � ð Þ 1=3 1 � 1; I0

ABCD 1⁄4 dACdBD � ð Þ 1=3 dABdCD:

(6)

Constitutive equations for deviatoric stress and pressure are13

S0 1⁄4 I0:C:e

0

þð Þ 1=3 I0:C:1 � ��e; (7)

p 1⁄4 B0=B0 ð Þ exp �B0 1⁄2 ð Þ� �e 1� � ð Þ 1=3 e

0 : C : 1: (8)

Here, B0 and B0 are the reference bulk modulus and pressure

derivative of the bulk modulus and C is the tensor of

second-order elastic constants referred to the (unrotated)

crystal frame. The colon denotes contraction over two pairs

of indices, e.g., C:e0 ð ÞAB 1⁄4 CABCDe0

CD. Pressure dependence

of shear elastic coefficients, implemented elsewhere for

cubic crystals,13 is omitted in Eq. (8) because of limited experimental data for the material of present interest.

The flow rule for slip is13

sk 1⁄4 r : s

k � mk � � 1⁄4 sk

0 c_

k �

� �� þ n � �=c_0

� �m

sgn c_

k � �: (9)

The resolved shear stress on slip system k is sk. Material parameters are initial and constant slip strength, sk

0, for each

slip system, reference strain rate, c_0, and rate sensitivity, m.

Constant n � c_0 provides a finite strength at zero strain rate.

B. RDX

Physical properties of RDX single crystals are listed in

Table I. The description applies to the a phase, the stable

polymorph for pressures under �3.8 GPa and temperatures

under � 480K.

Elastic properties are listed in Table II. Isentropic

second-order elastic constants15,16 and bulk modulus have

been converted to isothermal values at 295 K via the usual

thermodynamic formulae,17 incorporating anisotropic thermal expansion18 and specific heat.19 Voigt’s notation is

used: CABCD $ Cab, where Greek indices 1, 2, ..., 6. Voigt

(GV) and Reuss (GR) bounds17 on the effective shear modulus are also listed. Differences between Voigt and Reuss

bounds for bulk modulus B0, on the order 1–3%, are considered insignificant.

As is evident from Table II, reported values of secondorder elastic constants can vary substantially. Values

obtained using resonant ultrasonic (RUS) methods15 listed in

Table II are in reasonably close agreement with those

reported by other researchers using the same method.20 Values obtained using Brillouin scattering16 listed in Table II

are notably different, with particularly larger bulk stiffness

and shear stiffness in certain directions. Values obtained

using a third technique, impulsive stimulated thermal scattering,21 are similar to those obtained using RUS. Possible reasons for differences among measurements of elastic

constants of organic molecular crystals are discussed elsewhere.22 Values predicted using empirical atomic models23,24 also exhibit differences from those obtained in

experiments, though these predicted values tend to align

TABLE I. Structural and physical properties of RDX single crystals

(ambient).

Property Value Ref.

Space group Pbca 5

Crystal structure Orthorhombic

Lattice parameters [nm] a 1⁄4 1.3182 5

b 1⁄4 1.1574

c 1⁄4 1.0709

Mass density [g/cm3

] 1.806 16

TABLE II. Isothermal second-order elastic constants of RDX (converted

from ambient isentropic values).

Property Value15 Value16

C11 [GPa] 24.56 36.48

C22 [GPa] 18.85 24.49

C33 [GPa] 17.33 20.78

C12 [GPa] 7.61 0.90

C13 [GPa] 5.30 1.26

C23 [GPa] 5.24 8.16

C44 [GPa] 5.15 11.99

C55 [GPa] 4.06 2.72

C66 [GPa] 6.90 7.68

B0 [GPa] 10.5 11.2

GV [GPa] 6.06 9.26

GR [GPa] 5.72 6.40

063512-2 J. D. Clayton and R. Becker J. Appl. Phys. 111, 063512 (2012)
114

more closely with those from RUS than from Brillouin

scattering.

In this work, the two sets of elastic constants15,16 are

used, because these are the softest and stiffest reported experimental measurements, respectively; results obtained

using these two sets might be expected to bound the actual

response. Results obtained using these particular RUS constants15 are very close to those that would be obtained using

similar values.20,21

Pressure and temperature dependencies of second-order

elastic coefficients have been calculated using molecular dynamics;23 however, these predicted values have not been

validated using experiments, and some discrepancies exist

among calculated second-order elastic constants at room

temperature and experimental values.16 For this reason, in

the present study, the nonlinear elastic model only incorporates pressure dependence of the compressibility obtained

experimentally14 and not that of all elastic coefficients, specifically, B0 1⁄4 6.95.

Crystal plasticity theories incorporating higher-order

elastic constants have also been developed12,25 and might

offer an improved description of effects of volumetric and

shear deformations on tangent elastic moduli; such an

approach is not pursued here, because higher-order elastic

constants (e.g., 20 independent third-order constants for an

orthorhombic crystal17) are unknown for RDX. In a crystal

plasticity model of PETN,12 a tetragonal crystal, the Cauchy

relations were used to estimate unknown third-order elastic

constants.

Potential slip systems in RDX—as identified from hardness versus orientation profiles,5 indentation force versus

depth data,6,7 and residual surface impressions from indentation6

—are listed in Table III. Slip system geometry (Fig. 1)

is referred to a Cartesian system with axes (X1, X2, X3) parallel to lattice vectors (a, b, c). Listed initial slip system

strengths are upper bounds estimated from analysis of load

excursion data using the analytical Hertzian solution for frictionless spherical indentation into a semi-infinite, linear elastic, isotropic material.6 Other systems may become active

(and those listed may become inactive) for loading regimes

involving very different pressures, temperatures, and/or

strain rates; e.g., molecular dynamics simulations8 suggest

that partial dislocation loops may glide on (001)[010] during

shock loading at pressure in excess of � 1 GPa.

Strengths for various families of systems are varied

parametrically between physically reasonable bounds on the

order of the theoretical strength,6,7 where G 1⁄4 GR1⁄4 6.4

GPa.16 Because strengths of families of slip systems are not

known precisely a priori from experiments, shear strengths,

sk

0, are varied parametrically over the range listed in Table

IV. This range is physically descriptive of homogeneous

nucleation of glissile dislocation lines and loops.6,7,17 Other

parameters, which provide a nearly rate-independent

response, are also listed in Table IV.

III. INDENTATION

A. Boundary value problem

The constitutive model of Sec. II is implemented in the

ALE3D multi-physics code. Simulations of indentation are

performed using an implicit solver for static equilibrium.

The problem geometry mimics previous experimental

studies.6,7 A spherical diamond indenter of radius R 1⁄4 1.482

lm is used to indent a flat surface of a single crystal of RDX

of variable lattice orientation. Diamond is represented as an

isotropic nonlinear elastic material with B0 1⁄4 443 GPa,

G 1⁄4 538 GPa, and B0 1⁄4 4.0.

The substrate is represented by a right circular cylinder.26 The cylinder is assigned a height and radius of 2R; further increases in dimensions of the cylinder did not affect

results of interest. The indenter is modeled as a half-sphere.

Each body is discretized using eight-node hexahedral elements with selective reduced integration. The mesh of the

substrate is highly refined in the vicinity of contact beneath

the indenter, where stress fields are inhomogeneous, and the TABLE III. Slip systems in RDX single crystals (indentation).

System k Miller indices m0 s0

Max. strength

[GPa] Ref.

1 (021)[100] (0, 0.880, 0.475) [1, 0, 0] 0.885 5, 6

2 0ð Þ �21 [100] (0, –0.880, 0.475) [1, 0, 0]

3 (011)[100] (0, 0.679, 0.734) [1, 0, 0] 0.645 5, 6

4 0ð Þ �11 [100] (0, –0.679, 0.734) [1, 0, 0]

5 (010)[100] (0, 1, 0) [1, 0, 0] 0.885 6

6 (010)[001] (0, 1, 0) [0, 0, 1] 0.885 5, 6

FIG. 1. Slip systems in RDX (unit cell parameters not to scale).

TABLE IV. Crystal plasticity model parameters for RDX.

Property Value

sk

0 G=40 <

� sk

0 <

� G=10

m 5 � 10–3

c_0 10–2/s

n 10–7/s

063512-3 J. D. Clayton and R. Becker J. Appl. Phys. 111, 063512 (2012)
115

mesh coarsens progressively with distance from the initial

contact point. Simulations with smaller elements demonstrated that further increases in mesh refinement did not

affect results of interest.

During the loading phase, the upper face of the halfspherical indenter is assigned a constant (downward) velocity of D_ 1⁄4 10 nm/s in the laboratory X3-direction, leading to

strain rates similar to those of experiments.6 The lower face

of the cylinder is rigidly fixed, while the lateral sides (circumference) are traction-free. Indentation depth is denoted

by D; actual depth d of the tip of the sphere in contact with

the surface is monitored as an outcome of the solution. Only

for a rigid indenter would d 1⁄4 D. Indentation force P is the

sum of nodal forces along the upper face of the half sphere

acting in the direction of D, i.e., the sum of forces work conjugate to prescribed nodal velocities.

Contact between the indenter and substrate is assumed

frictionless, following previous studies that rely on analytical

solutions.6,7 Experimental measurements of dynamic friction

for RDX single crystals sliding on a glass substrate27 suggest

a friction coefficient on the order of unity for loads under 1 g

(�10 mN), wherein contact is characterized as elastic, and a

friction coefficient of 0.35 for loads in excess of 10 g,

wherein contact is characterized as plastic. The present indentation simulations consider a different geometry, smaller

system sizes (loads under 1 g), and slower sliding velocities

(on the order of 10 nm/s), so the reported experimental values for friction coefficients27 may not strictly apply here. In

additional simulations,28 it was found that differences in indentation force among cases invoking frictionless and sticking contact were insignificant. In some simulations,

unloading is also performed, whereby, after a peak depth is

attained, the upper face of the indenter is assigned an upward

velocity of D_ 1⁄4 10 nm/s until contact is released. If plastic

deformation has occurred, then some residual deformation

remains in the substrate upon unloading.

The present simulations enable direct quantification of

surface and subsurface slip on each system. In contrast, hardness or indentation experiments5–7 require substantial interpretation of data to deduce slip activity and do not provide a

quantitative measure of relative contributions of each slip

system to the overall strain field. In the aforementioned

experiments, visual observations of slip traces are restricted

to residual surface profiles, whereas simulations enable visualization of subsurface slip activity.

Simulations of indentation onto (001), (021), and (210)

planes are reported in Sec. III B. Let C(ijk) be the fourthorder matrix of elastic constants of the crystal oriented for

indentation into crystallographic plane (ijk). Let R(ijk) be the

corresponding rotation matrix. Then,

Cð Þ ijk

ABCD 1⁄4 Rð Þ ijk

AE Rð Þ ijk

BF Rð Þ ijk

CG Rð Þ ijk

DH CEFGH; (10)

TABLE V. Elastic stiffness and Young’s modulus of RDX oriented for indentation on planes (001), (021), and (210).

Plane [RUS15] Plane [Brillouin16]

Stiffness (001) (021) (210) (001) (021) (210)

C11 [GPa] 24.56 24.56 20.15 36.48 36.48 22.17

C22 [GPa] 18.85 16.78 17.33 24.49 24.94 20.78

C33 [GPa] 17.33 17.62 23.06 20.78 26.98 28.29

E11 [GPa] 20.85 20.85 16.72 36.40 36.40 19.12

E22 [GPa] 15.69 14.05 15.40 21.28 22.84 18.04

E33 [GPa] 15.40 14.19 19.31 18.04 25.55 23.61

FIG. 2. Indentation force vs applied displacement; model predictions obtained

using RUS (Ref. 15) and Brillouin (Ref. 16)

elastic constants compared to experiment

(Ref. 6): (a) indentation into (001); (b) indentation into (021); (c) indentation into

(210).

063512-4 J. D. Clayton and R. Becker J. Appl. Phys. 111, 063512 (2012)
116

where CEFGH corresponds to elastic constants of Table II. Values of elastic stiffness are listed in Table V. Perhaps most pertinent to indentation in the X3-direction, Cð Þ 210

33 > Cð Þ 02133 > Cð Þ 00133

for either set of constants. Note that Cð Þ ijk

33 from Brillouin

scattering16 always exceeds the corresponding value from

RUS.15 Components of compliance Sab are those of the inverse

of Cab. Young’s moduli are E111⁄4 1/S11, E221⁄4 1/S22, and

E331⁄4 1/S33. From Table V, Eð Þ ijk

aa from Brillouin scattering16

always exceeds the corresponding value from RUS15 for a 1⁄4 1,

2, and 3.

B. Model predictions

Predictions obtained using each set of elastic constants15,16 are compared directly for substrate orientations

(001), (021), and (210) in Fig. 2. In results marked as

“elastic”, slip is suppressed to permit assessment of the elasticity model in isolation and to permit deduction of the yield

point upon comparison with “elastic-plastic” results, wherein

the slip model of Sec. II is enabled. For the latter results

shown in Fig. 2, for each set of elastic constants, strength is

set to a constant value of sk

0 1⁄4 G/20 for all six slip systems.

Significantly closer agreement between model and

experiment6 is obtained from the elastic constants from

RUS15 for indentation onto (001) and (021) planes. The first

experimental data point in each figure corresponds to the

maximum depth at which the indentation process remains

elastically reversible, i.e., initiation of the first excursion

from a smooth force-displacement profile.6,7 Comparable accuracy is obtained from either set of elastic constants for indentation on (210) planes. The present simulations strongly

suggest that elastic constants obtained from RUS15,20 provide a more realistic representation of elastic stiffness during

nano-indentation than elastic constants obtained from Brillouin scattering,16 with the latter appearing too stiff. Results

shown in Fig. 2 invoke the geometrically nonlinear elastic

model with pressure-dependent bulk modulus; comparison

with results of additional calculations28 with B0 1⁄4 0 for RDX

demonstrated that the effect of nonlinear compressibility

becomes noticeable for D & 50 nm.

In simulations, the yield point (initiation of slip) can be

deduced as the indentation depth beyond which elastic and

elastic-plastic model predictions begin to differ. From Fig. 2,

the predicted yield point matches the experimental excursion

point reasonably well for each orientation when RUS elastic

constants15 are used in the model. In contrast, the indentation

depth at which yielding is predicted is premature for (001)

and (021) orientations when Brillouin constants16 are used.

This difference is presumably a result of attainment of larger

resolved shear stresses at a given depth of indentation when

stiffer constants are used.

Predicted forces exceed experimental values at larger indentation depths in each orientation and for both sets of elastic constants. Results corresponding to more compliant

elastic constants15 provide closer agreement to experimental

values than results corresponding to stiffer elastic constants.16 As noted elswhere,26,28 uncertainty in the true tip radius, R, of the indenter could lead to discrepancies between

simulations and experiments. Surface fractures and/or subsurface fractures could contribute to a loss of stiffness that

would be reflected only in the experimental data. RDX is

prone to cleavage fracture on planes (001), (010), (001),

(241), and 2ð Þ 41 .

2,4,6 Experimental data demonstrate nearly

horizontal steps in force versus displacement corresponding

to discrete shear discontinuities and/or fracture events that

are not readily resolved by a constant strength continuum

crystal plasticity model, such as the one formulated here. In

FIG. 3. Cumulative total slip contours for indentation to depth D 1⁄4 200 nm

using elastic constants from RUS (Ref. 15) and slip system strength G/20; a

slice along the centerline of the cylinder normal to the laboratory X1 axis is

shown: (a) indentation into (001)[X-axis normal to (100)]; (b) indentation

into (021)[X-axis normal to (100)]; (c) indentation into (210)[X-axis normal

to 2ð Þ 30 ].

063512-5 J. D. Clayton and R. Becker J. Appl. Phys. 111, 063512 (2012)
117

experiments,6 surface fractures are detected only at loads significantly exceeding yield excursion.

Details of slip system interactions and pressure dependence of shear strength are omitted in the present model. Indentation experiments have suggested the importance of

cross slip,6 and atomic modeling has noted that different slip

mechanisms may become important at high pressures.8

Incorporation of these effects into a more complex slip

model might provide closer agreement with experiment, e.g.,

if glide resistance were to decrease with pressure, since local

pressures under the indenter can achieve several GPa.28

Atomic modeling8,24 may provide insight into dependence of

slip resistance on pressure (and temperature, etc.) not available from experimental methods.

Model predictions of cumulative slip for indentation to a

depth of D 1⁄4 200 nm onto planes (001), (021), and (210) are

shown in Fig. 3. In each case, a slice along the centerline of

the cylinder normal to the laboratory X1 axis (i.e., X-axis) is

shown. For indentation on (001) and (021) planes, the X1

axis is normal to a (100) plane; for indentation on (210), the

X1 axis is approximately normal to a 2ð Þ �30 plane. Cumulative total slip, c, is defined as

c 1⁄4 X

k

ck 1⁄4 X

k

ð

c_

k �

� ��dt; (11)

with c

k the monotonically increasing cumulative slip on system k. Slip activity is greater for indentation into (021) and

(210) planes than for indentation on (001). For this particular

viewing plane, slip contours for indentation on (021) are

noticeably asymmetric. The wireframe mesh of the indenter

is drawn in each figure; in order to enable clear visualization

of heterogeneous slip distributions in the RDX, the mesh of

the substrate, which is considerably more refined than that of

the indenter, is not shown.

Tables VI and VII list maximum local cumulative slip

(i.e., the maximum value of c

k at any location in the RDX

substrate) at an indentation depth D 1⁄4 200 nm when various

elastic constants15,16 are implemented. Total slip, c, listed in

the bottom row of each table is not necessarily the sum of all

c

k listed in a given column, because the location in the substrate where total slip is maximum does not necessarily correspond to the location where each c

k is maximum.

Orientation (001) exhibits slip primarily on system 6. Significant activity of multiple slip systems is evident for indentation onto (021) and (210) planes. Trends are qualitatively

similar, regardless of choice of elastic constants. Cumulative

slip magnitudes are generally slightly larger in Table VII

than in Table VI, because larger stresses are attained at the

same depth of indentation when stiffer elastic constants are

prescribed.

Elastic-plastic simulations were also performed, wherein

the strength, sk

0, of one family of systems was set to G/20,

with that of all others set to G/10. Comparison of force versus depth profiles among these simulations enabled further

assessment of the most active slip systems for each crystal

orientation.28 Results were consistent with those in Tables

VI and VII: slip system 6 is dominant for indentation on

(001) and (021), while systems 1–5 are all important for indentation on (210).

Model predictions for uniform slip strengths, sk

0, of G/

10, G/20, and G/40 are compared with experimental data6 in

Fig. 4 for indentation on (210). Reduction of slip strength

from G/20 to G/40 enables closer agreement with experiment

at larger indentation depths, but also leads to premature initiation of slip and under-prediction of force at smaller depths

TABLE VII. Maximum local slip c

k at indentation depth of 200 nm for indentation on (001), (021), and (210) planes; Brillouin elastic constants.16

System k Plane (001) Plane (021) Plane (210)

1 0.01 0.19 0.41

2 0.01 0.00 0.41

3 0.06 0.13 0.28

4 0.06 0.00 0.28

5 0.01 0.10 0.52

6 0.35 0.51 0.26

all (c) 0.35 0.56 0.58

TABLE VI. Maximum local slip c

k at indentation depth of 200 nm for indentation on (001), (021), and (210) planes; RUS elastic constants.15

System k Plane (001) Plane (021) Plane (210)

1 0.01 0.19 0.40

2 0.01 0.00 0.40

3 0.08 0.16 0.29

4 0.08 0.00 0.29

5 0.00 0.10 0.51

6 0.31 0.41 0.16

all (c) 0.31 0.46 0.52

FIG. 4. Indentation force vs applied displacement for elastic-plastic indentation into

plane (210) with uniform slip strengths G/

10, G/20, and G/40 compared to experiment

(Ref. 6); model predictions obtained using

(a) RUS (Ref. 15) and (b) Brillouin (Ref. 16)

elastic constants.

063512-6 J. D. Clayton and R. Becker J. Appl. Phys. 111, 063512 (2012)
118

near experimental excursion points, i.e., near initial yield.

Similar trends were predicted28 for indentation onto (001)

and (021) planes when strength is increased or decreased

uniformly among slip systems.

Loading-unloading simulations of indentation on

(001), (021), and (210) planes were performed to enable

comparison with experimental observations. Force versus

displacement predictions are shown in Fig. 5 for a single

load-unload cycle, corresponding to loading to D � 200 nm

and then unloading to D 1⁄4 0. During much of the loading

phase, orientation (001) is slightly less stiff than orientations (021) and (210). However, at D � 200 nm, P is

largest for orientation (001). Recall that orientation (001)

demonstrates a lower elastic stiffness, but also less slip

activity. Hence, at larger indentation depths, increased slip

activity for orientations (021) and (210) lowers their

effective tangent stiffness and, hence, P below that of

orientation (001). Hysteresis is also substantially greater

for indentation on (021) and (210) than for (001), again

demonstrating less slip activity in the latter. Orientation

(210) demonstrates the most hysteresis (and greatest slip

activity) and the largest elastic stiffness for much of the

unloading phase.

Figure 6 shows total cumulative slip, c, on unloaded

surfaces of (001), (021), and (210) planes. In each case, the

maximum indentation depth prior to unloading is D � 200

nm. Non-circular contours are consistent with activity of

fewer than five geometrically independent systems necessary

to accommodate an arbitrary plastic strain field.6

Surface slip contours from individual systems were also

examined.28 Predicted surface slip activity results primarily

from system 6 for indentation on (001). Slight contributions

to the circular total slip trace in Fig. 6(a) are due to systems

3 and 4, i.e., f011g[100]. Surface slip activity is predominantly from system 6, with minor contributions from systems

1 and 3, for indentation on (021). Surface slip activity is predominantly from system 5, with contributions from systems

1, 2, and 6, for indentation on (210). Faint contributions

from systems 3 and 4 are also predicted for indentation on

(210).

Active slip planes during the loading history at the specimen surface, as predicted here, are compared with those

deduced from experimental surface impressions6 in Table

VIII. It is noted that, in simulations of indentation on (210),

FIG. 5. Indentation force vs applied displacement for a single load-unload

cycle on each of (001), (021), and (210) planes; elastic constants from RUS

(Ref. 15), uniform slip strength G/20.

FIG. 6. Residual total slip contours at the

surface for indentation to depth D � 200 nm

followed by unloading; results obtained

using elastic constants from RUS (Ref. 15)

and uniform slip system strength G/20: (a)

indentation into (001); (b) indentation into

(021); (c) indentation into (210).

063512-7 J. D. Clayton and R. Becker J. Appl. Phys. 111, 063512 (2012)
119

systems 3 and 4, i.e., f011g[100], demonstrate substantial

activity within the bulk material (i.e., beneath the surface),

but little cumulative slip at the surface. Though not presented

graphically in this paper, predicted residual bulk and surface

slip contours were also obtained using the relatively stiff

elastic constants from Brillouin scattering.16 Predicted slip

contours appeared qualitatively similar, regardless of choice

of elastic constants from RUS15 or Brillouin.16

IV. CONCLUSIONS

A nonlinear anisotropic elastic-plastic model has been

developed for RDX. The model accounts for orthorhombic

elastic constants, pressure-dependent compressibility, and

dislocation glide on up to six distinct slip systems.

Numerical simulations of spherical indentation on

(001), (021), and (210) planes of single crystals show significant influences of elastic anisotropy and nonlinearity on

force-displacement data. Model predictions for initial elastic response using constants measured with resonant ultrasound spectroscopy agree with experimental forcedisplacement data for indentation on (001), (021), and

(210) planes. Predictions using constants measured with

Brillouin scattering are in reasonable agreement with

experiments for indentation on (210), but are stiffer than

experiments for indentation on (001) and (021). Orientation

(001) is elastically most compliant, in agreement with

experiments.

Critical shear strengths associated with slip initiation

have been estimated as G/20, where G is a representative

elastic shear modulus. Initial yield points predicted by the

model are in close agreement with experimental load excursion data when elastic constants from resonant ultrasound

spectroscopy are used. Predictions of force for larger indentation depths, wherein predicted plastic slip is substantial,

tend to exceed experimental values, regardless of which set

of elastic constants is used. The constant strength (i.e., perfectly plastic) slip model implemented here is unable to replicate nearly horizontal steps in indentation force observed in

experiments. Such steps may correspond to discrete slip

events of width too fine to be captured by a conventional

continuum slip model. Dependencies of shear strength on

slip history and pressure have been omitted; incorporation of

such physics, for example, as suggested by atomic simulations, might provide improved agreement. Fractures

observed in experiments (at the surface) or not observed

(subsurface) would also explain the higher stiffness in simulations relative to experiments.

Simulations suggest that slip planes (010) and f011g

contain active systems for indentation on (001), with slip on

system (010)[001] dominating the inelastic response; experimental surface observations confirm that these, as well as

f021g slip planes, may also be active. Simulations suggest

that slip planes (010) and f021g, and to a lesser extent

f011g, are active at the specimen surface for indentation on

(210); these same planes have also been confirmed as active

in experiments. Simulations suggest that planes (010), (011),

and (021) contain active systems for indentation on (021);

particular slip planes active for this orientation have not been

reported from experiments.

The present results suggest that system (010)[001] provides the largest contribution to the inelastic material

response (i.e., post-yield force versus displacement curve)

for indentation on (001) and (021) planes, while five systems

f021g[100], f011g[100], and (010)[100] all contribute to

inelastic response for indentation on (210) planes. Plastic deformation and hysteresis are more extensive for indentation

on (021) and (210) than for indentation on (001). Since much

plastic deformation occurs in the bulk of the material, and

since different slip mechanisms may be prominent at the surface and in the bulk, the present results offer new insight into

inelastic mechanical behavior of RDX not available from experimental observations of residual surface topography

alone.

The model developed here, when used with elastic constants obtained from resonant ultrasonic methods, is

thought to provide an accurate representation of the nonlinear anisotropic response of RDX single crystals up to and

including the onset of slip. The present model is also

thought to provide a qualitatively reasonable depiction of

activity of different slip systems when a uniform and constant shear strength on the order of G/20 is prescribed.

Refinements of the model are needed to address any reduction in stiffness associated with discrete or highly localized

shear events or cleavage fractures observed at larger indentation depths.

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TABLE VIII. Significantly active slip planes at the specimen surface during

indentation of (001), (021), and (210) planes.

Indentation surface Present results Experiment6

(001) (010), f011g (010), f011g, f021g

(021) (010), (011), (021) Not reported

(210) (010), f011g, f021g (010), f011g, 0ð Þ 21

063512-8 J. D. Clayton and R. Becker J. Appl. Phys. 111, 063512 (2012)
120

18H. H. Cady, J. Chem. Eng. Data 17, 369 (1972). 19D. Bedrov, J. B. Hooper, G. D. Smith, and T. D. Sewell, J. Chem. Phys.

131, 034712 (2009). 20R. B. Schwarz, D. E. Hooks, J. J. Dick, J. I. Archuleta, and A. R. Martinez,

J. Appl. Phys. 98, 056106 (2005). 21B. Sun, J. M. Winey, N. Hemmi, Z. A. Dreger, K. A. Zimmerman, Y. M.

Gupta, D. H. Torchinsky, and K. A. Nelson, J. Appl. Phys. 104, 073517

(2008).

22L. L. Stevens, D. E. Hooks, and A. Migliori, J. Appl. Phys. 108, 053512

(2010).

23T. D. Sewell and C. M. Bennett, J. Appl. Phys. 88, 88 (2000). 24L. B. Munday, P. W. Chung, B. M. Rice, and S. D. Solares, J. Phys. Chem.

B 115, 4378 (2011). 25J. D. Clayton, Proc. R. Soc. Ser. A 465, 307 (2009). 26R. Nowak, T. Manninen, K. Heiskanen, T. Sekino, A. Hikasa, K. Niihara,

and T. Takagi, Appl. Phys. Lett. 83, 5214 (2003). 27J. K. A. Amuzu, B. J. Briscoe, and M. M. Chaudhri, J. Phys. D: Appl.

Phys. 9, 133 (1976). 28J. D. Clayton and R. C. Becker, U.S. Army Res. Lab. Tech. Report ARLTR-5864, 2012.

063512-9 J. D. Clayton and R. Becker J. Appl. Phys. 111, 063512 (2012)
121

Research@

Particle-Based Multiscale Coarse Graining 

with Density-Dependent Potentials: 

Application to Molecular Crystals 

(hexahydro-1,3,5-trinitro-s-triazine)

Sergei Izvekov, Peter W. Chung, and Betsy M. Rice

U.S. Army Research Laboratory, Aberdeen Proving Ground, MD

ABSTRACT

We describe the development of isotropic particle-based coarse-grain models 

for crystalline hexahydro-1,3,5-trinitro-s-triazine (RDX). The coarse graining 

employs the recently proposed multiscale coarse-graining (MS-CG) method, 

which is a particle-based force-matching approach for deriving free-energy 

effective interaction potentials. Though one-site and four-site coarse-grain (CG) 

models were parameterized from atomistic simulations of non-ordered (molten 

and ambient temperature amorphous) systems, the focus of the paper is a 

detailed study of the one-site model with a brief recourse to the four-site model. 

To improve the ability of the one-site model to be applied to crystalline phases 

at various pressures, it was found necessary to include explicit dependence 

on a particle density, and a new theory of local density-dependent MS-CG 

potentials is subsequently presented. The density-dependency is implemented 

through interpolation of MS-CG force fields derived at a preselected set of 

reference densities. The computationally economical procedure for obtaining 

the reference force fields starting from the interaction at ambient density 

is also described. The one-site MS-CG model adequately describes the 

atomistic lattice structure of α-RDX at ambient and high pressures, elastic 

and vibrational properties, pressure-volume curve up to P = 10 GPa, and the 

melting temperature. In the molten state, the model reproduces the correct 

pair structure at different pressures as well as higher order correlations. The 

potential of the MS-CG model is further evaluated in simulations of shocked 

crystalline RDX.

The Journal of Chemical Physics, Volume 135, Page 044112 (2011)
122

THE JOURNAL OF CHEMICAL PHYSICS 135, 044112 (2011)

Particle-based multiscale coarse graining with density-dependent

potentials: Application to molecular crystals

(hexahydro-1,3,5-trinitro-s-triazine)

Sergei Izvekov,a) Peter W. Chung, and Betsy M. Rice

U.S. Army Research Laboratory, Aberdeen Proving Ground, Maryland 21005, USA

(Received 21 January 2011; accepted 15 June 2011; published online 27 July 2011)

We describe the development of isotropic particle-based coarse-grain models for crystalline

hexahydro-1,3,5-trinitro-s-triazine (RDX). The coarse graining employs the recently proposed

multiscale coarse-graining (MS-CG) method, which is a particle-based force-matching approach for deriving free-energy effective interaction potentials. Though one-site and four-site

coarse-grain (CG) models were parameterized from atomistic simulations of non-ordered (molten

and ambient temperature amorphous) systems, the focus of the paper is a detailed study of the onesite model with a brief recourse to the four-site model. To improve the ability of the one-site model

to be applied to crystalline phases at various pressures, it was found necessary to include explicit

dependence on a particle density, and a new theory of local density-dependent MS-CG potentials is

subsequently presented. The density-dependency is implemented through interpolation of MS-CG

force fields derived at a preselected set of reference densities. The computationally economical procedure for obtaining the reference force fields starting from the interaction at ambient density is also

described. The one-site MS-CG model adequately describes the atomistic lattice structure of α-RDX

at ambient and high pressures, elastic and vibrational properties, pressure-volume curve up to P

= 10 GPa, and the melting temperature. In the molten state, the model reproduces the correct pair

structure at different pressures as well as higher order correlations. The potential of the MS-CG

model is further evaluated in simulations of shocked crystalline RDX. [doi:10.1063/1.3607603]

I. INTRODUCTION

Modeling of molecular crystals is becoming an important

tool to aid materials design in numerous applications ranging

from drugs, electronic devices, and energetic materials. Computer simulations using the molecular dynamics (MD) simulation method based on sophisticated atomistic force fields have

greatly advanced our understanding of dynamics in molecular solids on time and length scales difficult to readily access

through experimental techniques. However, the MD modeling of molecular crystalline materials at the atomistic level

encounters difficulties due to a significant disparity between

characteristic time scales for intramolecular and lattice movements as the former are governed by strong intramolecular

chemical bonds while the latter rely on the weak interactions between molecules, which are typically uncharged and

non-polar. The presence of fast intramolecular modes in the

molecular solids makes it nearly impossible to extend atomistic computer simulations to mesoscale using presently available computer power. Understanding behavior at such a scale

is necessary in many areas such as study of microstructured

and heterogeneous crystals, crystals with extended defects, or

response of materials to external mechanical stimuli (shock,

shear, etc.). The difficulties have prompted development of

new approaches for computer simulations at the mesoscale.

In recent years, particle-based coarse-graining approaches,

a)Author to whom correspondence should be addressed. Electronic mail:

sergiy.izvyekov@us.army.mil.

where groups of atoms are replaced with structureless interaction centers that interact through effective (coarse-grain) potentials, have grown increasingly popular. A large variety of

schemes for constructing such potentials has been reported in

the literature (for recent reviews see Ref. 1).

The development of coarse-grain (CG) potentials for

condensed-phase systems presents additional challenges

beyond the accuracy and validity of the underlying atomistic

potentials.1, 2 The first difficulty is a so-called “representability paradigm” as a general impossibility to design a CG

potential which is able to reproduce every property of the

atomistic ensemble even at a single thermodynamic state.3

For example, the thermodynamics of a CG ensemble of N

particles matches that of the reference atomistic ensemble if

the CG potential accurately approximates the atomistic free

energy W(RN ) in CG coordinates RN = (R1,..., RN ). However, at non-zero temperature T, the W(RN ) potential, which

can be referred to as the free-energy based CG potential, cannot reproduce the total average internal potential energy of the

system, U(RN ), as the W(RN ) contains both the potential energy and the entropic contributions associated with averaged

out atomistic coordinates as W(RN ) = U(RN ) − T S(RN ).

This feature is one prominent manifestation of the representability problem. CG potentials derived to match the

atomistic free-energy W(RN ) are preferable as they can

correctly reproduce the density correlations in the CG ensemble, and therefore the correct thermodynamics. Another

difficulty associated with the use of CG potentials is due

to poor state-point transferability as the potentials naturally

0021-9606/2011/135(4)/044112/17/$30.00 135, 044112-1

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123

044112-2 Izvekov, Chung, and Rice J. Chem. Phys. 135, 044112 (2011)

depend on the thermodynamic state at which they were

parametrized.4, 5 The CG modeling of molecular compounds

is further complicated by the fact that there is no “universal”

functional form to describe CG interactions as, for example,

a Lennard-Jones form to represent short-ranged dispersion

interactions between atoms. Moreover, the functional form

suitable to parameterize the CG interactions could not be

even determined a priori from physical principles and in most

cases a numerical representation is the only way to fit CG potentials accurately. Furthermore, CG potentials are essentially

many-body functions and must be pairwise decomposable for

ease of implementation and low computational expense.

Recently, an efficient bottom-up force-matching (FM)

based approach was developed which properly addresses

the above challenges. This method, the multiscale coarsegraining (MS-CG) method,5–9 is similar to conventional

structure inversion methods (e.g., Boltzmann inversion,

Schommers inversion,10 or Inverse Monte Carlo11 approaches) which are all based on Henderson’s uniqueness

theorem,12 in that it yields a pairwise decomposition of

the all-particle CG potential of mean force (PMF), but

better captures higher order correlations of the reference

ensemble. The MS-CG method can be naturally extended to

include dependence on thermodynamic state variables such

as density. Importantly, the MS-CG method yields a force

field which can be implemented directly into MD codes, as

the MD method does not require knowledge of the potential

to integrate the dynamics. This becomes an advantage if the

density dependency is introduced.

Thus far, the MS-CG method has been extensively

applied to soft-matter systems.9 Because the quality of

numerical MS-CG models depend on how well the intermolecular separations are sampled, the MS-CG approach

would be difficult to apply directly to crystals. The goal of

this research is to explore to what extent a FM model from

the disordered phase (e.g., liquid) is transferable to crystalline

structures. We discuss application of the MS-CG method to

coarse-grain intermolecular interactions in hexahydro-1,3,5-

trinitro-s-triazine (RDX), which is an important energetic

crystal widely used as a secondary explosive and in propellants. The primary model we investigate is a one-site model

in which every molecule of RDX in the crystal is represented

using a coarse-grained particle with a single interaction site.

However, on several occasions it is useful to also discuss and

compare the one-site results with those of a four-site model.

One of the major applications we sought for our one-site

MS-CG model is studying crystalline RDX under a shock

load. Therefore, to improve transferability to high pressures,

we introduce a computationally economical procedure for

constructing a density-dependent potential, which follows the

recipe outlined in Ref. 5. We also present a new theory for

local density dependent MS-CG potentials, which were previously introduced in our recent work5 at a rather ad hoc level.

We proceed as follows: In Sec. II, we describe the computational method and present new theory for local-density

dependent potentials. In Sec. III, we describe details of

application of the developed formalism to coarse graining

of crystalline RDX. We discuss the derivation of MS-CG

potentials, their validation, and then report applications of the

MS-CG model to study shocked RDX. Conclusions are given

in Sec. IV.

II. METHOD

A. MS-CG method

The MS-CG method is described elsewhere5, 8, 9 as a

FM-based approach for constructing free-energy based CG

potentials (e.g., Gibbs potential for NPT or Helmholtz potential for NVT ensembles). The MS-CG approach uses

the configurations and associated forces sampled from the

atomistic (we assume here canonical, NVT) ensemble of

n atoms partitioned into N groups to construct a pairwise additive approximation to the free-energy (Helmholtz)

potential,

Watm(RN ,N,V,T ) = −kT ln�δ(rn − RN )�, (1)

in the CG coordinates RN = (RI , I = 1, N), which describe

translational degrees of freedom (dof) of the groups. In Eq. (1)

rn = (RN , xP ) are canonical atomistic coordinates, where xP

denotes integrated out (intragroup) dof. The brackets �· · ·� denote the canonical average. In terms of coordinates xP and

their conjugated momenta pP

x , we can write �δ(rn − RN )� as

an ensemble average over xP ,

�δ(rn − RN )� ≡< 1 >xP = 1

Zn



dxP exp[−βUatm(RN , xP )],

(2)

where Uatm(RN , xP ) is the atomistic potential energy function; Zn is the system partition function excluding the kinetic

contribution; and β = 1/kT . The N!�δ(rn − RN )� = �(RN )

is an N-particle distribution function. The RI coordinate

can be any weighted average RI = 

i=1,sI cI iri



, where

i=1,sI cI i = 1 and sI is the number of atoms in the Ith

group. Commonly, the RI coordinate is associated with the

group’s center-of-mass (c.m.); i.e., cI i = mi/MI , where mi,

MI are atom and group masses. The dependence of Watm

on the number of intraparticle dof P and the volume VP

in xP space develops due to the ensemble average over xP

in Eq. (2). For a fixed number of intraparticle dof per CG

particle, the P and VP variables can be replaced with the

number of CG particles N and the volume of the CG system

V. Furthermore, the CG coordinates RN may be viewed

as thermodynamic state variables for the thermodynamic

potential Watm. Finally, the CG Hamiltonian which governs

the dynamics in CG system becomes

HCG(PN , RN ,N,V,T ) =



I=1,N

P2

I

2MI

+Watm(RN ,N,V,T ),

(3)

where PI is the momentum of the CG particle conjugated

to RI .

The MS-CG interaction is constructed through forcematching of the thermodynamic force field,

FPMF

I (RN ,N,V,T ) = −[∇RI Watm(RN ,N,V,T )]N ,V ,T .

(4)

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124

044112-3 Multiscale coarse graining of crystals J. Chem. Phys. 135, 044112 (2011)

The FPMF field is an ensemble average over the xP dof,

FPMF

I (RN ,N,V,T ) =



FCG

I (RN , xP )



xP

�1�xP

, (5)

of the atomistic instantaneous c.m. force field

FCG

I (RN , xP ) = 

i=1,sI

fi, (6)

where fi are instantaneous atomistic forces. Therefore,

the MS-CG interaction potential Watm [Eqs. (1) and (2)] can

be interpreted as an N-particle PMF. Now, define a function

GMSCG

I (RN , �) that has an additive decomposition into pairwise and central force fields

GMSCG

I (RN , �) = 

J �=I

gMSCG(RI J , �)nI J , (7)

where RI J = RI − RJ , RI J = |RI J |, nI J = RI J /RI J , and �

is a set of adjustable parameters. A least-squares approximation of FPMF

I can be obtained from GMSCG

I (RN , �0) by finding

the optimal set �0 ∈ � that minimizes5

χ(�) =

 

I=1,N

|FCG

I (RN , xP ) − GMSCGI (RN , �)|

2



. (8)

Then, the potential

W MSCG(RN , �) = 

I<J

wMSCG(RI J , �), (9)

where

wMSCG(R,�) =

 ∞

R

dR�gMSCG(R�, �), (10)

implements a pairwise decomposition of the Watm(RN ,

N,V,T ). In a practical implementation of the MS-CG

method, the gMSCG(R,�) function in Eq. (7) can be a cubic

spline function defined on a mesh {Rk} up to a preselected cutoff distance Rcut . The problem of minimizing χ(�) is then reduced to an overdetermined system of linear equations, which

is solved using a block-averaging scheme.5, 6 As shown in

Ref. 5, the MS-CG procedure eliminates part of the atomistic

free energy associated with the integrated-out dof, resulting

in a pressure deficiency in the CG ensemble. This deficiency

can be corrected using a virial constraint5, 6 in Eq. (8).

It is worthwhile to comment on the difference between

the MS-CG methodology and the RDF inversion methods,

which are based on the Henderson’s uniqueness theorem.12

The Henderson’s theorem postulates the existence of a unique

isotropic pairwise effective potential in a system of sametype particles with an arbitrary interaction which reproduces

the RDF of the ensemble. Such an effective potential, wH (R)

= wH (R,N,V, T ), can be found using the target RDF g(R)

as an input through an iterative scheme.10 The g(R) is related

to the two-particle PMF as

Watm(R2

) = −kT ln[�(R2)] = −kT ln[ρ2g(R12)], (11)

where �(R2) = 

V N−2 dRN−2�(RN )/(N − 2)! is a twoparticle distribution function and ρ is the particle’s density.

Similar to the MS-CG potential, the wH (R) represents an approximate pairwise decomposition of the all(N)-particle PMF

as Watm(RN ) ≈ 

I<J wH (RI J ) = WH (RN ) [cf. Eq. (9)]. It

can be shown that in the N-particle ensemble with the interactions described by the wH (R), the Watm(R2) and therefore

two-particle correlations are reproduced exactly. On the other

hand, the MS-CG approximation wMSCG(R) is not necessary

to reproduce the exact two-particle PMF and corresponding

RDF. Rather, the wMSCG(R) is biased to approximate the entire spectrum of the k-particle PMFs (k = 2, N), resulting in

better capturing of the many-particle correlations.

B. Density-dependent MS-CG potentials

As the W MSCG(RN , �) represents a pairwise approximation to the system free energy Watm(RN ,N,V,T ), it

is explicitly dependent on the ensemble thermodynamic

state variables N, V, and T, thus limiting the state-point

transferability3, 4, 13 of the model. In the homogeneous equilibrium system, such as bulk crystal or liquid, the Watm is a function of the global density ρG = N/V as Watm(RN , ρG,V,T )

or Watm(RN ,N,ρG, T ), depending on whether N or V is chosen as an independent thermodynamic variable.

Strictly speaking, as briefly discussed in Sec. II A, since

the canonical average in Eq. (2) is over xP , the Watm depends

on the total number of intraparticle dof P and volume VP of

the phase subspace xP sampled by the system and therefore on

the global density ρP ,G = P/VP . The ρP ,G can be replaced

with the density of CG particles ρG under the assumption

of a fixed number of intraparticle dof per CG particle and

if VP ∝ V . For the density-dependent potentials, it is convenient and correct to assume that the ρP ,G (or equivalently

ρG) is not determined by either the dimension of CG coordinate phase space N or the CG particle coordinates RN , and

is actually an independent thermodynamic variable. For similar reasons, we can regard VP (or equivalently V) as a state

variable independent of ρP ,G (ρG) and RN . We will assume

ρP ,G = ρG, VP = V , and use NG = ρGV to distinguish the

dimension of intraparticle phase space from the dimension N

of CG configuration RN .

The CG interaction can be viewed as an effective interaction between particles immersed into a “sea” of intraparticle

dof density that may vary. In such a picture, the “intraparticle”

chemical potential,

μ(RN ) =

∂Watm

∂NG



RN ,V ,T

= 1

V

∂Watm

∂ρG



RN ,V ,T

,

(12)

determines a response of the effective interaction in an ensemble of N CG particles due to a change in density of intraparticle dof under constant configuration of CG particles 

RN 

, V,

and T. Within this formalism, all thermodynamic equations

relating the Helmholtz free energy potential and other thermodynamic potentials, which are functions of state variables RN ,

N,V,T , remain valid. The above picture is analogous to the

effective electron mediated ion-ion interactions in a nucleielectron plasma, with electrons considered as “intraparticle”

dof, if the latter are well localized at the nuclei.

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125 044112-4 Izvekov, Chung, and Rice J. Chem. Phys. 135, 044112 (2011)

C. Global density dependency

Improved transferability with regard to both density and

temperature conditions can be achieved directly by incorporating dependence on a global density ρ = ρG in the MS-CG

potential.5 Following Ref. 5, the global density potential can

be implemented using a linear interpolation of the MS-CG

forces parameterized at several reference densities ρi as

gMSCG(R, ρ) = aρ × gMSCG(R, ρi+1)

+ (1 − aρ) × gMSCG(R, ρi),

aρ = ρ − ρi

ρi+1 − ρi

, ρ ∈ [ρi, ρi+1], (13)

where i is an index over a finite number of states selected

for the parameterization. The forces gMSCG(R, ρ) in Eq. (13)

can then be used directly in MD simulations. In contrast,

Eqs. (9) and (10) cannot be used to calculate the total potential

energy of the CG system with respect to any other reference

state of a different density (including a vacuum), as an assumption of a density change along system trajectories leads

to

GMSCG

I (RN , ρ) = −∇RI W MSCG(RN , ρ)

− [W MSCG(RN , ρ)]�

ρ∇RI ρ. (14)

However, if we assume that the force field GMSCG

I (RN , ρ)

can be uniquely integrated to a potential function (which

is a reasonable assumption even for the local density

approximation14), then the total energy is pairwise and can

still be determined using Eq. (9) with

wMSCG(R, ρ) =

 ∞

R

dR�gMSCG(R�, ρ(R�)), (15)

where ρ(R�

) = R�3

R3 ρ. Therefore, the total potential energy using Eqs. (9) and (15) is obtained by calculating the work by

the GMSCG

I (RN , ρ) force field as the system uniformly expands from the initial state with the density ρ, and thus represents the interparticle potential energy with respect to infinite

separation of the particles.

Importantly, for a piecewise interpolation of Eq. (13), the

potential Eq. (15) is continuously differentiable with respect

to density, and therefore the chemical potential μ [Eq. (12)]

is a continuous function of the density. This can be seen by

transforming the potential to an integral over ρ� = R�3

/R3ρ

as

wMSCG(R, ρ) = R

3ρ2/3

  ρ1

ρ

+



i=1,imax−1

 ρi+1

ρi

+

 0

ρimax



× dρ�

ρ�−1/3gMSCG(ρ�1/3ρ−1/3R, ρ�), (16)

and then easily verifying that for the linear interpolation

Eq. (13) [wMSCG(R, ρi + 0)]�

ρ = [wMSCG(R, ρi − 0)]�ρ, and

therefore wMSCG(R, ρ) is continuously differentiable with

respect to ρ.

D. Local density dependency

For simulations of macroscopically structured systems

such as those with multiple phases, under external fixed potential, or quasiequilibrium states such as within a shock

front, the potentials need to be dependent on local density

ρ = ρ(RN ). Under constant NVT conditions, the local density

refers to the number of particles within the elemental volume

in the configurational phase space. Since the number of particles within this elemental volume defined by the local density

is not a constant, the local density potential refers rather to the

grand-canonical (constant chemical potential-volume, μV T )

ensemble instead of NVT. This observation can be used to find

a representation of a local density potential for the NVT ensemble, which is consistent with the existence of global density variable. Recall that Watm(RN , ρG,V,T ) is considered

to be a function of the global density of intraparticle dof ρG

and corresponding volume as introduced in Sec. II A with

a respective interpretation of the chemical potential μ(RN )

Eq. (12). We will derive this relation for a homogeneous system that can be characterized by a spatially uniform μ(RN ),

meaning that μ is not a function of xP .

Using the Legendre transformation, the constant μV T

potential [see Eq. (12)] is introduced as

Watm(RN , μ, V , T ) = Watm(RN , ρG,V,T ) − μρGV.

(17)

The corresponding equation for the MS-CG pairwise

approximation Eqs. (9) and (10) is

W MSCG(RN , μ, V , T )

= 

I=1,N



W MSCG

I (RN , ρG,V,T ) − μ

N ρGV



= 

I=1,N



W MSCG

I (RN ,

μ

N ,V,T )



, (18)

where

W MSCG

I (RN ,ρ,V,T ) = 1

2



J �=I

wMSCG(RI J ,ρ,V,T ).

(19)

At this step, in order to proceed with defining a local

density, we have to make an assumption on the nature

of the dependence of W MSCG

I (RN ,ρ,V,T ) on ρ. Because

wMSCG(RI J ,ρ,V,T ) = 0 for RI J > Rcut, and therefore the

sum in Eq. (19) is over only those particles which are within

Rcut of the Ith site, we may postulate that

W MSCG

I (RN ,ρ,V,T ) = W MSCGI (RN , ρI ,V,T ), (20)

where we introduce the “local” density

ρI = NI /V Rcut, (21)

and where V Rcut = VI is the volume of sphere of radius Rcut

centered at the Ith particle, and NI is a number of intraparticle dof within the sphere. In other words, Eq. (20) assumes that the effective potential W MSCG

I experienced by the

Ith CG particle is sensitive to changes in its local density of

only those intraparticle dof which are associated with the CG

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126

044112-5 Multiscale coarse graining of crystals J. Chem. Phys. 135, 044112 (2011)

particles with which the Ith particle interacts. Note, that V

in the r.h.s. of Eq. (20) should not be replaced with VI because the W MSCG

I is an extensive function of V. Otherwise, if

W MSCG

I (RN , ρI ,V,T ) = W MSCGI (RN , ρI , VI , T ), the chemical potential μ(RN ) [Eq. (12)] will be a spatially varying function of intraparticle coordinates xP , thus contradicting the assumption that the system is homogeneous. Next, in applying

the Legendre transformation [Eq. (17)] to terms W MSCG

I in the

last sum in Eq. (18), we obtain

W MSCG(RN , μ, V , T ) = 

I=1,N

W MSCG

I (RN , ρI ,V,T )

− μ

N



I=1,N

ρIV. (22)

Now, we can introduce the “global” density variable as

ρG = 1

N



I=1,N

ρI , (23)

and again using the transformation in Eq. (17), we derive a relation between the global and local density MS-CG potentials

as

W MSCG(RN , ρG,V,T ) = W MSCG(RN , μ, V , T ) + μρGV

= 

I=1,N

W MSCG

I (RN , ρI ,V,T ).

(24)

It is easily seen that in the thermodynamic limit (N → ∞,

N/V = const), the definition in Eq. (23) is equivalent to

ρG = N/V .

Next, Eq. (24), with only the density dependence shown,

can be written explicitly in a pairwise form as

W MSCG(RN , ρG) = 

I<J

wMSCG(RI J , ρI , ρJ ), (25)

whereupon using Eq. (15) we may write

wMSCG(RI J , ρI , ρJ ) = 1

2

(wMSCG(RI J , ρI )

+ wMSCG(RI J , ρJ )). (26)

Note that similar to W MSCG and W MSCG

I , the wMSCG

(R, ρ,V ) should be considered as extensive functions of V.

Eqs. (25) and (26) represent a pairwise implementation of the

local density potential.

If we assume that the spatial variation of the density is

small,

ρI = ρG + δρI ,δρI /ρG � 1, (27)

then the potential in Eq. (26) can be approximated as

wMSCG(RI J , ρI , ρJ ) ∼= wMSCG

RI J , ρG +

1

2

(δρI + δρJ )



= wMSCG(RI J , ρI J ), (28)

where

ρI J = 1

2

(ρI + ρJ ). (29)

Equation (28) results in a computationally more efficient

scheme, as the potential wMSCG(R, ρ) requires the knowledge

of the density at a single point instead of two as in Eq. (26) and

can be used in simulations if the spatial variation of density on

length scales l ∼ Rcut is relatively small. In our derivations,

the RN are considered as independent thermodynamic state

variables, and therefore wMSCG(RI J , ρI ) and gMSCG(RI J , ρI )

are related by Eq. (10); then Eqs. (26) and (28) can be written

for forces. In particular, in the limit Eq. (27), we introduce the

density dependent pairwise force as

gMSCG(RI J , ρI , ρJ ) = gMSCG(RI J , ρI J ) (30)

which was introduced in our previous work in Ref. 5 on an ad

hoc level. An alternative to Eq. (29) could be

ρI J = max(ρI , ρJ ), (31)

which agrees with Eq. (29) to within O(δρI /ρG). To obtain

the potential with respect to a gas of non-interacting CG particles, the force in Eq. (30) should be integrated in accordance

to Eq. (15). Such a potential should be also utilized to calculate the system total energy in the simulations of inhomogeneous systems.

In closing, we would like to make several important conclusions on the applicability of the local density potentials

presented in Eqs. (21) and (25), and (26): (I) The local density

can be introduced if the free energy Watm is separable into a

sum of “local” terms Watm

I associated with each CG site. The

pairwise approximation gives a natural way [see Eqs. (10),

(15) and (19)] to construct such a separable representation.

In principle, under certain circumstances, the partitioning of

the free-energy potential into local terms (i.e., constructing

the local density approximation) could be accomplished for

systems with many-body interactions; (II) The notion of local

density involves all particles within the range of interaction.

Therefore, the local density becomes poorly defined for systems with long-ranged interactions; (III) It is generally incorrect to use the definition of local density as in Eq. (21) with

an arbitrarily small V Rcut .

It is instructive to compare the above formalism with the

formalism for many-body conservative potentials which are

free-energy based and density dependent, as developed by

Pagonabarraga and Frenkel.14, 15 Their postulates of the notion of local and global densities and their relation are similar

to those which we have derived more rigorously.

Within the approach developed by Allen and

Rutledge,13, 16 the local density potentials are derived by

inverting the RDFs [or two-particle PMFs, see Eq. (11)] at

local densities, which are related to the RDF at the global

density through consistency equations. Because the MS-CG

potentials represent a similar inversion of the N-particle PMF,

the method by Allen and Rutledge can be readily generalized

to the MS-CG potentials. Unfortunately, the consistency

equations of Allen and Rutledge are formulated in terms

of energy, whereas the present MS-CG method requires

forces. Therefore, the application of this technique to MS-CG

potentials has not been attempted.

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127 044112-6 Izvekov, Chung, and Rice J. Chem. Phys. 135, 044112 (2011)

E. Parameterization of density dependency

Ideally, gMSCG(R, ρi) in Eq. (13) should be obtained

from separate atomistic simulations at the appropriate thermodynamic conditions at each reference density ρi. For RDX

however, the slow intermolecular dynamics in dense states

precludes a sufficiently large sampling of atomistic configurations for the purpose of FM at P ≥ 2 GPa. For this reason,

we developed a less direct but computationally more efficient

scheme to construct a reference potential set {gMSCG(R, ρi)}

than that used in our earlier study for liquid nitromethane

(NM).5 The scheme is based on the difference between two

MS-CG potentials calculated for NM in accordance with

Eq. (10), one of which corresponds to gMSCG(R, ρi) derived

without a virial constraint and the other derived with a virial

constraint. The former MS-CG potential is pressure deficient

while the latter produces the correct pressure. In that study,

we showed that the difference between the two potentials is

close to a linearly decaying function. In the same study, we

observed that the density contribution can be described in the

CG potential by a linear term over a fairly broad range of densities. This suggests that for ρi > ρ0, where ρ0 is the ambient

density, the force term gMSCG(R, ρi) can be constructed as a

sum of ambient gMSCG(R, ρ0) and a constant term, where the

latter has the linearly decaying potential. We used the form

gMSCG(R, ρi) = gMSCG(R, ρ0) + gMSCG

i (R), (32)

where gMSCG(R, ρ0) was force-matched to the atomistic

simulation at the ambient density and the gMSCG

i (R) is

approximated as

gMSCG

i (R)=

⎧

⎪⎪⎨

⎪⎪⎩

gi, R ≤ Rcut − δR

gi sin�

(Rcut−R)

δR

π

2

�

, Rcut − δR < R ≤ Rcut

0, R>Rcut

.

(33)

In order to improve energy conservation in the dynamics

simulations, the gMSCG

i (R) force is smoothly attenuated to

zero within a δR distance to Rcut [second line in Eq. (33)].

In the past, a linear correction similar to Eq. (33) has been

used to improve pressure properties in structure-based coarsegraining frameworks.17

For each ρi > ρ0, the parameter gi was adjusted to reproduce the atomistic V − P isotherm of the atomistic α-RDX

bulk crystal as discussed below. As discussed later, varying

only the parameter gi at higher pressures is not always sufficient to obtain the target density while preserving the correct

lattice symmetry. To improve the match in the density, the

gMSCG

i (R) term was subjected to additional rescaling modification as

gMSCG

i (R, a) = 1

a

gMSCG

i

�1

a

R

�

(34)

with regard to an adjustable parameter, a.

Similar to Ref. 5, we will refer to the gMSCG(R, ρi) interactions as MS-CG/Pi, where Pi is the pressure which corresponds to ρi, while the corresponding density dependent

model introduced in accordance to Eq. (13) is referred to

as the MS-CG-D model. Within the MS-CG-D model, for

systems with a macroscopically inhomogeneous particle density, the local density Eqs. (21) and (26) should be used. This

scheme was used to evaluate the local density dependence of

the interactions in systems with strong density gradients such

as in shock fronts. In the systems in which the density gradients are deemed to be relatively small (e.g., simulations of

melting), we used the computationally simpler scheme from

Eqs. (21), (28), and (31). In equilibrium homogeneous simulations such as bulk phases, the density in the latter scheme

was replaced with the ρG = N/V as discussed in Secs. II B

and II C.

Practically, in the CG simulations, the gMSCG(R, ρi) were

stored as two-dimensional data arrays. In simulations with local density dependent potentials, the local density for each

particle was evaluated at each time step of a MD simulation

by counting the particles within a sphere of radius Rcut centered at the particle. All local densities were then rescaled uniformly to satisfy the consistency equation Eq. (23) with the

instant global density ρG = N/V as estimated from a current system volume. For the system sizes used in the study,

Eq. (23) is well satisfied, so that the rescaling of local densities is not necessary. The pairwise forces were then evaluated through the interpolation scheme in Eq. (13) and the

total system energy was updated using Eq. (15) if necessary.

Therefore, in the simulations of inhomogeneous systems, the

interaction between the CG particles changes at each step of

the dynamics following the density changes.

III. RESULTS AND DISCUSSION

A. Reference atomistic simulations and details

of coarse graining

Atomistic MD simulations of the RDX systems were performed using the all-atom flexible force field developed by

Smith and Bharadwaj (SB model).18, 19 The SB model accurately predicts many physical properties of crystalline RDX,

including its crystal structure, elastic properties, and thermal

expansion.20 This model was also employed to simulate RDX

under shock conditions.20 We will also use data from all-atom

simulations using the Sorescu-Rice-Thompson (SRT) force

field,21, 22 which is rigid body, and also the force field by Boyd

et al.23 for comparison with the MS-CG-D results.

The coarsening of the RDX molecule into one- and foursite models is depicted in Fig. 1. As discussed in the Introduction, the present version of the MS-CG method cannot use reference data sampled from simulations of ordered

(e.g. crystalline) systems due to the piecewise functional

form of the pairwise terms gMSCG(RI J , �) [Eq. (7)]. Therefore, two noncrystalline systems were used to obtain reference trajectory and force data. The first was molten RDX at

T = 550 K, and the second was disordered RDX obtained

by annealing the molten RDX to T = 350 K. Electrostatic

interactions were treated with the Ewald method24 and the

short-ranged interactions were cut off at 1.0 nm. The simulations were carried out at P = 0 GPa and a cubic supercell containing 144 molecules was used. Periodic boundary conditions were imposed in all directions. The constant

isobaric-isothermal (NPT) conditions were simulated using

Melchionna’s modification of the Hoover algorithm, in which
128

044112-7 Multiscale coarse graining of crystals J. Chem. Phys. 135, 044112 (2011)

O

N

C H RDX 

One-site model

NO2 

C3N3

Four-site model

FIG. 1. Coarse graining of the RDX molecule into one-site and four-site

representations.

the equations of motion couple a Nosé-Hoover thermostat

and a barostat,25 with relaxation times of 0.5 and 3 ps,

respectively.25

First, the two systems were equilibrated under constant

NPT conditions for 10 and 30 ns, respectively. The configurations, velocities, and forces were then sampled from constant NVT simulations and these data were then used in the

subsequent CG FM procedure. The value of the instantaneous

virial was recorded for each configuration during the NVT

stage and later imposed through the virial constraint during

the force-matching.5 The 550 K system was simulated for

30 ns with configurations recorded at 2 ps intervals; the 350 K

system was simulated for 60 ns with configurations recorded

at 4 ps intervals, thus providing a total of 15 000 configurations for each system to be used in the FM procedure. The

sampled data were then transformed into the CG coordinates

of interest (either molecular or group c.m.) and merged such

that each high temperature configuration was followed by the

lower temperature one.

The merged 350 K and 550 K state point data (350/

550 K sampling) were used to parameterize the onesite MS-CG force field at P = 0 GPa (MS-CG/0 model). The

MS-CG/0 was served as a base interaction in the procedure

outlined in Sec. II E to build the MS-CG/Pi,Pi > 0, force

fields and then the MS-CG-D model as described in details in

Sec. III B and III C. For a comparison, we also force-matched

the one-site model from only “pure” state point 550 K data.

Two four-site models, which are non-density dependent,

were developed from the 350K/550 K and only 550 K data,

respectively. The MS-CG forces were determined using

the block averaging scheme5, 6 with the block size of 10

on a radial mesh with bin size of 0.1 a.u. The cutoff Rcut

= 29 a.u. (15.35 Å) for one-site models and Rcut = 23 a.u.

(12.17 Å) for four-site models were used.

The atomistic and CG MD simulations were performed

using the DL_POLY 2.20 simulation package,26 which was

modified to enable density dependency of the interactions in

accordance with Eqs. (13) and (15).

B. MS-CG potentials without density dependency

The one-site effective force as yielded by the forcematching to mixed 350/550 K and pure (molten) 550 K

state points at ambient pressure are displayed in Fig. 2(a).

The corresponding one-site MS-CG/0 potentials are shown in

Fig. 2(b). The MS-CG potential cannot be determined at short

0 0.2 0.4 0.6 0.8 1 1.2 1.4

R , nm

-0.25

0

0.5

1

1.5

2

F(R), 10-2 a.u.

550 K

350 K / 550 K

(a)

Wall

Exp

0 0.2 0.4 0.6 0.8 1 1.2 1.4

R, nm

-20

0

20

40

60

80

100

120

 W(R), kJ/mol

P = 0 GPa

P = 0.5 GPa

P = 1 GPa

P = 2 GPa

P = 3 GPa

P = 5 GPa

P = 10 GPa

P = 0.0 , T = 550 K

(b)

 Rcore, exp

 Rcore, wall

FIG. 2. Panel (a): One-site MS-CG force profiles as obtained from forcematching to 550 K (thin solid) and 350/550 K (thick solid) samplings. For

350/550 K force (MS-CG/0 model), possible choices of core interactions corresponding to core distances outlined in panel (b) in vertical lines are shown

in dash-dash-dotted (hard-wall core) and dashed (exponential core). Panel

(b): One-site MS-CG potentials at different pressures and T = 300 K. Vertical dashed lines mark possible selections of the core distance. Shown are

potentials at P (GPa): 0 (filled circles), 0.5 (filled squares), 1 (filled triangles), 2 (empty circles), 3 (empty squares), 5 (empty triangles), 10 (stars).

The P = 0 GPa potential corresponds to 350 K/550 K force in panel (a).

Dotted line shows the potential corresponding to 550 K force in panel (a). On

both panels dot-dashed line outlines the atomistic c.m.-c.m. RDF in arbitrary

units.

distances R<Rcore (referred to as the core region), since this

region is not sampled in the reference simulation; therefore,

the MS-CG potential must be extrapolated in this region. The

exponential core (shown in Fig. 2) typically produces slightly

better results (e.g. RDFs) at ambient conditions. However, at

shocked states the hard-wall potential (also shown in Fig. 2)

performed slightly better and, therefore, was chosen for the

calculations presented here.

A more portable representation of the 350/550 K MSCG/0 force can be obtained through a least-squares fit of the

tabulated spline data by powers of 1/R

g(R) = nmax

n=2

An/Rn. (35)

A better fit can be obtained with the more rapidly convergent

Chebyshev series

g(R) = A0/2 +nmax

n=1

AnT n(R ̄), (36)

where T n(R ̄) is the Chebyshev polynomial of degree n and

R ̄ = (2/R − 1/Rcore − 1/Rcut)/(1/Rcore − 1/Rcut).

(37)

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129

044112-8 Izvekov, Chung, and Rice J. Chem. Phys. 135, 044112 (2011)

TABLE I. Coefficients An of the least-squares fit of the one-site MS-CG

force at P = 0 GPa using the polynomial equation (35) (second column) and

Chebyshev series equations (36) and (37) (third column), with power of the

term n shown in the first column. Atomic units for force and distance were

used. For the Chebyshev series, the following radii were used in Eq. (37):

Rcore = 8.5 a.u., Rcut = 29.011 a.u. The expansions were switched to zero

using the linear switching function as in Eq. (33) with δR = 4.0 a.u. for a

polynomial expansion and with δR = 2.0 a.u. for a Chebyshev expansion.

The fourth column shows the values of parameters for the density dependent

contribution Eq. (33). At R<Rcore = 8.5 a.u., the force extrapolated as discussed in Sec. III B. A cutoff of 1.535 nm must be applied to the expansions.

n An, P ol (a.u.) An, Cheb (a.u.) gi, a (a.u.)

0 0.0 0.295090311504×10−02 P = 0.5 GPa:

1 0.0 0.187126981336×10−02 0.7745×10-4, 1.0

2 −164.77839357004 −0.249111348590×10−04 P = 1 GPa:

3 39914.936788713 −0.620876162674×10−03 1.8×10−4, 1.0

4 −3312200.9193869 0.143777893709×10−03 P = 2 GPa:

5 140675713.66804 −0.445005999642×10−04 3.6×10−4, 0.9784

6 −3525421465.5670 0.429487649160×10−04 P = 3 GPa:

7 55168999115.916 −0.212373560733×10−03 5.4×10−4, 0.9679

8 −543866611821.71 0.257218649672×10−04 P = 5 GPa:

9 3253068550427.9 0.244979561419×10−03 8.6×10−4, 0.9534

10 −10456759286650 −0.761237529621×10−04 P = 10 GPa:

11 11504453156792 −0.952354112853×10−04 1.88×10−3, 0.7468

12 10486861064266 0.163080035111×10−04

13 4483868093300.0 0.389889598497×10−04

14 1370873909039.2 −0.605523294055×10−04

The nmax = 14 was selected for both expansions

Eqs. (35) and (36). The summary of polynomial representations is given in Table I. To improve energy conservation in

the MD simulation, the polynomials were smoothed at Rcut

as described in the caption to the table. In both cases, slight

differences in properties may result from simulations using

these analytical representations when compared to results

using the tabulated MS-CG potentials and forces.

C. Density dependent one-site MS-CG model

Next, the global density dependency to the MS-CG/0

model was added in accordance with the procedure outlined

in Sec. II E. The density dependence was introduced by parameterizing the set {gMSCG(R, ρi)} to reproduce the V − P

isotherm for the reference atomistic system at preselected

pressures. Because P(V ) = −∂W/∂V , where W is the total

free energy of the system, the ability of the MS-CG potentials to reproduce the atomistic P(V ) ensures a correct change

of the free energy W under a change in pressure (and, therefore, density) as well as the system’s compressibility. The latter determines the magnitude of particle density fluctuations,

and therefore thermodynamics of the particle ensembles including those under constant volume conditions. To reproduce

the atomistic ensemble thermodynamics, the density fluctuations in the CG ensemble must match those in the atomistic

system. Conclusively, the correct compressibility and, thus,

density dependency of the CG interactions is important to

have even in constant volume simulations. The compressibility can be viewed as a measure of the model transferability to

different density states in the vicinity of the reference state.

Furthermore, because MS-CG/Pi potentials represent a pairwise decomposition of the free energy of reference atomistic

system in the CG coordinates, the correspondingly parameterized MS-CG-D model should yield the correct free energetics

of the CG ensemble. The reference system was chosen to be

the ambient crystal.

Under ambient thermodynamic conditions the most stable crystalline phase of RDX is the α-polymorph, in which

RDX adopts an orthorhombic unit cell containing eight

molecules in the space group P bca.

27 The parameters gi

and a in Eqs. (33) and (34) were optimized to reproduce the

density of the bulk α-RDX from constant stress-temperature

(NsT) simulations at pressures 0.5, 1, 2 3, 5, 10 GPa and at

T = 300 K. Because of the large cutoff of the MS-CG interactions (15.35 Å), the CG simulations used a supercell composed of 4 × 6 × 6 unit cells of α-RDX. At P ≥ 2 GPa, the

target ρi could not be achieved by merely varying the repulsion parameter gi. At small gi, the crystal collapses into a

structure of a different symmetry than that of the experimental crystal. At sufficiently large gi, the target crystal structure

is stable but its density is too low. For this reason, fitting was

done in two steps. In the first step, the repulsion gi was decreased to the value for which the structure is still stable but

would undergo change with a further decrease of gi. Next, the

gMSCG

i (R) [Eq. (33)] was rescaled according to Eq. (34) with

the factor a being adjusted to achieve the target ρi. The corresponding potentials from the resulting forces gMSCG(R, ρi)

[Eq. (10)] are depicted in Fig. 2(b). The MS-CG potentials

and forces used in all simulations reported herein are in tabular form and can be obtained upon request. The simulations

with local density potentials were performed as described at

the end of Sec. II E. Because the calculation of the system total energy using Eq. (15) is computationally expensive, it was

performed only when the total energy was required (e.g., to

calculate thermodynamic quantities).

D. Crystal structure

We compared the lattice structure of the molecular c.m.

of bulk α-RDX with the crystal structure by the MS-CG-D

model at T = 4.5 K. The atomistic MD supercell was

a 2 × 3 × 3 block of crystallographic unit cells (144

molecules) of α-RDX. The MS-CG-D simulations were

performed using a 4 × 6 × 6 supercell and under constant

NsT ensemble conditions and using global density dependent

potentials as discussed in Sec II E. Remarkably, the relaxed

MS-CG-D supercell maintained orthorhombic symmetry with

the CG site positions being close to the molecular c.m. within

the corresponding atomistic supercell. The striking similarity

can be seen from Fig. 3, in which the atomistic unit cell

and the corresponding rectangular block from the MS-CG-D

supercell are superimposed. For purposes of comparison, we

will denote this latter rectangular block of eight sites as the

MS-CG “unit cell”. The unit cell geometries and corresponding densities from the atomistic and MS-CG-D models are

compared in Table II. The differences between the atomistic

and MS-CG-D lattices can be better seen in Fig. 4, and can

be characterized as small shifts of certain crystallographic

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130

044112-9 Multiscale coarse graining of crystals J. Chem. Phys. 135, 044112 (2011)

[010] 

[100] 

[001] 

FIG. 3. Comparison of RDX unit cells calculated at T = 4.5 K using atomistic and MS-CG-D models. For the atomistic model unit cell, red balls denote

the molecular c.m. locations superimposed upon the ball-and-stick representations of the molecules, whereas green balls denote the CG interaction sites.

planes. An analysis of the MS-CG lattice indicates that it is

of a higher symmetry, that of hexagonal close-packed (hcp)

A3-type (Pearson symbol hP2, space group P63/mmc). The

hP2 lattice has two sites in the unit cell and coordination

number of 12 [see Fig. 5(c)]. The highest packing efficiencies

are found in hcp and fcc lattices. However, the observed hP2

lattice is not a structure of highest density for the MS-CG/0

potential (without density dependency). In computer experiments exploring the compressibility of this crystal, we found

that under external stimuli (e.g., stress), the hP2 lattice described by the MS-CG/0 potential irreversibly transforms into

a higher density structure with crystalline domains that differ

from hP2 symmetry as shown in the insert to Fig. 4(a). The

new packing has a higher density by 3% as compared to the

hP2 lattice. This is one artifact of the MS-CG model. Similar

domains appear at the solid-liquid interface in the MS-CG-D

simulations as discussed below [Sec. III H and Fig. 7(a)].

Unfortunately, the four-site model is unable to produce

a correct crystal symmetry beyond approximately 30 K. The

CG crystal configuration evolves into a state with significant

orientational disorder regardless of the initial configuration.

However, the density remains close to the atomistic value. A

strong anisotropy of the site-site interactions within the crysTABLE II. Lattice parameters a, b, c and density ρ of the RDX crystal at

different T and P = 0 GPa from atomistic and MS-CG-D models.

a(nm) b(nm) c(nm) ρ(kg/m3)

T = 4.7 K:

atm (SB) 13.36 11.33 10.34 1885

MS-CG-D 13.93 12.06 10.10 1739

T = 298 K:

atm (SB) 13.46 11.53 10.53 1805

atm (SRT)a 13.40 11.80 10.73 1740

MS-CG-D 13.96 12.09 10.13 1726

Exptb 13.18 11.57 10.71 1806

aT = 300 K (Refs. 22 and 34). bT = 300 K (Ref. 27).

(d)

(a) 

[010] 

[100] 

[001] 

(b) 

[010] 

[100] 

[001] 

(c)

[010] 

[100] 

[001] 

[001] 

[010] 

[100] 

FIG. 4. Comparison of the α-RDX crystal lattice structure from simulations

using the atomistic (red balls denote molecular c.m.) and MS-CG-D (green

balls) models. Panel (a): [001]-view of lattice from the MS-CG-D simulation;

Panel (b): Same view as in panel (a) for the atomistic model; Panel (c): [100]-

view of superimposed lattices; Panel (d): [010]-view. In panels (a), (b), and

(c) dotted and dashed lines outline location of planes in the atomistic lattice,

which correspond to planes shown in dotted and solid lines for MS-CG-D

lattice. In panel (b), atomistic planes (dashed) move to new locations in MSCG-D lattice (solid) in panel (a). In panel (c), the two distinct planes in the

atomistic representation (dashed) collapse into a single plane in the MS-CG

representation (solid line).

tal arrangement, likely between the ring (C3N3) and NO2

sites (see Fig. 1), which cannot be represented adequately

by spherically isotropic interactions within the four-site description, is probably one reason for the non-transferability of

the four-site model to the crystalline structures. In this sense,

the one-site RDX model is advantageous as it lacks orientational dof and, thus, does not suffer from orientational melting

due to an anisotropy of molecular-molecular interactions being averaged within the one-site description. As a remedy, a

finer coarse graining, possibly of the ring group, may be exercised to build the CG models having orientational dof, which

are capable of reproducing the atomistic crystal symmetry of

the RDX.

E. Pairwise structure in molten phase

Radial distribution functions (RDFs), which map out pair

correlations, is an important test for the quality of the freeenergy based CG model. As shown in Fig. 5(a), all MS-CG/0

models produce liquid RDFs in satisfactory agreement with

the respective atomistic c.m. RDF. The one-site MS-CG/0

model matched to 550 K sampling naturally gave a better

RDF as compared to the 350/550 K potential. For P > 0 GPa,

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131 044112-10 Izvekov, Chung, and Rice J. Chem. Phys. 135, 044112 (2011)

0 0.2 0.4 0.6 0.8 1 1.2 1.4

R , nm

0

1

2

3

4

5

RDF

c.m.-c.m. , Atm

c.m.-c.m. , 1-site MS-CG , 350/550 K

c.m.-c.m. , 1-site MS-CG , 550 K

C3N3-C3N3 , Atm

C3N3-NO2 , Atm

NO2-NO2 , Atm

C3N3-C3N3 , 4-site MS-CG , 550 K

C3N3-NO2 , 4-site MS-CG , 550 K

NO2-NO2 , 4-site MS-CG , 550 K

(a)

0.4 0.6 0.8 1 1.2 1.4

R , nm

0

1

2

3

RDF

 0 GPa

0.5 GPa

 1 GPa

 2 GPa

 5 GPa

Atm

MS-CG-D (b)

0.4 0.6 0.8 1

R , nm

0

4

8

12

16

RDF , 

Nc

(c) Atm

MS-CG-D

RDF

Nc

FIG. 5. Panel (a): c.m. RDFs from atomistic (solid), MS-CG 350/550 K (dotted), and 550 K (dashed) simulations of liquid RDX at T = 550 K using onesite (black) and four-site (color) MS-CG models. Panel (b): c.m. RDFs from

atomistic (solid) and one-site MS-CG-D (dashed) simulations at different P.

Shown are P (GPa): 0 (cyan), 0.5 (black), 1 (blue), 2 (green), and 5 (red).

Panel (c): c.m. RDFs from atomistic (thick solid) and MS-CG-D (thin solid)

simulations of crystalline RDX at T = 300 K. The corresponding running

coordination number Nc shown in dashed lines.

we employed the ad hoc procedure of Eq. (33) to construct the

reference gMSCG(R, ρi) [Eq. (13)] and thus, the results may

not yield the same result as that produced by direct application of the force-matching. Therefore, we have tested the validity of the MS-CG-D model produced using Eqs. (8), (33),

and (34) by comparing the predicted liquid RDFs under different pressures with the atomistic simulations [both shown in

Fig. 5(b)]. As is evident, the MS-CG-D model describes adequately the pressure dependence of the liquid structure. For

example, the first maximum in the atomistic RDF decreases

in intensity by 0.18 and shifts by 0.041 nm to shorter separations as the pressure increases from ambient to 2 GPa, in

satisfactory agreement with the respective values of 0.17 and

0.040 nm from the MS-CG-D model. At P = 5 GPa, structure appears in the first peak in the atomistic RDF reflecting

an increased molecular order in the first coordination shell.

For liquid NM, the one-site MS-CG model, which was forcematched explicitly to a liquid phase under 5 GPa, was able

to reproduce accurately similar increases in structure in the

first coordination shell under high pressures.5 The MS-CGD model using the ad hoc density dependency is not able to

capture such changes in the effective interaction. However,

the MS-CG-D RDF exhibits a step-like structure in the front

slope, which gradually develops as more pressure is applied.

This structure vaguely resembles the shape of the front slope

of the atomistic RDF and the similarity becomes more pronounced with increasing pressure. These observations suggest

that the gi used in Eq. (33) may reflect relevant pressure effects in the effective intermolecular interactions. Finally, in

Fig. 5(c), we present the RDFs from crystal simulations. The

features in the crystal MS-CG RDF are consistent with those

of close-packed structure. In particular, the coordination is 12.

The atomistic lattice has 12 molecules within a shell of the

size of the MS-CG first coordination shell (0.8 nm), however,

the density is structured within this shell.

In Fig. 5(a), we also present a comparison of RDFs for

the four-site model [Fig. 1] force-matched to the 550 K data.

As seen, the four-site model yields the structure, which is in

very good agreement with its atomistic counterpart.

F. Many-body effects

The RDFs provide a measure of the correlation between

pairs of particles only, whereas an orientational order parameter can provide a measure of many-body correlations in a

system. Therefore, the ability of the MS-CG-D model to capture the many-body correlations can be determined by comparing the distributions of an orientational order parameter

with those generated from an atomistic model. While several

forms of orientational order parameters are available,28 we

have chosen a form similar to that introduced to study tetrahedral close-range ordering in liquid water.29 For the I th site,

we introduce the instantaneous order parameter as

ψI = 1 − 1

Nc

3

8



K=1,Nc

(cos θIK − 1)2, (38)

where the θIK is the smallest angle formed by the vector RIK

joining the I th site and the Kth closest site from the preselected coordination shell of Nc sites in the liquid configuration and the analogous vector in the reference (crystal) configuration. In the perfect (i.e., reference crystal) alignment,

�ψI � = 1 where �...� denotes ensemble averaging. If the arrangement is random, as in an ideal gas, then �ψI � = 0. The

averaging ψI among the I site and the J sites in its coordination shell yields an averaged local order parameter

ψ ̄ I = 1

Nc + 1



ψI +

J

ψJ



. (39)
132

044112-11 Multiscale coarse graining of crystals J. Chem. Phys. 135, 044112 (2011)

0.5 0.6 0.7 0.8 0.9 1

ψI

0

5

10

15

20

25

P(

ψI

)

Atm

MS-CG-D , Atm reference config 

MS-CG-D , CG reference config

Nc = 3

Nc

 = 7

Nc

 = 12

FIG. 6. Distributions of the order parameter Eqs. (38) and (39) with different coordination number Nc for atomistic (thick) and MS-CG-D (thin) models. The MS-CG-D results shown for reference crystal structure {RIK } from

atomistic (solid) and MS-CG-D (dashed) models.

The order parameter �ψ ̄ I � is a convenient parameter for

monitoring (orientational) amorphization of the crystal. The

global order parameter ψ ̄ is the average of ψ ̄ I over all sites in

the system. The level of agreement between the atomistic and

MS-CG-D distributions P(ψ ̄ I ) is a measure of how well the

local Nc-body correlations are captured by the MS-CG models. In partially ordered phases such as in a liquid, the exact

shape of the P(ψ ̄ I ) is determined by the set of reference vectors {RIK} as well by the size of the coordination shell Nc.

Fig. 6 shows the atomistic and MS-CG-D P(ψ ̄ I ) for different Nc and different reference structures. The atomistic distributions were calculated with the {RIK} for the optimized

atomistic lattice, while the MS-CG-D P(ψ ̄ I ) were calculated

for both the atomistic and MS-CG-D {RIK}. The atomistic

and MS-CG-D distributions agree remarkably well if the same

(atomistic) reference structure is used, thus affirming the ability of MS-CG-D model to capture the higher-than-pairwise

correlations in liquid state.

G. Melting

We determined the melting temperature Tmelt of the

CG crystal from the solid-liquid coexistence (two-phase)

method.30 The melting point of the MS-CG crystal was

estimated from a solid-liquid mixture with a planar interface

as shown in Fig. 7(a) and using local density dependent potentials as elaborated in Sec. II E. The interface was prepared

by joining liquid and solid slabs of material simulated at

480 K. The sample contained 9206 sites. The interface moves

into the solid phase if T >Tmelt and moves in the opposite

direction or remains stationary if T <Tmelt . The dynamics

of the liquid-solid interface was monitored using temporal

profiles of ψ ̄ (z), which is calculated by averaging ψ ̄ I within

a slab of a preselected width centered at z (see Fig. 11, where

similar profiles shown for the shocked sample). The Tmelt

was determined from a series of NsT simulations at different

temperatures starting with T = 480 K, and increasing the

300 350 400 450 500 550 600

T , K

0.94

0.95

0.96

0.97

0.98

0.99

1

ψ

20 40 60 80

t , ps

0.94

0.96

0.98

1

550 K

450 K

(b)

(a)

FIG. 7. Panel (a): Snapshots of the liquid-solid interface from simulation

of slab melting at T = 490 K, at which close-packed structure formed at

the interface (upper), and more common configuration (bottom). In the upper snapshot, the closed-packed domains of different from hP2 structure are

contoured by dotted line. Such domains are artifact of the model reflecting

limited transferability of the model to heterogeneous structures such as interfaces. These domains are similar to those shown in inset to Fig. 4(a) and discussed in the text. Panel (b): Equilibrium global order parameter, �,  ̄ Eq. (39)

with Nc = 12 as a function of T for crystal. Insert shows � ̄ (t) for crystal at

T = 450 and 550 K.

temperature in increments of 5 K. The melting temperature

was obtained (Tmelt = 490 K) and is in reasonable agreement

with the experimental value of 462 K. The atomistic SRT

model used with a flexible intramolecular potential energy

function predicts Tmelt = 439 K.31

The ψ ̄ (T ) curve for the bulk crystal is shown in Fig. 7(b).

As seen the crystal needs to be superheated to T = 510

− 520 K before it starts melting. Above this range, it takes

about 40 ps for the sample to melt [Fig. 7(b) inset]. In atomistic simulations of α-RDX, the crystal was superheated to

T = 510 K before melting was observed.32 Full orientational

disorder was reached at T = 540 K and translational freedom

at T = 580 K. In the MS-CG-D simulation of the perfect

crystal, domains of closed packed symmetry different from

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133

044112-12 Izvekov, Chung, and Rice J. Chem. Phys. 135, 044112 (2011)

hP2, which are similar to those shown in Fig. 4(a), first appear. The occurrence of such domains is responsible for the

finite slope in the ψ ̄ (T ) curve in the 510 − 520 K interval.

Such domains were occasionally observed at the liquid-solid

interface at T >Tmelt as contoured in Fig. 7(a). This is an

artifact of the model.

H. Elastic properties

The isothermal Young’s moduli E1,2,3 and shear moduli

G23, G31, G12 were evaluated for the MS-CG-D model at several thermodynamic state points using the direct strain-stress

method. For orthorhombic symmetry as in our case, there are

six independent elements of the strain tensor (εk, k = 1, 6).

Within the direct strain-stress approach, we first derive the

stress-strain curve from the NVT MD simulations for each εk.

The stress-strain curve was determined by uniformly deforming the simulation cell initially in a stress-free state, which has

basis vectors h0 = (a0

i ), to a new geometry h = (ai). The new

h is selected to have a strain tensor [εij ] = 1

2 ([hh−1

0 ]

T hh0 −

I), where I is a unity matrix, with only one preselected

εk �= 0 [ε1, ε2, ε3 = ε11, ε22, ε33, ε4, ε5, ε6 = 2ε23, 2ε31, 2ε12

in Voigt notation]. The stress σ for the deformed structure is

then measured from the NVT simulation. The residual stress

σ − σ0, where σ0 is the NVT equilibrium value for the h0

cell, represents a point on the stress-strain curve. For each εk,

five strain increments from 0.0 to 0.015 were applied. The

elements of the stiffness matrix (elastic constants) were then

evaluated from the slope at zero of a quadratic least-squares

fit of the stress-strain points. The elastic moduli and Poisson

ratios were then determined from the elastic constants using

the Voigt average, which gives an upper bound estimate for

shear moduli.

The calculated elastic moduli are reported in Table III.

The experimental data were obtained from acoustic measurements by Haussuhl et al.33 The atomistic simulation data

are compiled from the SRT and SB models.34, 35 At all state

points, the MS-CG-D results are in reasonable agreement

with the atomistic data, with overall better agreement with

the SB model (from which the MS-CG-D potential was

derived). One of the most prominent differences between the

MS-CG-D and atomistic results is the degree of anisotropy

among Young’s moduli. The MS-CG-D results at T = 298 K

are almost isotropic, whereas both SB and SRT models

demonstrate anisotropy in the moduli. There are two major

reasons for this discrepancy. The first is due to the use of

virial data5 (



i=1,3 σii vs. V) to match the MS-CG-D model,

resulting in a more “isotropic” σ for the MS-CG-D crystal.

Indeed, at ambient conditions, the MS-CG/0 model (without

a density dependency) yields a very low E3 = 0.7 GPa, while

the MS-CG-D model predicts the E3 as a largest component

(10-15% larger than the E1 and E2). Interestingly, the SB and

SRT models yield different ordering of Young’s moduli. For

the SRT model, E2 > E3 is in agreement with experiment,

while the SB model yields an inverted ordering, which is

carried over to the MS-CG-D model. We note that the SRT

model only describes intermolecular interactions, and the

simulations were performed within the rigid-molecule

TABLE III. Young’s (E) and shear (G) moduli at different T (K) and P

(GPa). Units are GPa.

E1 E2 E3 G23 G31 G12

T = 218, P = 0:

atm (SRT)a 27.3 27.0 17.4 9.6 6.0 9.1

MS-CG-D 24.2 20.5 22.9 8.1 8.0 4.7

T = 298, P = 0:

atm (SB) 19.3 16.2 18.5 7.6 4.9 3.0

atm (SRT)a 24.2 21.1 15.4 8.4 5.3 7.6

MS-CG-D 19.7 19.4 21.9 7.2 7.1 4.1

Expt 20.9 16.0 15.6 5.2 4.1 6.9

T = 304, P = 1:

atm (SRT)a 29.8 32.1 26.6 12.7 8.0 11.6

MS-CG-D 25.5 26.8 28.5 8.1 8.0 4.9

aReference 34.

approximation. The second reason, which in part relates to

the first one, is due to a different structure of atomistic and

MS-CG-D lattices as discussed in Sec. III D. The MS-CG-D

lattice belongs to a more symmetrical crystalline space

group that in turn contributes to the absence of a significant

anisotropy. The isothermal bulk modulus BT = 13.0 GPa was

obtained from a fit to the P − V curve [see Fig. 8(a)] using

the Murnaghan equation.36 The SB35 and MS-CG-D values of

BT are identical. If the isotropic relation BT = E/3(1 − 2ν)

is assumed, where E is the arithmetic average of Young’s

moduli, the MS-CG-D model yields the average Poisson ratio

ν = 0.24, which compares well with the SB model prediction

(ν = 0.27). Importantly, the MS-CG-D model predicts the

shear moduli in reasonable agreement with the reference

atomistic data. With regard to the shear moduli, the most noticeable qualitative difference is the degree of anisotropy. The

MS-CG-D G23 and G31 moduli have similar values, while the

atomistic models predict more anisotropy for these shears.

Most likely, this feature is due to the same reasons as the

behavior of Young’s moduli discussed above. With regard to

comparison with experiment, the SRT model results correlate

with directional response, while the SB and MS-CG-models

differ. Finally, it is notable that the MS-CG-D model exhibits

the expected decrease in moduli with increasing temperature

and increase in moduli with increased pressure.

I. Thermal properties

The MS-CG-D model accurately reconstructs the ambient V − P isotherm of crystalline RDX as displayed in

Fig. 8(a), and is most certainly a consequence of the way

the model was parameterized as described in Sec. II E.

For the molten state, the MS-CG-D V − P isotherm is

noticeably stiffer at P > 0.5 GPa. This discrepancy could

be due to the behavior of the �Uconf − P isotherm [shown

in Fig. 8(b)]. For the molten phase, the MS-CG-D model

predicts an increase in �Uconf with increasing pressure for

P < 1 GPa, whereas the atomistic model shows a decrease.

However, at higher pressures the shapes of the MS-CG-D

and atomistic �Uconf − P curves visually are similar. The

poorly predicted V − P and �Uconf − P isotherms in the

molten phase by the MS-CG-D model suggest the importance

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134 044112-13 Multiscale coarse graining of crystals J. Chem. Phys. 135, 044112 (2011)

0 2 4 6 8 10

P , GPa

0.7

0.75

0.8

0.85

0.9

0.95

1

V/V0

Atm , Cryst

MS-CG-D , Cryst

Atm , Liq

MS-CG-D , Liq

(a)

0 2 4 6 8 10

P , GPa

0

20

40

60

80

100

Uconf- U conf

0 , kJ/mol

Atm , Cryst

MS-CG-D , Cryst

Atm , Liq

MS-CG-D , Liq

(b)

FIG. 8. Comparison of isotherms for crystalline RDX at T = 300 K

(crosses) and molten RDX at T = 550 K (circles) from atomistic (solid)

and MS-CG-D (dashed) models. Panel (a): V/V0 − P isotherm, where V0 is

ambient volume; Panel (b) �Uconf − P isotherm, where �Uconf = Uconf

− Uconf

0 is the change in the total internal potential energy with respect to

the ambient value Uconf

0 .

of structural rearrangements of the molecules as pressure

increases, which cannot be captured by a MS-CG-D model

derived from crystal P − V data.

The significantly degraded thermal expansion properties

of the MS-CG crystal could be attributed to the use of P − V

data at a single temperature (300 K) to parameterize the

density dependency. At T = 300 K and P = 0 GPa, the

MS-CG-D model yields a coefficient of constant stress

thermal expansion αV = 22.3 × 10−6 1/K, 7.3 times lower

than the atomistic value of 163.8 × 10−6 1/K35 and the

experimental value of 193.4 × 10−6 1/K.37 The lower αV

from the MS-CG-D model is a direct consequence of the fact

that the temperature transferability of the MS-CG-D model

is inherited from its density dependency, while in reality the

T enters the effective interactions as an independent thermodynamic parameter. We calculated a constant volume specific

heat capacity for the MS-CG-D model using the fluctuation

formula38 as cV = 97 J/kg · K, which is about an order of

magnitude lower than the experimental value of 1088 J/kg · K.

The ratio (cP − cV )/cV = α2

V BT T /ρcV is about 0.01. In the

atomistic model cV is dominated by the excitations of the

vibrational modes of the RDX molecules; removing them

in the CG description results in cV being severely underestimated. The ambient value of the Grüneisen coefficient

� = BT αV /ρcV , which is a measure of overall anharmonicity of the crystal, is estimated for the MS-CG-D crystal as

� = 1.726. For the atomistic model assuming the experimental cV , � = 1.084 and the experimental value is � = 1.103.

The similar values of Grüneisen parameter obtained by the

atomistic and MS-CG-D descriptions indicate roughly the

same level of anharmonicity. The inclusion of molecular

structure to the MS-CG model, such as in the four-site representation, will improve the heat capacity of the system, and

therefore the equation of state which governs quasiequilibrium phenomena such as shock propagation [see Sec. III K].

J. Vibrational spectrum

In Fig. 9 the vibrational density of states (VDOS) for the

bulk crystal at T = 300 K calculated with the MS-CG-D and

atomistic models are compared. The VDOS was determined

using the power spectrum of the velocity autocorrelation

function (VACF) VACF(t) = �v(t)v(0)�. For the atomistic

system, the VDOS was obtained from the all-atom and

c.m. VACFs. The lattice part of the atomistic c.m. VDOS is

expected to be dominated by the optical modes. For a better

comparison, the atomistic VDOSs in Fig. 9 are normalized

on the integral intensity up to 50 cm−1, and the MS-CG-D

VDOS rescaled. The part of the experimental spectrum39 of

α-RDX, which is due mainly to optical lattice vibrations,

extends to about 150 cm−1 [regions (I) and (II) in Fig. 9]. The

0 50 100 150 200 250

 λ , cm -1

arbtr units

(I)

Lattice Modes:

- c.m. transl

- c.m. rot

(II)

Lattice Modes with

some NO2

 rotations:

- c.m. rot

- c.m. transl

- CNNO

Atm

Atm , c.m. 

MS-CG -D

(a)

0 500 1000 1500 2000 2500 3000

 λ , cm -1

0

1

2

3

4

5

E(λ) , 10−3

(b)

Lattice Modes

present

FIG. 9. Panel (a): All-atom VDOS (dashed), atomistic VDOS of c.m. (thin

solid), and MS-CG-D (thick solid) in arbitrary units. The vertical dashed lines

mark the parts of spectrum: (I) Only lattice modes associated with movements

and rotations of molecular c.m.; (II) Mix of c.m. lattice modes with wagging

and torsional modes of nitro groups. Panel (b): All-atom VDOS with portion

of spectrum shown in panel (a) marked with a vertical solid line. The vertical dashed line (150 cm−1) marks the low frequency region in which lattice

modes are present.
135 044112-14 Izvekov, Chung, and Rice J. Chem. Phys. 135, 044112 (2011)

major contributor to the c.m. lattice modes are translations

and rotations of the molecules; however, low-frequency intramolecular vibrations may contribute.23 The low-frequency

boundary of the spectrum is determined by the size of the

supercell. The atomistic c.m. VDOS is nearly all confined to

the < 150 cm−1 region, indicating that the SB model captures

adequately the lattice part of the spectrum. In the VDOS

region (I), which in the experimental spectrum extends to

about 60 cm−1, all vibrations can be assigned to the lattice

modes. In region (II) (60-150 cm−1 according to experimental

data), the lattice modes are mixed with the wagging and

torsional motions of the nitro groups.23 Outside the regions

(I), (II), all modes originate from intramolecular vibrations.

As seen from Fig. 9(a), the all-atom and c.m. VDOSs start

deviating at approximately 50 cm−1 due to the contribution

of intramolecular modes in the all-atom VDOS. Outside

region (II) of the c.m. VDOS, there is a weak tail stretching to

about 200 cm−1, which may correspond in part to a cluster of

vibrational modes centered at 180 cm−1 in which (according

to other studies) significant motion of nitro groups is accompanied with c.m. movements. In the atomistic model of RDX

by Boyd et al.,

40 the upper boundary of the lattice spectrum

is aligned at 190 cm−1. The VDOS from the MS-CG-D

model occupies the frequency region below 35 cm−1 and

is composed of three well defined peaks39 at 18, 24, and

27 cm−1. The first major peak exhibits splitting into three

levels. The MS-CG-D VDOS possibly corresponds to a cluster of five lowest laying normal modes, which occupy the 18

− 29.2 cm−1 range in the experimental spectrum. The modes

from region (I), which are not reproduced by the MS-CG-D

model, can be attributed to the lattice vibrations due to the

molecular rotational dof, which are projected out in the CG

description.

K. Shocked RDX

As a first step in the evaluation of the performance of the

MS-CG-D model in simulating shocked RDX, we calculated

thermodynamic properties of the explosive on its RankineHugoniot curve,

Hg(T,V ) = 0 = E − E0 − 1

2

(P + P0)(V0 − V ), (40)

which describes the locus of thermodynamic states with

the specific volume, internal energy, and pressure (V

= 1/ρ , E, P) accessible by shock loading from an initial thermodynamic state (V0 = 1/ρ0, E0, P0). The RankineHugoniot states under hydrostatic compression were calculated using the procedure described by Erpenbeck41 from

multiple NsT simulations integrated for 50 ps. We used

the 4 × 6 × 6 supercell and T0 = 300 K, P0 = 1 bar. The

E was calculated using Eqs. (9) and (15). The Hugoniot in

the P − V/V0 plane is shown in Fig. 10(a) along with

shock data by Olinger et al.,

42 and in the P − T plane in

Fig. 10(b). The MS-CG-D P − T locus is much steeper than

its atomistic counterpart due to the smaller number of dof

(by a factor 21) and low cV (referred to in Sec. III I). On

the Hugoniot locus, the MS-CG-D model goes into the in

elastic regime at about Pmelt = 0.45 GPa as it partially melts.

(b)

01234 5 6 7 8 9

P , GPa

300

350

400

450

T , K

Atm

500

1000

1500 MS-CG-D

Tmelt(P) , a

KK= 2.78

012345

P , GPa

0.8

0.85

0.9

0.95

1

V/V0

Atm

MS-CG-D

Expt , Olinger

(a)

FIG. 10. Panel (a): Hugoniot curves of RDX from the atomistic model (circles), the MG-CG-D model (triangles), and experiment (solid line). Panel

(b): Pressure vs. temperature for the calculated shock Hugoniot of RDX from

atomistic (circles) and MS-CG-D (squares) models. Dashed line depicts the

MS-CG-D melting curve using the Kraut-Kennedy relation as discussed in

the text. Right T-axis is for the atomistic model and left T-axis is for the

MS-CG-D model.

The melting temperature can be related linearly to a volume change �V using the Kraut-Kennedy relation43 Tmelt

= T 0

melt(1 + aKK�V /V0), with the constant aKK = 2(�

− 1

3 ). For the MS-CG-D model, aKK = 2.78 and the resulting melting curve Tmelt(P) is depicted by the dashed line in

Fig. 10(b). The values of aKK estimated from experimental

data are typically aKK = 1.54 − 1.66.44 The MS-CG-D estimate of aKK suggests that the CG crystal will melt at fairly

moderate shock loads above Pmelt = 0.44 GPa, which is

close to the observations from the Hugoniot calculations. For

the SB model, the elastic limit in uniaxial compression along

the [100] direction is reached at about 8 GPa.35 For a shock

along the [001] direction, elastic behavior was observed to

12 GPa.19 The temperature can be related to the specific internal energy E and V as T (V,E) = T0( V0

V )

� + (E−ES (V ))

cV . The

first term describes isentropic compression and contributes to

a temperature increase of about 25 K/GPa for the MS-CG-D

model and 15 K/GPa for the atomistic model. Because the

MS-CG-D model underestimates the atomistic E − ES(V )

by the specific rotational and vibrational kinetic energies

[Fig. 8(b)], the low cV (as discussed in Sec. III I) is a reason for superheating of the shocked CG crystal. The superheated shocked states caused by the abnormally low cV is an

artifact generic to the mesoparticle modeling and in principal

can be eliminated by extending the model to include the energy exchange between the CG particles and their internal dof.

Such a formalism requires the modeling of the intraparticle

Dldd 08 M2012 t1286316739Rditibtibjt tAIP liihthtt//ji/bt/ihtdii
136

044112-15 Multiscale coarse graining of crystals J. Chem. Phys. 135, 044112 (2011)

2

4

6

8

10

ρmol , nm

-3

Atm

MS-CG-D

0

1

2

3

T , 103 K

Tatm , K

-10 0 10 20

X , nm

0.8

0.9

1

ΨX

(b)

3

4

5

6

7

ρmol , nm

-3

Atm

MS-CG-D

0

100

200

300

400

T , K

Tatm , K

-10 0 10 20

X , nm

0.8

0.9

1

1.1

ΨX

(a)

FIG. 11. The profiles of (top frame) molecular density ρmol, (middle frame)

effective translational temperature Teff , atomistic temperature Tatm (filled

circles), and (bottom frame) global order parameter ψ ̄X [Eq. (39)] in shocked

RDX from atomistic (thick) and MS-CG-D (thin) simulations. Panel (a):

Vp = 1 km/s; Panel (b): Vp = 3 km/s. The MS-CG-D ρmol curve is shifted up

by difference �ρ0 = 0.215 in the values of atomistic and MS-CG-D molecular density of crystal at T = 4.5 K for a better comparison.

specific heat capacity ci(T,P) as a function of temperature

and pressure. The ci(T,P) function has to be chosen to restore correct atomistic Hugoniot locus in P − T plane. Several models for such mesoparticle modeling have been proposed in the past, which include the constant energy DPD

method45 and mesodynamics with an internal thermostat.46

In Fig. 11 the wave profiles for a planar shock along

the [100] direction for two particle velocities Vp = 1 and

3 km/s are presented. The shock wave was generated by impacting a flyer plate and its target as described elsewhere.47

The initial velocities of the flyer and the target were 5

3Vp

and −1

3Vp. The atomistic translational and CG temperature Ttransl is shown as an effective temperature Teff

= Ttransl/Natm, where Natm = 21 is the number of atoms

in the RDX molecule. For the CG model, the Ttransl is

a physical temperature. We have also plotted the atomistic temperature Tatm to illustrate the degree of energy

transfer from the shock load into the molecular dof. For

the atomistic system, we have calculated stress profiles

P and P[100] using the scheme of Lindahl et al.48 The

P at the shock front estimated from the stress profile is

7.5 GPa for Vp = 1 km/s, with about 20% due to intramolecular contributions and 35 GPa for Vp = 3 km/s with about

50% due to intramolecular stresses. The sufficiently close

agreement of the atomistic and MS-CG-D Teff profiles for Vp

= 1 km/s suggests that the atomistic Hugoniot with T considered as Teff is adequately reproduced by the MS-CG-D

model for shock strengths up to 7-8 GPa. For example, if we

assume that the final equilibrium state of the shock wave is

that denoted by the vertical dashed line in Fig. 11(a), then the

slopes for the Hugoniot T − P with T = Teff are estimated as

1555 K/GPa for the atomistic model and 1558 K/GPa for the

MS-CG-D model. A comparison between the Tatm and Teff

for the atomistic models in Fig. 11(a) and 11(b) provides insight into the degree of energy transfer from molecular translational dof excited upon shock impact into intramolecular vibrational modes as a function of impact energy. For Vp = 1

km/s, such energy transfer is small, whereas energy transfer from the molecular translational dof to the intramolecular vibrational modes behind the shock front is more efficient

for the higher velocity impact. The atomistic profile ψ ̄X for

Vp = 1 km/s indicates an absence of orientational melting:

the molecules for the most part relax to their original orientations upon shock wave passage, while for Vp = 3 km/s

melting is evident with significant kinetic energy imparted to

the intramolecular vibrations. Increased numbers of molecular collisions in the liquid state could explain the larger degree of vibrational heating seen for the Vp = 3 km/s impact.

Finally, we have also carried out a simulation in which the

shock load is 0.3 km/s; no melting was observed in the MSCG-D system.

IV. CONCLUSIONS

We discussed the application of the MS-CG approach,

which is based on force-matching of the atomistic ensemble, to develop particle-based CG models of molecular crystals. The particular application was the development of a

transferable one-site CG model for RDX suitable for simulating shocked crystals. The imposed requirements on the

model transferability in a broad range of thermodynamic conditions and across the phases was accomplished by introducing a density-dependent potential. The density-dependent

model was constructed in two steps. In the first step, the MSCG potential was derived explicitly through force-matching

of data sampled from simulations of molten and amorphous

RDX states at ambient pressure. The derived potential predicted an orthorhombic (hP2-A3) symmetry of the crystal

lattice, which was very close to the native geometry of the
137 044112-16 Izvekov, Chung, and Rice J. Chem. Phys. 135, 044112 (2011)

atomistic lattice, and within 5% of the density predicted by

the atomistic model. In the next step, the density-dependency

of the MS-CG potential was implemented through interpolation of potentials parameterized at several reference pressures

up to 10 GPa. In contrast to our earlier work on nitromethane,

the reference potentials were constructed using an empirical

force term fitted to the crystalline V − P isotherm. The form

of the density dependent force contribution was close to the

distance-independent term as suggested by our earlier works

on MS-CG modeling of nitromethane. This new procedure requires much less computational efforts than the direct application of the force-matching. The resultant model was able to

reproduce a wide range of the structural and thermodynamic

properties of crystalline and liquid states of RDX, including

pair (RDF) and many-body (orientational ordering) correlations, elastic and vibrational properties, and isotherms. As a

result of coarse graining of the intramolecular dof, the MSCG model predicts a specific heat that is more than 10 times

lower than that of the atomistic value. This deficiency in the

model is manifested in a significant deviation from the atomistic result in the Hugoniot P − T locus. The Hugoniot data

indicate that the MS-CG RDX crystal may melt under unrealistically low (< 0.5 GPa) shock loads, reflecting the importance of the intramolecular dof in the shock event. We can

conclude that the application of the MS-CG model to perform

realistic simulations of a shocked RDX crystal would require

the introduction of an intraparticle equation-of-state and explicit heat exchange between particles and intraparticle heat

reservoirs, such as that in the constant energy DPD method.45

To date, numerous types of CG models exist.2 All

classes of coarsening approaches, including structure-based

and top-down techniques, may result in different coarsegrain representations of the same system, each potentially

yielding different properties. Such observations exemplify

the nonuniqueness of coarse graining and the problems

surrounding the representability paradigm. The advantage of

bottom-up approaches, such as MS-CG or structure-based

methods, is the greater thermodynamic consistency with the

atomistic models as these methods systematically project the

atomistic microscopic statistical information to mesoscales. If

the properties of interest are primarily dependent on pairwise

correlations (e.g., RDFs) the structure-based methods are

quite useful. On the other hand, the MS-CG models by design

are capable of better capturing the many-body correlations as

compared to structure-based models. This makes them more

amenable to high ordered systems such as crystals. Therefore

the MS-CG models would be advantageous in simulations of

those properties in which many-body effects are important.

The reported isotropic one-site model of RDX opens

up several possibilities to extend quasimolecular modeling

of crystalline explosives to mesoscales, which are still not

accessible with existing all-atomistic models. For instance, a

knowledge of detailed molecular dynamics at the mesoscale

is of crucial importance for understanding how the sensitivity

of solid explosives is affected by the presence of extended

defects such as voids which, when shocked, lead to hotspots

via void collapse.49 Due to the low computational cost of

the one-site model of RDX, its prospective application may

include simulations of the hydrodynamic regime of void

collapse. Namely, under conditions of large voids, when the

void collapse takes much longer than the shock to traverse

the void, the void collapse dynamics evolves into the hydrodynamic regime, which is characterized by the formation of

a molecular jet that emerges from upstream wall. The jetting

may lead to a hot spot formation through enhanced focusing

of kinetic energy onto the downstream side. The detailed

dynamics of the jetting mechanism for realistic crystalline

systems is still poorly understood and in the past was studied

using highly idealized or unrealistically small systems.50

Thus, computational approaches that make the mesoscale

representation of materials tractable are needed. This effort is

a significant step towards realistic modeling at mesoscales.

ACKNOWLEDGMENTS

The authors wish to thank Dr. John Brennan and

Ms. Sarah Hamdan for helpful comments. This research was

supported by the DoD High Performance Computing Modernization Program Software Application Institute for Multiscale Reactive Modeling of Insensitive Munitions. Computing

support was provided by the DoD Supercomputer Resource

Center.

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139

Research@

Laser-Induced Plasma Chemistry of the 

Explosive RDX with Various 

Metallic Nanoparticles

Jennifer L. Gottfried

U.S. Army Research Laboratory, Aberdeen Proving Ground, MD

ABSTRACT

The feasibility of exploiting plasma chemistry to study the chemical 

reactions between metallic nanoparticles and molecular explosives such as 

cyclotrimethylenetrinitramine (RDX) has been demonstrated. This method, 

based on laser-induced breakdown spectroscopy, involves the production of 

nanoparticles in a laser-induced plasma and the simultaneous observation 

of time-resolved atomic and molecular emission characteristic of the species 

involved in the intermediate chemical reactions of the nanoenergetic material 

in the plasma. Using this method, it has been confirmed that the presence 

of aluminum promotes the ejection process of carbon from the intermediate 

products of RDX. The time evolution of species formation, the effects of laser 

pulse energy, and the effects of trace metal content on the chemical reactions 

were also studied.

Applied Optics, Volume 51, Number 7, pages B13-B21 (2012)
140

Laser-induced plasma chemistry of the explosive RDX

with various metallic nanoparticles

Jennifer L. Gottfried

U.S. Army Research Laboratory, RDRL-WML-B, Aberdeen Proving Ground,

Maryland, 21005, USA (jennifer.gottfried@us.army.mil)

Received 27 September 2011; revised 10 January 2012; accepted 13 January 2012;

posted 13 January 2012 (Doc. ID 155400); published 26 January 2012

The feasibility of exploiting plasma chemistry to study the chemical reactions between metallic nanoparticles and molecular explosives such as cyclotrimethylenetrinitramine (RDX) has been demonstrated.

This method, based on laser-induced breakdown spectroscopy, involves the production of nanoparticles in

a laser-induced plasma and the simultaneous observation of time-resolved atomic and molecular emission characteristic of the species involved in the intermediate chemical reactions of the nanoenergetic

material in the plasma. Using this method, it has been confirmed that the presence of aluminum promotes the ejection process of carbon from the intermediate products of RDX. The time evolution of species

formation, the effects of laser pulse energy, and the effects of trace metal content on the chemical

reactions were also studied.

OCIS codes: 140.3450, 300.6365.

1. Introduction

It is well known that two-component explosives consisting of metal particle fuels and oxidizers can produce more than twice the energy of high performance

molecular explosives alone. Aluminum powder (typically ∼50 μm average diameter) is frequently added

to explosives and propellants to improve their performance. The addition of metallic aluminum results in

a considerable increase in the heat of explosion and

higher temperatures since alumina (Al2O3) has a

high heat of formation [1]. Other fuels used in

heat-producing pyrotechnics include Ti, Mg, Ni, Zr,

and Be [2]. Recently there has been considerable interest in using nanoparticle fuels, which offer the

possibility of faster energy release, more efficient

combustion, and controllable explosive performance

(via particle size or passivation layer properties)

[3,4]. The development of nanoenergetic materials

has been limited by the lack of fundamental understanding regarding the chemical dynamics involved.

The need to understand the chemical mechanisms

of combustion, thermal explosion, and detonation is

essential in order to develop more efficient explosives

and propellants. In addition, the chemical products

produced during detonation can negatively affect

device performance, so an understanding of the

chemical reactions involved during the decomposition of component mixtures is crucial. In a recent

study, Song et al. investigated the formation of C

and AlO in an aluminized-cyclotrimethylenetrinitramine (RDX) shock tube explosion [5]. Both unreacted

carbon (C, C2) and AlO decrease rocket performance.

By monitoring the C2 and AlO emission, they observed that increasing the quantity of Al particles

(∼4.5 μm diameter) added to the RDX resulted in increased C2 emission. They subsequently studied the

effect of nanoaluminum on the detonation of RDX in

a similar experimental setup [6].

Here, we demonstrate an alternate approach for

studying the chemical reactions of molecular explosives and metallic nanoparticles. This approach,

based on laser-induced breakdown spectroscopy

(LIBS), involves generating the metallic nanoparticles via laser ablation and studying the chemical

reactions that occur in the laser-induced plasma by

monitoring the time-resolved emission spectra. In

the past decade, laser ablation has increasingly been

used to produce metallic nanoparticles. The size and

distribution of the nanoparticles can be controlled by

using sequential laser pulses, varying repetition

rates, laser fluence, wavelength, and pulse width,

1 March 2012 / Vol. 51, No. 7 / APPLIED OPTICS B13
141

and choice of carrier gas (air, argon, nitrogen, etc.)

[7–10]. LIBS is an analytical technique that has been

widely investigated over the past several decades for

both qualitative and quantitative material analysis

[11]; previous applications include the detection of

explosives [12]. It is only recently that a few groups

have started investigating the chemical reactions

involving explosive materials that occur during the

lifetime of the laser-induced plasma (tens to hundreds of microseconds) and the subsequent effect

on the LIBS signature [13–18].

The potential advantages to this approach to

studying nanoenergetic materials include (1) little

or no sample preparation is needed (the laser directly

ablates the explosive and substrate material), (2) the

intermediate chemical reactions of the nanoenergetic

material can be studied on a small scale (microgram

quantities), eliminating the need for a shock tube or

other explosive containment apparatus, (3) the properties of the laser (pulse energy, wavelength, pulse

duration) can be tuned to control the size and distribution of the particles formed, and (4) time-resolved

relative concentrations of a large number of atomic

and molecular species can be tracked simultaneously. With this approach, any type of material

can be ablated with the laser and its plasma chemistry studied as long as the laser energy exceeds the

breakdown threshold.

In this work, the laser ablation of thin film residues of RDX on various metal substrates (generating

nanoenergetic particles) has been combined with the

observation of time-resolved optical emission from

the resulting high-temperature plasma. The relative

concentrations of C, C2, CN, H, N, and O (as well as

any other metal substrate-related atomic, molecular,

or ionic species emitting from 200–940 nm) were

tracked during the lifetime of the plasma (typically

<1 ms). Several experiments designed to understand

the time dependence of species formation in the plasma, to determine the effect of bulk and trace metals

on the chemistry of RDX, and to study the effects of

laser pulse energy were performed.

2. Experimental Methods

Colleagues at the U.S. Army Research Laboratory

provided Class 1 (<850 μm particle diameter), military-grade RDX. Double pulse LIBS spectra of a thin

layer of RDX crushed on the surface of a 1-mm-thick

Al substrate (Sigma-Aldrich, 99.999%) were obtained

using two Big Sky CFR400 lasers (1064 nm, 225 mJ

per laser) focused on the sample surface with a 10 cm

focal length convex lens. The interpulse delay between the two laser pulses (Δt  2 μs) was selected

to minimize the amount of atmospheric contribution

to the O, N, and H emission signals. A pierced mirror

was used to collect the spatially integrated plasma

emission in a 600 μm fiber optic. The fiber was inserted into a Catalina Scientific echelle spectrometer

(SE200) through a 25 μm pinhole (200–1000 nm). An

Apogee (AP2Ep) ICCD (gain  600) served as the

detector. The gate delay (tdelay) was varied from 0

to 10 μs while the integration time (tint  1 μs) was

held constant at the minimum value for the detector. Ten single-shot spectra at each gate delay were

recorded.

To study the laser-induced plasma chemistry, a

second double pulse LIBS setup was employed. For

this experiment, two Nd:YAG lasers (Continuum

Surelite, 1064 nm, maximum ≃420 mJ per laser)

were focused onto the sample surface with a 10 cm

lens. The plasma emission was directed into a

400 μm fiber optic using a pierced mirror setup.

An echelle spectrometer (Catalina Scientifc EMU65) was paired with an electron multiplying CCD

(EMCCD) detector (Andor iXon, gain  2) to collect

the LIBS spectra (200–1000 nm) under an argon flow

with the following timing parameters: Δt  2 μs,

tint  50 μs, and tdelay  1 μs. Both detectors [intensified CCD (ICCD) and EMCCD] were corrected

for their spectral responses using a calibrated deuterium tungsten–halogen light source. For the samples used to study the plasma chemistry, the RDX

was suspended in acetonitrile, which enables fairly

reproducible, semiquantitative sample deposition.

Typical RDX concentrations of ∼10 mg∕ml were applied to the substrate surfaces with multiple deposits

from a 10 μl syringe, resulting in residue concentrations of ∼2.5 μg∕mm2; approximately 13 μg of RDX

was sampled with each laser shot. The substrates

used for this experiment were high-purity metals

(≥1 mm thick) obtained from Sigma-Aldrich: Al

(99.999%), Cu (99.999%), Ni (99.98%), Sn

(99.998%), and Ti (99.998%). Fifty single-shot spectra of RDX residue on each of the pure metal substrates were obtained. Spectra of RDX residue on

an Al2O3 sample from our laboratory and Al alloy

standard reference material (SRM) purchased from

the National Institutes of Standards and Technology

(NIST SRM 1256b, alloy 380; NIST SRM 1259,

alloy 7075; NIST SRM 1715, alloy 5182) were also

acquired.

Additional experiments performed using the Continuum double pulse laser setup include the sampling of mixtures of RDX with <75 μm diameter Al

powder (Sigma-Aldrich). The Al powder was added

to 8 mg of RDX in several weight ratios (0∶1, 1∶1,

2∶1, and 4∶1). The mixtures were then spread onto

a 6.5 cm2 area on the surfaces of the pure metal substrates Al, Cu, Ni, Sn, and Ti. Fifteen single-shot

spectra were acquired under argon for each sample.

Unlike the pure RDX residue, which tends to stick

easily to substrate surfaces, a significant portion of

the RDX/Al mixtures was blown off the substrate

by the shock wave from each laser shot. An exhaust

port was placed near the point of laser ablation to ensure none of the explosive mixtures were inhaled or

otherwise dispersed into the laboratory.

The dependence of the emission spectra on laser

pulse energy was studied using the Continuum

double pulse laser with the echelle/EMCCD spectrometer. As before, the RDX was suspended in acetonitrile and deposited on the pure Al substrate

B14 APPLIED OPTICS / Vol. 51, No. 7 / 1 March 2012
142

(∼2.5 μg∕mm2). Twenty spectra of both the pure substrate and the RDX residue were acquired under

argon using a single 210 mJ laser pulse, a double

210 mJ pulse (total energy 420 mJ), a single

420 mJ pulse, and a double 420 mJ pulse (total energy 840 mJ).

3. Results and Discussion

A. Time-Resolved Emission Spectroscopy

Background-corrected peak emission intensities for

the atomic and molecular species relevant to RDX

were tracked over the first 10 μs of the laser-induced

plasma lifetime (Fig. 1). Although the emission spectra collected with a 0 μs gate delay (and a 1 μs gate

width) contain the background continuum characteristic of bremsstrahlung emission (which persists for

the first several hundred nanoseconds), discrete

emission lines were also observed. After subtracting

the average background signal near the relevant

emission line, the peak intensities for the C

(247.856 nm), H (656.285 nm), N (746.831 nm),

and O (777.194∕777.417 nm) lines were normalized

to the total emission intensities of selected atomic

and molecular species, including CN (388.340 nm),

C2 (516.520 nm), and AlO (484.220 nm). Because

the RDX residue was applied to an aluminum substrate, Al I and Al II emission features were observed

in addition to AlO (formed by the reaction of the Al

from the substrate and O liberated from the RDX).

The cause of the initial increase in the normalized

O intensity over the first several microseconds is unclear, but may be the result of slower dissociation reactions involving the release of atomic O. The O

(after 2 μs), N, and H emission from the RDX slowly

decays over time, while the C emission is relatively

constant for the first 10 μs of the plasma lifetime. The

energy required to promote C into the excited state

(which subsequently emits radiation at 247.856 nm

as the electron relaxes to a lower energy level) is

much lower than that for the O, N, and H lines

(7.68 eV compared to 10.7 , 12.0, and 12.1 eV, respectively) [19] and is therefore less affected by the

decreasing thermal energy in the plasma at later

times. Because the double pulse LIBS technique

[20] was used to collect the data, only a minimal

amount of air was entrained into the plasma and

most of the O, N, and H signals were from the

RDX residue (and potentially contaminants on the

aluminum surface).

As previously reported for other materials with

much lower laser energies [21–23], the molecular

species exhibit very different time-dependent behavior in the laser-induced plasma. Of the three

molecules observed in the LIBS spectra of RDX on

aluminum, only the CN is present at very early times

in the plasma lifetime since some CN fragments are

formed by the initial laser ablation of the RDX

molecule (Figure 2). These initial CN fragments

are dissociated into atomic C and N atoms by collisions in the high-temperature plasma [13]. The ratio

of singly ionized Al (466 nm) to neutral Al (309 nm) is

also shown in Fig. 2. The excitation temperatures

calculated based on the Boltzmann two-line method

using the Al I lines (309 and 394 nm) follow the same

trend as the ionization ratio—both decrease as the

plasma cools. The calculated temperatures drop from

∼11; 000 K at a gate delay of 0 μs to around 3700 K at

a gate delay of 10 μs. As the plasma begins to cool,

CN, C2 and AlO form from chemical reactions, e.g.,

Eqs. (1–4) [13], Eq. (5) [24], and Eq. (6) [25], resulting

in an increase in molecular emission at times >2 μs:

C  N  M → CN  M; (1)

C  N2 → CN  N; (2)

CO  C → C2  O; (3)

CN  C → C2  N; (4)

Al  O  M → AlO  M; (5)

Al  CO2 → AlO  CO. (6)

The formation of molecular species during the lifetime of the laser-induced plasma is clear evidence

Fig. 1. (Color online) Peak atomic emission intensities as a function of gate delay (gate width 1 μs). Error bars represent 95%

confidence intervals.

Fig. 2. (Color online) CN, C2, and AlO emission (shown as a function of gate delay with a gate width  1 μs) increases as the plasma cools and becomes less ionized. Error bars represent 95%

confidence intervals.

1 March 2012 / Vol. 51, No. 7 / APPLIED OPTICS B15
143

that chemistry is occurring in the highly excited,

nonequilibrium environment. While most LIBS practitioners therefore limit the detector gate width in order to minimize changes in the emission intensities

(at the expense of sensitivity), our goal was to determine if useful information about the species involved

in the plasma chemistry can be obtained using longer

gate widths. While using longer gate widths results

in stronger overall emission intensities, changes in

emission intensities resulting from chemical reactions in the plasma (as well as changes in the plasma

temperature) also affect the relative emission intensities for the different species in the plasma over the

integrated plasma lifetime.

B. Plasma Chemistry

Selected regions of spectra (double pulse LIBS under

argon) from the five high-purity substrates with and

without RDX are shown in Figure 3. The increase in

the C, H, N, O, and CN emission with the presence of

RDX (dotted red traces) is clearly visible. In this experiment very little air was entrained in the laserinduced plasma, so most of the nitrogen used in

the formation of CN originates from the RDX. The

highest excitation temperatures for the RDX/metal

plasmas (calculated using the Ar I lines at 750.387

and 751.465 nm) correspond to Ti > Al > Cu > Sn >

Ni (Table 1). The Ti substrate plasma had the greatest increase in temperature with the addition of

RDX, while the plasma for the Sn substrate remained essentially unchanged. The higher plasma

temperature for the blank Sn implies less laser energy coupled to the sample and more energy went

into heating the plasma. In this case, there would

be fewer ablated Sn atoms in the plasma to affect

the chemistry.

The extent of ionization for each RDX/metal plasma

(represented by the ion/neutral ratio) followed almost

the same trend as the temperature, Ti > Al > Cu >

Ni > Sn (Fig. 4). The ion/neutral ratios are drastically

different among the substrates. Al and Ti have the

lowest ionization potentials of the five substrates

(5.986 and 6.828 eV, respectively, compared to

7.726 eV for Cu, 7.640 eV for Ni, and 7.344 eV for

Sn) [26], and their plasmas have the highest ion/

neutral ratios. The ionization of the Al plasma is substantially decreased when RDX residue is added to

the substrate, while the ionization in the Ti and Cu

plasmas increased. The ionization of the Ni and Sn

substrate plasmas showed little change with the

addition of RDX residue. These properties of the substrate will significantly affect the chemical reactions

occurring in the laser-induced plasma.

Figure 5 shows three contour plots created by plotting the relative emission intensities of species within the laser-induced plasma (normalized to the Ar

line at 763.511 nm) generated by ablation of RDX

residue on the five substrates. These plots illustrate

the variations in the plasma chemistry caused by the

different substrates. Figure 5(a) shows that, as the C

intensity in the plasma increases, the CN intensity

increases, as well. The CN intensity is essentially independent of the atomic N content, however, which

indicates that Eq. (2) is the dominant mechanism

for formation of CN in the plasma, rather than

Eq. (1). This observation is consistent with previous

results (e.g., [23]). The O intensity increases as the N

increases, while increasing C content results in decreasing O [Fig. 5(b)]. This result can be explained

by the formation of CO and CO2 in the plasma via

the reactions

C  O  M → CO  M; (7)

CO  O → CO2. (8)

Equations (7) and (8) are important reactions in the

decomposition of RDX, since the formation of gaseous

Fig. 3. (Color online) Selected regions of LIBS spectra for (a) Al, (b) Cu, (c) Ni, (d) Sn, and (e) Ti with (dotted red) and without (solid black)

RDX residue.

Table 1. Calculated Excitation Temperatures for Metal Substrates

With and Without RDX Residue

Temperature (K)

Substrate Blank RDX Residue % Change

Al 9594  251 12195  796 27%

Cu 9369  142 11802  611 26%

Ni 9839  177 11335  488 15%

Sn 11639  398 11434  585 −1.8%

Ti 9123  176 16834  1295 85%

B16 APPLIED OPTICS / Vol. 51, No. 7 / 1 March 2012
144

products such as CO and CO2 is highly exothermic

and results in the liberation of large quantities of

energy upon detonation [27]. Figure 5(c) shows that,

as the CN intensity increases, the O intensity decreases. This decrease in O content in the presence

of CN can be explained by the reaction

CN  O → CO  N. (9)

The concentration of atomic O in the laser-induced

plasma is, therefore, highest in conditions of low C

and low CN.

The assumption for the preceding analysis is that

the minor differences in plasma temperature between substrates with RDX residue on them (with

the exception of Ti, see Table 1) are less important

than the effects of the ablated metal on the plasma

chemistry. Since the subsequent analysis agrees with

previous studies, we believe this assumption is reasonable. For a more quantitative approach to this

analysis, inclusion of the temperature and ground

state populations would need to be considered. In addition, the plasma chemistry is extremely complex,

and there may be further complications not considered in this simplified approach.

To further understand the effect of metals on the

chemical reactions of RDX in the laser-induced plasma, Al powder was mixed with the RDX in varying

concentrations (0∶1, 1∶1, 2∶1, and 4∶1) and applied

to the five substrates. Al is rapidly oxidized by oxygen (O, O2), CO2, H2O, and possibly other species,

such as NO, N2O, and NO2, to form AlO before reaching its final equilibrium state, alumina (Al2O3) [28].

Because of the high heat of formation for Al2O3, the

addition of Al to explosive formulations results in a

considerable increase in the heat of explosion. In

these experiments, the laser-induced plasma temperature for each substrate increased by as much

as 55% with the addition of Al powder to the RDX

(Fig. 6). The Ti was the only substrate that showed

an initial decrease in plasma temperature with Al

(compared to the pure RDX residue).

As the concentration of Al powder was increased,

the substrate emission lines decreased (due to increased residue surface coverage) and the Al emission lines increased. While the C, H, N, H, and CN

emission lines decreased with increasing Al (possibly

because less RDX was being sampled in the plasma),

the C2 and AlO increased on all substrates (Fig. 7).

This result is confirmed by the conclusions of

Song et al. [5], who observed the increase in C2

and AlO emission with increasing Al content during

an aluminized-RDX explosion in a shock tube. The

higher the concentration of Al in the laser-induced

plasma, the more O the Al scavenges to form AlO,

and the less O available to react with the C to form

CO. This results in an increase in condensed phase C

aggregate (soot) and C2. Although the increase in

Fig. 4. Ion/neutral intensity ratios for the five high-purity substrates with and without RDX residue. The Al I 669.867 nm

and Al II 466.306 nm lines were used to calculate ion/neutral ratios

for Al since the 309.271 nm line was saturated. Other emission lines used for calculating ion/neutral ratios included: Cu I

296.116 nm and Cu II 227.626 nm, Ni I 343.356 nm and Ni II

243.789 nm, Sn I 333.062 nm and Sn II 533.236 nm, Ti I

373.890 nm and Ti II 462.928 nm. Error bars represent 95%

confidence intervals.

Fig. 5. (Color online) Contour plots for the (a) N, C, CN, (b) N, C,

O, and (c) CN, N, O emission intensities of RDX on five different

metal substrates (Al, Cu, Ni, Sn, and Ti).

1 March 2012 / Vol. 51, No. 7 / APPLIED OPTICS B17
145

plasma temperature is more significant in this set of

experiments, the accumulation of soot with increasing Al content has been observed in the previous

experiment [5], substantiating our assumption that

the increase in C2 emission results from chemical

reactions in the plasma.

C. Laser Pulse Energy Dependence

Conventional formulations of energetic materials

containing Al particles to improve the performance

of explosives and propellants use particles with a

mean diameter of ∼50 μm, but the reactivity of Al

nanoparticles has been shown to increase with decreasing size [29]. Previous experiments have demonstrated that ablation of an Al target with a pulsed

laser produces Al nanoparticles with a mean primary

particle size of 5–500 nm (depending on laser pulse

duration and fluence as well as carrier gas type and

pressure) [7,9,30]. It has also been demonstrated that

using higher laser pulse energies increases the number of smaller nanoparticles produced [10].

A comparison of the LIBS spectra of RDX on Al obtained using different pulse energy schemes (ranging

from 210 to 840 mJ) showed that higher laser pulse

energy results in both increased overall spectral

emission and relative AlO emission (Fig. 8); the smaller Al particle sizes produced by the high laser pulse

energy are more reactive (due to the higher relative

surface area) and, therefore, are oxidized more

quickly. While the slight increase in plasma temperature observed with increasing laser pulse energy affects the observed emission intensities, the standard

deviation for the calculated temperatures was only

505 K, so we would expect this effect to be minimal.

The double laser pulses result in stronger RDX

spectra (compared to single pulse spectra of the same

total pulse energy), but weaker AlO emission, since

less O is entrained in the laser-induced plasma (as

seen by comparing the O intensity from the blank

Al at 420 mJ versus 2 × 210 mJ in Fig. 8). The relative C intensity increases with increasing laser pulse

energy. Unlike with the AlO, there is no decrease in C

emission between a single 420 mJ pulse and a double

2 × 210 mJ pulse. Despite the strong C signal, no C2

was observed at any laser pulse energy. RDX does not

contain any C─C bonds, so no C2 fragments are

formed by the initial laser ablation (Fig. 2). Without

the presence of additional Al to scavenge the O, the O

in the plasma reacts with C (Eq. 7), preventing the

formation of significant amounts of C2 [15].

D. Effect of Trace Elements on RDX Chemistry

In addition to studying the interaction of pure metal

substrates and Al additive with RDX, we also investigated the effect of minor impurities or trace elements on the chemical reactions occurring in the

laser-induced plasma. LIBS spectra of a series of primarily aluminum samples were obtained, including

an Al2O3 substrate and three NIST standard reference material Al alloys. Table 2 lists the concentrations of the trace elements for each sample (when

known), divided into two categories: those observed

in the double pulse LIBS spectra in argon, and those

that were not observed. Only the Al2O3 substrate

spectra contained significant O emission lines, indicating that the atmospheric contributions to the

LIBS signal were negligible (none of the blank substrates had measurable N lines). All five substrates

have C and H lines, even after multiple cleaning

shots to remove surface contamination.

Fig. 6. (Color online) Excitation temperatures for RDX/Al mixture residues on Al, Cu, Ni, Sn, and Ti substrates calculated using

Al I lines at 308.215 and 396.152 nm. Error bars represent 95%

confidence intervals.

Fig. 7. (Color online) As the concentration of Al in the RDX/Al

mixture was increased, the C2 (top) and AlO (bottom) emission

from the laser-induced plasma increased on all five metal substrates, including the Al and Cu substrates.

B18 APPLIED OPTICS / Vol. 51, No. 7 / 1 March 2012
146

RDX residue was applied to the five aluminum

substrates and Fig. 9 shows the C emission intensities (normalized to the Ar line at 763.511 nm) observed in the LIBS spectra. The high-purity

aluminum substrate, with the fewest trace elements,

gave the strongest C signal from the RDX residue.

The presence of more Al in the plasma increases

the C signal by reducing the amount of free O available to react with C via Eq. (7). However, much of this

increase in C emission intensity can be attributed to

the higher plasma temperature with the pure Al substrate (12;086  356 K). The four Al alloys, on the

other hand, produce nearly identical plasma temperatures when the RDX residue is present (within

the standard deviation of the shot-to-shot variation).

For example, alloy 5182 (with RDX) results in a plasma temperature of 10;565  145 K, compared to

10;608  243 K for alloy 7075—yet the increase in

C emission signal is nearly fourfold for alloy 7075.

Differences in the plasma chemistry resulting from

the trace elements present in the aluminum alloys

are clearly evident. Alloy 7075 has the highest Zn

concentration, and the highest C signal of the three

alloys. Alloy 5182, despite having the highest Al concentration of the three alloys, has the highest Mg

concentration. Since Zn has a significantly lower

oxidation potential (0.76 V) than Mg (2.37 V) [26],

it makes sense that alloy 7075 would have the highest C signal among the alloys and alloy 5182 would

have the lowest—the presence of easily oxidizable

species will decrease the amount of O available to react with the C. Alloy 380 also contains a significant

amount of Si, which also has a low oxidation potential (0.91 V) [26]. Alumina has the highest O content,

which would typically favor the formation of CO∕CO2

(i.e. less C emission); however, this effect may be offset by its higher aluminum content than the other

three alloy substrates.

4. Conclusions

Although, at first glance, an explosive detonation and

a laser-induced plasma may not seem to have much in

Fig. 8. Emission intensity for AlO (top) and O (bottom) obtained

from RDX residue on Al with a single 210 mJ pulse, a 420 mJ

pulse, and double pulse excitation (2 × 210 and 2 × 420 mJ). Error

bars represent 95% confidence intervals.

Table 2. Trace Metal Content (Including Elements Observed in LIBS Spectra) for Five Different Aluminum Samples

Minor Components

Sample % Al Observed Not Observed

Alumina (Al2O3) unknown Si, Mg, Fe, Ti, Sr, Cr, Pb, Cu Mn, Zn, Ni, Sn, Be

Alloy 380 82.99 9.362% Si, 3.478% Cu, 1.011% Zn, 0.865% Fe,

0.3857% Mn, 0.4135% Ni, 0.0188% Sr, 0.877% Ti,

0.0572% Cr, 0.0637% Mg, 0.1075% Pb, 0.0212% V

0.35% Sn

Alloy 7075 89.76 5.44% Zn, 2.48% Mg, 1.60% Cu, 0.025% Be, 0.173% Cr,

0.18% Si, 0.205% Fe, 0.079% Mn, 0.063% Ni

Alloy 5182 94.58 4.474% Mg, 0.3753% Mn, 0.1553% Si,

0.0494% Cu, 0.199% Fe

0.034% Cr, 0.015% Pb, 0.0195% Ni, 0.0002% Sr,

0.0335% Ti, 0.0174% V, 0.0505% Zn

Aluminum 99.999 Si, Cu, Fe Mn, Zn, Sn, Pb, Ni, Sr, Ti, Cr, V, Be

Fig. 9. Emission intensities for the C line from the LIBS spectra

of RDX residue on five different Al samples. Error bars represent

95% confidence intervals.

1 March 2012 / Vol. 51, No. 7 / APPLIED OPTICS B19
147

common, they are both very high-temperature (thousands of Kelvin), high pressure (up to ∼105 atm) environments subjected to a shock wave resulting in the

rapid decomposition of the explosive into atomic constituents and subsequent exothermic formation of

gaseous products. Understanding the chemistry that

occurs in the reaction front of an explosive detonation

is extremely difficult due to the short time scale and

violent release of energy in the form of heat, sound,

light, and blast. In addition, full-scale detonation testing is quite expensive. The ability to study the chemical reactions of explosives in a small-scale laboratory

environment that does not require extensive safety

precautions would greatly benefit the development

of improved explosive formulations. The use of optical

emission spectroscopy to study the chemical reactions

of laser-generated nanoenergetic materials in laserinduced plasmas is extremely promising. By monitoring the emission intensity of different reactant species as a function of time, a better understanding of

the chemistry of metallic nanoparticles and molecular

explosives at high temperatures can be achieved,

eventually enabling the development of explosive formulations with higher explosive power and fewer

harmful byproducts.

Time-resolved, broadband emission of chemical

species involved in the reaction of RDX and metallic

nanoparticles in a laser-induced plasma has been observed. Using the methodology presented here for

investigating the chemical processes involved in nanoenergetic material reactions, we observed the increase in C2 emission resulting from an increase in

Al powder additive, as confirmed by the observations

of Song et al. [5] in a shock tube explosion. The time

evolution of species formation, the effects of laser

pulse energy, and the effects of trace metal content

on chemical reactions were also studied. In addition

to matrix effects induced by the interaction of the laser with the metal substrate (affecting the amount of

material ablated, plasma temperatures, electron

densities, etc.), we have shown that differences in

the plasma chemistry of RDX with ablated metals

significantly alters the LIBS spectra of RDX.

In a recent kinetic modeling study by Ma and

Dagdigian [23], the effect of Al particles ablated from

the substrate surface on the plasma chemistry was

neglected based on assumptions as to the entrainment of the ablated Al in the laser-induced plasma

[16]. Their initial experimental results [16], which

formed the basis for this assumption, were obtained

under very different experimental conditions than

the present study (a single laser pulse of 20 mJ at

355 nm on a thin Al foil was used). In addition,

the effect of the presence of organic residue on the

laser–material interaction was not considered. Their

resulting kinetic model significantly underestimated

the molecular emission of C2 (and CN) from the plasma [23]. Our current experimental results showing

that increasing the Al content in the plasma results

in increased C2 emission as a result of the preferential oxidation of Al over C suggest that including the

effects of Al in the kinetic model may improve the

results.

Future improvements to the current experimental

setup should include the incorporation of on-line

monitoring of particle size distribution [10] in order

to directly correlate the plasma chemistry to particle

size effects, and the monitoring of mid-IR emission

from the plasma to obtain information on the concentration of important species such as CO, CO2, NO,

and NO2 in the plasma [31]. Experiments are underway to study the particle sizes produced by laser

ablation of RDX. Additional experiments to be performed include tuning the laser properties, such as

pulse duration and wavelength, to adjust the size

and size distribution of the ablated particles, and exploring other explosive formulations, such as TNT/Al

and Al/Teflon.

The author thanks Dr. Frank C. De Lucia, Jr. for

numerous thoughtful discussions relating to this

work.

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Biographies of ARL Authors

Jan L. Allen is a chemist with expertise in materials and solid state chemistry. Dr. Allen’s interest is focused on Li-ion battery 

materials, with particular emphasis on electrode and electrolyte materials. He received his B.S. in chemistry from Truman State 

University in Missouri in 1989. He continued his education at Northwestern University and earned a Ph.D. in inorganic chemistry 

in 1993, under the direction of Professor Ken Poeppelmeier. Following postdoctoral research at the University of Caen in Caen, 

France, with Professor Bernard Raveau, and at the University of Kentucky with Professor Peter Eklund, Allen worked at the Gillette 

Research Institute, the corporate laboratory of the Gillette Company in Gaithersburg, Maryland. In 2000, Allen joined ARL. Dr. 

Allen received the Department of the Army Research & Development Achievement Award for Technical Excellence in 2002 and 

2011. 

Radhakrishnan Balu is a researcher in computational material science, with a research focus on energetic materials 

and nanomaterials used in electronics and optical devices. He has been working at ARL since 2007 and uses state-of-theart quantum mechanics-based calculations leveraging the high performance computing infrastructure for research. He has 

over 20 publications, including peer-reviewed journal papers. He has five degrees in science and engineering, with a Ph.D. 

in computational chemistry from the University of Maryland Baltimore County, an M.S. in biotechnology from Johns Hopkins 

University, and a B.S. in computer science from the Indian Institute of Science Bangalore.

Richard Becker received a B.S. and M.S. in mechanical engineering from the University of Pittsburgh. His M.S. thesis work 

on micromechanical modeling of ductile void growth was continued at Brown University, where he received his Ph.D. in 

engineering. Dr. Becker accepted a position at Alcoa, pursuing modeling of crystallographic texture evolution and anisotropy, 

thermo-mechanical process modeling, and constitutive model development. In 1999, he went to Lawrence Livermore National 

Laboratory, where he worked on algorithm and code development, and multi-scale model development and validation for strength 

and fracture models under the Advanced Supercomputing Initiative. Dr. Becker joined ARL in 2009, where he has been focusing 

on multi-scale modeling, material model development, and model implementation in large-scale hydrocodes. Dr. Becker is on the 

editorial board of several journals and is a member of the American Academy of Mechanics, the American Society of Mechanical 

Engineers, and the Sigma Xi Scientific Research Society.

Jennifer Ciezak-Jenkins is a research physicist within the Energetic Materials Science Branch. Her research examines the 

properties of materials under thermomechanical extremes using optical spectroscopy, x-ray and neutron probes, and diamond 

anvil cell techniques. Dr. Ciezak-Jenkins began her research career in molecular spectroscopy and electronic structure theory 

and obtained an NRC postdoctoral fellowship with ARL and Dr. Betsy Rice in 2004. Dr. Ciezak-Jenkins attended St. Bonaventure 

University, where she studied chemistry and physics (B.S., 2001). She did her graduate work in chemical physics at Syracuse 

University (Ph.D. 2004). After an NRC post-doctoral fellowship with Drs. Betsy Rice and Samuel Trevino (2004-2006), she joined 

ARL as a research physicist in 2006. She is the recipient of five ARL Director’s Research Initiative (DRI) grants concerning 

polymeric nitrogen, structural bond energy release, ultrahard phases of carbon, and low energy nuclear reactions, and is a 

member of the American Physical Society, the American Chemical Society, and the Neutron Scattering Society of America. 

Peter Chung is a Team Leader in the Simulation Sciences Branch, Computational Sciences Division of CISD. He holds a Ph.D. 

in mechanical engineering from the University of Minnesota and a B.S. in aerospace engineering from the University of Virginia. 

After a one-year NRC Post-doc, he joined the civil service at ARL in 2001. He is the recipient of the 2010 Army Science Conference 

Best Paper Co-Award, 2008 ICCES Outstanding Young Investigator Award in Computation, 2004 US RDECOM Outstanding Young 

Scientist, and 2003 Army Superior Civilian Service Award. He has served as mentor for 15 current and former post-docs and 

students, and has participated on eight Ph.D. and M.S. degree committees. He currently leads a research team in areas related 

to computational basic research of multiscale science, materials, physics, chemistry, and mechanics. He has published 38 

archival journal papers and 42 conference papers and presentations, and has given nine invited talks.

Wayne A. Churaman received B.S. and M.S. degrees in electrical engineering from the University of Maryland, College Park, in 

2005 and 2010, respectively, with a specific focus on microelectronics and microrobotic platforms. He has been working as a 

MEMS Engineer for the ARL since 2005. His interests include experimental analysis of nanoenergetic materials, integration of 

nanoenergetic materials with MEMS, and MEMS sensor design and fabrication. He is currently pursuing a Ph.D. in mechanical 

engineering at the University of Maryland, College Park, developing nanoenergetic porous silicon as a novel actuator to enhance 

mobility of microrobotic platforms.
150

John D. Clayton has been on ARL’s Technical Staff at APG, Maryland, since 2003. Dr. Clayton served at APG as a NRC Postdoctoral 

Fellow; he was previously a Visiting Scientist, performing graduate research at Sandia National Laboratories in Livermore, 

California (2000), and he performed undergraduate research at Worcester Polytechnic Institute, Massachusetts (1997). Dr. 

Clayton is a member of several professional associations, including the American Ceramic Society, the American Academy of 

Mechanics, the American Physical Society, the American Society of Mechanical Engineers, the Society of Engineeering Science, 

and the U.S. Association for Computational Mechanics. His work has garnered him the ARL Award for Publication (2011), as well 

as multiple DRI Awards, one of which (“Phase Field Modeling”) was rated as one of the Top 3 DRI projects in 2010.

Arthur von Cresce holds a Ph.D. in materials science and engineering from the University of Maryland, College Park, and joined 

ARL in January 2012 as a materials engineer. His primary research focus is on polymer science, the synthesis of organic 

polymers, and their manipulation and characterization. He currently performs organic synthesis as it relates to Li-ion batteries, 

and making additive and solvent compounds. He also performs a variety of electrochemical characterizations of Li-ion battery 

systems, as well as materials characterizations such as FTIR, XPS, and mass spectrometry. 

Luke J. Currano received B.S., M.S., and Ph.D. degrees in mechanical engineering from the University of Maryland, College 

Park, in 2000, 2002, and 2010, respectively. He has worked as a MEMS Engineer for ARL since 2001. He is currently Team 

Leader of the MEMS and Nanoenergetic Devices team at ARL. His research interests include integration of energetic materials 

with MEMS, high force/displacement MEMS actuators, acceleration switches, no-power sensors, and biologically inspired 

MEMS devices.

Jennifer L. Gottfried received a B.S. in chemistry from Ohio Northern University in May 2000, an M.S. in physical chemistry from 

the University of Chicago in August 2001, and a Ph.D. in physical chemistry from the University of Chicago in August 2005. She 

joined ARL in September 2005 as a postdoctoral fellow and was hired as a Physical Scientist in May 2008. In 2009, Dr. Gottfried 

was the Gold medal winner for Rookie Employee of the Year—Technical Scientific and Program Support (Excellence in Federal 

Career Award). She was a co-recipient of a 2009 Research and Development Award for “Stand-off Detection of Explosives,” coinvestigator on one of the top three FY07 DRI projects, and principal investigator for one of the top four FY09 DRI projects. Dr. 

Gottfried has contributed to a double-pulse LIBS review paper and a book chapter on explosives detection, and has written more 

than 29 additional peer-reviewed papers. 

Sergiy Izvyekov is a Physical Scientist who joined the Weapons and Materials Research Directorate at ARL in April 2009. Dr. 

Izvekov received a Ph.D. in Physics from National University of Kiev (Kiev, Ukraine) in 1992. He has authored 62 articles in 

peer-reviewed journals (over 1700 citations), four book chapters, and numerous presentations at national and international 

conferences. Additionally, he received the 27th Army Science Conference Best Paper Award in “Advanced High Performance 

Computing in Physical Sciences and Engineering.” Dr. Izvekov is a member of the American Chemical Society.

T. Richard Jow currently leads the ARL team developing advanced energy storage materials for high power density and high 

energy density electrochemical and electrostatic energy storage devices. Dr. Jow has been the technical lead of the DOE-ARL 

Interagency Agreements on Li-ion electrolytes and batteries project since 2000. He also managed Army Capacitor ManTech 

Development Program for pulsed power applications (2004-2008). Dr. Jow received his B.S. and M.S. in physics from Tsing Hua 

University, Taiwan, China, and received his Ph.D. in materials science and engineering from Northwestern University in 1977. 

Dr. Jow joined ARL in 1989. He received the Allied-Signal Inventors Award in 1987, the ARL Technical Achievement Award in 

Science in 1996, and the Army R&D Achievement Award in 1996, 1999, 2001, 2002 and 2011. Dr. Jow has also authored and 

co-authored over 130 journal publications and received 27 U.S. patents.

William D. Mattson is a Research Physicist in the Weapons and Materials Research Directorate at ARL. Dr. Mattson obtained his 

B. S. in physics and computer science at the University of Maryland Baltimore County in 1996. He further obtained his Ph.D. in 

physics at the University of Illinois at Urbana-Champaign in 2003, under the direction of Professor Richard Martin. Dr. Mattson 

has worked for the ARL since 1993, and on his current work in the simulation of energetic materials since 1994. Dr. Mattson 

has more than 30 publications, and more than 70 presentations. Dr. Mattson developed and serves as the PI for theory of 

the Disruptive Energetics ARL Mission program. Dr. Mattson has also participated in developing scalable software for classical 

atomistic molecular dynamics simulations under several DoD HPCMP programs, and currently serves as the Computational 

Technology Lead for Computationals Chemistry and Material Science. 
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Christopher Morris joined ARL in 2007 as a DARPA-sponsored fellow under the S.M.A.R.T. fellowship program, and was hired a 

government employee in 2009. He earned a Ph.D. in electrical engineering from the University of Washington in 2007, following 

a two-year stint at a biotechnology startup company, and an M.S. in mechanical engineering in 2000, also from the University 

of Washington. Current research interests include the use of stored chemical or mechanical energy for novel high-force and 

large-displacement actuators; fuze and initiation technologies; and new packaging and assembly methods to enable complete, 

functional, and self-contained microsystems. He has authored or co-authored over 45 peer-reviewed journal and conference 

papers, with a recent paper on RFID-integrated battery-free MEMS actuators highlighted as being in the top 0.5% of all IEEE 

publications for that month.

Lynn B. Munday works in ARL’s Simulation Sciences Branch, Computational Sciences Division. He holds a Ph.D. from the 

University of Maryland, an M.S. from the University of New Mexico, and a B.S. from Montana State University, all in mechanical 

engineering. He joined the staff of ARL in late 2011 after working an intern. His research is in computational solid mechanics. 

He has published three archival journal papers, and four conference papers and presentations.

Jeffrey A. Read received his Ph.D. in chemistry from the University of Michigan in 1996, under Dr. Anthony Francis, where he 

studied the spectroscopy of metal chalcogenides and their intercalates. From 1993 to 1997, he conducted applied research 

on new materials and electrolytes for Li-ion polymer batteries at Ultralife Batteries, with emphasis on cycle life and storage at 

elevated temperatures. From 1997 to 1998 he worked on advanced development of rechargeable Zn-Air batteries at AER Energy 

Resources, with particular focus on cycle life and failure analysis. Dr. Read began at ARL in 1999 and has focused on primary 

battery research for Soldier applications. Chemistries studied include Li/MnO2, Li/CFx, and Li/Air batteries. Dr. Read is Team 

Leader on the group looking to the next generation of primary and rechargeable batteries, including Li/S and rechargeable 

batteries using ceramic conducting membranes. 

Betsy M. Rice is a research chemist and Leader of the Multiscale Reactive Modeling Team in the Energetic Materials Science 

Branch, Weapons and Materials Research Directorate of ARL. She is responsible for initiating, planning, and personally 

performing research investigations to determine microscopic details of the physical and chemical processes of materials of 

interest to the Army, including energetic materials and materials related to armor/armaments. Dr. Rice’s expertise lies in the 

theoretical chemistry areas of classical molecular simulation and quantum mechanical molecular characterization directed 

toward advanced modeling of materials that are critical components of several DoD weapons and S&T mission areas. Dr. Rice 

began her tenure with the Army as a National Research Council Postdoctoral Fellow at the Ballistic Research Laboratory in 1989, 

and was hired in 1990. Dr. Rice received a B.S. Degree in chemistry from Cameron University, Lawton, Oklahoma, in 1984, and 

was awarded a Ph.D. in chemistry at Oklahoma State University in 1987. Dr. Rice also holds the position of Senior Scientist at the 

High Performance Computing Software Application Institute for Multiscale Reactive Modeling of Insensitive Munitions. Dr. Rice 

was named as an ARL Fellow in 2000.

Xiaoming Ren is a Senior Research Chemist, focusing on alkaline fuel cell technology for portable and stationary applications, 

and developing a lithium air battery. He came to ARL in 2010, after experience as a Post-Doctoral Research Fellow at the 

University of Texas at Dallas and the Los Alamos National Laboratory, as well years of experience in private industry. Dr. Ren 

received a B.S. from Chengdu University of Science and Technology in 1984, and a Ph.D. from the Memorial University of 

Newfoundland in 1993.

Dat T. Tran obtained his B.S. and Ph.D., both with honors, in chemistry from SUNY Binghamton in 1998 and 2005, respectively. 

He pursued postdoctoral studies in the area of bio-inspired materials synthesis at Cornell University during 2006-2007. He joined 

ARL in June 2008 as a research chemist. His areas of research include synthesis and characterization of inorganic materials for 

liquid phase desulfurization of logistic JP-8 fuel and catalysts materials for Army fuel cell/battery applications. In his research 

areas, Dr. Tran has authored 20 peer-reviewed journal papers and several proceedings papers and technical reports. His research 

work has often been cited by other research workers in the field. Dr. Tran is an active member of the American Chemical Society.
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Jeffrey Wolfenstine is currently working at ARL in the area of Li-ion batteries for both Soldier and vehicle applications. Prior to 

joining ARL, he was a professor of mechanical engineering at the University of California Irvine investigating the high temperature 

mechanical behavior of composites, intermetallics, and ceramic materials. He is an internationally recognized scientist in the 

field of materials science, with main emphasis on processing, characterization, and properties of structural and energy storage 

materials. He has published over 100 journal papers with a citation index over 1,000, and his research is referenced in textbooks. 

He received outstanding paper of the year for one journal, two Army R&D awards, and one best paper at an Army science 

conference, and has two patents and five patent disclosures.

Kang Xu is a Research Chemist whose areas of interest include energy storage materials, interface chemistry in electrochemical 

devices, and bio-inspired synthesis of materials. Dr. Xu received a B.S. in organic chemistry from Southwest University, an M.S. 

in polymer chemistry from Institute of Chemical Physics, Academy of Sciences, and a Ph.D. from Arizona State University. He 

joined ARL in 1997, and has been cited since then for his technical excellence with Department of Army R&D Awards (1999, 

2001, 2002 and 2011), Best Paper Award at Army Science Conference (2008), the Citation for Leadership Excellence in DDR&E 

Wearable Power Prize Competition (2008), ARL Publication (2005) and Science Awards (2011), respectively. He serves as 

a frequent panelist for industry, DOE, Electrochemical Society, and other international academic communities. He has 114 

publications in peer-reviewed journals, and has edited/co-authored three books and book chapters. His work has over 3,800 

citations in literature with an h-index of 34.

Sheng S. Zhang is a research chemist in Electrochemistry Branch of ARL’s Sensors and Electron Devices Directorate. He received 

a Ph.D. in physical chemistry from the University of Science and Technology Beijing, China, in 1993. After experience in academy 

and industry, he joined ARL as a postdoctoral associate in 1997, and became an employee in 2002. His research activities 

focus on advanced materials for electrochemical energy storage and conversion devices, including rechargeable Li-ion batteries, 

lithium batteries, metal-air batteries, and electrochemical capacitors. He investigates high energy density batteries beyond 

Li-ion, including Li/CFx battery, Li-air battery, and lithium-sulfur batteries. Dr. Zhang has authored 117 peer-referred journal 

papers, published 21 patents/applications, and edited a review book. His publications have received an h-index of 34 with over 

3200 citations. He received two Army R&D Achievement Awards in 2000 and 2001, and serves as an editorial board member 

for International Journal Electrochemistry, Journal of Energy Storage and Conversion, The Open Electrochemistry Journal, and 

Dataset Papers in Materials Science.
Approved for public release; distribution unlimited • June 2012