A Strong Energy Portfolio for a Strong America 

Energy efficiency and clean, renewable energy will mean a stronger economy, a cleaner environment, 

and greater energy independence for America. Working with a wide array of state, community, industry, and 

university partners, the U.S. Department of Energy’s Office of Energy Efficiency and Renewable Energy 

invests in a diverse portfolio of energy technologies.

For more information contact:

EERE Information Center

1-877-EERE-INF (1-877-337-3463)

www.eere.energy.gov
U.S. Department of Energy 

Office of FreedomCAR and Vehicle Technologies 

1000 Independence Avenue S.W. 

Washington, DC 20585-0121 

FY 2005 

Progress Report for Heavy Vehicle Propulsion Materials 

Energy Efficiency and Renewable Energy 

Office of FreedomCAR and Vehicle Technologies 

Advanced Materials Technologies 

Edward Wall Program Manager, OFCVT 

Rogelio Sullivan Advanced Materials Technologies Team Leader 

James Eberhardt Technology Development Area Specialist 

May 2006 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

CONTENTS 

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

2. MATERIALS FOR FUEL SYSTEMS.................................................................................................. 5 

A. Low-Cost, High-Toughness Ceramics ............................................................................................ 5

3. MATERIALS FOR EXHAUST AFTERTREATMENT..................................................................... 9 

A. Materials for Exhaust Aftertreatment.............................................................................................. 9 

B. Development of Materials Analysis Tools for Studying NOx Adsorber Catalysts 

(CRADA No. ORNL-02-0659 with Cummins, Inc.) ...................................................................... 17 

C. Development of NOx Sensors for Heavy Vehicle Applications...................................................... 25 

D. Ultra-High Resolution Electron Microscopy for Characterization of Catalyst 

Microstructures and Deactivation Mechanisms .............................................................................. 31 

E. Microstructural Changes in NOx Trap Materials under Lean and Rich Conditions 

at High Temperatures...................................................................................................................... 39 

F. Aftertreatment Catalysts Materials Research .................................................................................. 47 

G. Catalysis by First Principles............................................................................................................ 49 

H. Durability of Diesel Particulate Filters (CRADA with Cummins, Inc.) ......................................... 61 

4. MATERIALS FOR AIR HANDLING, HOT SECTION, AND STRUCTURAL 

COMPONENTS...................................................................................................................................... 69 

A. High-Temperature Advanced Materials for Lightweight Valve Train Components ...................... 69 

B. Engineered Surfaces for Diesel Engine Components...................................................................... 77 

C. Cermet Composites for Wear Applications..................................................................................... 87 

D. Processing and Characterization of Structural and Functional Materials for Heavy-Vehicle 

Applications .................................................................................................................................... 89 

E. NDE Development for Valve-Train Components for Diesel Engines ............................................ 97 

F. Durability of Diesel Engine Component Materials......................................................................... 103

G. Life Prediction of Diesel Engine Components................................................................................ 109 

H. Low-Cost Manufacturing of Precision Diesel Engine Components ............................................... 115 

I. Cost-Effective Machining of Titanium Materials ........................................................................... 123 

J. Advanced Cast Austenitic Stainless Steels for High-Temperature Components ............................ 131 

K. Viscoplastic Twinning in an Ultrafine TiAl/Ti3Al Laminate Composite........................................ 139 

L. Ti-6Al-4V Billet Feedstock Manufacture and Evaluation .............................................................. 147

M. High-Density Infrared Surface Treatment of Materials for Heavy-Duty Vehicles ......................... 151 

N. High-Temperature Aluminum Alloys ............................................................................................. 155 

O. Application of Titanium Alloys for Heavy-Duty Vehicles ............................................................. 163 

P. Mechanical Behavior of Ceramic Materials for Heavy-Duty Diesel Engines ................................ 169 

Q. Powder Processing of Nanostructured Alloys Produced by Machining ......................................... 179 

R. Deformation Processes for the Next-Generation Ceramics............................................................. 185 

S. Development of Titanium Component Applications in Heavy-Duty Diesel Engines..................... 191 

T. High-Speed Machining of Titanium ............................................................................................... 199 

iii 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

5. TEST AND MATERIALS STANDARDS ............................................................................................ 207

A. Rolling Contact Fatigue .................................................................................................................. 207 

B. Implementing Agreement for a Programme of Research and Development on Advanced 

Materials for Transportation Applications ...................................................................................... 215 

C. Mechanical Property Test Development ......................................................................................... 223 

D. An Integrated Surface Modification of Engineering Materials for Heavy

Vehicle Applications....................................................................................................................... 231 

Appendix A. ACRONYMS AND ABBREVIATIONS ........................................................................... 237 

iv 
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

1. INTRODUCTION 

Heavy Vehicle Propulsion Materials 

Advanced materials are an enabling technology for fuel-efficient heavy-vehicle truck engines. The Heavy 

Vehicle Propulsion Materials Project is organized around the following technology issues: fuel systems; 

exhaust aftertreatment; air handling, hot section, and structural components; and standards. 

Materials for Fuel Systems 

Currently, the fuel system represents a significant portion of the cost of a heavy-duty diesel engine. Enabling 

materials and cost-effective precision manufacturing processes are instrumental in developing improved fuel 

injection systems. In addition to new and improved materials, improved manufacturing and inspection 

methods for the injector components are being developed. 

Manufacturing technology for nickel aluminide–titanium carbide cermet fuel system components continued 

this year in coordinated efforts at Oak Ridge National Laboratory (ORNL), CoorsTek, Inc., and Cummins, 

Inc. ORNL supplied a large pilot-scale batch of processed powder mixtures with alternate Ni3Al precursors to 

CoorsTek (a parts supplier) for injection molding of test components and completed work with CoorsTek for 

scale-up of the manufacturing process. Sintered parts were supplied to Cummins Engine for rig testing of 

machined components. 

A novel “pin-on-twin” scuffing test has been developed to evaluate fuel injector materials in diesel fuel and 

low-sulfur fuel environments in a simulated fuel system environment. It is in beneficial use for evaluating 

fuel system materials. The system is available to users of the High Temperature Materials Laboratory 

(HTML) at ORNL, as well as to the Propulsion Materials Program. 

Materials for Exhaust Aftertreatment 

The reduction of nitrogen oxides (NOx) and particulate emissions is critically important to the Office of 

FreedomCAR and Vehicle Technologies (OFCVT) program and is highly materials dependent. The U.S. 

Department of Energy (DOE) goals for improved efficiency of heavy vehicles are greatly complicated by 

engine design and exhaust aftertreatment technologies designed to meet the mandatory U.S. Environmental 

Protection Agency (EPA) emission regulations for 2007 and 2010. Materials and systems research is being 

conducted to minimize the potentially negative effects of emission-reduction technologies on fuel economy 

and to result in cleaner and more efficient engines. 

The durability of exhaust aftertreatment systems in heavy vehicles is a concern. A lifetime of at least 500,000 

miles is expected, and a 1,000,000-mile lifetime is desired (compared with 100,000 miles for automobiles). 

Exposure of the aftertreatment systems to high temperatures, vibration, erosion, and chemical attack by 

species in the oil and fuel results in degradation of performance. The effects of exposure in service on the 

microstructure and microchemistry of the aftertreatment systems are being characterized, and this research 

may lead to the development of more durable systems. The development of advanced NOx sensors is being 

pursued to facilitate optimal engine and aftertreatment control strategies. 

Accomplishments this year included significant progress in a collaborative Ford–ORNL program intended to 

facilitate deployment of a NOx trap for lean diesel or gasoline exhaust by (1) investigating materials issues 

related to deterioration of NOx trap performance with aging as a result of thermal and sulfation-desulfation 

cycles and (2) investigating materials that are robust under the lean NOx trap operating conditions. The latter 

objective includes the synthesis of new materials. Accomplishments of this effort include completing a study 

1 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

of microstructural changes in a series of model catalysts during aging under lean and rich conditions at 

500°C; beginning the updating of the ex-situ reactor to enable transmission electron microscopy (TEM) 

sampling under lean, rich, or stoichiometric conditions as well as lean-rich cycles; and equipping a new 

synthesis laboratory for the preparation of NOx trap materials. 

A new agreement was initiated entitled “Catalyst Materials by Design.” The effort is led by a team that 

includes ORNL, the University of Notre Dame, and the University of California at Davis. The objective is to 

search for durable emissions catalysts by means of an integrated approach using computational modeling, 

synthesis and processing, characterization via high-resolution scanning TEM/TEM, and testing of catalyst 

materials. Computational modeling was carried out using density-functional-theory–based first-principles 

calculations. A series of platinum and rhenium (Re) nanoclusters, Re nanoclusters, and carbonylated Re 

clusters supported on commercial alumina were synthesized and characterized with respect to the size and 

chemistry of atomic clusters and the performance of the clusters in contact with carbon monoxide, NOx, and 

hydrocarbons. 

Research conducted at the High Temperature Materials Laboratory at ORNL is focused on the development 

and utilization of new capabilities and techniques for ultra-high-resolution scanning TEM/TEM to 

characterize the microstructures of catalytic materials of interest for reducing NOx emissions in diesel and 

automotive exhaust systems. This research aims to relate the effects of reaction conditions on the changes in 

morphology of heavy metal species on “real” catalyst support materials, typically oxides. During this period, 

the morphologies of tri-rhenium carbonyl clusters on gamma-alumina support material were characterized at 

the atomic level to demonstrate for the first time the unambiguous imaging of clusters of known structure. 

Also, the nature of platinum catalyst structures on mixed-oxide support materials was characterized, including 

the imaging of platinum species dispersed as single atoms and in precrystalline clusters known as “rafts.” 

Exhaust aftertreatment materials projects are ongoing at both Caterpillar and Cummins. These projects are 

significant in that they represent a departure from an earlier culture in which the diesel engine companies 

relied heavily on catalyst manufacturers to provide the needed technologies and did not actively participate in 

the development of catalyst materials. These diesel engine manufacturers are actively collaborating with 

catalyst suppliers in the development of improved catalyst materials, and they are contributing to the 

development of a fundamental understanding of catalyst performance that is important to both suppliers and 

users of catalyst systems. In addition to the in-house efforts at Caterpillar and Cummins, Cummins and 

ORNL collaborated, through cooperative research and development agreements, in characterizing lab- and 

engine-tested catalysts via X-ray diffraction, spectroscopy, and microscopy. They also collaborated in 

characterizing diesel particulate filters and developing probabilistic design tools to predict the useful lifetimes 

of the filters. 

Ford Motor Co. and ORNL are collaborating on the development of a NOx sensor that can be used in systems 

for on-board remediation of diesel engine exhaust. The sensor should have an operating temperature of 600– 

700°C and be able to measure NOx concentrations from 1 to 1500 ppm at oxygen levels from 5 to 20 vol %. 

Prototype sensing elements are fabricated by patterning electronically conductive and catalytic layers onto 

oxygen-ion–conducting substrates. The sensing elements are characterized for NOx response, oxygen 

sensitivity, and response time. Accomplishments this year included developing sensing elements that display 

near “total NOx” behavior and are stable in simulated long-term service. 

Materials for Air Handling, Hot Section, and Structural Applications 

Engine design strategies for meeting EPA emission requirements have resulted in the need for significantly 

higher turbocharger boost. The higher boost requirements result in higher heat of compression and greater 

thermal and fatigue loads on turbocharger components. 

2 
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Caterpillar began a new project in 2003 to design and fabricate a cost-competitive diesel engine turbocharger, 

using lightweight titanium materials, that provides a reduction in both fuel consumption and transient 

emissions. Caterpillar has designed a series turbocharger for use on the C15 engine platform. This 

turbocharger consists of one turbo wheel and two compressor wheels that are attached to a single rotating 

drive shaft. This compact design will replace the two-turbocharger system that is presently installed in the 

C15 engine. Titanium aluminide will be used for the turbo wheel, and one of the compressor wheels will be 

made from a titanium alloy. The project was completed this year. Highlights include successfully developing 

joining technology for the rotor and shaft, evaluating the physical and mechanical properties of commercially 

available materials, and accessing the commercial viability of the advanced turbocharger. 

Caterpillar and ORNL won a 2003 R&D 100 Award for the development of CF8C-Plus cast stainless steel. 

The new high-temperature stainless steel may have near-term applications in diesel engine exhaust manifolds 

and turbocharger housings. Highlights include demonstrating that CF8C-Plus cast stainless steel is far 

superior to SiMo cast iron at 600°C and above. CF8C-Plus also shows strength, creep resistance, and aging 

resistance benefits compared with other commercial cast stainless steels being considered for heavy-duty 

diesel turbocharger housings at 750°C. Significant progress has been made in commercializing the alloy: a 

6700-lb gas-turbine end-cover was cast from CF8C-Plus at MetalTek. MetalTek has cast more than 30,000 lb 

for various applications. 

Caterpillar, in collaboration with Argonne National Laboratory and ORNL, has a project to design and 

fabricate prototype engine valves from silicon nitride and TiAl materials that are 30% lighter than steel valves, 

provide a 200% increase in service lifetime, and potentially increase fuel efficiency in advanced engines by 

10%. A probabilistic design approach was developed for the high-hardness valve materials. The friction 

welding of TiAl valve heads and Ti-6V-4Al valve stems was successfully optimized. The effects of surface 

finish on the performance of silicon nitride valves were evaluated, and the results indicated that valves with 

good surface finish performed well in bench tests. 

This year, in preparation for an engine test to be conducted in FY 2006, test valves were procured and 

characterized for dimensional and surface properties and via laser-scattering for surface or sub-surface defects. 

In addition, an evaluation of the machinability of TiAl was initiated, along with a study of friction, wear, and 

corrosion of the alloy. A Caterpillar C15 ACERTTM engine was donated to the National Transportation 

Research Center for future component testing. 

Caterpillar is also developing innovative approaches to thermal barrier and wear-resistant coatings for engines. 

Durability issues for thermal sprayed coatings, particularly thermal barrier coatings, remain the major 

technical challenge to their implementation in new engine designs. New approaches to coating design and 

fabrication are being developed to aid in overcoming this technical hurdle. A highlight from this year is the 

coating of pistons for a homogeneous charge compression ignition system with amorphous steel for singlecylinder engine testing. 

The feasibility of reducing the weight of a heavy-duty engine by substituting a lighter material for the cast 

iron engine block and cylinder head was evaluated by ORNL and Ricardo, Inc. Finite element analysis was 

used to compute the stresses, temperatures, and fatigue safety factors of a 5.9-L diesel engine, run at the 

maximum design power, for three lightweight casting alloy systems: titanium, aluminum, and magnesium. 

The analysis indicated the feasibility of simple material substitutions in all cases, with the exception that the 

titanium alloy cylinder head will require inserts to reduce the temperature in the valve bridge area. Reducing 

the weight of the entire engine by up to 33% for the magnesium alloy and 20% for the aluminum alloy, and 

up-rating the power of the engine by 50% while reducing the weight by 15% for the titanium alloy, were 

predicted to be feasible. 

3 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Efforts in cost-effective manufacturing were carried out by ORNL, the University of Michigan, and Purdue 

University. Technology for machining difficult materials, such as titanium alloys, was developed in the 

ORNL/University of Michigan collaboration. ORNL characterized the grinding characteristics of TiC/Ni3Al 

matrix cermets that are promising candidates for diesel engine fuel injector plungers and wear parts, and 

investigated the possibility of using cermets as cutting tool materials for high-speed titanium machining. The 

University of Michigan and ORNL participated in a consortium of manufacturers led by Third Wave Systems, 

Inc., to develop finite-element-analysis–based machining modeling software and to acquire new machining 

process modeling capabilities. Purdue University is investigating the consolidation of low-cost machining 

chips to produce nanocrystalline components with high strength and hardness. 

Materials and Testing Standards 

OFCVT has an International Energy Agency “Implementing Agreement for a Programme of Research and 

Development on Advanced Materials for Transportation Applications.” Accomplishments this year include 

the development by NIST of a new design principle for surface texture design to enable friction reduction 

under boundary lubrication conditions. 

NIST also leads a project to develop standard testing methods for advanced materials, primarily ceramics. 

During this year, NIST revised, refined, and improved several current ASTM and ISO standards. 

The rolling contact fatigue (RCF) effort is led by ORNL. Its objectives include the characterization of the 

RCF performance of ceramics and tribological coatings; determination of the effects of subsurface damage, 

microstructure, material properties, and contact stress on RCF performance; and correlation of RCF 

performance measured by different internationally-used RCF test techniques. Highlights this year include 

completion of a summary report on RCF test methods and result interpretations used in Germany, Japan, 

United Kingdom, and the United States and the 

development of a method to evaluate and discriminate the elastic properties of ball bearings in situ using 

resonance ultrasound spectroscopy. 

Rogelio Sullivan James J. Eberhardt 

Team Leader, Materials Technologies Technology Development Area Specialist 

Office of FreedomCAR and Vehicle And Chief Scientist 

 Technologies Office of FreedomCAR and Vehicle 

Energy Efficiency and Renewable Energy Technologies 

Energy Efficiency and Renewable Energy 

4 
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

2. MATERIALS FOR FUEL SYSTEMS 

A. Low-Cost, High-Toughness Ceramics 

T. N. Tiegs, F. C. Montgomery, and J. O. Kiggans 

Oak Ridge National Laboratory 

P.O. Box 2008, MS-6087 

Oak Ridge, TN 37831-6087 

(865) 574-5173; fax: (865) 574-4357; e-mail: tiegstn@ornl.gov 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Prime Contract No: DE-AC05-00OR22725 

Objectives 

• Demonstrate that significant improvement in the reliability of structural ceramics for advanced diesel engine 

applications could be attained if the critical fracture toughness (KIc) were increased without strength 

degradation. 

• Confirm that cost is a major factor in determining the applicability of new materials in engine components. 

• Develop high-toughness materials that are also low cost. 

Approach 

• Show that TiC-Ni3Al composites have a combination of superior physical properties and mechanical behavior 

using conventional powder processing methods. 

• Develop TiC-based composites with 40–60 vol % Ni3Al because they have expansion characteristics very close 

to those for steel. 

• Collaborate development effort with CoorsTek, Inc., and Cummins Engine on processing scale-up and engine 

testing. 

• Determine initial mechanical properties of ceramic-particulate-reinforced TiAl composites and make 

assessment of application in heavy vehicle systems. 

Accomplishments 

• Supplied large pilot-scale batch of processed powder mixtures with alternate Ni3Al precursors to CoorsTek for 

injection molding of test components. 

• Completed work with CoorsTek, Inc., (a parts supplier) for scale-up of the processing and to supply them with 

pilot plant scale quantities of powder mixtures for injection molding trials. 

• Supplied sintered parts produced in conjunction with CoorsTek to Cummins Engine Co. for rig testing of 

machined components. 

5 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Future Direction 

• End the project with the transfer of the TiC-Ni3Al cermet technology to industry. 

Introduction 

TiC-Ni3Al composites are under development 

for application in diesel engines because of desirable 

physical and mechanical properties. For these applications, the Ni3Al volume contents are on the order 

of 30 to 50 vol % to match the thermal expansion of 

steel. Typically, flexural strengths greater than 

1000 MPa up to 800°C and fracture toughnesses 

higher than 15 MPa m are obtained for the 

composites. The composites are densified by liquid 

phase sintering, and most of the early work used 

gas-atomized Ni3Al particles with fine TiC powders. 

Later work was done using nickel and NiAl powders 

(along with the TiC) to form Ni3Al by an in situ 

reaction during sintering. Over the last few years the 

in-situ reaction process was developed significantly 

because it produced high mechanical properties and 

developed a fine TiC grain size. The finer grain sizes 

were favored because of better wear resistance. 

The fabrication techniques and equipment 

employed in production are very similar to those for 

the fabrication of WC-Co hard metals, and thus, the 

processing costs are well established. However, 

when the economics of producing the TiC-Ni3Al 

composites was examined, a significant cost was 

associated with the use of the NiAl precursor 

powder (about 55% of the total raw material cost). 

Because the costs of the starting raw materials can 

be a significant fraction of the total cost of a component, alternative materials for fabricating the cermets 

are of interest. Part of the reason for the high cost of 

the NiAl is that it is produced only as a specialty 

powder at the present time. The development effort 

is being done in collaboration with CoorsTek, Inc. 

Results 

Alternate Precursors for Ni3Al Formation 

All previous work used a combination of nickel 

and NiAl for an in-situ reaction to form the Ni3Al. A 

new NiAl commercial source has been identified 

and powder procured. Characterization of the powder by x-ray diffraction indicated that it was 

stoiciometric and had a pattern identical to earlier 

material from a different supplier. Particle size 

analysis showed the mean particle size was ~9.8 μm, 

which was similar to the previous material. 

However, the size distribution was broader than the 

earlier NiAl. 

A large batch (3 kg) of TiC-50 vol % Ni3Al 

powder was processed with the new NiAl powder 

and shipped to CoorTek for injection molding (IM) 

trials. Test parts were fabricated and shipped back to 

Oak Ridge National Laboratory (ORNL). Initial 

sintering results indicated the densification behavior 

was similar to the earlier composites, but slightly 

less: 6.05 g/cm3

 vs 6.11 g/cm3. 

Based on those results, new IM tooling was 

designed and fabricated at CoorsTek. Several large 

batches (>15 kg total) of TiC-50 vol % Ni3Al powder were milled, homogeneously mixed, and shipped 

to CoorsTek. Parts with the new tooling have been 

IM and returned to ORNL. The specimens are currently awaiting sintering when the furnace is 

repaired. 

While the initial sintering results indicated the 

densification behavior was similar to the earlier 

composites, flexural strength measurements revealed 

a rather low strength in comparison to the previous 

material from a different NiAl manufacturer. The 

comparison was 945 ± 150 MPa for the earlier materials and 600 ± 37 MPa for the newest batch. Scanning electron microscopy (SEM) examination of a 

fracture surface showed numerous agglomerates 

≤100 μm in size. Energy dispersive x-ray analysis 

(EDAX) indicated they were NiAl particles. 

Evidently the new batch of NiAl powder contained 

larger particles (even though the powder was listed 

as <44 μm). 

A modified milling procedure was instituted 

with the NiAl, and particle size analysis indicated all 

particles were <10 μm. Sintering of test parts with 

the modified powder was improved, and densities of 

~6.11 g/cm3 were obtained. This density is equal or 

better than the previous high-strength materials. A 

test batch (~3 kg) of powder with the modified 

milling procedure has been sent to CoorsTek for IM. 

Parts and test MOR bars have been sintered at 

ORNL and returned to CoorsTek for testing. Sintering behavior of these injection molded materials was 

similar to previous cermets, and the microstructure 

showed no unusual defects. Consequently, a large 

batch (>12 kg total) was fabricated and sent to 

CoorsTek for IM of test parts. 

6 
Heavy Vehicle Propulsion Materials 

Assessment of Ceramic-Particulate-Reinforced 

TiAl Composites 

TiAl offers significant improvements in weightto-strength ratios compared to most advanced alloys. 

Such materials could have applications in heavy 

vehicle propulsion systems. Ceramic-metal 

composites that utilize TiAl as the bonding phase 

potentially could show increased fracture strength, 

higher elastic modulus, better creep resistance, and 

higher wear resistance than the base alloy. An 

exploratory study was initiated to fabricate ceramicparticulate-reinforced TiAl composites and 

determine the mechanical properties. Initial samples 

were fabricated by hot-pressing to readily obtain 

high-density specimens suitable for testing. An 

assessment was made to determine if the mechanical 

properties are suitable for further development for 

heavy vehicle applications. 

Several compositions were fabricated including: 

TiC-50 vol % TiAl, TiN-50 vol % TiAl, and TiC-67 

vol % TiAl. The TiAl was an alloy with the 

composition, Ti-48 Al-2 Cr-2 Nb, which has been 

shown to have excellent properties. The powder 

mixtures were made using similar methods to those 

used for the TiC-Ni3Al cermets. Hot-pressing 

produced samples with closed porosity and densities 

of 94–95% theoretical density. Microstructural 

examination revealed good wetting and distribution 

of the TiC and TiAl particles. No obvious core-rim 

structure (like those observed with TiC-Ni3Al 

cermets) was observed. However, the TiC did 

appear to have reacted with the TiAl to produce a 

reaction zone surrounding remnant TiC particles. 

EDAX showed these zones to contain Ti, Al, Nb, 

Cr, and C. X-ray diffraction (XRD) results indicated 

the formation of Ti3AlC2, which confirmed a 

reaction occurred between the TiAl and the TiC. 

In contrast to the TiC, the TiAl did not appear to 

have wet the TiN particles very well. A reaction 

between the TiAl and the TiN was also evident, and 

EDAX indicated a metallic phase of Ti-Al-Nb-Cr 

and a second phase of Ti-Al-N. XRD showed the 

formation of AlTi2N, indicating a reaction between 

FY 2005 Progress Report 

the starting components. Indent hardness determinations showed the hardness was lower than similar 

ceramic particle volume loadings in Ni3Al-based 

cermets. For example, the hardness of the TiC-50 

vol % TiAl, TiN-50 vol % TiAl, and TiC-67 vol % 

TiAl composites was 4.0–4.6 GPa, 5.7–5.8 GPa, and 

4.2–4.3 GPa, respectively. For comparable TiC-50 

vol % Ni3Al cermets, the hardness values would be 

>8 GPa. 

Because of the reactions between the TiAl alloy 

and the TiC or TiN reinforcements, the prospects for 

sintering materials and obtaining good mechanical 

properties is not encouraging. However, other 

ceramic particles, such as TiB2, ZrB2, HfB2, and 

Al2O3, should be compatible in the TiAl alloy matrix 

and provide reinforcement. 

Conclusions 

Pilot-scale batches of TiC-Ni3Al composite 

mixtures with alternate Ni3Al precursors were 

produced and supplied to CoorsTek for IM of test 

components. The prospects for sintering TiAl-based 

materials and obtaining good mechanical properties 

is not encouraging. However, other ceramic 

particles, such as TiB2, ZrB2, HfB2, and Al2O3, 

should be compatible in the TiAl alloy matrix and 

provide reinforcement. 

Presentations and Publications 

T. N. Tiegs, “Aluminide Bonded Ceramic 

Composites,” at the American Ceramic Society 

Conference on Advanced Ceramics and Composites, 

Cocoa Beach, Florida, January 23–27, 2005. 

T. N. Tiegs, F. C. Montgomery, P. A. 

Menchhofer, and P. F. Becher, “Aluminide Bonded 

Ceramic Composites,” to be published in Ceram. 

Eng. Sci. Proc., Am. Ceram. Soc., Westerville, Ohio 

(2005). 

T. N. Tiegs, “High Toughness Materials— 

TiC-Ni3Al Cermets,” at the Heavy Vehicles 

Materials Peer Review, Oak Ridge, Tennessee, 

September 13–15, 2005. 

7 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

3. MATERIALS FOR EXHAUST AFTERTREATMENT 

A. Materials for Exhaust Aftertreatment 

Corey Shannon, Herbert DaCosta, Julie Faas, Matt Stefanick, Svetlana Zemskova, Christie Ragle, Ronald 

Silver, Paul Park, and Craig Habeger 

Caterpillar, Inc. 

P.O. Box 1875 

Peoria, IL 61656 

(309) 578-4065; fax (309) 578-2953; e-mail: park_paul_w@cat.com 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: Caterpillar, Inc., Peoria, Illinois 

Objectives 

• Develop advanced materials applied for diesel engine aftertreatment systems that will comply with future 

emission regulations. 

• Assess impact of phosphorus on oxidation, lean-NOx, NOx adsorber (LNT) and urea-SCR (selective catalytic 

reduction) catalysts. 

• Identify diesel particulate filter (DPF) substrate materials that have high durability and filtration efficiency with 

minimum impact on engine efficiency. 

Approach 

• Assess impact of phosphorous on aftertreatment systems by comparing fresh and aged catalyst performance 

using a fuel burner and a bench test system. 

• LNT: Evaluate phosphorus impact on NO oxidation, NOx storage and reduction functions. 

• Urea-SCR: Evaluate phosphorus impact on zeolite and vanadia based catalysts using temperature and NH3/NOx 

ratio sweeps. 

• Oxidation catalyst: Develop a correlation between bench and field in terms of deactivation mechanisms. 

• Use previously developed test protocol to evaluate filtration efficiency and pressure drop of catalyzed and 

uncatalyzed DPF samples from various suppliers. 

Accomplishments 

• Completed lean-NOx project (unfinished tasks carried over from last year). 

⎯ Sulfur impacts on catalytic performance more severely than phosphorus. 

⎯ Paraffins contributed lean-NOx catalysis; however, benzene and mono-cyclo-paraffin were not active 

reductants, whereas aromatics and dicyclo-paraffin inhibited the catalytic reaction. 

9 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

• Identified the cause of LNT deactivation due to phosphorus exposure. 

⎯ NOx storage capacity function was affected first and more severely than NO oxidation and NOx reduction. 

⎯ Phosphorus impacted the NOx storage function more at low NOx slip than at high NOx slip conditions. 

⎯ A model was used to estimate the impact of phosphorus deactivation on fuel penalty for the LNT. 

• Identified the cause of vanadia- and zeolite-based SCR catalysts deactivation due to phosphorus exposure. 

⎯ Phosphorus similarly impacted both NH3 and NO conversions. 

⎯ Zeolite catalyst showed higher phosphorus tolerance initially, but both catalysts had similar deactivation 

after extended exposure. 

⎯ Speculate that phosphorus blocked NH3 adsorption sites, which prevented NO reduction. 

• Found that phosphorus impacted LNTs more than SCRs, but presented a challenge for both catalyst 

technologies. 

• Identified a promising DPF substrate material with filtration efficiency larger than 95% and pressure drop 

significantly lower than conventional cordierite material. 

Future Direction 

• Capture intellectual property and identify a partner for commercialization (lean-NOx project). 

• Assess impact of combination of phosphorus, sulfur, and thermal deactivation on oxicat, LNT, and SCR 

catalysts. 

• Explore mitigation strategies for phosphorus, sulfur, and thermal deactivation phenomena. 

• Evaluate the effect of thermal deactivation on filtration efficiency of new DPF substrate materials. 

Introduction 

The objective of this project is to develop and 

evaluate materials that will be utilized in aftertreatment systems for diesel engine applications to comply with future government emission regulations. 

The materials include catalysts for NOx abatement 

and filtration media for particulate matter control. 

This year’s focus is on the determination of the 

effect of phosphorus deactivation on the performance of LNTs, SCR catalysts and oxidation catalysts, and on identifying diesel particulate filter 

(DPF) substrate materials that have high durability 

and filtration efficiency with minimum impact on 

engine efficiency. 

Approach 

Catalyst Deactivation 

The study was focused on identifying the impact 

of phosphorus exposure on catalyst formulations by 

testing the performance of fresh and aged catalysts. 

The aged catalysts were prepared by exposing to 

phosphorus in a diesel fuel burner using a diesel fuel 

doped with tricresyl phosphate that mimicked the 

effect of phosphorus from combusted lube oil. Using 

x-ray florescence (XRF), we analyzed the amount of 

phosphorus adsorbed as a result of this exposure. 

PM Trap 

A developed test protocol was used to evaluate 

filtration efficiency and pressure drop of catalyzed 

and uncatalyzed DPF samples from various suppliers. In order to test DPF core samples that are supplied in sizes larger than 1 in. × 3 in., we designed 

and built a new chemical reactor with the capability 

to hold samples up to 2 in. × 6 in. 

Results 

Catalyst Deactivation 

Lean-NOx Catalyst Durability: Durability studies using sulfur and phosphorus aging have been 

performed on the Lean-NOx catalyst developed from 

the previous study. Phosphorus was introduced to 

the sample using a diesel fuel burner burning 

phosphorus-blended zero-sulfur fuel, and the NOx 

reduction performance was tested after each exposure increment. NOx reduction of the catalyst maintained more than 40% NOx conversion after 

reaching 30,000 h engine-equivalent (C15 engine) 

10 
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

phosphorus exposure (Figure 1). Phosphorus content Lean-NOx Catalyst Sensitivity to Diesel Fuel 

of the aged samples was quantified using XRF. 

Sulfur was introduced by SO2 exposure directly 

using a bench test system, monitoring the NOx 

reduction performance as a function of time-onstream with SO2. Sulfur content of the exposed 

samples has been quantified using a LECO CS-200 

Carbon/Sulfur Determinator. NOx reduction performance of the catalyst was reduced to ~30% NOx 

reduction after 100-h engine-equivalent (C15 

engine) sulfur exposure. 

The negative impact of sulfur exposure was 

more pronounced than that of phosphorus exposure, 

both in magnitude and rate of degradation. Understanding the effects of sulfur and phosphorus provided insight into the development of catalyst materials and led to mitigation options best suited for the 

intended environment. 

Components: A reductant study over the selected 

combination catalyst has demonstrated much lower 

activity when diesel fuel was used as a reductant 

compared to the desired reductant (example: 

dodecane). The influence of diesel fuel components 

on NOx reduction performance has been investigated 

using a bench test system to determine the source of 

the activity loss (test condition: 500 ppm NOx, 

3000 ppm liquid hydrocarbons as C1, 7% H2O, 9% 

O2, 50 ppm propene, 500 ppm CO, and 8% CO2 with 

35,000 h–1 space velocity). Mixtures were prepared 

such as 80% of dodecane and 20% of various 

organic hydrocarbons, including paraffin, isoparaffin, aromatic, diaromatic, and combinations thereof. 

Isoparaffins (max. 74% NOx reduction) showed NOx 

reduction performance at a similar level to dodecane 

(77%) (Figure 2). Monoaromatics and variations on 

Figure 1. The effect of phosphorus and sulfur on the developed lean-NOx catalyst performance based 

on C15 Caterpillar engine equivalent hours exposure. 

Figure 2. The effect of diesel component on lean-NOx catalyst performance 

(20% of diesel components mixed with dodecane). 

11 
FY 2005 Progress Report 

paraffins had a neutral effect on the NOx reduction 

performance. Among the types of reductants tested, 

diaromatics strongly hindered the performance of 

the combination catalyst (48%), also shifting the 

operation temperature to a higher temperature range. 

Understanding the effects of fuel components provided insight into the development of catalyst materials as well as a fuel processor to obtain the proper 

reductant from diesel fuel to supply to the catalyst 

for maximum performance. 

NOx Adsorber Catalyst Durability: A series of 

model LNTs were prepared by coating 400-cpsi cordierite monolith cores with γ-Al2O3 washcoat. The 

cores were then dipped in cerium acetate solution, 

dried, then calcined at 550oC for 5 h to give a loading of 20 wt % CeO2. The process was repeated with 

a barium acetate solution to give a loading of 

11 wt % BaO, and then with platinum/rhenium 

nitrate solution, resulting in a loading of 5 wt % 

platinum and 0.6 wt % rhenium. The finished 

catalysts were reduced in 5% H2 at 350oC for 2 h. 

The finished catalysts were tested for NOx 

conversion using a laboratory reactor with 1-in. 

diameter by 3-in.-long catalyst cores at a space 

velocity of 55,000 h–1. Each sample was degreened 

at 275oC for 20 min in a reducing gas (2.5% H2, 

2.5% CO, 5% CO2, 5% H2O, balanced with N2), 

then switched to an oxidizing mix (12% O2, 5% 

CO2, 5% H2O, balanced with N2). The 200-ppm NOx 

(10% NO2) was added to load the samples with NOx. 

After saturating the samples with NOx at 275oC, the 

LNTs were ramped to 425oC in the oxidizing mix 

without the NOx, regenerated in the reducing gas, 

and reloaded with NOx. After determining the initial 

performance of the model LNT, the sample was 

exposed to 3.6 g/L of phosphorus using a diesel fuel 

burner with phosphorus-doped fuel at 250oC. The 

aged sample was then reanalyzed for NOx 

conversion in the lab reactor. The aging process was 

repeated to increase phosphorus exposure to 10 g/L. 

The NO oxidation to NO2, NOx storage, and 

NOx reduction function of the NOx adsorbers was 

tested after loading with 3.6 and 10 g/L phosphorus, 

with and without a blank monolith in front as a 

phosphorus getter. Results indicate that phosphorus 

degrades LNTs. The NO oxidation function was not 

affected by the exposure of 3.6 g/L phosphorus; 

however, the NO to NO2 conversion dropped from 

45% to about 35% after exposure to 10 g/L 

phosphorus. 

Heavy Vehicle Propulsion Materials 

In contrast to the relatively stable NO oxidation 

to NO2 function, a significant impact was observed 

on the catalyst NOx storage capacity even on exposure to 3.6 g/L of phosphorus. Figure 3 shows that at 

20% NO slip, the moles of NOx stored per mole of 

available barium in the LNT dropped by 65%. 

Furthermore at a constant storage capacity of 

0.05 moles NOx/mole barium, the NOx slip increased 

from 6% on a fresh LNT to 30% after 3.6 g/L 

phosphorous, and to 90% after exposure to 10 g/L 

phosphorous. 

The reduction of NO to N2 under the rich condition was negatively affected by the loss of NOx storage capacity. Figure 4 shows the NOx storage 

capacity at 80% slip of a series of fresh LNTs 

Figure 3. The effect of phosphorus exposure on LNT 

NOx storage for three different phosphorus 

exposures. 

Figure 4. The effect of phosphorus exposure on LNT 

reduction of NOx in a rich gas stream. 

12 
Heavy Vehicle Propulsion Materials 

plotted against the amount of NOx conversion. A 

straight line was fit to the data with a correlation 

coefficient of 0.968. Exposing the catalysts to 

3.6 g/L phosphorus reduced the storage capacity and 

thus decreased the NOx reduction. However, testing 

the LNTs in the opposite direction from which they 

were exposed to phosphorus revealed a significant 

boost in NOx conversion from 0.0695 to 0.1014 g/L. 

When the LNT cores are exposed to phosphorus, the 

phosphorus tends to deposit in a nonuniform profile 

with the greatest amount of phosphorus in the front 

of the core and less phosphorus in the rear. Any NOx 

introduced to the front of this aged sample cannot be 

adsorbed until it gets far enough into the core where 

phosphorus is not blocking the adsorption sites. 

Upon switching to rich gas, the front section of the 

aged LNT has no NOx to release, and thus no reduction occurs. If NOx containing gas is introduced at 

the rear of the core, NOx begins storing immediately 

inside the new entrance of the sample because of the 

low phosphorus part of the sample. In this case, 

when the system is swept by rich gas, the upstream 

section of the LNT has stored NOx to release, and 

the rhenium in the downstream (high phosphorus 

part of the sample) can reduce the NOx to N2. Thus 

the rhenium in the LNT appears to remain active for 

NO reduction even if no NOx is stored due to the 

phosphorus-blocking adsorption sites. 

Results from this first phase of study show that 

phosphorus deactivation may become a significant 

source of deactivation as durability requirements 

become increasingly stringent. For LNTs, the NOx 

storage capacity function is affected first and most 

strongly. 

SCR Catalyst durability: Zeolite- and vanadiabased SCR catalysts were provided by a catalyst 

supplier. The samples were degreened according to 

the company’s recommended procedure prior to 

testing. Both sets of samples were evaluated using a 

temperature ramp from 100 to 550°C in lean conditions at an NH3/NOx ratio of 1.11. After determining 

the initial activity of the SCR catalysts, samples 

were exposed to 3.6 g/L of phosphorus using the 

previously described diesel fuel burner with 

phosphorus-doped fuel at 250°C. The aged sample 

was then re-analyzed for NOx conversion in the lab 

reactor. The aging process was repeated to increase 

phosphorus exposure to 10 g/L. Figure 5 shows the 

results over the fresh and 3.6-g/L phosphorusexposed vanadia-based samples. 

FY 2005 Progress Report 

Figure 5. The effect of 3.6-g/L phosphorus exposure on 

NH3 and NOx conversion over vanadia SCR. 

The results shown in Figure 5 indicate the 

vanadia-based SCR catalyst is strongly impacted by 

3.6-g/L phosphorus exposure, especially at temperatures above 300o

C. Both NOx and NH3 conversions are impacted, and to the same degree. As the 

temperature is increased above 450o

C, the NH3 oxidation reaction becomes significant over this catalyst, and thus the NOx conversion decreases due to 

ammonia reacting to become NO. In other results 

(not shown), the rate of ammonia oxidation appears 

to increase relative to the rate of the SCR reaction 

with phosphorus exposure. As the phosphorus 

begins to block active sites, the first sites lost will be 

those involving adjacent pairs of vanadia for NOx 

reduction, and the other vanadia sites will be available for a longer period of phosphorus exposure. 

Since the SCR reaction requires two adjacent vanadia sites and ammonia oxidation only requires one 

site, this could explain the observed effect on relative rates (NH3 oxidation vs NOx reduction). 

The same tests were run with the zeolite-based 

SCR. Figure 6 shows the results over the fresh and 

3.6-g/L phosphorus exposed zeolite-based catalysts. 

Compared to Figure 5, the zeolite appears to be less 

affected by this level of phosphorus than the vanadia 

catalyst. The effect of phosphorus seems to be relatively constant across a range of temperatures in this 

case. Both the NOx and NH3 conversions are 

affected to a similar degree. Unlike the vanadia 

sample, the zeolite catalyst does not show a significant drop in NOx conversion at temperatures above 

450o

C due to ammonia oxidation. 

13 
FY 2005 Progress Report 

Figure 6. The effect of 3.6-g/L phosphorus exposure on 

NH3 and NOx conversion over zeolite SCR. 

The results from tests of fresh and aged samples 

of both types were then compared using an 

Arrhenius type plot assuming pseudo first-order 

kinetics. The results in Figure 7 show that similar 

kinetics for the SCR reaction occur on both catalyst 

types. 

Below 400°C, the impact of phosphorus on the 

reaction rate is lower because a significant portion of 

the SCR reaction takes place in the bulk, where the 

phosphorus does not penetrate. At temperatures 

above 400°C more and more of the reaction relatively occurs on the top surface of the catalyst, 

which is also where most of the phosphorus is 

deposited. 

Figure 7. Kinetics of the SCR reaction over fresh and 

aged catalysts. 

Heavy Vehicle Propulsion Materials 

The differences between fresh and aged curves 

of the same type formulation are typical of what is 

observed when a physical blockage of active sites 

leads to activity loss. Results over the vanadia sample show that it loses significant activity on exposure 

to 3.6 g/L phosphorus, but it shows a smaller decline 

as the exposure is increased to 10 g/L. In contrast, 

the zeolite is not affected significantly by exposure 

to 3.6 g /L phosphorus, but the activity significantly 

decreases as the exposure increases to 10 g/L. Interestingly, increasing the temperature after phosphorus 

exposure seems to help the zeolite sample regain 

some of its activity but hurts the vanadia sample, 

possibly due to increased ammonia oxidation over 

the latter. 

Phosphorus Adsorption 

The amount of phosphorus that each type of 

formulation adsorbed after being exposed to 10 g/L 

phosphorus was determined by XRF. Results are 

shown in Figure 8. 

The total amount of phosphorus adsorbed did 

not differ significantly from one catalyst type to the 

other, and the total amount adsorbed was about 15% 

of the phosphorus to which the sample was exposed. 

Similar deposition profiles were also determined for 

all the types of NOx reduction catalysts tested. Out 

of the total amount of phosphorus that was adsorbed 

by a sample, most of it (65%) adsorbed on the front 

third of the core, 25% was adsorbed by the middle 

third, and 10% or less was adsorbed by the bottom 

third of the sample. 

Figure 8. XRF analysis of phosphorus adsorption on 

NOx reduction technologies. 

14 
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

PM Trap 

Various catalyzed and uncatalyzed DPF materials were obtained from suppliers to screen filtration 

efficiency and backpressure. Filtration efficiency 

was determined using a Scanning Mobility Particle 

Sizer (SMPS), and backpressure was measured 

using a pressure transducer. 

Figure 9 shows the filtration efficiencies of both 

the new DPF substrate and cordierite. Both graphs 

are averages of several fresh uncatalyzed DPF 

samples. The data shows that the filtration efficiencies of both materials are greater than 95%. 

Figure 10 shows the results of pressure drop as a 

function of particulate loading time. The arrows 

notes backpressure measurements bypassing the 

DPF samples. The measurements were taken before 

and after loading the sample with particulate matter 

to account for variations as a function of particulates 

produced by the fuel burner. The data show that the 100 

new DPF substrate yields substantially lower presFigure 10. Pressure drop as a function of particulateloading time for both Cordierite and for the 

new material substrate. 

sure drop during particulate-loading than Cordierite 

and for the new material substrate. 

% Efficiency (#/c

m

3) 

98 

Figure 11 shows filtration efficiency of the new 96 

DPF substrate material. The filtration efficiency of 

the catalyzed samples maintain above 98% through­ 94 

out the particulate-loading process. 

To examine the effect of catalyst coating on the 92 

pressure drop of the new substrate material, the 

pressure drop was measured over both catalyzed and 90 

uncatalyzed samples (Figure 12). Uncatalyzed sam- 0 10 20 30 40 50 60 

ple shows lower pressure drop within 20 min of the Time (min) 

test; however, the pressure drop significantly 

increases as a function of particulate loading time. Figure 11. Filtration efficiency of the new catalyzed 

The catalyzed sample shows opposite pressure drop DPF substrate material. 

Figure 9. Filtration efficiency of fresh uncatalyzed samples of the new material and of cordierite (average of 

several samples). 

15 
Pressure Drop (kPa) 

FY 2005 Progress Report

1.2 

1.0

Uncatalyzed

Catalyzed 0.8 

0.6 

0.4 

0.2 

0 10 20 30 40 50 

Time (min) 

Figure 12. Pressure drop for catalyzed and uncatalyzed 

fresh samples of the new material substrate.

profile. The sample yields higher backpressure in the 

early stages of particulate loading, but the pressure 

drop rate decreases in the late stages of particulate 

loading possibly because the coated catalyst 

enhances regenerating soot collected in the sample. 

Conclusions

Lean-NOx Catalyst 

The project was completed with the assessments 

of catalyst durability and diesel fuel component 

effect on the developed catalyst formulation. The 

catalyst is relatively durable to phosphorus deactivation but vulnerable to sulfur deactivation. The 

mitigation plan was developed to minimize the 

sulfur deactivation. The technology is ready to 

transfer to an organization specialized in retrofit 

applications. 

Catalyst Durability 

Results from the catalyst deactivation study

showed that phosphorus exposure by itself can lead 

to loss of catalyst activity. The loss of activity 

Heavy Vehicle Propulsion Materials 

appears to be tied to a physical blockage of surface 

active sites by the phosphorus compound. Most of 

the phosphorus adsorbs on the inlet side of the catalyst, whereas the adsorbed phosphorus concentration 

is lower at the outlet. For LNTs, the phosphorus 

blocks the adsorption of NO2 on storage sites. For 

SCR, the phosphorus prevents NH3 from adsorbing. 

The impact of phosphorus is greater on LNT activity

than it is for SCR, but the resulting loss of activity

can lead to a loss of fuel economy in both cases. 

PM Trap 

A promising material was identified that provides filtration efficiency larger than 95% with low 

pressure drop. In addition, catalyst coating did not 

present adverse effect on pressure drop. 

Publications and Presentations

R. G. Silver, M. Stefanick, and A. Alletag, 

“Impact of Phosphorus on NOx Aftertreatment 

Catalysts,” poster presentation at 2005 DEER 

meeting, Chicago, Illinois (2005). 

R. G. Silver, M. Stefanick, J. M. Faas, and 

S. Zemskova, “Impact of Phosphorus Exposure on 

Lean NOx Trap Performance,” presented at the 19th

Annual Meeting of the North American Catalysis 

Society, Philadelphia (2005). 

P. Park and C. Boyer, “Effect of SO2 on the 

Activity over Ag/γ-Al2O3 Catalysts for NOx Reduction in Lean Conditions,” presented at the 19th

Annual Meeting of the North American Catalysis 

Society, Philadelphia (2005). 

P. Park and C. Boyer, “Effect of SO2 on the 

Activity over Ag/γ-Al2O3 Catalysts for NOx Reduction in Lean Conditions,” Appl. Catal. B., 59, 27–34, 

2005. 

16 
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

B. Development of Materials Analysis Tools for Studying NOx Adsorber 

Catalysts (CRADA No. ORNL-02-0659 with Cummins, Inc. ) 

Thomas Watkins, Larry Allard, Doug Blom, Michael Lance, and Harry Meyer 

Oak Ridge National Laboratory 

P.O. Box 2008 

Oak Ridge, TN 37831-6064 

(865) 574-2046; fax: (865) 574-3940; e-mail: watkinstr@ornl.gov 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Objective 

• Produce a quantitative understanding of the processing and in-service effects on NOx adsorber catalyst 

technology leading to an exhaust aftertreatment system with improved catalyst performance capable of meeting 

the 2007 emission requirements. 

Approach 

• Characterize lab-engine tested samples with X-ray diffraction (XRD), spectroscopy, and microscopy. 

• Correlate findings with Cummins data and experience. 

Accomplishments 

• Evaluated gradient formation on a macro scale of active elements on a catalyst as a function of catalyst history 

and operating conditions. 

• Supported continued characterizations of new materials from various stages of the catalyst’s life cycle. 

• Begun determination of the soot and ash distribution as a function of macroscopic position within the NSR 

catalyst support brick. 

Future Direction 

• Evaluation of thermal degradation of a NOx stoichiometric ratio (NSR) catalyst as a function of macroscopic 

position within the catalyst support brick. 

• Continue the determination of the soot and ash distribution as a function of macroscopic position within the 

NSR catalyst support brick. 

Introduction will need to integrate aftertreatment with engine 

control systems. Currently, no commercial off-the- To meet the 2007 emission requirements for shelf technologies are available to meet these diesel exhaust, aftertreatment in diesel engines may standards. Consequently, Cummins, Inc. is working be necessary. The technology necessary for 2007 to understand the basic science necessary to 

17 
FY 2005 Progress Report 

effectively utilize these catalyst systems. Oak Ridge 

National Laboratory (ORNL) is assisting with the 

materials characterization effort. 

Base metal oxides (BMOs) are major components in current SOx and NOx adsorbers that Cummins seeks to use in NOx adsorber catalyst systems. 

The function of these adsorbers is to collect oxysulfur (SOx) and surface nitrite/nitrate (NOx) species. 

These species are to be released from these surface 

sites during regeneration, where the adsorber BMO 

is either heated to some critical temperature or 

exposed to a reducing or reactant atmosphere. Sulfur 

poisoning of adsorber catalysts is a major problem 

that must be resolved for BMO-based emission 

reduction technologies to become commercially 

viable. 

Approach 

In general, the crystal structure, morphology, 

phase distribution, particle size, and surface species 

of catalytically active materials supplied by 

Cummins will be characterized using XRD, Raman 

spectroscopy, and electron microscopy. These materials will come from all stages of the catalyst’s life 

cycle: raw materials, as-calcined, sulfated, regenerated, etc. Both ORNL and Cummins personnel have 

participated in this work. 

Samples 

Briefly, Cummins provided various series of 

“core” samples taken from cordierite “bricks” containing platinum, γ-alumina and BaCO3 or platinum 

and γ-alumina. The general features of each series 

will be given below in the text as needed. The 

brick/core samples consisted of a series with varied 

amounts of platinum loading, a sample that had lost 

performance after numerous desulfations, and a 

sample loaded with sulfur. 

Results 

The following is a summary of the work from 

October 2004 to September 2005. 

X-ray 

X-ray diffraction (XRD) was employed to 

understand the crystalline nature of the samples. For 

the unfamiliar, the following analogy can be applied: 

Heavy Vehicle Propulsion Materials 

As a fingerprint identifies a person, so a diffraction 

pattern identifies a crystalline material. 

A model system was analyzed and consisted of 

platinum particles dispersed on a γ-alumina (Al2O3) 

on a cordierite (Mg2Al4Si5O18, the core/brick 

material) substrate. This series of samples were 

“redoxed pulsed” to remove nitrogen from the oxidation catalyst. Here, redox pulsing entails going 

from a rich/reducing fuel mixture to a lean/oxidizing 

fuel mixture and back. Samples A, B, C and D were 

designated the baseline, 6 h at 350°C, 6 hours at 

350°C-low volume, and 6 h at 450°C, respectively. 

The conditions for sample D were regarded as the 

most severe and would be the most likely to cause Pt 

crystallite growth. Figure 1 compares the XRD 

pattern of sample D to the as-received γ-alumina 

powder. The powder has a high surface area, and the 

very broad peaks of the powder indicate a small 

crystallite size, 3 nm. A crystallite is a region of 

coherent diffraction and can be thought of as a subgrain, which is a very small perfect crystal. The 

narrower alumina peaks of sample D, indicate that 

the crystallite size has increased and that the surface 

area has likely decreased. Figure 2 shows that the γalumina has transformed to θ-alumina and indicates 

that sample has experienced temperatures in the 727 

to 1027°C range.1

 Note the platinum and θ-alumina 

superimpose. Figure 3(a) shows the experimental 

set-up used to collect platinum crystallite information as a function of macroscopic length along the 

sample. As was inferred above, the x-ray crystallite 

size is related to the x-ray peak broadening. Typically, the size of a crystallite is smaller than the 

grain size of a material as observed on the scanning 

electron microscopy (SEM). Significant contributions to x-ray peak broadening occur when the 

crystallite size is less than 100 nm. The lower bound 

is approximately 2 nm. Figure 3 shows the crystallite size data for samples A, B, C, and D. The crystallite size did not vary systematically along the 

length. As will be shown later, the size does not 

agree with that found in the microscopy section. 

This is likely due to the aforementioned peak 

superposition. Thus, the crystallite sizes given may 

more accurately reflect the size of the θ-alumina. 

18 
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 1. The diffraction patterns for the as-received γ-alumina powder and sample D (redox pulsed 

for 6 h at 450°C). 

Figure 2. The diffraction pattern for sample D (redox pulsed for 6 h at 450°C). 

19
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

(a) (b) 

Figure 3. (a) Silicon wafer masks allow measurement of data every 10 mm. (b) Platinum/alumina crystallite size as a 

function of length along the sample. 

Raman Spectroscopy 

Raman spectroscopy measures the characteristic 

vibrational energy levels of molecules and crystals 

and so is very sensitive to any changes in bonding, 

stoichiometry, and phase/symmetry. 

Soot oxidation has been a critical issue for 

regeneration of diesel particulate filters (DPFs). The 

design of the regeneration strategy and the durability 

requirement for DPFs depend on an understanding 

of soot burning. Raman spectroscopy has been used 

to monitor surface configuration changes. Under 

thermal aging, engine soot undergoes a transformation from a mostly amorphous carbon structure 

into more orderly polyaromatic structures, which 

exhibit resistance toward oxidation. Certain 

neutralization products, such as thio-phosphorous 

and phosphorous derivatives, can also alter the 

thermal behavior of soot through interactions 

between ash- and soot-adsorbed species. The 

incomplete regeneration can be attributed to a 

buildup of a physical barrier that prevents oxygen 

and/or NO2 from diffusing effectively onto the 

carbon surface. 

Raman spectroscopy offers a unique way to 

characterize the long-range interactions and the 

order-disorder effects on carbonaceous materials. 

Both the first and the second-order features can be 

seen with Raman. Pristine graphite crystals show 

two first-order Raman features located at 42 and 

1578 cm–1. The 1578 cm–1 band, with a full width at 

half-maximum bandwidth of 16 cm–1 (see Figure 4), 

is referred to as the graphitic band (G). Grinding the 

graphite crystal generates an additional band at 

1335 cm–1, which can be assigned as the disordered 

band (D). Additional grinding gives an increase in 

the intensity and bandwidth of the disordered band. 

The appearance of the disordered band can be 

attributed to the disorder relaxation where the wavevector selection rules are broken by lattice discontinuities and structural defects. 

The Raman spectra of the fresh and the artificially matured samples of A1C are shown in 

Figure 4. After heating at 950°C, additional bands 

begin to show around 680 ~ 1000 cm–1, which can 

be assigned to metal oxides. When samples are 

heated at high temperatures, spectral interference 

20 
Heavy Vehicle Propulsion Materials 

Figure 4. Raman spectra of A1C engine soot aged at 

various temperatures. 

from ash components with the carbon bands begins 

to become pronounced. As the maturation temperature is increased, the graphitic band becomes 

sharper, and the intensity is enhanced. The transition 

frequency of the disordered band is also shifted to a 

slightly lower frequency. Similar trends are also 

seen in other samples. All these spectral trends 

suggest a structural change of the sample during 

thermal maturation. 

The frequency ratios of the graphitic to disorder 

peak (R = fG/fD) measured for all three samples are 

plotted in Figure 5. All three samples follow a general trend with an increase of R with the maturation 

temperature. The increase of R for soot maturation is 

from 0.37 (<650°C) to 0.48 (>850°C). The change 

shows little statistical significance at temperatures 

Figure 5. The correlation between Raman band ratio 

(R = fG/fD) and temperature. 

FY 2005 Progress Report 

below 700°C. Above 750°C, a rapid enhancement is 

observed, and the curve levels off beyond 850°C. 

The major change seems to occur around 700°C. 

This temperature is consistent with most young coal 

samples where their maturation scale changes dramatically under artificial pyrolysis around 700°C. 

This result shows that Raman spectroscopy can be 

used to monitor the degree of long-range order in 

soot samples, which is a key parameter in determining the regeneration processing of DPFs. More 

research is planned to better understand the kinetics 

of soot burning and DPF aging. 

Microscopy 

Scanning transmission electron microscopy 

(STEM) and electron probe microanalysis (EPMA) 

techniques were used to continue characterization of 

the microstructure of NOx trap catalysts and oxidation catalysts. In the previous report, the microstructure and chemistry of samples from Reactor II 

(platinum loading ~2 wt %) at different positions 

along the reactor were studied, and the changes in 

morphology of platinum particles as well as the distribution of elements such as the trapping component within the washcoat were elucidated. Subsequently, a variety of both NOx trap catalysts and 

oxidation catalysts were examined to quantify the 

precious metal particle size distributions both as a 

function of aging as well as macroscopic position in 

the reactor. In addition, the location and structure of 

the BMO trapping component of the NOx trap 

systems were determined. Some example results and 

main conclusions from the microscopy experiments 

are given below. 

The platinum particles of a fresh Pt/Ba/ 

γ-alumina model catalyst sample were examined. 

The sample had a target loading of 1.6 g/in.3

 and 

was exposed to 500°C in its preparation. The 

majority of platinum particles were observed to be 

between 1 and 2 nm in size, as can be seen in 

Figure 6. The HA-ADF STEM image has bright 

contrast for heavy elements, while lighter elements 

appear darker. The particles were too fine for quantification via image processing. “Spectral imaging” 

is an advanced technique for elemental analysis in 

electron beam instruments whereby a full EDS 

spectrum from 0–20 keV is collected at every individual pixel location. The collected data can be 

manipulated off-line to generate 1-dimensional and 

21
FY 2005 Progress Report 

Figure 6. HA-ADF (Z-contrast) STEM image of 

platinum particles on alumina washcoat in 

sample PBA1. The majority of the particles 

are 1–2 nm in diameter in this micrograph. 

Heavy Vehicle Propulsion Materials 

2-dimensional elemental line scans and maps. Figure 

7 shows both the HA-ADF STEM image and an 

extracted barium X-ray map for the sample. The Xray map conclusively shows the spatial distribution 

of the barium-containing phase in this sample, and 

comparison with the high angle-annular dark field 

(HA-ADF) image allows for the identification of 

barium containing areas relative to the morphology 

and brightness seen in the HA-ADF image. Samples 

containing Pt/BMO/γ-alumina were also examined. 

Spectral imaging identified the trapping component 

as BMO-based. As seen in Figure 8, the BMO 

phase is widely distributed but not as easily identified from just the HA-ADF images. 

Unlike previous specimens, sample D (see X-ray 

section above) came as a much longer core 

Figure 7. (a) HA-ADF STEM image. (b) Barium X-ray map extracted from spectral image. Barium containing 

phase corresponds to the large bright area in (a). 

Figure 8. (a) HA-ADF STEM image. (b) BMO X-ray map. There has been some uncompensated shift 

between the reference image and the map. The BMO phase is widely distributed, but not easily 

identified in the HA-ADF image alone (a). 

22 
Heavy Vehicle Propulsion Materials 

with the request to analyze the platinum particle size 

distribution down the length of the monolith. Sample 

D1 was taken near the smooth/flat upstream face of 

the monolith [see left side of inset in Figure 3(b)] 

and was closer to the engine. Sample D2 was from 

an area 7 cm away along the length of the monolith, 

near the rough/jagged side. Figure 9 shows the 

comparison of the platinum particle size 

distributions for samples D1 and D2. Sample D1 

exhibits a very typical Gaussian distribution of particle sizes, while sample D2 has a bimodal distribution with both a population of small catalyst parti 

cles and large catalyst particles. For 757 particles in 

Figure 9. Comparison of platinum particle size 

distributions from two different areas on the 

same monolith. Each distribution is based on 

ten images each. 

FY 2005 Progress Report 

D1, the median and mean particle sizes were 6.82 

and 6.76 nm, respectively. For 1125 particles in D2, 

the median and mean particle sizes were 5.87 and 

5.48 nm, respectively. The higher temperatures that 

the platinum particles at location D1 experienced 

appear to have coarsened the particle slightly. 

References 

1. R-S. Zhou and R. L. Snyder, “Structures and 

Transformation Mechanisms of the η, γ and θ Transition Aluminas,” Acta Cryst., B47, 617–30 (1991). 

Presentations and Publications 

T. R. Watkins, L. Allard, D. Blom, M. J. Lance, 

H. Meyer, L. Walker, C. Narula, R. D. England, 

B. Epling, H. Fang, and T. Yonushonis, 

“Development of Materials Analysis Tools for 

Studying NOx Adsorber Catalysts,” [Revised] 

presented at 11th Diesel Engine Emission Reduction 

(DEER) Conference, Chicago, Illinois, August 24, 

2005. 

T. R. Watkins, L. Allard, D. Blom, M. J. Lance, 

C. Narula, H. Meyer, R. D. England, H. Fang, 

T. Yonushonis, and H. Hess, “Catalyst Characterization, CRADA No. ORNL-02-0659 with Cummins, 

Inc.,” presented at the Heavy Vehicle Materials 

Program Review, Oak Ridge National Laboratory, 

Oak Ridge, Tennessee, September 14, 2005. 

H. L. Fang, M. J. Lance, and G. S. Painter, 

“Raman Characterization of Pt-support Interactions 

on Metal Oxides,” in preparation. 

23

______________________ 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

C. Development of NOx Sensors for Heavy Vehicle Applications 

Timothy R. Armstrong, David L. West, Fred C. Montgomery, Jaco H. Visser,* and David J. Kubinski* 

Oak Ridge National Laboratory 

P.O. Box 2008, Mail Stop 6186 

Oak Ridge, TN 37831-6186 

(865) 574-7996; fax: (865) 241-0112; e-mail: armstrongt@ornl.gov 

*Scientific Research Laboratory, Ford Motor Co. 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Prime DOE Contract No.: DE-AC05-00OR22725 

Subcontractor: Ford Motor Company, Detroit, Michigan 

Objectives 

• Develop NOx sensors for remediation and monitoring of diesel engines. 

• Develop sensors that have an operating temperature of 500–700oC and are able to measure NOx concentrations 

from ∼10–1500 ppm at oxygen levels from 5 to 20 vol %. 

• Ensure that sensors are selective for either NO or NO2, or able to measure “total NOx” ([NO] + [NO2]) since 

“NOx” refers to mixtures of NO and NO2. 

Approach 

• Fabricate prototype sensing elements by screen printing electrode layers onto oxygen-ion conducting [typically 

yttria stabilized zirconia (YSZ)] substrates. 

• Operated elements either in a “non-Nernstian” mode (where the output is a voltage) or under direct current 

electrical bias. 

• Characterize sensor in bench setting and in NG engine exhaust for NOx sensor response, sensitivity to varying 

[O2], and recovery/response kinetics. 

• Develop breadboard based sensor for exhaust gas sensing. 

Accomplishments 

• Developed sensing elements displaying near “total NOx” behavior and stable in simulated long-term service. 

• Developed sensing elements that are capable of repeatedly detecting [NOx] in the range of ~10 ppm and can 

operate at temperatures as high as 750oC with ±1 ppm sensitivity. 

• Developed novel techniques for applying electrical stimulus to combat drift of the sensor background output. 

25 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Future Direction 

• Evaluate the stability and performance of these sensing elements in real exhaust streams to quantify and correct 

for drift. 

• Characterize the cross-sensitivity to other species (e.g., CO) that may be present in real exhaust streams. 

• Clarify both the sensing mechanism and the cause(s) for sensor drift. 

Introduction 

The primary pollutants (excluding the greenhouse gas CO2) from the combustion of low-sulfur 

fuels are CO, hydrocarbons (HC), and oxides of 

nitrogen (“NOx,” a mixture of NO and NO2). Sparkignited, direct-injection (SIDI) auto and truck 

engines employ a three-way catalyst (TWC) that can 

remove all three of these pollutants. However, the 

TWC is only effective over a narrow range of [O2] 

in the exhaust, losing its effectiveness for NOx 

removal at high partial pressures of O2. This means 

that the currently employed TWC is not effective for 

NOx remediation of the exhaust from compressionignited (diesel) engines, as these are O2-rich 

(typically 5–20 vol % O2). 

In this current absence of a “lean-NOx” catalyst, 

three technologies are proposed to meet the challenge of NOx remediation of diesel exhausts: “LeanNOx” traps (LNTs), selective catalytic reduction 

(SCR) with hydrocarbons, and SCR with urea. All of 

these technologies will require on-board NOx 

sensors, to control trap regeneration (LNT) or 

reagent injection (SCR). 

In this project we are developing NOx sensing 

elements that can be used in diesel engine exhausts. 

The minimum requirement for this application is an 

operating temperature of 500–700oC and the ability 

to measure 1–1500 ppm NOx. It should be emphasized that “NOx” refers to mixtures of NO and NO2; 

therefore, full NOx characterization of the exhaust 

may require two of the following three concentrations to be measured: [NO], [NO2], [NOx] (= [NO] 

+ [NO2]). This project is a cooperative research and 

development agreement with Ford Motor Company. 

Approach 

The approach selected is to build a sensing platform on yttria-stabilized zirconia (YSZ). YSZ, an 

oxygen-ion conducting material, is currently widely 

used in [O2] sensors for gasoline engines. This 

approach is not original, because two NOx sensors 

based on YSZ are near commercialization and have 

been proven by the engine manufacturers association 

not to be adequate to measure NOx less than 

50 ppm. In addition, both are relatively complex 

with multicavity constructions. This construction is 

required because in the first design2 preliminary O2 

removal from the sampled exhaust is required before 

the NOx is measured amperometrically. In the 

second design3 the sampled exhaust passes over a 

conversion electrode in order to convert the NOx 

into NO2 before measurement in the sensing cavity. 

The lion’s share of work to date has focused on 

the sensing element portion of the sensor. This is the 

portion of the sensor that responds to the presence of 

NOx. The reason for adopting this approach is that 

the working characteristics of the sensing element 

will dictate many aspects of sensor design (operating 

temperature, need for catalytic overlayers, etc.). 

To fabricate sensing elements, electrodes are 

screen-printed and fired onto YSZ substrates 

(Figure 1). These elements are then tested for NOx 

response using a furnace to simulate elevated temperature service and a commercial gas mixing unit to 

deliver blends of N2, O2, NOx (either NO or NO2), 

and other species as required. 

Results 

Much of the work on this project has been 

devoted to all-oxide “total NOx” sensing elements. 

These sensing elements, a sample of which is 

depicted in Figure 1, consist of compositionally 

identical oxide electrodes on a YSZ substrate. These 

sensing elements are attractive because the use of a 

single electrode material will reduce production 

cost, and more importantly, their “total NOx” (similar response to NO and NO2) sensing behavior 

26
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 1. Prototype sensing element. 

would enable measurement of [NOx], regardless of 

the [NO]:[NO2] ratio in the exhaust gas. This obviates the necessity to perform any NOx conversion of 

the type described in Ref. 3. 

Several electrode compositions have been 

evaluated in order to develop a composition with 

large NOx response magnitudes. The best of these, 

when operated at 600oC and 7 vol % O2, had 

responses to 77 ppm NO and NO2 of –37%* and 

–40%, respectively. However, long-term (>200-h) 

lab-scale testing of this composition revealed that it 

was unstable in exhaust conditions, and thus the 

development of new compositions, stable in exhaust 

conditions, has been carried out. 

The NOx sensing performance of two promising 

new compositions is shown in Figure 2. It can be 

seen that the NOx responses are similar in magnitude 

over the concentration range 20–190 ppm, and 

efforts are ongoing to narrow the response between 

[NO] and [NO2]. 

Figure 3 shows X-ray diffraction patterns from 

extensively tested sensing elements made with old 

electrode composition and the new composition. The 

new composition is clearly more stable as shown by 

the lack of decomposition products. 

The data in Figure 2 were collected at 600oC. It 

may be desirable to operate the sensing element at 

higher temperatures, both in order to improve the 

response/recovery time and prevent condensation on 

the sensing element. The higher temperature NO 

*This is the change in measured voltage when NOx is 

present. 

Figure 2. NOx responses of two new electrode 

compositions. 

Figure 3 . X-ray diffraction patterns of used elements 

made with (a) original and (b) new 

compositions. 

sensing performance of these sensing elements is 

illustrated in Figure 4, which shows sensing element 

performance at 750oC. Two aspects of the sensing 

performance in Figure 4 differ from that observed in 

Figure 2. First, the magnitude of the changes due to 

27 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 5. Response to identical changes in [NO] 

and [O2] after 0, 10, and 70 h of 

continuous operation. 

Figure 4. Sensing performance at 750oC. In (a) [NO] is 

stepped through the sequence 0, 77, 49. 31, 20, 

10, 0 ppm at the indicated [O2]. Computed 

changes in voltage are shown in (b). 

NO are much smaller. Second, the dependence of 

the response on [NO] is linear. Note also that the 

response decreases with increasing [O2]. This means 

that operation at these temperatures will most likely 

require lowering of the [O2] in the sampled exhaust 

(by means of a pumping cell). 

These sensing elements are electrochemical in 

nature, and electrochemical sensing elements are 

prone to drift. Drift is frequently observed because 

electrochemical sensors usually rely on charge transfer at interfaces, and these interfaces will age with 

use of the sensor. 

Sample drift behavior observed with these 

sensing elements is illustrated in Figure 5. Shown 

there are the element responses to identical programmed changes in [NO] and [O2] after 0, 10, and 

70 h of continuous operation. The background 

resistance of the element is monotonically increasing, but the responses to the gas concentration 

changes are repeatable. 

To combat this drift, we have been investigating 

alternate methods of applying an electrical stimulus 

to the sensing element. One such example is shown 

in Figure 6. 

Figure 6. Element performance with alternating 

application of positive and negative 

electrical stimulus. 

The electrical stimulus was alternated between 

positive and negative values (only the negative 

values are shown in Figure 6). During every third 

negative excursion, the [NO] concentration was 

varied systematically as indicated in Figure 6. Operated in this manner, the sensing element is capable 

of repeatedly detecting [NO] changes on the order of 

10 ppm over more than 10 h of continuous 

28
Heavy Vehicle Propulsion Materials 

operation. These are thought to be the demands that 

monitoring of a lean NOx trap would place on a NOx 

sensor. 

Conclusions 

Sensing elements capable of “total NOx” sensing behavior are under development. During the 

course of this project the following accomplishments 

have been achieved: 

• Sensing elements displaying near “total NOx” 

behavior and stable in simulated long-term 

service have been developed. 

• NO sensing elements have been demonstrated. 

• Sensing elements are capable of repeatedly 

detecting [NOx] in the range of ~10 ppm and 

can operate at temperatures as high as 750°C 

with ±1 ppm sensitivity. 

• Sensing elements stable for more than 1000 h in 

oxygen and steam containing environments have 

been developed. 

The challenges ahead are fourfold: (1) the 

stability and sensitivity of these elements must be 

examined in real exhaust streams, (2) the crosssensitivity to other species that may be present in the 

exhaust must be characterized, (3) techniques for 

combating drift must be introduced and optimized, 

and (4) the sensing mechanism of these elements 

needs to be elucidated. 

This last item is important as ease of fabrication 

and cost may dictate changes in materials or geometry during the incorporation of these elements into 

sensors suitable for any engine application. Understanding of the sensing mechanism will be critical in 

predicting how such changes might affect the sensor 

performance. 

References 

1. K. B. J. Schnelle and C. A. Brown, “NOx 

control,” in Air pollution control technology 

handbook, CRC Press, 2001, pp. 241–255. 

FY 2005 Progress Report 

2. N. Kato, H. Kurachi, and Y. Hamada, “Thick 

film ZrO2 sensor for the measurement of low NOx 

concentration,” SAE Technical Paper 980170, 1998. 

3. A. Kunimoto, M. Hasei, Y. Yan, Y. Gao, 

T. Ono, and Y. Nakanouchi, “New total-NOx sensor 

based on mixed potential for automobiles,” SAE 

Technical Paper 1999-01-1280, 1999. 

Patents 

F. C. Montgomery, D. L. West, T. R. 

Armstrong, and L. C. Maxey, “NOx Sensing 

Devices Having Conductive Oxide Electrodes,” 

ORNL Docket 1348C. 

Presentations and Publications 

1. D. L. West, F. C. Montgomery, and T. R. 

Armstrong, “Total NOx sensing elements with 

compositionally identical oxide electrodes,” 

accepted by Journal of the Electrochemical 

Society, September 2005. 

2. D. L. West, F. C. Montgomery, and T. R. 

Armstrong, “NO-selective NOx sensing 

elements for combustion exhausts,” accepted by 

Sensors and Actuators B, February 2005. 

3. D. L. West,* F. C. Montgomery, and T. R. 

Armstrong, “All-oxide “total NOx” sensing 

elements,” 207th Meeting of the Electrochemical 

Society, 2005. 

4. D. L. West, F. C. Montgomery, and T. R. 

Armstrong, “DC electrical-biased, all-oxide 

NOx sensing elements for use at 873 K,” 29th 

International Cocoa Beach Conference on 

Advanced Ceramics and Composites, 2005. 

5. D. L. West, F. C. Montgomery, and T. R. 

Armstrong, “High-T NOx sensing elements 

using conductive oxides and Pt,” Proceedings of 

ICEF: Engines for Mobile, Marine, Rail, Power 

Generation and Stationary Applications, 2004. 

29 

____________________ 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

D. Ultra-High Resolution Electron Microscopy for Characterization of Catalyst 

Microstructures and Deactivation Mechanisms 

L. F. Allard, D. A. Blom, D. W. Coffey, C. K. Narula, and M. A. O’Keefe* 

Oak Ridge National Laboratory 

P.O. Box 2008, MS-6064 

Oak Ridge, TN 37831-6064 

*Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, CA 94720 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Objectives 

• Develop and utilize new capabilities and techniques for ultra-high resolution transmission electron microscopy 

(UHR-TEM) to characterize the microstructures of catalytic materials of interest for reduction of NOx

emissions in diesel and automotive exhaust systems. 

• Relate the effects of reaction conditions on the changes in morphology of heavy metal species on “real” catalyst 

support materials (typically oxides). 

• Develop capability to experimentally study the effects of experimental reaction treatments on TEM samples of 

NOx trap catalyst materials before and after reactions, using the High-Temperature Materials Laboratory 

(HTML) “ex-situ” catalyst reactor with a specimen holder designed for use in the aberration-corrected electron 

microscope (ACEM). 

Approach 

• Utilize Oak Ridge National Laboratory ACEM to characterize for the first time the atomic morphology of 

heavy-metal species such as platinum and rhenium on oxide support materials via techniques of annular darkfield (ADF) or “Z-contrast” imaging. 

• Utilize dedicated scanning transmission electron microscope (STEM) instrument and techniques of spectrum 

imaging to characterize the structure and chemistry of a series of lean NOx trap (LNT) and diesel particulate 

NOx reduction (DPNR) filter samples being studied by Ford Research Laboratory colleagues. 

Accomplishments 

• Characterized the morphologies of tri-rhenium carbonyl clusters on gamma-alumina support material, at the 

atomic level to demonstrate for the first time the unambiguous imaging of clusters of known structure. 

• Characterized the nature of platinum catalyst structures on mixed-oxide support materials, including the 

imaging of platinum species dispersed as single atoms and in precrystalline clusters known as “rafts.” 

• Performed initial analysis of the structures of several manufacturers’ NOx trap and DPNR filter materials, in the 

as-manufactured condition and after an initial “de-greening” process. 

31 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

• Demonstrated the potential for the Z-contrast atomic imaging process to allow determination of the precise 

shapes of nanoparticles and clusters. 

• Developed and fabricated a new specimen rod and modifications to the HTML ex-situ reactor system to permit 

catalyst reaction studies using the ACEM. 

Future Direction 

• Focus on the use of the ACEM to characterize the location of sulfur species on NOx reduction catalyst 

materials, to enable characterization of the mechanisms of sulfur poisoning of these materials. This will involve 

not only imaging of the catalysts at the single-atom level, but also the use of the electron spectroscopy 

capabilities on the ACEM to obtain chemical information at the atomic level. 

• Continue to work on the critical problem of determination of precise shapes of catalyst nanoparticles to determine the most important characteristics for control of catalytic performance. 

• Extend new imaging and spectroscopy techniques to include studies of bi-metallic catalyst particles (e.g., goldpalladium). 

Technical Progress 

Imaging of ultrafine atomic clusters of heavy 

metals on low atomic number supports is expected 

to be significantly enhanced with the advent of 

research on the new aberration-corrected 

STEM/TEM electron microscope (JEOL 2200FSAC “ACEM”). In this report, we illustrate the 

potential for imaging atomic species down to the 

single atom level, with samples of platinum on thin 

carbon films, platinum on and rhenium on alumina. 

The latter samples are part of a specimen suite comprising primarily tri-rhenium carbonyl clusters on 

gamma alumina, provided by Prof. B. C. Gates and 

students at UCalDavis. Details of the changes in 

cluster morphology related to decarbonylation 

experiments in an ex-situ reactor system will be 

reported by Narula, Moses, and Allard in a separate 

quarterly. Platinum on carbon samples were prepared by sputter deposition from a platinum target 

using argon ions, in our ion-sputtering unit. The 

ultra-thin carbon films were prepared by controlled 

evaporation of carbon onto NaCl substrates, then 

floating films (nominally 10Å to 15Å thick) on 

water, and picking the thin films up on holey carbon 

support films. The thin carbon is essentially not 

imaged in the ADF imaging mode of the ACEM, 

allowing only platinum atoms and clusters to be 

revealed in bright contrast. 

The platinum on carbon experiment was conducted using two specimens. One sample was prepared using dual-gun sputtering for 20 s under 

standard beam energy conditions. The second 

sample was prepared using identical beam 

conditions, but only a single gun, giving nominally 

one-half the amount of platinum species on the 

carbon film surface. The ACEM was adjusted to 

achieve a fully corrected electron probe, using an 

illuminating aperture that calculated to produce a 

beam diameter of just about 1 Å. 

Platinum/carbon imaging results: Figures 1 and 

2 show the ADF images of the 1-gun and 2-gun 

depositions, respectively. Single atoms and “rafts” 

of atoms having 2–3 layers are seen in both samples. 

It is clear that there is a direct correlation between 

the deposition parameters and the relative amounts 

of platinum observed on the respective samples. 

Also note that these images were acquired prior to 

achieving fully stable operation on the instrument, 

and also while using a nonideal aperture size in the 

probe-forming lens system, so that no better than 

1-Å resolution was expected. The line profile shown 

over a single platinum atom (raw data) indicates just 

over 1Å in diameter, as expected. The impressive 

result from this experiment is that we expect to be 

able to image single atoms of a heavy element with a 

resolution significantly lower than the expected 

0.7-Å ultimate resolution of the microscope. A second observation (not illustrated here) is that the 

atoms and clusters are mobile on the surface of the 

carbon film and continually change position between 

successive image acquisitions. This phenomenon is 

not unexpected and has been observed in other sample types observed to date, including the rhenium on 

alumina samples described below. It is important to 

be able to (eventually) determine the relationship 

between images recorded at highest resolution on 

32
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 1. (a) Platinum sputtered onto thin carbon film, showing single atoms, clusters, and “rafts.” (b) A 

second sputter coating of a carbon film with platinum, accumulating about double the total 

platinum relative to the first coating. Rafts are larger and fewer clusters are present, but no discrete 

crysallinity is evident even in the larger rafts. 

Figure 2. (a) BF image of thin plate of g-Al2O3, decorated with rhenium atoms. The phase contrast from the 

alumina obscures unambiguous determination of location of rhenium species. (b) Same area in 

ADF image, showing rhenium species in bright contrast. Single atom at A, apparent dimmer at B, 

clear trimer at C, small raft of atoms at D. 

the ACEM and the original distribution of atoms and for providing a reliable quantification of cluster 

clusters on as-prepared catalyst specimens. Correla- morphologies and distributions. 

tion of imaged morphologies with results from other Rhenium/g-Al2O3 imaging results: Figure 2 is a 

measurement techniques, again as being done in the bright-field (BF) STEM image of a thin γ-Al2O3 

work with UCalDavis, may prove the most valuable flake, showing high phase contrast and an inability 

33 
FY 2005 Progress Report 

to unambiguously distinguish the precise location of 

rhenium atoms and tri-rhenium clusters. The ADF 

image of the exact same sample area is shown in 

Figure 3, where the rhenium species are clearly seen 

in bright contrast. These images were also acquired 

with a probe diameter of about 1 Å. A single atom 

(A), an apparent dimer (B) and a tri-rhenium cluster 

(C) are seen (the raw digital images were smoothed 

using a 3 × 3 filter). Interestingly, one atom of the 

“dimer” is brighter in contrast than the other; an 

intensity profile (Figure 4) shows double the intensity in the bright atom, which is consistent with the 

intensity which would result from two atoms, one 

above the other, in the trimer arrangement shown in 

the cartoon. It is therefore likely that the dimer is 

actually a trimer seen precisely on edge, a configuration possible if the trimer were attached to a ledge 

on the alumina support just a few Ångströms in 

height. The trimer at C shows equal intensity for the 

individual atoms and is therefore a classical trirhenium cluster. The atom-atom spacing in trimer C 

is slightly different than the spacings in a nearby 

Heavy Vehicle Propulsion Materials 

cluster D (2.0 Å vs 2.3 Å), suggesting that the rhenium atoms (perhaps still carbonylated) are influenced by the crystallography of the γ-alumina crystal 

lattice. Direct observations of atoms and clusters 

such as illustrated here, made possible by the new 

generation of aberration-corrected electron microscopes, and coupled with molecular simulations, 

offer the potential for a deeper understanding of the 

changes in morphology of heavy metal species in 

catalyst systems with aging treatments. 

Characterization of fine particle morphology: 

The direct relationship in the ADF image between 

number of atoms in a column and the intensity of the 

column was shown clearly in the example of the trirhenium cluster on edge above. In our FY 2004 

annual report we also demonstrated this effect for an 

ideal TEM image of a platinum cube-octahedron. 

Here we show a bright-field (BF) image of a 7-nm 

gold particle supported on a 10-nm-thick amorphous 

carbon film, processed to extract the so-called “exit 

surface wave” image, at a resolution in the 1-Å 

range. The problem with the BF images is that the 

Figure 3. Analysis of the dimer at B shows one “atom” is twice the intensity of 

the adjacent, where the adjacent atom intensity profile is precisely the 

height of the single atom at A. This suggests the dimer is actually a 

trimer tethered to a ledge in the alumina surface and is oriented edge-on 

to the electron beam, as indicated by the cartoon. 

34
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 4. (a) Reconstructed BF Exit Surface Wave image of 7-nm gold particle, with intensity profile inset. 

The profile shows discrete steps in intensity column by column, but the carbon background makes 

interpretation ambiguous. (b) Direct ADF image from STEM, with intensity profile A showing 

number of atoms column by column. 

phase contrast from the carbon film makes a significant contribution to the contrast of the atomic columns in the particle. However, in ADF images, 

because of its very low scattering power relative to 

the high atomic number of the heavy metal particle, 

the carbon film contributes essentially no signal to 

the image. Figure 4 compares the gold particle on 

carbon in BF [Figure 4(a)] with the gold particle on 

carbon [Figure 4(b)] in the ADF aberrationcorrected STEM image. Because of the phase contrast due to the carbon film in the BF image, the 

interpretation of discrete intensity jumps column to 

column are ambiguous. However, a line profile 

along a typical row of atoms in the gold particle 

ADF image shows the discrete jumps from column 

to column. One possible interpretation of the number 

of atoms in the individual columns is given on the 

profile inset in Figure 4(b). These are the most 

promising results obtained to date that show the possibility for ADF imaging in the ACEM to allow full 

determination of the precise shape of nanometersized catalyst particles. Further work is needed to 

work out image processing methods to allow automatic analysis of particle shapes from ADF images. 

Development of new reactor specimen rod: To 

conduct catalyst reaction experiments involving 

examination of catalyst powder samples in the 

ACEM before and after reaction in the HTML exsitu catalyst reactor system, a new specimen rod for 

the ACEM had to be designed and fabricated. The 

new rod was built with the capability to protect the 

catalyst specimen from atmospheric exposure during 

transfer of the sample into the microscope, then back 

to the reactor for sequential reaction experiments. 

Figure 5 shows the reactor rod [Figure 5(a)] and the 

rod tip [Figure 5(b–d)]. The tip contains a recess to 

allow a small disc that would hold a specimen grid 

to be installed, as shown in Figure 5(c). By manipulating the control knob at the end of the rod, thin 

sheets of phosphor-bronze spring stock can be slid 

above and below the specimen grid, allowing the 

grid to be encapsulated in an inert gas (with a metalto-metal seal) during the transfer. Experiments utilizing the new capability on the ACEM for reaction 

studies will be detailed in forthcoming quarterly 

reports. 

35 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 5. (a) Specimen rod designed for use with the ACEM to permit catalyst reaction studies in the 

HTML ex-situ reactor system. Tip details shown below. (b) Tip of specimen rod with disc to hold 

specimen grid adjacent. (c) Rod tip with disc in place. (d) Rod manipulated to position thin metal 

sheets above and below the specimen (metal-to-metal seal), allowing encapsulation of the sample 

in an inert gas. 

Analysis of Supplier Catalysts 

In partnership with colleagues at Ford Research 

Laboratory and the National Transportation 

Research Center (NTRC), catalyst materials including lean NOx trap (LNT) and diesel particulate NOx 

reduction (DPNR) materials have been characterized 

in a program to understand the structures and 

behavior of catalysts being used in vehicles overseas 

(Europe). The LNT materials were supplied by 

Umicore Co. [their GDI (gasoline direct injection) 

catalyst] and by Toyota (their Avensis D-CAT, or 

Diesel Clean Advanced Technology catalyst). 

Toyota also supplied DPNR material for their 

Avensis vehicle. These materials were characterized 

initially in the “fresh” or as-received condition, with 

general microstructure and chemistry provided by 

electron microprobe analysis, and further highresolution imaging and microanalysis using the 

Hitachi TEM and dedicated STEM instruments in 

the HTML, as well as the ACEM for ultra-high 

resolution imaging. 

A full presentation of the results of the supplier 

catalyst studies is beyond the scope of this report, 

but a few interesting observations are shown as 

examples. The first observation using the electron 

microprobe in backscattered electron imaging (BSE) 

mode showed the Umicore LNT catalyst bulk 

structure to be a standard single washcoat deposition 

on a porous cordierite monolith having channels 

with square cross section. Figure 6(a) shows a typical BSE image of a corner of washcoat, where several blocky aggregate phases are evident (in BSE 

mode, lighter contrast reflects higher average atomic 

number of the phase, so baria would be lighter in 

contrast than alumina, for example). The Toyota 

D-CAT LNT structure was similar (not shown here), 

except a cordierite monolith having channels with 

hexagonal cross section was employed. 

36
Heavy Vehicle Propulsion Materials 

Figure 6. BSE image of corner of a channel in Umicore 

LNT catalyst monolith, showing single-layer 

washcoat composed of blocky aggregates of 

constituents including alumina, baria-alumina 

and ceria-zirconia. 

In contrast to these more standard materials, the 

Toyota DPNR catalyst utilized a more highly porous 

monolith structure, and the washcoat material was 

much finer grained and infiltrated within the pore 

structure, with none remaining on the channel walls, 

as illustrated in Figure 6(b) and 6(c). The DPNR 

structure proved difficult to thin for TEM work, so 

FY 2005 Progress Report 

results of the study of that material will be reported 

in a future quarterly. The LNT catalysts were prepared for TEM analysis by standard ion milling 

techniques, so thin sections of the washcoat were 

available for high-resolution imaging (see Figure 7). 

The Fresh Umicore LNT structure comprised 

alumina, baria-alumina, and ceria-zirconia (primarily ceria) phases, with platinum as the primary 

heavy-metal catalytic species. Interestingly, both the 

alumina and baria-alumina phases contained magnesia (MgO) as a significant constituent (at the 

10 wt % level), which is not typical of similar catalysts from other manufacturers. The role of MgO is 

not yet clear. 

The Fresh Umicore material showed a dispersion of platinum clusters and ultra-fine particles that 

were best imaged with the ACEM, as particles 

ranged from the 1- to 1.5-nm size range (widely dispersed) down to uniformly dispersed clusters at the 

near single atom level. After a “de-greening” treatment (standard practice for LNT catalysts), the 

platinum dispersion was observed in the Hitachi 

HD-2000 STEM with ADF imaging to form discrete 

particles in the 2- to 10-nm size range. Images at 

roughly equivalent magnifications from the ACEM 

[Figure 8(a)] and the dedicated STEM [Figure 8(b)] 

illustrate this particle growth. The advantages of 

Figure 7. (a) Low-magnification image of polished cross section of Toyota DPNR catalyst structure, 

showing infusion of washcoat throughout pores of the cordierite monolith. (b) DPNR structure at 

higher magnification (compare to Figure 6). Washcoat (light contrast) is very fine grained relative 

to the aggregate structure of the LNT washcoat seen in Figure 6. Black areas are embedding 

media, and grey areas are cordierite. 

37 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 8. (a) High-resolution ACEM ADF image of 

ultra-fine platinum clusters level in “Fresh” 

Umicore LNT catalyst, in area of alumina 

phase. (b) ADF image from HD-2000 STEM 

showing particle after a “de-greening” 

treatment. 

imaging heavy-metal particles in the dark-field 

mode, both with the dedicated STEM for standard 

resolution imaging and with the new aberration-corrected instrument are clear. The effective change in 

platinum particle morphologies with further treatments in a benchtop reactor should serve to permit 

further understanding of the degradation in performance of these LNT catalysts, and will be reported in 

later quarterlies. 

References 

1. M. A. O’Keefe, E. C. Nelson, and L. F. 

Allard, “Focal Series Reconstruction of Nanoparticle 

Exit Surface Wave,” Microscopy and Microanalysis, 

2003. 

Presentations and Publications 

L. F. Allard and C. K. Narula, “Electron 

Microscopy of LNT Materials: Microstructural 

Changes with Aging,” Presentation at 2005 

CLEERS Workshop, Dearborn, Michigan, May 

2005. 

V. A. Bhirud, M. J. Moses, D. A. Blom, L. F. 

Allard, T. Aoki, S. Mishina, C. K. Narula, and B. C. 

Gates, “Alumina-supported Tri-rhenium Clusters 

Revealed by Aberration-corrected Electron 

Microscopy,” Microscopy and Microanalysis (2005) 

(First Place Award winner in Physical Sciences 

competition). 

C. K. Narula. L. F. Allard, M. J. Moses, Y. Xu, 

W. Shelton, W. F. Schneider, G. Graham, and 

J. Hoard, “Lean NOx Traps—Microstructural 

Studies of Real World and Model Catalysts,” DEER 

2005, Palmer Hilton, Chicago, August 21–25, 2005. 

M. A. O’Keefe, L. F. Allard, and D. A. Blom, 

“HRTEM Imaging of Atoms at Sub-Ångström 

Resolution,” Journal of Electron Microscopy, 

Special Issue in Memory of Prof. John M. Cowley, 

p. 1–12, August 25, 2005. 

M. A. O’Keefe, L. F. Allard, and D. A. Blom, 

“Resolution Quality and Atom Positions in SubÅngström Electron Microscopy,” Microscopy and 

Microanalysis, in press. 

38
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

E. Microstructural Changes in NOx Trap Materials under Lean and Rich 

Conditions at High Temperatures 

C. K. Narula, M. J. Moses, and L. F. Allard 

Oak Ridge National Laboratory 

P.O. Box 2008, MS 6116 

Oak Ridge, TN 37831-6116 

(865)574-8445; e-mail: narulack@ornl.gov 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Objective 

• Facilitate the development of nitrogen oxides (NOx) trap catalysts for lean gasoline or diesel exhaust treatment 

by 

— Investigating materials issues related to deterioration of NOx traps performance upon aging as a result of 

thermal and sulfation-desulfation cycles. 

— Designing new lean NOx trap (LNT) materials that are robust under the LNT operating conditions. 

Approach 

• Study microstructural changes in model LNT catalyst systems under normal operating conditions and 

accelerated aging using simulated diesel exhaust in a reactor system to understand the causes of failure. 

• Correlate microstructural changes with performance by evaluating model LNT catalyst systems on a bench-top 

reactor 

• Design new LNT catalysts that are robust under operating and aging conditions. Rapidly screen new catalysts 

for microstructural changes using an ex-situ reactor system and detailed evaluation on a bench-top reactor. 

Accomplishments 

• Tested the updated ex-situ reactor with simulated stoichiometric, lean, rich, and lean/rich cycling exhaust 

capability to treat transmission electron microscopy (TEM) samples by reproducing the results of the old 

reactor. 

• Completed a study of microstructural changes in the model LNT catalyst, 2%Pt-98%[10%CeO2-ZrO2­

90%(2%La2O3-98%BaO•6Al2O3)], and 2%Pt/γ-Al2O3 using simulated diesel exhaust (without sulfur oxides) 

under the following conditions: 

⎯ 60-s/5-s lean/rich cycling at 500°C/4 h for 2%Pt-98%[10%CeO2-ZrO2-90%(2%La2O3-98%BaO•6Al2O3) 

⎯ 240-s/60-s lean/rich cycling at 500°C/4 h for the model catalysts and 2%Pt/γ-Al2O3 

⎯ Accelerated aging at 700°C for 16 h using 60-s/5-s lean/rich cycling and 240-s/60-s lean/rich cycling for 

2%Pt-98%[10%CeO2-ZrO2-90%(2%La2O3-98%BaO•6Al2O3) 

39 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

• Carried out a preliminary study of microstructural changes under accelerated aging conditions (700°C, 16 h) 

using simulated exhaust (with SO2) at 240-s/60-s lean/rich cycling for 2%Pt-98%[10%CeO2-ZrO2­

90%(2%La2O3-98%BaO•6Al2O3). 

• Neared completion of the assembly of the bench-top flow reactor and installation of software. We will initiate 

experiments to correlate microstructural changes in LNT materials to LNT performance. 

Future Direction 

• Initiate correlation of microstructural changes with LNT performance by testing model catalysts on a bench-top 

reactor for performance. 

• Guided by the results of correlation of microstructural changes with performance, design new LNT catalyst 

systems and carry out systematic evaluation. 

Introduction 

The introduction of diesel-engine–based heavyduty trucks and passenger vehicles depends on the 

successful development of a strategy to treat NOx 

emissions. There presently is no catalyst or combination of catalysts that can convert NOx into inert 

gases under oxidizing conditions over a complete 

range of exhaust temperatures. Among NOx treatment strategies, LNTs are the most likely candidates 

for early deployment because they are consumertransparent (no action is needed on the part of 

consumers) and they can be system-integrated into 

current vehicle control strategies.1 

NOx traps collect engine-out NOx during lean 

operation and treat it during short rich operation 

cycles.2

] The NOx traps can be considered to be 

derived from commercial three-way catalysts 

(TWCs) installed to treat stoichiometric emissions 

from engines operating at air-fuel ratios of ~14. 

Therefore, the basic components of NOx traps are 

identical to TWCs. The advanced version of a TWC 

is a two-layer system on a honeycomb substrate with 

the inner layer based on platinum-alumina and the 

outer layer on rhodium-ceria-zirconia. The NOx 

traps derived from advanced TWCs are identical to 

these with the exception of a high baria content (the 

upper limit being close to 20%) in the alumina layer. 

Fresh NOx traps work very well, but they cannot 

sustain their high efficiency over the lifetimes of 

vehicles. The gradual and persistent deterioration in 

the performance of commercial TWCs is quite well 

known.3 The performance deterioration in NOx traps 

is believed to be caused by aging due to high-temperature operation and sulfation-desulfation cycles 

necessitated by the sulfur oxides (SOx) in the emissions from the oxidation of sulfur in fuel. The 

formation of barium aluminate is also considered to 

be a cause of performance deterioration since barium 

aluminate forms on thermal aging and is an inefficient NOx adsorber. Until last year, the aging-related 

microstructural changes in precious metals components were not available in open literature. In our 

previous reports, we summarized our studies of fresh 

and aged supplier samples that clearly show changes 

that occur on aging under various operating conditions. The prominent changes are sintering and 

migration of precious metals and migration of 

barium, leading to reduced precious metal-adsorber 

surface area available for NOx adsorption in lean 

cycles. 

In order to design a thermally durable NOx trap, 

there is a need to understand the changes in the 

microstructures of materials that occur during various modes of operation (lean, rich, and lean/rich 

cycles). This information can form the basis for 

selection and design of new NOx trap materials that 

can resist deterioration under normal operation. The 

first goal of the project is to determine if one or all 

of the microstructural changes take place during 

lean, rich, or lean/rich cycles. 

The tasks to achieve this goal are as follows: 

• Complete microstructural characterization of 

fresh and thermally aged NOx trap materials to 

determine the species formed as a result of 

aging. 

• Complete microstructural characterization of 

fresh NOx trap materials after exposure to 

simulated lean, rich, and lean/rich cycling conditions with and without SOx to determine the 

species formed during lean cycles. 

• Complete microstructural characterization of 

fresh NOx trap materials after exposure to 

40 
Heavy Vehicle Propulsion Materials 

accelerated lean/rich aging conditions to determine the species formed as a result of elevated 

temperatures. 

• Correlate microstructural changes with the performance of LNTs. 

The second goal of the project is to investigate 

and design new catalyst materials that can withstand 

NOx trap operating conditions without undergoing 

detrimental structural changes. The results from the 

first goal will provide insights into changes that 

occur in NOx trap materials at a microstructural level 

upon extended exposure to NOx trap operating conditions, enabling selection and design of materials 

for the second goal. 

Approach 

Our approach involves a detailed study of 

microstructural changes in model LNT catalyst 

systems using an updated ex-situ reactor under 

normal operating conditions, and accelerated aging 

using simulated diesel exhaust in a reactor system to 

understand the causes of failure and the deterioration 

mechanism. 

The next step in this process is establishment of 

a correlation between microstructural changes and 

the performance of model LNT catalyst systems 

using a bench-top reactor. 

Finally, we plan to design new LNT catalysts 

that are robust under operating and aging conditions. 

In this protocol, we will use microstructural changes 

as a rapid screening tool for new catalysts and will 

carry out detailed evaluation using a bench-top 

reactor. 

Results 

Our previous studies of the microstructural 

changes that occurred in a supplier LNT system 

(based on Pt/BaO–Al2O3 and CeO2–ZrO2 materials) 

upon aging on (1) a pulsator at Ford Motor 

Company (2) a dynamometer at Ford, and 

(3) vehicles in gasoline direct-injection sparkignition engines in Europe show the following 

results. 

After pulsator aging, lean and rich aged samples 

showed that the sintering of platinum particles 

occurs during aging and barium migrates into the 

ceria-zirconia layer. Both of these factors reduce the 

platinum-barium oxide surface area where NOx 

adsorption and reduction takes place during lean and 

FY 2005 Progress Report 

rich cycles, respectively. The stoichiometric aging 

also leads to the migration of barium into the ceriazirocnia layer, but the sintering of platinum is less 

severe. 

The dyno-aged samples showed extensive 

sintering of platinum and its migration in the ceriazirconia layer. The sintering of rhodium was also 

observed. The migration of barium into the ceriazirconia and the sintering of the precious metal 

component could explain the deterioration of 

performance. 

The analysis of samples evaluated on a vehicle 

after 32,000 km and 80 km showed that the bulk of 

the sintering of precious metals occurred in the 

stages of on-vehicle aging. 

These results primarily show precious metal 

sintering and barium migration to be causes of performance deterioration in NOx trap materials in 

gasoline engine exhaust conditions. 

While the laboratory aging protocols for diesel 

engines were in development, we compiled data 

showing early-stage changes in model NOx trap 

materials under diesel conditions. We also reported 

the synthesis and characterization of three catalyst 

systems, i.e. (Pt/Al2O3), (2%Pt-98%[10%CeO2­

ZrO2-90%(2%La2O3-98%BaO.6Al2O3)]), and 

(2%Pt, 5%MnO2-93%[10%CeO2-ZrO2­

90%(2%La2O3-98% BaO.6Al2O3)]). Now that the 

laboratory aging protocols are in public domain, we 

have carried out a preliminary study of microstructural changes in LNT catalyst materials under the 

accelerated aging conditions. The results are summarized in the following paragraphs. 

Microstructural Changes in Model Catalysts 

The TEM samples of a model LNT catalyst, 

2%Pt-98%[10%CeO2-ZrO2-90%(2%La2O3­

98%BaO•6Al2O3)] and a standard 2%Pt/γ-Al2O3 

were treated under lean/rich cycling at operating 

(500oC) and accelerated aging (700oC) conditions 

using a simulated diesel exhaust on the updated exsitu reactor (Figure 1). The microstructural changes 

in the samples were analyzed by scanning TEM 

(STEM). The simulated diesel exhaust contained 

CO, CO2, H2, C3H6, NO, H2O, and O2 with N2 as a 

balance gas. The SO2 was added to the simulated 

exhaust stream in some accelerated aging studies. 

All tests were run at a gas flow of 100 cc/min. A 

summary of the results is presented in Tables 1 

and 2. 

41 
___ 

___ 

___ ___ 

FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 1. Ex-situ reactor set-up. 

Table 1. Platinum particle size change under lean/rich cycle treatment at 500°C/4 h using 

simulated exhaust without SO2 

Catalyst Fresh sample Lean/rich diesel aging 

60 s/5 s 240 s/60 s 

2%Pt/γ-Al2O3 0.9–2.3 nm 0.7–2.7 nm 

2%Pt-98%[10%CeO2-ZrO2-90%(2%La2O3-98% 

BaO•6Al2O3)] 0.7–2.2 nm 0.7–2.8 nm 0.7–2.1 nm 

Table 2. Platinum particle size change under lean/rich cycle treatments at 700°C using 

simulated exhaust 

Lean/rich diesel exhaust 

Catalyst Aging time 60 s/5 s 240 s/60 s 

w/o SO2 w/o SO2 w/ SO2 

Fresh 0.6–2.9 nm 0.9–2.5 nm 1.0–2.1 nm 

4h 0.5–5.0 nm 0.6–4.4 nm 1.0–3.7 nm 

2%Pt-98%[10%CeO2-ZrO2­

90%(2%La2O3-98% 8h 0.6–5.4 nm 0.8–5.3 nm 0.8–4.6 nm 

BaO•6Al2O3)] 

12h 0.6–5.9 nm 0.7–4.7 nm 

16h 0.7–4.7 nm 

42 
Heavy Vehicle Propulsion Materials 

Operating conditions—lean/rich cycling at 

500°C on ex-situ reactor: After 4 h of treatment in 

a simulated exhaust flow without SO2, no platinum 

sintering was observed regardless of lean/rich pulse 

time (Figure 2). 

• 60-s/5-s lean/rich cycling: 2%Pt-98% 

[10%CeO2-ZrO2-90%(2%La2O3­

98%BaO•6Al2O3)]: STEM of the fresh sample 

showed platinum particles in the 0.7–2.2 nm 

range. The electron microscopy images show no 

significant growth in particle size, 0.7–2.7 nm, 

after aging. 

• 240-s/60-s lean/rich cycling: 2%Pt-98% 

[10%CeO2-ZrO2-90%(2%La2O3­

98%BaO•6Al2O3)]: STEM images of the fresh 

sample showed platinum particles in the 

0.7–2.2 nm range. The particles showed no 

increase in particle size range after aging. STEM 

images of the fresh 2%Pt/γ-Al2O3 showed 

platinum particles in the 0.9–2.3 nm range, and 

no growth in particle size was seen after 4 h. 

Accelerated aging—lean/rich cycling at 

700°C: The accelerated aging of a model catalyst, 

2%Pt-98%[10%CeO2-ZrO2-90%(2%La2O3-

FY 2005 Progress Report 

98%BaO•6Al2O3)], was carried out using simulated 

exhaust without SO2 under 60-s lean/5-s rich pulses 

or 240-s/60-s rich pulses, and no significant difference was noted. Interestingly, platinum sintering 

was evident after only 4 h of sample treatment under 

accelerated aging conditions. A gradual growth in 

platinum particles in the 2.0–5.0 nm range was seen 

after each additional 4 h of treatment. Addition of 

SO2 to the simulated exhaust did not lead to any 

additional noticeable platinum sintering. Figures 3 

and 4 depict the results of the accelerated aging 

study at 240 s/60 s lean/rich cycling with and without SO2, respectively. 

60-s/5-s lean/rich accelerated aging treatment 

using simulated exhaust without SO2: 

• STEM images of the fresh sample showed platinum particles in the 0.6–2.9 nm range. 

• After 4 h of aging, sintering of the particles 

could be seen in the STEM images by the 

appearance of particles in the 2.0–5.0 nm range. 

• After each of the remaining four increments, 

STEM images showed further platinum sintering 

with a gradual increase in the number of 

2.0–5.0 nm particles. 

Figure 2. (a) Particle size distribution of multiple areas of a model LNT catalyst, 2%Pt-98% 

[10%CeO2-ZrO2-90%(2%La2O3-98%BaO•6Al2O3)], after lean/rich cycling and 

60-s/5-s pulse using simulated exhaust without SO2. (b) A STEM image of one area 

recorded in (a). (c) Particle size distribution of multiple areas of a model LNT catalyst 

after lean/rich aging at a cycling time of 240 s/60 s. and (d) A STEM image of one area 

recorded in (c). 

43 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 3. These graphs show platinum particle size distribution of the model LNT 

catalyst, 2%Pt-98%[10%CeO2-ZrO2-90%(2%La2O3-98%BaO•6Al2O3)], 

fresh and after accelerated aging at 700°C in 4-h increments for 16 h and 

240-s/60-s lean/rich pulse using simulated diesel exhaust without SO2. 

The same four areas of the sample were used for each distribution. 

44 
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 4. These graphs show platinum particle size distribution of the model LNT 

catalyst, 2%Pt-98%[10%CeO2-ZrO2-90%(2%La2O3-98%BaO•6Al2O3)], fresh 

and after accelerated aging at 700°C in 4-h increments for 8 h and 240-s/60-s 

lean/rich pulse using simulated diesel exhaust with SO2. The same three areas of 

the sample were used for each distribution. 

240-s/60-s lean/rich accelerated aging treatment using simulated exhaust without SO2: 

STEM images of the fresh sample showed platinum 

particles in the 0.6–2.5 nm range. 

• After 4 h of aging, sintering of the particles 

could be seen in the STEM images by the 

appearance of particles in the 2.0–5.0 nm range. 

• After each of the remaining 4-h increments, 

STEM images showed further platinum sintering 

with a gradual increase in the number of 2.0–5.0 

nm particles. 

240-s/60-s lean/rich accelerated aging treatment using simulated exhaust with SO2: 

• STEM images of the fresh sample showed platinum particles in the 1.0–2.1 nm range. 

• After 4 h of aging, sintering of the particles 

could be seen in the STEM images by the 

appearance of particles in the 2.0–4.0 nm range. 

• After another 4 h of aging, STEM images 

showed further platinum sintering with a gradual 

increase in the number of 2.0–5.0 nm particles. 

Modifications to Ex-situ Reactor 

Final modifications to the ex-situ reactor were 

completed this year to include the ability to cycle 

through lean and rich conditions. This involved the 

use of a bench-top flow furnace and the installation 

of additional gas-flow controllers and solenoid 

valves that could be controlled through a computer 

using LabView software. A picture of the modified 

ex-situ reactor setup is shown in Figure 1. 

Assembly of Bench-Top Flow Reactor 

The assembly of the bench-top flow reactor and 

installation of software is nearing completion. We 

will initiate experiments to correlate microstructural 

changes in LNT materials to LNT performance. 

Conclusions 

Our results for microstructural changes clearly 

show that LNT materials are relatively stable under 

normal operating conditions of diesel engines. Practically no sintering of precious metals is observed 

that could cause performance deterioration. However, our accelerated aging studies demonstrate that 

45 
FY 2005 Progress Report 

rapid sintering of precious metals can take place, 

leading to performance deterioration. Interestingly, 

SOx appears to place an indirect role in the performance deterioration of LNTs. Since there is no significant difference in precious metal sintering during 

accelerated aging using simulated exhaust with or 

without sulfur, it would appear that SOx plays no 

role in precious metal sintering, but hightemperature treatment is required only for desulfation. Thus if there were no SOx in diesel exhaust, 

there would be no need for high-temperature desulfation and the LNT would be quite durable. 

Our next goals are to finish accelerated aging 

tests in the presence of SO2 and to start evaluating 

new model catalyst systems for their performance 

and correlate them to microstructural changes. 

References 

1. N. Miyoshi, S. Matsumoto, K. Katoh, 

T. Tanaka, J. Harada, N. Takahashi, K. Yokota, 

M. Sigiura, and K. Kasahra, “Development of New 

Concept Three-Way Catalyst for Automotive LeanBurn Engines,” paper 95-0809, Society of Automotive Engineers, Warrendale, Pennsylvania, 

February 11, 1995. 

2. N. Takahashi, H. Shinjoh, T. Iijima, 

T. Suzuki, K. Yamazaki, K. Yokota, H. Suzuki, 

N. Miyoshi, S. Matsumoto, T. Tanizawa, S. Tateishi, 

and K. Kasahara, “The New Concept 3-Way 

Heavy Vehicle Propulsion Materials 

Catalyst for Automotive Lean-Burn Engine: NOx 

Storage and Reduction Catalyst,” Catalysis Today, 

27, 63–69 (1996). 

3. H. S. Gandhi, G. W. Graham, and R. W. 

McCabe, “Automotive Exhaust Catalysis,” J. Cat., 

216, 433–442 (2003). 

Presentations and Publications 

C. K. Narula, W. A. Shelton, L. F. Allard, 

Y. Xu, M. Moses, W. Schneider, and B. Gates, 

“Combining Theory and Experiments in Studies of 

Structural Changes in LNT Materials,” presented at 

the Eighth DOE Crosscut Workshop on Lean 

Emissions Reduction Simulation, CLEERS 2005, 

University of Michigan–Dearborn, Dearborn, 

Michigan, May 17–19, 2005. 

C. K. Narula, L. F. Allard, M. Moses, Y. Xu, 

W. Shelton, W. F. Schneider, G. Graham, and 

J. Hoard, “Lean NOx Traps—Microstructural 

Studies of Real World and Model Catalysts,” DEER 

2005, Chicago, August 21–25, 2005. 

C. K. Narula, C. S. Daw, J. W. Hoard, and 

T. Hammer, “Materials Issues Related to Catalysts 

for Treatment of Diesel Exhaust,” Int. J. Appl. 

Ceram. Tech., 2, 452 (2005). 

C. K. Narula, M. Moses, and L. F. Allard, 

“Ex-situ Reactor Studies of Microstructural Changes 

in NOx Traps under Operating Conditions,” J. Phys. 

Chem. (in preparation). 

46 
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

F. Aftertreatment Catalysts Materials Research 

Matthew Henrichsen and Randy Stafford 

Cummins, Inc. 

1900 McKinley Ave. 

Columbus, IN 47203 

(812) 377-3279; e-mail: randy.j.stafford@cummins.com 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Prime Contract No: DE-AC05-00OR22725 

Subcontractor: Cummins, Inc., Columbus, Indiana 

Objective 

• Develop knowledge and tools necessary to satisfy the U.S. Environmental Protection Agency (EPA) 2007 and 

2010 diesel emissions regulations while minimizing the total cost of ownership. 

Approach 

• Develop reactors and test methods for understanding the fundamental behavior of catalysts. 

• Evaluate novel base metal oxide catalysts for oxidation of diesel soot with conventional and microwave 

heating. 

• Improve understanding of sulfation and desulfation of commercial NOx reduction catalysts through reactor 

studies. 

Accomplishment 

• Completed program final report. 

Future Direction 

• Continue sulfation and desulfation studies with private funding. 

• Continue the search for low-temperature soot oxidation catalysts. 

Introduction 

The 2007 and 2010 EPA regulations on particulate and NOx emissions from diesel engines are significantly lower than the current regulatory standards. These regulations are seen as aftertreatmentforcing by the diesel industry. 

NOx traps (also referred to as NOx storage and 

release catalysts) are a primary path toward compliance with the 2007 and 2010 EPA regulations. 

However, the poisoning effects of sulfur reduce the 

efficiency of NOx traps to capture and destroy NOx 

emitted from diesel engines. Even though sulfur 

levels will be reduced in future diesel fuels, any 

amount of sulfur in fuel will be detrimental because 

the effect of sulfur on NOx traps is cumulative. NOx 

traps can be desulfated by high-temperature reduction. Desulfation, however, can require extended 

time at high temperature. If the desulfation process 

47 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

requires 10 min for completion for light-duty 

applications, then desulfation will be impractical 

because the average time between engine start and 

stop for a light-duty vehicle is less than 10 min. 

Catalysts commonly used for soot oxidation in 

catalyzed soot filters are based on platinum group 

metals and cerium oxides. The use of platinum 

group metals significantly increases the cost of 

catalyst systems. Base metal oxide catalysts, while 

typically less active than platinum group metal 

catalysts, are much less expensive. There have been 

reports in the literature that microwave heating 

accelerates soot oxidation in the presence of some 

base metal oxide catalyst more than conventionally 

heating to a similar temperature accelerates oxidation. If true, this could allow very low temperature 

regeneration of soot filters, significantly reducing 

the stress on the system and the cost of regeneration. 

Approach 

No experimental work was conducted in 

FY 2005. The program final report was written and 

submitted to close the program. 

Results 

The program final report was submitted in 

February 2005 to close this program. 

Summary and Conclusions 

A test method was developed that allows a relatively inexpensive mass spectrometer to differentiate 

between molecules of similar mass and different 

chemical behavior. Reactors were developed to 

allow comparison of chemical kinetics with microwave and conventional heating. A reactor for artificial aging of multiple catalyst cores was designed. 

Catalyzed soot oxidation with conventional and 

microwave heating was explored. Data from direct 

comparisons provided no reasonable basis to think 

that microwave heating significantly improves 

catalytic soot oxidation over conventional heating at 

a similar temperature. Though very low soot oxidation rates were found with some base metal oxide 

catalysts, these catalysts are very hygroscopic and 

unlikely to be successful in practical application. 

One commercial NOx adsorber catalyst was 

evaluated in reactor studies. While other commercial 

catalysts are expected to have similar qualitative 

behavior, temperatures required for desulfation will 

be higher. The temperature at which desulfation 

rates become reasonable is fairly sharply defined. 

While desulfation at 500°C removed nearly 80% of 

the sulfur, desulfation at 475°C failed to remove 

even 20%. As expected, much of the sulfur removal 

occurs during the early portion of a desulfation 

treatment, while a full performance recovery is 

dependent on removing most of the sulfur. 

48
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

G. Catalysis by First Principles 

C. K. Narula, L. F. Allard, V. Bhirud, B. Gates, M. Moses, W. Schneider, W. Shelton, Y. Xu 

Oak Ridge National Laboratory 

P.O. Box 2008, MS 6116 

Oak Ridge, TN 37831-6116 

(865) 574-8445, fax: (865) 576-6298; e-mail: narulack@ornl.govl 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Prime Contract No: DE-AC05-00OR22725 

Objective 

Search for durable emission catalysts [lean Nox (nitrogen oxides) trap (LNT), three-way catalyst (TWC), OC] from 

a protocol based on 

• an integrated approach between computational modeling and experimental development 

• design and testing of new catalyst materials to rapidly identify the key physiochemical parameters necessary for 

improving the catalytic efficiency of these materials. 

Approach 

• Theoretical modeling 

— Density functional theory-based first-principles calculations 

— Optimization of platinum (Pt) clusters supported on alumina 

— Interaction of CO, NOx, and hydrocarbons (HCs) with Pt clusters supported on alumina 

• Experimental system 

— Synthesize Pt and rhenium (Re) nanoclusters on morphologically diverse alumina supports 

— Evaluate systems for CO, NOx, and HC oxidation 

Accomplishments 

• Completed the synthesis of a series of Pt and Re clusters supported on commercial alumina. 

— The microstructural studies of freshly prepared Pt/alumina (by impregnation) show it to contain 

10–20 atom clusters. Thermal treatment leads to larger particles. 

— The microstructural studies of freshly prepared Re/alumina (by impregnation) show it to contain 

10–20 atom clusters. We prepared smaller clusters of 1–3 atoms from the decarbonylation of 

H3Re3(CO)12/γ−Al2O3. Thermal treatment leads to larger particles. 

• Synthesized alumina sol-gel and molecular sieve support materials and began efforts to increase the thermal 

stability of the molecular sieves. 

• Carried out modeling of isolated Pt clusters. Theory tells us that the following are true for Pt nanoclusters: 

— Reactivity (adsorption) shows strong size dependence. 

49 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

— Smaller clusters are much more prone to oxidation. 

— Oxidized clusters adopt 1-dimensional or 2-dimensional structures. 

— In an O2 atmosphere, higher T or lower pO2 leads to lower extent of oxidation and makes the reduction of 

oxidized Pt clusters more facile Æ Pt oxides becoming better oxidizing agent. 

— Lower oxide phases are more prevalent for larger clusters. 

• Equipped the synthesis laboratory to enable preparation of catalyst materials under inert conditions. A benchtop 

catalyst testing reactor has been installed (funded by a project from DOE Office of Industrial Technologies) and 

is being tested before catalyst evaluation experiments start. 

Future Direction 

• Understand the structures of nanoclusters on support. Our results show that nano-clusters are 10–15 atoms. 

— Develop theoretical models to understand cluster oxidation state (oxidized, reduced, in equilibrium), 

dependence on cluster size, and the kinetics of oxidation. 

— Conduct experimental studies on the structures of these systems using X-ray absorption fine structure 

(EXAFS) and X-ray absorption near-edge structure (XANES). 

• Support interaction with clusters and its impact on the structure and reactivity: 

— cationic or zero-valent metals or both 

— cationic metals at the metal–support interface 

• Analyze the reactivity of the clusters 

— Conduct theoretical and experimental studies on structure-reactivity correlation toward CO, NOx, and HC 

oxidation studies 

— Determine the impact of cluster size on catalytic activity 

• Translate these results into the design of durable supported catalysts for LNTs (and other systems such as TWC, 

OC for diesel) 

Introduction 

This research focuses on an integrated approach 

between computational modeling and experimental 

development, design, and testing of new catalyst 

materials, which we believe will rapidly identify the 

key physiochemical parameters necessary for 

improving the catalytic efficiency of these materials. 

The typical solid catalyst consists of nanoparticles on porous supports. The development of new 

catalytic materials is still dominated by trial-anderror methods, even though the experimental and 

theoretical bases for their characterization have 

improved dramatically in recent years. Although it 

has been successful, the empirical development of 

catalytic materials is time-consuming and expensive 

and brings no guarantee of success. Part of the difficulty is that most catalytic materials are highly nonuniform and complex, and most characterization 

methods provide only average structural data. Now, 

with improved capabilities for synthesis of nearly 

uniform catalysts—which offer the prospects of high 

selectivity as well as susceptibility to incisive characterization combined with state-of-the science 

characterization methods, including those that allow 

imaging of individual catalytic sites—we have a 

compelling opportunity to markedly accelerate the 

advancement of the science and technology of 

catalysis. 

Computational approaches, on the other hand, 

have been limited to examining processes and phenomena using models that had been much simplified 

compared with real materials. This limitation was 

mainly a consequence of limitations in computer 

hardware and in the development of sophisticated 

algorithms that are computationally efficient. In 

particular, experimental catalysis has not benefited 

from the recent advances in high-performance computing that enable more realistic simulations 

(empirical and first-principles) of large ensemble 

atoms, including the local environment of a catalyst 

site in heterogeneous catalysis. These types of 

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Heavy Vehicle Propulsion Materials 

simulations, when combined with incisive microscopic and spectroscopic characterization of catalysts, can lead to a much deeper understanding of the 

reaction chemistry that is difficult to decipher from 

experimental work alone. 

Thus a protocol to systematically find the optimum catalyst can be developed that combines the 

power of theory and experiment for atomistic design 

of catalytically active sites and can translate the 

fundamental insights gained directly to a complete 

catalyst system that can be technically deployed 

(Figure 1). 

Although it is no doubt computationally challenging, the study of surface, nanometer-sized, metal 

clusters may be accomplished by merging state-ofthe-art, density-functional-based electronic-structure 

techniques and molecular-dynamic techniques. 

These techniques provide accurate energetics, force, 

and electronic information. Theoretical work must 

be based on electronic-structure methods, as 

opposed to more empirical-based techniques, so as 

to provide realistic energetics and direct electronic 

information. 

A computationally complex system, in principle, 

will be a model of a simple catalyst that can be 

synthesized and evaluated in the laboratory. It is 

important to point out that such a system for an 

experimentalist will be an idealized simple model 

catalyst system that will probably model a “realworld” catalyst. 

Thus it is conceivable that a “computationally 

complex but experimentally simple” system can be 

examined by both theoretical models and experimental work to forecast improvements to obtain 

optimum catalyst systems. 

Approach 

The theoretical modeling is based on densityfunctional theory (DFT) studies of Pt clusters to 

understand the relationship between cluster size, 

structure, composition, and reactivity. This, coupled 

with first-principles thermodynamics, provides 

insights into the effects of oxidizing atmosphere 

Pt Pt/Al Pt Pt/A 2O3 l

2O3 TWC, LNT

FY 2005 Progress Report 

(O2)—finite (T, pO2), structure, composition, and 

redox potential on particle size. The results provide 

guidance for investigations of larger/supported 

clusters. The next steps of this work will model 

supported clusters, based on input from experimental results, and their interactions with CO, NOx, and 

HC. 

Experimentally, we have synthesized a series of 

Pt and Re clusters and particles supported on morphologically different γ–aluminas. The microstructural characterization reveals structural differences 

in 1–3 atom and 10–20 atom clusters. We will 

evaluate these catalysts for their CO, NOx, and HC 

oxidation capability. Guided by the theoretical 

models and experimental results, we will synthesize 

a new set of catalyst materials with higher durability 

under operating conditions. 

Results 

Over the last year, theoretical studies have been 

carried out on the oxidation behavior of gas-phase Pt 

nanoparticles and metallic clusters. On the experimental side, a range of γ–Al2O3 supported Pt and Re 

clusters have been synthesized and characterized. 

Alumina sol-gels and molecular sieves have also 

been synthesized to use as supports. The results are 

summarized in the following subsections. 

Alumina Support Materials 

We synthesized sol-gel alumina by the reported 

procedure from the hydrolysis of Al(sec-C4H9O)3, 

nitric acid peptization, and gelation. The material 

gives an amorphous X-ray diffraction between 5 and 

70° 2θ. The alumina sol surface area calculated from 

a multipoint BET (Brunauer, Emmett, and Teller) 

measurement without degassing is 170.82 m2/g. 

Scanning electron microscopy (SEM) images show 

some macro-pores on the surface, but no ordered 

structure can be seen (Figure 2). 

For the synthesis of alumina molecular sieves, 

we employed a structure-templated method using 

CombinatorialDurable 

Candidate 

Materials 

Candidate

Materials

Catalyst

Screening Materials creening Materials

Theoretical ModelsExperimental Models

Figure 1. Theory and experiment, combined with simulation, can translate to a deployable catalyst system. 

51 

x 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 2. SEM images of sol-gel alumina. 

cetyltrimethylammonium bromide, reported by values are about half of the reported BET surface 

Deng et al.1 X-ray diffraction (XRD) of the alumina areas (~470 m2/g). Efforts are in progress to determolecular sieves and SEM images (Figure 3) con- mine an accurate BET surface area before underfirm the reproducibility of this synthetic route. Since taking further modifications to the synthesis 

the rate of stirring has a significant impact on the procedure. 

morphological structure of alumina molecular Since the alumina molecular sieve structure 

sieves, we carried out the synthesis using a magnetic collapses at elevated temperatures, we are exploring 

stirrer and used stirring rates of 64 and 128 rpm. The synthetic routes to incorporate stabilizers into the 

BET surface areas of 128-rpm and 64-rpm samples alumina molecular sieve structure. Employing the 

were 208 m2/g and 234 m2/g, respectively. These Pinnavaia method, we added lanthanum (La) 

Figure 3. SEM images of the Al2O3 molecular sieves prepared by stirring at 64 rpm (a) shows the 

surface openings of the macropores (channel openings) and the channels that extend 

through the material. (b) is a close-up of the channel opening at an angle. (c) is closeup of macropore channels and some of the micropores within the channels can be 

discerned (d) shows an enlargement of the micropores within the channels from (c). 

52
Heavy Vehicle Propulsion Materials 

nitrate during the synthesis of alumina molecular 

sieves. This method did not lead to improved thermal durability of alumina molecular sieves. The 

powder XRD (Figure 4) of 2 wt % La-doped 

alumina molecular sieves, after calcining at 500°C/ 

5.5 h in flowing air, exhibits no diffraction peaks, 

suggesting that the materials are amorphous 

materials. Further calcining at 700°C/5.5h and 

FY 2005 Progress Report 

900°C/5.5 h led to low-intensity broad peaks in 

powder XRD that were assigned to aluminum oxide. 

The SEM of La-impregnated samples resembled that 

of pure alumina molecular sieves after initial 

calcining at 500°C/5.5h (Figure 5). The impact of 

calcining La-impregnated alumina molecular sieves 

for 5.5 h at 700°C after the initial 500°C/ 5.5 h can 

be seen in the SEM (Figure 6). The structures 

Figure 4. X-ray diffraction patterns of Al2O3 molecular sieves synthesized at 64 rpm, doped with 2 wt % La and 

calcined under flowing air at (a) 500°C/5.5 h, followed by annealing at (b) 700°C/5.5 h, and finished with a 

final annealing at (c) 900°C/5.5 h. The diffraction patterns on the left are without background subtraction and 

those on the right are with background subtraction. Diffraction lines of aluminum oxide are of 

JCPDF # 75-0921. 

Figure 5. SEM images of the Al2O3 molecular sieves impregnated with 2% La and calcined at 

500°C/5.5h. The images show the surface openings of the macropores (channel 

openings) and the channels that extend through the material. 

53 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

of many porous channels were maintained; however, electron microscope (ACEM) in the HAADF-STEM 

the coverage and uniformity of the channels have 

diminished. This suggests that La did not get 

incorporated in the molecular sieve structure. As a 

next step, we will explore our method of employing 

heterometallic alkoxides as precursors for the 

incorporation of La in alumina molecular sieves. 

Platinum-Alumina System 

While synthesis and characterization of sol-gel 

and molecular sieve alumina are in progress, we 

have synthesized and characterized Pt/Al2O3 materials using commercial alumina. We carried out the 

synthesis of nanocluster Pt/commercial γ–alumina 

by impregnation of γ–alumina with H2PtCl6•6H2O 

followed by calcination in air. The scanning 

transmission electron microscopy (STEM) of 2% 

Pt/γ–Al2O3 (Figure 7) shows approximately 1-nm 

particles with a narrow distribution centered in the 

1–1.5-nm range (Figure 7B). High-magnification 

images obtained using the aberration-corrected 

mode indicate particles of 10–20 atoms and also 

show some smaller platinum clusters (Figure 7c 

and d). 

Thermal treatment of this material leads to 

~10-nm or larger Pt clusters supported on γ–alumina 

with the Pt particle size depending on the temperature and duration of thermal treatment. In order to 

prepare smaller clusters, a possible approach is to 

synthesize platinum carbonyl clusters on alumina. A 

variety of Pt carbonyl clusters {[Pt15(CO)30]2−, 

[Pt6(CO)12]2−, [Pt9(CO)18]2−} on MgO are known.2 

A synthesis of such clusters on a variety of aluminas 

(commercial γ–alumina, sol-gel alumina2 and 

molecular sieve alumina1 will be challenging, but a 

simple decarbonylation will furnish Pt clusters of 

fewer than 10 atoms. 

The theoretical efforts, which are based on 

DFT, have so far been focused on the effect of 

cluster size and the impact of oxidation on the 

properties—including structure, composition, and 

Figure 6. SEM images of the Al2O3 molecular sieves impregnated with 2% La and calcined first 

at 500°C/5.5h followed by further calcination at 700°C/5.5h (a) shows the surface 

openings of the macropores (channel openings) (b) shows a piece of alumina with 

channels that extend through the material (c) shows some channels that do not appear 

to extend through the entire piece of alumina (d) show a piece of alumina with a 

honeycomb structure. 

54
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 7. (a) STEM images of the 2% Pt/γ-Al2O3 sample and (b) the particles size 

distribution. (c and d) ACEM HAADF-STEM images of 2% Pt/γ-Al2O3. 

reactivity—of isolated Pt clusters. We have 

concluded the investigation of Ptx nanoclusters 

(x = 1, 2, 3, 4, 5, and 10) using the DFT codes VASP 

and Dacapo. The main results are summarized in 

Figure 6. 

A salient feature is that the oxidation of the Pt 

clusters, either partial or complete, is more exothermic (more negative oxidation energy) than the oxidation of bulk Pt to the bulk oxides. The oxidation 

energies display strongly non-linear dependence on 

cluster size and generally become more exothermic 

as the size decreases. This means that smaller Pt 

clusters are much more prone to oxidation than 

larger Pt clusters or bulk Pt. Complete oxidation of a 

Ptx cluster to PtxO2x is favored energetically over 

partial oxidation to PtxOx, regardless of the cluster 

size. 

Upon oxidation, the structure of Pt clusters also 

changes substantially. Whereas the original Pt clusters are generally 3-dimensional, the oxidized clusters prefer flatter, 1- or 2-dimensional shapes. The 

PtxOx clusters form closed loops, in which the Pt 

atoms adopt the shapes of triangles, quadrangles, 

pentagons, etc., with each Pt atom coordinated to 

two O atoms. The PtxO2x series, on the other hand, 

form open, bent linear chains, in which the middle 

Pt atoms are each coordinated to four O atoms. The 

structure of the oxidized Pt clusters is also very 

different from the 3-dimensional matrices of the 

bulk oxides (Figure 8). 

Bader charge analysis of the oxidized Pt nanoclusters reveals another difference between the 

oxidized Pt nanoclusters and the corresponding bulk 

oxides: the Pt-O bonds in the clusters have considerably less ionic character than those in the bulk 

oxides. 

Both cluster size and oxidation affect the reactivity of the Pt clusters. We studied the adsorption 

energies of O, O2, CO, and CO2 on the Pt1–5 clusters 

and found considerable variation with respect to 

cluster size. In general, adsorption on smaller clusters is more exothermic. The adsorption energies of 

O, O2, and CO2 exhibit a minimum of around x = 4 

before approaching the level on an extended Pt 

55 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Pt3O3

Pt3

Pt3O4

Pt3O8 NPt=3

Pt3O6

Pt3O3

Pt3

Pt3O4

Pt3O8 NPt=3

Pt3O6

Figure 8. (Left) Oxidation energies of the PtxOx (filled circles) and PtxO2x (open circles) clusters. Black and grey 

dashed lines indicate bulk PtO (–0.54 eV) and β-PtO2 (–1.58 eV) formation energies, respectively. The stickand-ball representations of the pure Pt and Pt oxide clusters are included. Red spheres represent O atoms, and 

grey spheres represent Pt atoms. The stoichiometries of Pt:O = 1:1 and 1:2 represent, respectively, partially 

and fully oxidized Pt systems. (Right) A temperature-O2 pressure phase diagram of Pt clusters from firstprinciples thermodynamic model. 

surface. CO adsorption grows monotonically weaker 

with increasing cluster size. The adsorption energies 

of O and CO are found to be weaker on several 

PtxOx clusters than on the corresponding pure Ptx 

clusters, suggesting that oxidation probably weakens 

the adsorption of atoms and molecules in general. 

Since DFT results are formally obtained at 0 K 

and in vacuum, we developed a first-principles 

thermodynamic model in order to incorporate the 

effects of an oxygen atmosphere with finite temperature and pressure on the composition of the Pt 

clusters (Figure 8). A temperature–O2 pressure 

phase diagram has been constructed for each of the 

Pt1–3 clusters. The results agree with the general 

physical insight that higher temperature or lower 

pressure favors lower extent of oxidation. Under 

conditions of interest to this project—e.g., diesel 

engine exhaust, which is generally found in the 

400–600 K range and has an O2 partial pressure of 

~0.1 bar—the extent of oxidation is found to 

decrease from a dioxide to a monoxide 

stoichiometry with increasing cluster size. No submonoxides are seen in any of the three phase 

diagrams. Therefore, we expect larger clusters to 

exist primarily as monoxides under similar 

conditions. 

In summarizing the theoretical findings, we note 

that both size and oxidation profoundly affect the 

structure, composition, and reactivity of Pt clusters, 

at least in the size range that we have studied so far. 

In the second year of this project, we will tackle the 

effect of another variable on the properties of Pt 

clusters, namely the metal oxide support that is an 

integral part of many heterogeneous catalysts, 

including automotive applications (e.g., TWC and 

LNT). We will study the reactivity of larger clusters, 

perhaps with sizes in the range of 10–20 Pt atoms, 

which are more accessible to current laboratory 

techniques. Our results so far have provided many 

hints that will facilitate the modeling of larger clusters and the investigation of support effects. 

Theoretical study also suggests that the structure 

of experimentally observed 10–20 atom Pt-clusters 

is different from that of bulk metal and in all likelihood has some oxygen incorporated. Three-dimensional structure mapping and EXAFS studies are in 

progress to determine detailed cluster structure. 

Rhenium-Alumina System 

Since Pt-carbonyl cluster chemistry on alumina 

is not well developed, we chose Re-carbonyl clusters 

supported on alumina as surrogates for Pt-alumina 

clusters of fewer than 10 atoms. We initiated this 

work by successfully synthesizing H3Re3(CO)12 by 

using literature methods and confirmed its formation 

by infrared spectroscopy.

4 In the previous report, we 

described the decarbonylation and characterization 

of decarbonylated clusters by EXAFS and XANES 

spectroscopy. The EXAFS parameters indicated 

complete decarbonylation of the Re. The data indicate the formation of tri-nuclear Re rafts, Re3. The 

decarbonylated samples, modeled as tri-rhenium 

56
Heavy Vehicle Propulsion Materials 

rafts, are shown by the XANES data to incorporate 

electron-deficient (cationic) Re. On the basis of a 

comparison of the observed XANES signature with 

various XANES of various reference materials, we 

tentatively identify the oxidation state of Re to be 

+4. This statement is rough and in need of theoretical assessment. 

The H3Re3(CO)12/γ–Al2O3 sample was imaged 

on the ACEM in the HAADF-STEM mode. The 

intact supported tri-rhenium rafts indicated by infrared spectra, EXAFS, and XANES were confirmed 

by atomic-resolution STEM images (Figure 9 and 

inset). In addition to tri-rhenium rafts, atomic rhenium, dimers, and larger polyatomic clusters were 

also atomically resolved. STEM images indicate that 

although slightly larger clusters were present, the 

sample was predominately clusters comprising 

1–3 Re atoms. 

STEM images of H3Re3(CO)12/γ–Al2O3 treated 

in H2 at 400°C to decarbonylate the clusters were 

also obtained using the ACEM. Images at high magnification indicated uniform clusters with diameters 

of approximately 0.4–0.8 nm. The instability of the 

microscope at the time of use hindered efforts to 

capture images of atomic resolution. Rhenium has a 

radius of 1.37 Å, which suggests that the clusters are 

composed of 1–3 Re atoms. 

Decarbonylation of H3Re3(CO)12/γ–Al2O3 was 

also done by treatment in helium at 400°C. STEM 

FY 2005 Progress Report 

images show metallic particles in the 1.4–3.9 nm 

range (Figures 10 and 11). Some small Re clusters 

were also present, but because of the 3-dimensional 

structure of the metallic nanoparticles, imaging was 

difficult. 

Interestingly, Re/γ–Al2O3 synthesized by ReCl5 

impregnation on commercial γ–Al2O3 and subsequt 

calcining led to small (1–1.5 nm) clusters. 

Microstructural characterization on the ACEM 

shows two types of areas. In one type of area 

(Figure 12), 10–20-atom clusters of essentially 1 nm 

were observed that were reminiscent of our Pt/γ– 

Al2O3 materials (Figure 8). While these clusters are 

not completely random in the atomic arrangement, a 

3-dimensional metallic order is not observed. The 

second area (Figure 13) shows smaller, less uniform 

clusters and single-atom structures. Neither area 

shows significant structure change as a result of time 

under the beam, suggesting a stabilizing interaction 

between Re and the support. The inconsistent loading may be attributed to the high Re loading. 

Conclusions 

We have prepared and characterized Pt and Re 

clusters of different sizes and properties (metallic 

and non-metallic) on commercial γ–Al2O3 ranging 

from single atoms to 2-nm 3-dimensional nanoparticles. Two other alumina supports with different 

structural properties, sol-gel and molecular sieves, 

Figure 9. ACEM HAADF-STEM images of 1 wt % Re/γ-Al2O3. (a) Image of H3Re3(CO)12/γ-Al2O3; the 

inset at the bottom right is an enlargement of the tri-rhenium raft located in the circle. (b) An 

enlargement of the boxed in area of (a). Traces 1 and 2 correspond to the particle intensity. Graphs 

indicate one atom in trace of particle (1) and three atoms in trace of particle (2). (c) Front and side 

views of Re3 cluster, tethered to a ledge parallel to the beam direction. 

57 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 10. (a–b) show aberration corrected STEM-HAADF images of two 

different Re/γ-Al2O3 particles prepared by decarbonylating 

[H3Re3(CO)12]/γ-Al2O3 in He at 400°C, and (c) shows the Re 

particle size distribution. 

Figure 11. The high magnification ACEM images show the 3D metallic structure of 

the Re particles prepared by decarbonylating [H3Re3(CO)12]/γ-Al2O3 in He 

at 400°C. 

58
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 12. These ACEM images Re/γ-Al2O3 prepared by impregnation. The 

images were taken consecutively showing no movement of the 

clusters. 

Figure 13. This high magnification ACEM image shows Re atoms 

and small clusters of Re/γ-Al2O3 prepared by 

impregnation method. 

are being prepared and will soon be ready for Pt and supported on the three different alumina supports, 

Re loading. Computational calculations and simula- obtain and compare their activity as CO oxidation 

tion have shown the oxidation behavior of small Pt catalysts, and use the experimental data as the inputs 

clusters and their energetically favored structures. for the theoretical models. 

Our next goals are to finish structural 

characterization and preparation of the clusters 

59 
FY 2005 Progress Report 

Presentations and Publications 

C. K. Narula, presentation at the Eighth DOE 

Crosscut Workshop on Lean Emissions Reduction 

Simulation, Dearborn, Michigan, May 17–19, 2005. 

L. F. Allard, presentation at the annual meeting 

of the Microscopy Society of America. 

C. K. Narula, presentation at DEER 2005, 

Chicago, August 21–25, 2005. 

Bruce Gates and W. Schneider, presentations at 

the North American Catalysis Society meeting, 

Philadelphia, May 22–27, 2005. 

V. A. Bhirud, C. Narula, and B. C. Gates, 

“γ-Al2O3 Supported Tri-rhenium Rafts: 

Spectroscopic and Microscopic Characterization,” in 

Proceedings of the 19th North American Catalysis 

Society Meeting, Philadelphia, May 22–27, 2005. 

Y. Xu, W. A. Shelton, and W. F. Schneider, 

“Effect of Particle Size on the Oxidizability of 

Platinum Clusters,” submitted to J. Phys. Chem. A. 

Y. Xu, W. A. Shelton, and W. F. Schneider, 

“Nanoscale Effects in the Reactivity of Pt Clusters 

towards CO oxidation,” in Proceedings of the 19th 

North American Catalysis Society Meeting, 

Philadelphia, May 22–27, 2005. 

Y. Xu, W. A. Shelton, and W. F. Schneider, 

“Theoretical Studies Based on Post-Hartree-Fock 

and DFT Methods,” in Synthesis and Applications of 

Heavy Vehicle Propulsion Materials 

Oxide Nanoparticles and Nanostructures, ed. J. A. 

Rodriguez, John Wiley and Sons; in preparation. 

V. A. Bhirud, M. J. Moses, D. A. Blom, L. F. 

Allard, Jr., T. Aoki, S. Mishina, C. K. Narula, and 

B. C. Gates, “Alumina-supported Tri-rhenium 

Clusters Visible by Aberration-Corrected Dark-field 

STEM,” Microscopy and Microanalysis 2005, 

Honolulu, July 31–August 4, 2005. 

C. K. Narula, W. A. Shelton, L. F. Allard, 

Y. Xu, M. Moses, W. Schneider, and B. Gates, 

“Combining Theory and Experiments in Studies of 

Structural Changes in LNT Materials,” Eighth DOE 

Crosscut Workshop on Lean Emissions Reduction 

Simulation, University of Michigan–Dearborn, 

Dearborn, Michigan, May 17–19, 2005. 

References 

1. W. Deng, M. W. Toepke, and B. H. Shanks, 

Adv. Func. Mater. 13, 61 (2003). 

2. B. C. Gates, Chem. Rev. 95, 511 (1995). 

3. C. K. Narula, S. R. Nakouzi, R. Wu, C. T. 

Goralski, and L. F. Allard, AIChE Journal 47, 744 

(2001). 

4. M. A. Andrews, S. W. Kirtley, and H. D. 

Kaesz, Inorg. Synth. 17, 16 (1977). 

5. A. S. Fung, P. A. Tooley, M. J. Kelley, D. C. 

Koningsberger, and B. C. Gates, J. Phys. Chem. 95, 

225 (1991). 

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Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

H. Durability of Diesel Particulate Filters (CRADA with Cummins, Inc.) 

Edgar Lara-Curzio 

Oak Ridge National Laboratory 

P.O. Box 2008, MS-6069 

Oak Ridge, TN 37831-6069 

(865) 574-1749; fax: (865) 574-6098; e-mail: laracurzioe@ornl.gov 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Objective 

• Identify and implement test techniques to characterize the physical and mechanical properties of ceramic diesel 

particulate filters (DPFs) and develop analysis tools for predicting their reliability and durability. 

Approach 

• Identify and implement test techniques to determine the physical and mechanical properties of DPF ceramic 

substrates. 

• Investigate the applicability of probabilistic design tools to DPF ceramic substrates. In particular, the 

applicability of the Ceramic Analysis and Reliability Evaluation of Structures (CARES) code to predict the 

reliability of DPF ceramic substrates will be investigated. 

Accomplishments 

• Developed a procedure to prepare double-torsion test specimens of porous cordierite. 

• Determined the fracture toughness and slow-crack-growth behavior of porous cordierite as a function of 

temperature and environment. 

• Carried out a round-robin testing program to determine the precision of 4-point bending test methods to 

determine the flexural strength of DPF ceramic substrates. 

Future Direction 

• Implement probabilistic design tools to predict the reliability and durability of DPF ceramic substrates when 

subjected to arbitrary thermomechanical histories. 

Introduction 

The pollution emitted by diesel engines 

contributes greatly to the nation’s air quality 

problems. Even with more stringent environmental 

regulations set to take effect in 2007, existing trucks 

and buses will continue to emit nitrogen oxides 

(NOx) and particulate matter (PM), both of which 

contribute to serious public health problems.1 

Fortunately, there are several technologies designed 

to reduce pollution from existing trucks and buses, 

such as DPFs. A DPF is a ceramic device that 

collects PM in the exhaust stream. The high 

temperature of the exhaust heats the ceramic 

61 
FY 2005 Progress Report 

structure and allows the particles inside to break 

down (or oxidize) into less harmful components. 

DPFs reduce emissions of PM, hydrocarbons, and 

CO by 60 to 90%. 

Most DPFs consist of a ceramic honeycomb 

with hundreds of cell passages partitioned by walls 

(Figure 1). Each cell passage has a square cell 

opening at one end and is closed at the other end so 

that the cell passages are alternately closed at each 

end. This structure forces the exhaust gases through 

the porous, thin ceramic honeycomb walls. When 

the gases carrying the carbon particles flow through 

the fine pores of the walls, the carbon particles are 

filtered out. High porosity values, in the range of 

60%, heighten filtration efficiency to more than 90% 

while reducing gas-flow resistance to avoid affecting 

the engine performance. 

The process of diesel PM collection begins as 

soon as the engine is started and continues while the 

engine is operating. As the carbon particles are 

Figure 1. Corning’s cordierite-based DPFs. 

Heavy Vehicle Propulsion Materials 

collected on the ceramic walls, the backpressure of 

the system increases. This problem is alleviated by 

burning the trapped PM, through a catalytic reaction 

using exhaust gas heat at 400°C or more, into CO2 

and water vapor that pass through the filter. This 

process, called regeneration, results in a cleaner 

filter. The regeneration process is dependent upon 

exhaust temperature, oxygen, NOx content, time, and 

PM levels. 

The key to the successful application of DPFs is 

to reliably regenerate the filter (e.g., to burn the PM 

that the filter continues to trap or collect). 

Traditionally, combustion of soot is done in an 

oxygen atmosphere (air). In air, soot will burn at 

about 500°C. However, this is not a typical 

operating temperature for diesel engine exhaust. 

Therefore, to burn soot in air, an active system—i.e., 

one that increases the temperature of the exhaust 

using some external heat source—is required. But if 

an active system is not carefully controlled, or if too 

much PM collects on the filter walls, the filters can 

experience an “uncontrolled burn” where the 

temperature increases to 600°C or more, resulting in 

damage to the filter element. 

The objective of this project is to develop and 

implement methodologies to predict the reliability of 

DPFs. A useful conceptual model for this purpose is 

the reliability bathtub curve, which describes 

reliability-related phenomena of a component over 

its life cycle.2

 A schematic of the reliability bathtub 

curve is depicted in Figure 2. It consists of three 

stages: the infant mortality phase is characterized by 

premature failures due to improper manufacturing or 

assembly, poor workmanship, or defects introduced 

during processing. The second stage of the curve 

corresponds to the useful life of the component and 

Figure 2. Reliability bathtub curves. 

62
Heavy Vehicle Propulsion Materials 

is characterized by a constant failure rate. In this 

regime, failures are typically associated with random, excessive loads. If sufficiently high safety 

factors are used during the design process, the magnitude of this failure rate should be negligible. The 

third stage of the bathtub curve is known as the 

wear-out phase, where the failure rate increases with 

time as a result of aging phenomena. Aging phenomena include thermal and mechanical fatigue, 

corrosion, creep deformation and environmentally 

assisted crack growth. The time at the onset of wearout is often regarded as the useful life of the 

component. 

The reliability bathtub curve can be used as a 

descriptor of how the failure rate of DPFs evolves 

over time. Infancy failures of DPFs could be related 

to manufacturing or process defects (large pores, 

inclusions, cracking) or defects introduced during 

assembly. Failures of DPFs during their useful life 

will be dictated by the intersection between the 

spectrum of thermomechanical loads and the distribution of DPF strengths. Such failures could result 

from excessive vibration, for example, or unwanted 

thermal excursions during transients associated with 

regeneration. Wear-out and degradation failures of 

DPFs could be associated with the growth of microcracks assisted by thermal fatigue, and/or by the 

chemistry of the environment, and/or by chemical 

and microstructural changes in the material due to 

long-term exposure to elevated temperatures in the 

exhaust environment. 

The objective of this project is to develop life 

prediction methods for DPFs and, by using those 

methods, to design durable, reliable DPFs. 

Approach 

Designing DPFs that are durable and reliable 

poses significant challenges. For example, the 

porosity of DPFs, which allows them to remove PM 

from the exhaust gas stream, has a deleterious effect 

on their mechanical strength. This is important 

because DPFs will experience demanding thermomechanical conditions during service. These include, 

for example, thermal shock resulting from rapid 

heating/cooling and stresses that arise from 

temperature gradients. The approach that will be 

followed in this project includes identifying and 

implementing test techniques for the nondestructive 

FY 2005 Progress Report 

evaluation of ceramic substrates to assess their 

integrity. Their physical and mechanical properties 

and the mechanisms responsible for their degradation in the various stages of the bathtub curve also 

will be determined and identified. These properties 

include thermal expansion, thermal conductivity, 

heat capacity, density, porosity, elastic properties, 

strength, fracture toughness, and resistance to crack 

growth at ambient and elevated temperatures, in air 

and in relevant environments. 

The information generated will be used in turn 

to implement probabilistic design tools. In particular, the applicability of the CARES code3

 to predict 

the reliability of DPF ceramic substrates will be 

investigated. Such probabilistic design methodologies are based on a combination of experimentally 

determined strength data, stress analyses of the 

component using a finite-element analysis, and 

selection of appropriate failure criteria. The durability (service life) of the component can also be 

predicted using this framework by considering the 

mechanisms that are responsible for the degradation 

of material strength, such as slow crack growth or 

creep. While these methodologies have been successfully applied to design structural components 

using dense ceramics, their applicability to the 

analysis and design of porous ceramic components 

has not been validated. 

Results 

During FY 2005, a procedure was developed to 

prepare double-torsion test specimens to determine 

the fracture toughness and slow-crack-growth 

behavior of porous cordierite. The double-torsion 

test configuration (Figure 3) consists of symmetric 

four-point loading around a crack or a notch on one 

end of a rectangular plate. One feature of this loading configuration is that the stress intensity factor is, 

at a first approximation, independent of crack length 

for a range of crack lengths. This feature makes 

double-torsion testing ideally suited for the evaluation of opaque and non-reflective plate-like test 

specimens where direct crack length measurements 

could be difficult to make. The double-torsion test 

method is also attractive for slow-crack-growth 

studies because of the relative stability of crack 

extension, in contrast to other testing configurations 

(e.g., SENB, CT). 

63
FY 2005 Progress Report 

Precrack 

Load, P Specimen 

t 

S 

Notch 

L 

Sm 

Figure 3. Schematic of the double-torsion test 

configuration. 

The preparation of double-torsion test specimens 

of porous cordierite, which poses challenges because 

the thickness of the DPF cell walls is of the order of 

300 μm, was accomplished following dry 

cutting/grinding operations. Sections from a DPF 

with a cell size of 2.25 × 2.25 mm and approximately 50% porosity were sliced using a diamond 

cut-off blade. Then subsequent material removal was 

accomplished using a diamond grinding wheel (150 

grit, 4000 rpm, 25 μm down feed). To ensure 

parallelism and flatness of the surfaces of the test 

specimen, a special fixture was fabricated using a 

Heavy Vehicle Propulsion Materials 

precision vice and gage blocks. Test specimens were 

prepared with the notch aligned either parallel or 

perpendicular to the orientation of the walls, and the 

final thickness of the test specimens was 280 μm. 

Figure 4 illustrates the sample preparation process. 

The fracture toughness of porous cordierite test 

specimens was determined by loading pre-cracked 

test specimens to failure at a constant crosshead displacement rate of 0.01 mm/s. The fracture toughness 

was calculated from the peak load using Equation 1: 

⎡3 1( + ν)⎤

1 / 2

KIC = PSm⎢ ⎥ , (1)

⎣ St4

ξ ⎦ 

where P is the peak load, S is the width of the test 

specimens, Sm is the moment arm, t is the specimen 

thickness, ν is Poisson’s ratio, and ξ is a finite beam 

thickness correction factor given by 

⎛ πS ⎞ 

ξ = 1− 1.26 ⎛ ⎜ t ⎞ ⎟ + 2.4 ⎛ ⎜ t ⎞ ⎟e

⎜

⎝

− 2t ⎟

⎠ .

 (2) ⎝ S ⎠ ⎝ S ⎠ 

Figure 5 shows a cordierite test specimen after a 

fracture toughness test, illustrating the trajectory of 

Figure 4. Schematic of procedure followed for the preparation of double-torsion test specimens. 

64
Heavy Vehicle Propulsion Materials 

Figure 5. Photograph of double-torsion test specimen 

after fracture toughness testing. 

the crack, which propagated along the major axis of 

the test specimen. In Figure 5, the notch is located 

on the right of the test specimen. It was found that 

the propagation of the crack was significantly influenced by the microstructure of the material, as 

illustrated by the scanning electron micrograph in 

Figure 6. To better characterize the interaction 

between the microstructure and crack propagation, 

exploratory work was carried out to determine the 

feasibility of imaging the propagation of the crack 

FY 2005 Progress Report 

through a double-torsion test specimen using X-ray 

computerized tomography (CT) scans. In addition to 

providing information on the role of the microstructure in the resistance of the material to crack 

propagation, this work could lead to the implementation of nondestructive evaluation techniques. 

Figure 7 shows a reconstructed image of a porous 

cordierite test specimen. Additional work will be 

pursued in this area in FY 2006. The fracture toughness of porous cordierite was found to decrease with 

temperature from a value of 0.45 ± 0.02 MPa m at 

20°C to 0.36 ± 0.07 MPa m at 800°C. The plot in 

Figure 8 summarizes those results. 

During FY 2005, work was initiated to investigate the slow-crack-growth behavior of porous 

cordierite using the load relaxation version of the 

double-torsion test specimen. According to this test 

procedure, a pre-cracked test specimen is loaded at a 

constant crosshead displacement rate that is lower 

than the load associated with fracture toughness. 

Then the crosshead is arrested and the load is 

monitored as a function of time. From the rate of 

load relaxation, which results from the increase in 

(a) 

(b) 

Figure 6. Scanning electron micrograph illustrating the influence of microstructure on crack propagation: (a) low 

magnification and (b) high magnification. 

65
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 7. Reconstructed image of porous cordierite test 

specimen from X-ray CT scan. 

Figure 8. Fracture toughness of porous cordierite as a 

function of temperature. 

compliance of the test specimen, it is possible to 

determine the rate of crack growth using Equation 3: 

da Pi ⎡ D ⎤⎛ dP ⎞

dt =−

P2 ⎢

⎣

ai +

B ⎥

⎦⎝

⎜ dt ⎠

⎟ , (3) 

where Pi is the initial load, ai is the initial crack 

length, and D and B are experimental constants 

associated with the load train compliance. 

In parallel with these efforts, a series of dynamic 

fatigue tests were also initiated in FY 2005 

to determine the slow-crack-growth behavior of 

porous cordierite according to ASTM C1465-00, 

“Standard Test Method for Determination of Slow 

Crack Growth Parameters of Advanced Ceramics by 

Constant Stress-Rate Flexural Testing at Elevated 

Temperatures.” According to this test method, the 

flexural strength of honeycomb test specimens is 

determined as a function of applied stress rate. The 

strength degradation exhibited with decreasing 

applied stress rate is the basis of this test method, 

which enables the evaluation of slow-crack-growth 

parameters of a material. 

This dual approach to the characterization of the 

slow-crack-growth behavior of porous cordierite is 

necessary to distinguish between material behavior, 

which will be obtained through the double torsion 

testing of test specimens obtained from the walls of 

cellular structures, and structural behavior, which 

will be obtained through dynamic fatigue flexural 

testing of beam test specimens with cellular structure. The results from these two sets of tests will 

provide the kinetics for crack growth to enable the 

prediction of the service life of porous cordierite 

DPFs using CARES. The slow-crack-growth 

behavior of porous cordierite and porous cordierite 

DPFs with cellular structure will be determined in 

air, water-vapor, a reducing environment, and 

simulated diesel exhaust environments. 

In FY 2005, a round-robin testing program was 

organized between Oak Ridge National Laboratory 

(ORNL) and Cummins to determine the precision in 

the determination of the 4-point flexural strength of 

cordierite DPFs. Each laboratory was responsible for 

testing 40 test specimens with nominal dimensions 

of 100 × 26 × 13.5 mm. Tests were carried out at 

ambient conditions under a constant crosshead 

displacement of 0.35 mm/min, using inner and outer 

spans of 19 and 89 mm. The test results, which were 

analyzed using Weibull statistics, are presented in 

Figure 9. It was found that the average and standard 

deviations of all tests were 6.73 and 0.83 MPa, 

respectively, and that the repeatability and reproducibility standard deviations were 0.35 and 

0.9 MPa, respectively. The Weibull modulus, which 

was determined according to the maximum likelihood method, was found to be 24.1. The discrepancies in the characteristic strength values were 

associated with procedures used to determine the 

dimensions of the test specimens. 

66
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 9. Weibull plot summarizing flexural strength results from roundrobin testing program. (U: ORNL data; {: Cummins data). 

Summary 

A procedure was developed to prepare test 

specimens from cellular DPFs to determine the 

fracture toughness of porous cordierite by the 

double-torsion test method. Using this test method, 

the fracture toughness of porous cordierite was 

found to decrease with increasing temperature from 

0.45 ± 0.02 MPa m at 20°C to 0.36 ± 0.07 

MPa m at 800°C. Work was initiated to determine 

the slow-crack-growth behavior of porous cordierite 

and cellular structures of porous cordierite, using the 

load relaxation variation of the double-torsion test 

method and the constant stress-rate flexural test 

method according to ASTM C1465, respectively. 

The results from these tests will enable prediction of 

the service life of porous cordierite DPFs. A roundrobin testing program was completed to assess precision in the determination of the 4-point flexural 

strength of cordierite DPFs. It was found that the 

average and standard deviations of all tests were 

6.73 and 0.83 MPa, respectively and that the repeatability and reproducibility standard deviations were 

0.35 and 0.9 MPa, respectively. The Weibull 

modulus was found to be 24.1 

References 

1. United States Environmental Protection 

Agency, Office of Transportation and Air Quality 

EPA420-F-03-017, June 2003. “Questions and 

Answers on Using a Diesel Particulate Matter Filter 

in Heavy-Duty Trucks and Buses.” 

2. G. S. Wasserman, Reliability Verification, 

Testing and Analysis in Engineering Design, Marcel 

Dekker, New York, 2003. 

3. N. N. Nemeth, L. P. Powers, L. A. Janosik, 

and J. P. Gyekenyesi, “Ceramics Analysis and 

Reliability Evaluation of Structures Life Prediction 

Program: Users and Programmers Manual,” NASAGRC, Life Prediction Branch (1993). 

Collaborators 

Amit Shyam, H-T. Lin, Randy Parten, and 

Claire Luttrell. 

Presentations and Publications 

A. Shyam and E. Lara-Curzio, “The Doubletorsion Testing Technique for the Determination of 

Fracture Toughness and Slow-Crack Growth 

Behavior of Materials: A Review,” accepted for 

67
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

publication in the Journal of Materials Science 

(2005). 

T. Yonushonis and E. Lara-Curzio, “Durability 

of Particulate Filters—PM 10461,” presented at the 

FY 2005 Heavy Vehicle Materials Program Review, 

Oak Ridge National Laboratory, Oak Ridge, 

Tennessee, September 14, 2005. 

68
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

4. MATERIALS FOR AIR HANDLING, HOT SECTION, AND 

STRUCTURAL COMPONENTS 

A. High-Temperature Advanced Materials for Lightweight Valve Train 

Components 

Jeremy S. Trethewey, Principal Investigator 

John Grassi, David Longanbach, and Chad Logan 

Caterpillar, Inc. 

Tech Center 854-80 TC-E 

14009 Old Galena Rd. 

Mossville, IL 61552 

(309) 578-0056; fax: (309) 578-2953; e-mail: Trethewey_Jeremy_S@cat.com 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract Number: DE-AC05-00OR22725 

Subcontractor: Caterpillar, Inc., Peoria, Illinois 

Objective 

• Evaluate prospective lightweight valve materials and determine feasibility of implementation in heavy-duty 

diesel and natural gas applications. 

Approach 

• Design, procure, and evaluate γ-phase titanium aluminide (TiAl) and silicon nitride (Si3N4) heavy-duty engine 

valves. 

• Demonstrate performance and durability advantages by performing 1000-h durability test on valves in a 

Caterpillar G3406 natural gas engine. 

• Develop finite-element analysis (FEA)-based life prediction model (NASA CARES/Life) that will generate life 

prediction curves for advanced valve materials. 

• Evaluate the thermal, chemical and mechanical properties of TiAl and Si3N4 that are critical to on-engine valve 

performance and durability. 

• Investigate manufacturing considerations (i.e., machinability and quality inspection technique) of these novel 

materials. 

Accomplishments 

• Procured 12 TiAl and 45 Si3N4 heavy-duty engine valves. 

• Performed pretest characterization on all valves [laser-scatter nondestructive evaluation (NDE), dimensional 

check with coordinate measuring machine, surface roughness]. 

• Completed preliminary life prediction model for Si3N4. 

69 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

• Kicked-off studies to examine friction, wear, corrosion, and machinability characteristics of TiAl. 

• Donated a Caterpillar C15 ACERTTM engine to the National Transportation Research Center (NTRC) for future 

component testing. 

Future Direction 

• Complete current G3406 engine test on TiAl and Si3N4 valves; 

⎯ perform laser scatter NDE scans to track crack/defect propagation;

⎯ perform fast fracture on specimens extracted from the valves to determine retained strength; 

⎯ correlate results with life prediction model. 

• Procure additional TiAl valve blanks for second round of engine tests; will use life prediction model as a design 

tool to optimize valve geometry. 

• Refine NASA/CARES Life prediction model to include non-axisymmetric deformation modes, enable TiAl 

modeling, and evaluate Probabilistic Design System (PDS) methodology. 

• Extend friction, wear, corrosion, and machinability studies of TiAl. 

Introduction 

Valve train components in heavy-duty engines 

operate under high stresses, at elevated temperatures, and in severely corrosive environments. 

Structural ceramics and emerging intermetallic 

materials are highly corrosion and oxidation resistant and possess high strength and hardness at elevated temperatures. These properties are expected to 

allow higher engine operating temperatures, lower 

wear, and enhanced reliability. In addition, the 

lighter weight of these materials (~50% of production alloys) will lead to lower reciprocating valve 

train mass that could improve fuel efficiency. 

Over the past decade, the automotive engine 

industry has demonstrated moderate success with 

TiAl as a valve material. Eylon et al.1

 demonstrated 

a 2% fuel savings on a Chevrolet Corvette that ran 

for 25,000 km with TiAl valves. This fuel efficiency 

increase is attributed to the decrease in valve train 

parasitic frictional losses. In another study with TiAl 

valves, Maki et al.2

 demonstrated a 1000-rpm 

increase in over-speed performance on a Nissan 

VRT35 engine. This capability would allow more 

efficient engines to operate in more demanding 

regimes. The current research and development program at Caterpillar, Inc., extends this body of 

knowledge from the automotive community and 

examines TiAl and Si3N4 in the context of heavyduty diesel and natural gas engine environments. 

The valve train material development effort will 

provide the materials, design, manufacturing, and 

economic information necessary to bring these new 

materials and technologies to commercial realization. With this information, component designs will 

be optimized using probabilistic lifetime prediction 

models, and validated in rig bench tests and shortand long-term engine tests. After establishing proofof-concept with valves, this design approach will be 

applied to other components made from hightemperature materials. 

Approach 

Because the broad goal of this project is to 

evaluate TiAl and Si3N4 materials for use in valve 

applications, many perspectives must be considered. 

As such, there are several facets to this project. The 

primary thrust toward component validation is 

achieved by fabricating TiAl and Si3N4 valves and 

testing them on a Caterpillar G3406 (in-line 6 

cylinder) natural gas engine. Previous annual reports 

documented the design process used to modify the 

current baseline metallic valve to accommodate the 

limited ductility of ceramic and intermetallic 

materials.3

 For example, the head dimensions were 

changed to reduce the stress concentration experienced in the fillet radius. Silicon nitride and TiAl 

valve blanks were formed, rough machined, friction 

welded to a Ti-6Al-4V shaft (TiAl only), and finish 

machined. 

A parallel effort to develop a design tool for 

high-strength, brittle materials is also being pursued. 

A FEA-based life prediction code (NASA 

CARES/Life) is used to evaluate component design 

and generate life prediction curves. This approach 

70
Heavy Vehicle Propulsion Materials 

uses material properties and thermo-mechanical 

boundary conditions as the primary inputs. A 

probability-based code then calculates the accumulated damage through the life of the valve and 

creates a life prediction curve. 

A variety of other studies contribute to the 

understanding of the behavior of TiAl and Si3N4 and 

how they may perform in an engine environment. 

Several studies examine the fundamental properties 

of the materials: tensile strength, creep strength, 

thermo-physical properties, friction coefficient, wear 

resistance, oxidation resistance, and corrosion resistance. Two other studies examine the manufacturability of the materials: the effect of machining damage on fatigue resistance and a NDE technique to 

evaluate surface quality.

4 

Results 

Valve fabrication 

Twelve TiAl and forty-five Si3N4 valves were 

fabricated and are under final preparations before 

engine testing. Figure 1 shows the steps of the fabrication process for a TiAl valve. Once the design was 

finalized, the geometry was sent to the suppliers for 

blank fabrication. The valve is then rough machined 

and friction welded to a Ti-6Al-4V valve. This 

FY 2005 Progress Report 

friction weld procedure is performed for three 

reasons: (1) cost effectiveness, (2) difficulty in 

maintaining the straightness specification on 

castings with high aspect ratio and (3) hightemperature strength is not required in stem. The 

final procedure, finish machining, is required to 

minimize the probability of a brittle failure caused 

by surface defects. 

As was noted in this program’s previous annual 

report5

 a supplier oversight resulted in incorrect 

surface finish that resulted in an early life failure on 

an impact rig. Subsequently, the valves were 

returned to the supplier for corrective finish 

machining. It is clear from the laser scatter NDE 

scans (Figure 2) that the majority of the machining 

damage was removed after the second finish 

machining procedure. 

Additional preengine test characterization was 

performed on all valves. A coordinate measuring 

machine (CMM) was used to ensure that the valves 

met the dimensional specifications. Also, a profilometry scan was performed on several critical sections of the valves to verify that finish machining 

achieved the necessary surface roughness. It was 

confirmed that the surface roughness specification 

was met for nearly all of the measured regions on all 

of the valves. 

Figure 1. Critical steps of the fabrication process for a TiAl valve. Si3N4 valves are subject to every step 

except the joining of the head to stem because they are formed as a single piece. 

71 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 2. Laser scatter NDE scans of the same Si3N4 valve after poor finish machining (a) and after 

correct finish machining (b). Note that the machining defects (white streaks) are minimized 

after the valve is finish machined properly. 

Valve engine testing 

Arrangements have been made to perform 

engine tests at the National Transportation Research 

Center (NTRC) located in Oak Ridge, Tennessee. 

These tests will be executed with the intention to 

compare valve performance and reliability of 

advanced materials to current production materials. 

The valves will be tested on a Caterpillar G3406 

natural gas genset platform. A G3406 cylinder head 

was outfitted with 32 thermocouples embedded in 

4 intake and 4 exhaust valve seat inserts. Pressure 

transducers were installed in each of the six cylinders, the intake manifold and exhaust manifold. In 

addition, several custom valve train components 

have been designed to accommodate the modified 

valve design (i.e. thicker head, change in keeper 

geometry). Custom seat inserts, valve guides, keeper 

locks and bridges have been procured and are ready 

for installation in the G3406 engine. 

A test plan has been generated to evaluate TiAl, 

Si3N4, and production valves performance in 

parallel. The material responses that will be studied 

are as follows: wear rates will be evaluated via 

profilometry and micrographic analysis; crack 

initiation and propagation will be evaluated using 

laser scatter and ultrasonic NDE; deposition 

formation will be evaluated via chemical analysis. 

The valves will be evaluated at various time 

intervals while accumulating 1000 h of engine test 

exposure. 

As previously mentioned, the valves have been 

finish machined and are ready for on-engine evaluation (Figure 3). However, there must be an acceptable familiarity with the performance of these novel 

valve materials prior to evaluation on engine tests. 

Thus, a set of two Si3N4 intake valves and two TiAl 

exhaust valves is currently accumulating time on an 

impact test rig (20 h to date) to assess early life wear 

and crack propagation. 

Valve modeling 

To rigorously evaluate the performance of a 

component made from a brittle material, it is insightful to consider the component’s lifetime as a function of thermal stresses, mechanical stresses, and 

the cyclic nature of these loads. With Caterpillar’s 

input, Jadaan Engineering & Consulting has been 

developing a lifetime prediction model, using NASA 

72
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 3. Finish valves ready for engine testing. 

CARES/Life code, for a Si3N4 exhaust valve sub- • Mechanical boundary conditions as a function of 

jected to the typical stresses seen in a G3406 engine crank angle [Figure 4(b)] 

environment. The inputs to this code follow: • Thermal boundary conditions 

• FEA model to establish the valve geometry • Mechanical properties of the valve material 

[Figure 4(a)] 

Figure 4. Inputs for the probabilistic life prediction model. (a) FEA valve model. (b) Mechanical 

boundary conditions as a function of crank angle. 

73 
FY 2005 Progress Report 

These inputs were supplied to Jadaan Engineering & Consulting for a G3406 valve experiencing 

representative operating conditions. A FEA model 

was then generated, and a transient reliability analysis was performed. Subsequent to applying the 

mechanical and thermal boundary conditions onto 

the valve geometry, Si3N4 material properties are 

assigned to the elements. The mechanical load spectrum [Figure 4(b)] was discretized into a finite number of steps, and the stress state was calculated for 

each set of loads. Using the calculated stress fields, a 

computational transient reliability analysis is performed. In this approach, the material is susceptible 

to slow crack growth (fatigue) and the probability of 

failure depends on both the number of cycles and 

load history sequence. The probability of failure 

during a given time step is directly impacted by 

those in the previous time steps. Hence, because of 

fatigue damage the probability of failure will 

increase with time as damage accumulates with 

additional load cycles. The generated life prediction 

curve is shown in Figure 5. According to this curve, 

at a safety factor of three, there is <0.0001% probability of failure after 10 years of continuous operation! Although these results may seem to be extraordinarily promising, they must be interpreted with 

caution. Note that four loads that have been modeled 

are not the only loads present; they have only served 

Figure 5. Life prediction curve generated from the preliminary probabilistic model. Safety factor 

represents the multiplication factor between 

the defined mechanical boundary conditions 

and the value used to calculate the probability 

of failure (i.e., if tip pressure was defined as 

10 MPa; the set of data with safety factor = 5 

used a tip pressure of 50 MPa). 

Heavy Vehicle Propulsion Materials 

to explore the model’s capability. Further refinement 

of the valve’s loading conditions must be made to 

model the valve’s true operating conditions (i.e., 

non-axisymmetric bending loads). It is expected that 

when more rigorous boundary conditions are applied 

that the probability of failure will increase compared 

to this preliminary model. 

In addition to refining the model’s boundary 

conditions (i.e., implementing a 3-D model in place 

of the current axisymmetric model), several new 

capabilities will be added to the model. The model 

will be extended to enable the modeling of TiAl. 

Titanium aluminide will be modeled as a material 

with a 3-parameter Weibull distribution.6

 Also, the 

ANSYS/PDS (Probabilistic Design System) will be 

employed to evaluate the effect that model parameter variation (of e.g., dimensional tolerance, external 

loading, characteristic strength) has on the life prediction result. 

Material characterization 

Investigation of basic material properties and 

manufacturing considerations supplement the engine 

test and modeling effort. For instance, fundamental 

thermo-physical and mechanical data were collected 

on two varieties of TiAl and several varieties of 

Si3N4. The machinability of TiAl and Si3N4 and the 

friction coefficients of these materials are two subjects of particular interest. 

Machining processes account for a significant 

portion of the cost of valve components. This is 

especially true for high-temperature materials such 

as intermetallics and ceramics because these 

materials show a greater sensitivity to surface 

features than metallic alloys. Thus, it is necessary to 

identify the most economical method of machining 

that achieves the required specifications (i.e., 

geometric tolerances, surface roughness). A series of 

experiments are planned that will investigate the 

capability of three machining techniques (turning, 

grinding, laser-assisted machining) on TiAl test 

specimens. The surface finish achieved by each 

technique will be quantified based on microscopy 

and profilometry measurements. These test 

specimens will then be subjected to axial stepped 

fatigue tests at room temperature and engine 

operating temperatures to evaluate the effect of 

machining damage and temperature on the fatigue 

performance of TiAl. These fatigue tests are to be 

performed at Oak Ridge National Laboratory. 

74
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Friction is a fundamental phenomenon in the 

valve train system. It behaves in both a detrimental 

manner (e.g., wear of valve seat contact) and beneficial manner (e.g., traction that allows retention of 

the rotocoil). There is currently no comprehensive 

database that captures the friction behavior of valve 

materials and their peripheral components. To 

address this, a matrix of experiments has been 

developed that will explore the interaction of various 

valve materials (TiAl, Si3N4 and austenitic steel) 

with the materials that they contact (seat insert, 

valve guide, keeper and bridge). These tests will be 

performed on an in-house Cameron-Plint test apparatus in an environment that matches the engine 

operating conditions as closely as possible. Results 

from these tests will aid modeling efforts and highlight the consequences of changing valve materials 

from the well-known production steel to an 

advanced material. 

Conclusions 

Lightweight, high-strength, high-temperature 

materials are enablers for advanced combustion concepts and fuel-efficient engine designs. This effort 

aims to examine two prospective materials, TiAl and 

Si3N4, in the context of heavy-duty diesel and natural gas engine environments. In this year progress 

has been made on several fronts: (1) toward conducting engine tests on these valve materials, 

(2) development of a probabilistic life prediction 

model, and (3) compilation of material properties of 

TiAl and Si3N4 that are critical to implementation of 

these materials in the extreme environment of a 

heavy-duty engine. 

Twelve TiAl and forty-five Si3N4 valves have 

completed the fabrication process. Their incorrect 

surface finish provided by the machining supplier 

was rectified. The valves have been subjected to 

several preengine test characterization assessments: 

CMM, profilometry and laser-scatter NDE. The first 

set of valves has accumulated 20 h on an impact rig 

and has demonstrated acceptable performance. Upon 

successful completion of the impact rig test, the 

valves will be tested on a Caterpillar G3406 engine 

at the NTRC. In addition to fabricating custom 

components that will enable this engine to accept the 

new modified valve design, the engine has been 

instrumented to capture the performance of each 

cylinder. 

Significant progress has been made toward 

developing a NASA CARES/Life prediction model. 

A preliminary model has been developed, and a life 

prediction curve has been generated. It was speculated that this preliminary curve significantly underpredicts the probability of failure due to incomplete 

representation of the boundary conditions. Model 

enhancements are currently underway and updated 

results will be correlated to the engine performance 

of the valve hardware. 

Thermo-physical properties (e.g., coefficient of 

thermal expansion, thermal conductivity, density) 

and mechanical properties (e.g., ultimate tensile 

strength, Young’s modulus, ductility) of TiAl and 

Si3N4 have been collected. Properties specific to 

implementation of these advanced materials in a 

valve environment are also under investigation (e.g., 

machinability, frictional coefficients, corrosion 

resistance). 

This diverse effort will help to equip design 

engineers with the information required to implement advanced lightweight materials in a valve train 

application. This program has made significant progress toward this goal in FY 2005 and is poised to 

significantly advance the case for TiAl and/or Si3N4 

as a future valvetrain material in FY 2006. 

References 

1. D. Eylon et al., “Development of permanentmold cast TiAl automotive valves,” Intermetallics, 

6, 1998. 

2. K. Maki et al., “Development of a high-performance TiAl exhaust valve,” SAE Paper 960303, 

Warrendale, Pennsylvania, 1996. 

3. M. J. Andrews and J. Chapa-Cabrera, “Hightemperature advanced materials for lightweight 

valve train components,” Heavy Vehicle Propulsion 

Materials 2003 Annual Progress Report, Energy 

Efficiency and Renewable Energy Office of 

FreedomCAR and Vehicle Technologies, March 

2004. 

4. J. G. Sun, “NDE development for ceramic 

valves for diesel engines,” Heavy Vehicle Propulsion Materials 2004 Annual Progress Report, 

Energy Efficiency and Renewable Energy Office of 

FreedomCAR and Vehicle Technologies, May 2005. 

5. J. S. Trethewey et al., “High-temperature 

advanced materials for lightweight valve train 

components,” Heavy Vehicle Propulsion Materials 

2004 Annual Progress Report, Energy Efficiency 

75 
FY 2005 Progress Report 

and Renewable Energy Office of FreedomCAR and 

Vehicle Technologies, May 2005. 

6. Biery et al., “Use of Weibull statistics to 

quantify property variability in TiAl alloys,” Metallurgical and Materials Transactions A, 33A, 

3127–3136 (October 2002). 

Presentations and Publications 

J. M. Zhang, J. G. Sun, M. J. Andrews, A. 

Ramesh, J. S. Trethewey, and D. M. Longanbach, 

“Characterization of subsurface defects in ceramic 

rods by laser scattering and fractography,” presented 

at the 32nd Annual Review of Progress in Quantitative Nondestructive Evaluation, Bowdoin College, 

Brunswick, Maine, July 31–August 5, 2005, and 

paper to be published in conference proceedings. 

Heavy Vehicle Propulsion Materials 

J. G. Sun, J. M. Zhang, J. S. Trethewey, J. A. 

Grassi, M. P. Fletcher, and M. J. Andrews, 

“Nondestructive evaluation of machining and benchtest damage in silicon nitride ceramic valves,” paper 

presented at the American Ceramic Society’s 29th 

Ann. Intl. Cocoa Beach Conf. & Exposition on 

Advanced Ceramics & Composites, Cocoa Beach, 

Florida, January 23–28, 2005, and to be published in 

conference proceedings. 

J. S. Trethewey, H.-T. Lin, and J. G. Sun, 

“Lightweight valve train materials,” presented at 

Department of Energy Merit Review, Oak Ridge 

National Laboratory, September 15, 2005. 

76
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

B. Engineered Surfaces for Diesel Engine Components 

M. Brad Beardsley 

Caterpillar, Inc. 

Technical Center Bldg. E/854 

P.O. Box 1875 

Peoria, IL 61656-1875 

(309) 578-8514; fax: (309) 578-2953; e-mail:Beardsley_M_Brad@cat.com 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: Caterpillar, Inc., Peoria, Illinois 

Objectives 

• Develop coating processes to apply amorphous steels and quasicrystal materials to engine pistons and cylinder 

heads for use as thermal barrier coatings in homogenous charged combustion ignition (HCCI) engine designs. 

• Quantify the diffusion stability of new quasicrystalline material at engine operating temperatures. 

Approach 

• Use high-velocity oxygen fueled (HVOF) processes to apply amorphous and quasicrystalline materials to 

engine piston and cylinder head. 

• Use microprobe analysis of diffusion couples to determine the stability of quasicrystalline and bond coating 

couples at engine operation temperatures. 

Accomplishments 

• Coated HCCI pistons with amorphous steel for single cylinder engine testing. 

• Measured diffusion profiles for the quasicrystalline material and three different bond coating materials. 

Future Direction 

• Modify the thickness and uniformity of the amorphous steel coating if needed after initial engine testing. 

• Develop graded coating structures for the quasicrystalline material based upon the diffusion profiles developed. 

Introduction 

Two new materials, amorphous steels and quasicrystalline materials, are investigated as thermal 

barrier coatings. Both materials are capable of being 

sprayed using high-velocity oxygen fuel (HVOF) 

processes that produce denser coatings with higher 

strengths than plasma spraying. The denser coating 

structure of HVOF coatings eliminates the need for 

sealing of the coating porosity and the higher 

strength enhances the coating durability. Amorphous 

steel coatings are reported to have low thermal 

conductivity as well as high hardness and toughness, 

making them attractive as thermal barrier coatings.1,2 

77 
FY 2005 Progress Report 

The quasicrystal being studied, Al71Co13Cr8Fe8, has 

been shown to be capable of 1100°C operation in a 

turbine engine.3

 Application of thermal barrier 

coatings (TBCs) to new homogeneous charged 

combustion ignition (HCCI) engine components can 

result in lower fuel consumption and lower engine 

component temperatures similar to that shown in 

standard diesel engine configurations (Figure 1). 

Approach 

Amorphous Steel Coatings 

The SHS717 amorphous steel produced by The 

NanoSteel Company, Maitland, Florida, was 

selected for use as a TBC. Thermal diffusivity of 

HVOF sprayed SHS717 shows that the amorphous 

Heavy Vehicle Propulsion Materials 

structure is retained up to approximately 700oC 

(Figure 2). New HCCI engine designs have surface 

temperatures below this range allowing the use of 

the amorphous coatings. New piston designs were 

procured, and coating processes were developed to 

apply up to 3.5-mm coating to the piston bowl. 

The piston components were coated using a 

kerosene-fueled HVOF torch manufactured by 

Praxair TAFA. The geometry of the part is shown in 

Figure 3. Because the geometry of the pistons consists of surfaces that are up to 60° from normal, they 

represent a challenge to coat uniformly. It is also 

necessary to maintain the torch orientation as close 

to 90° to the surface to obtain the highest coating 

density. 

Figure 1. High-density (sealed) TBCs can reduce fuel consumption and provide temperature reduction 

needed for HCCI engine components. 

78
Heavy Vehicle Propulsion Materials 

Figure 2. Exposure to temperatures above glass 

transition temperature (700°C) causes an 

increase in thermal diffusivity of amorphous 

steel. 

Figure 3. Coated piston geometry. 

Two types of robotic motion were investigated 

for coating deposition: (1) single radial 

[Figure 4(a)], and (2) raster combined with a short 

radial motion to cover the 60° angled part 

[Figure 4(b)]. The raster type motion is the easiest to 

program to achieve a uniform coating thickness but 

results in a high wastage of coating material due to 

the torch traversing off the part for much of the 

cycle. The radial motion is more efficient in material 

usage (lower cost) but difficult to program as the 

speed control of the robot is uncertain due to the 

high acceleration/deceleration rates required. As the 

circumference of the spray footprint depends on the 

position with respect to the center, the robot speed 

was adjusted accordingly to achieve a uniform 

FY 2005 Progress Report 

deposition rate. Initially, practice piston geometries 

were used to develop the coating. The pistons were 

rotated in a lathe for either robotic motion with the 

rotation speed at 1700 rpm. 

The raster type motion provides an easier 

method to control the deposition rate because each 

pass over the part is done at a constant speed. However, the raster type motion results in a part of the 

surface being sprayed under 30° angle rather than 

90°. Also, the stand-off distance is not held constant 

and varies by half of the bowl depth. In total, 20 

practice pistons were coated using the radial type 

motion and 2 using the raster type motion. Final two 

pistons coated for engine testing used the raster type 

motion as this provides the most uniform coating. 

Further work to achieve coating uniformity with the 

radial motion is required and will be pursued if the 

initial engine testing is successful. 

Quasicrystalline Materials 

Microprobe analysis of diffusion couples is 

being used to assess the thermal stability of the 

quasicrystalline coatings for use as thermal barrier 

coatings. HVOF coatings of the Al71Co13Cr8Fe8 

material were previously produced for diffusion 

couple testing: (1) the quasicrystal and steel substrate, (2) the quasicrystal, a Ni-17Cr-6Al-0.5Y 

bond coat, and steel substrate, (3) the quasicrystal, a 

Ni-31Cr-11Al-0.6Y bond coat, and steel substrate, 

and (4) the quasicrystal, a Fe-26Cr-8Al-0.4Y bond 

coat, and steel substrate. The diffusion couples were 

produced by spraying a steel sample, 12.5 mm 

diameter by 19 mm, with 0.5-mm-thick layers of the 

bond coatings and quasicrystal. These samples were 

then sealed in evacuated quartz tubes and held at 

temperatures of 500, 700, and 900°C for 25, 100, 

and 500 h. The samples were water quenched after 

annealing, and samples were prepared for examination by optical microscopy, scanning electron 

microscopy (SEM), and microprobe analysis. 

Results 

Amorphous Steel Coatings 

The microstructure of the coatings was evaluated using optical microscopy. The microstructure of 

the sample deposited using the raster motion 

79 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 4. Coating motions used (a) continuous radial motion (top) and 

(b) two-stage raster motion (bottom). 

(Figure 5) was compared to the microstructure produced using the radial motion (Figure 6). The coating deposited using the raster motion exhibits higher 

porosity in the radius area at the bottom of the bowl. 

This is caused by the turbulence created by the bowl 

geometry. The coating deposited using the radial 

motion exhibits higher density that is more uniform 

along the radius of the piston. However, thickness 

uniformity in coatings formed using a radial motion 

remains a challenge. Pistons coated using the raster 

motion were supplied for engine testing but have not 

yet been tested. 

Quasicrystalline Materials 

Microprobe analysis of the diffusion couple 

samples shows high mobility of aluminum into both 

the nickel- and iron-based bond coatings in only 

25 h at 700oC (Figures 7 and 8). Cobalt shows 

higher mobility into the iron based bond coating 

than the nickel. Further analysis of the diffusion 

zones at various temperatures and times is under 

way to provide a better understanding of the diffusion behavior of various species. 

80
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 5. Microstructure of the practice piston 9 sprayed using a raster type motion showing higher porosity in the 

radius area of the bowl. 

81 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 6. Microstructure of the practice piston 20 sprayed using a radial type motion showing lower 

porosity in the radius area of the bowl. 

82
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 7. Microprobe results for the quasicrystal and nickel-based bond 

coating showing high mobility of the aluminum into the bond 

coating. 

83 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 8. Microprobe results for the quasicrystal and iron-based bond coating 

showing mobility of the cobalt in the bond coating. 

Conclusions 

Amorphous Steel Coatings 

The microstructures of the coatings applied to 

engine test pistons can be improved using the radial 

motion, but the coating thickness uniformity will 

need to be improved. Successful engine testing of 

the initial coated pistons will provide the emphasis 

to further develop these coatings. 

Quasicrystalline Materials 

The reaction zone between the quasicrystalline 

and bond coatings of iron and nickel show high 

mobility of the aluminum at 700oC. Because the 

interface of the bond coating and quasicrystalline 

material will be below 500°C, this should not preclude their use as a thermal barrier coating system. 

Graded designs may be limited by this reactivity, 

and further analysis of the diffusion kinetics will be 

done to verify the temperatures of use for the quasicrystal and bond coatings combinations. 

84
Heavy Vehicle Propulsion Materials 

References 

1. D. J. Branagan, W. D. Swank, D. C. 

Haggard, and J. R. Fincke, “Wear Resistant 

Amorphous and Nanocomposite Steel Coatings,” 

Met. & Mat. Transactions A, 32A, 2615–2621 

(2001). 

2. D. Shin, F. Gitzhofer, and C. Moreau, 

“Development of Metal Based Thermal Barrier 

Coatings (MBTBCs) for Low Heat Rejection Diesel 

FY 2005 Progress Report 

Engines,” Proceedings of the 2005 International 

Thermal Spray Conference, 915–919 (2005). 

3. A. Sanchez, F. J. Garcia, J. M. Algaba, 

J. Alvarez, P. Valles, M. C. Garcia-Poggio, and 

A. Aguero, “Application of Quasicrystalline 

Materials as Thermal Barriers in Aeronautics and 

Future Perspectives of Use for These Materials,” 

Mat. Res. Soc. Symp. Proc., 553, 447–458 (1999). 

85 

___________________ 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

C. Cermet Composites for Wear Applications 

M. Anderson, T .M. Yonushonis, R. J. Stafford, M. Henrichsen, D .J. Gust, B. Seegmiller,* and J. Stephan* 

Cummins, Inc. 

Columbus, IN 47202-3005

(812) 377-3279; fax: (812) 377-7226; e-mail: randy.j.stafford@cummins.com 

*CoorsTek, Golden, Colorada 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Prime Contract No.: DE-AC05-00OR22725 

Subcontractor: Cummins, Inc., Columbus, Indiana 

Objectives 

• Optimize injection molding binder composition and sintering cycles to achieve 99+% density. This will include 

the determination of the shrink rate during sintering for prototype tool construction and demonstration. 

• Test prototype components in intended application on Cummins’ fuel system test rigs to show viability of 

material in an application. 

Approach 

• Verify properties, microstructure, and shrinkage in test bars manufactured from new source of NiAl powder in 

injection molded powder composition. Ensure that the properties are the same as for the previous NiAl powder. 

• Test prototypes of powder injection molded near net shape components of TiC/Ni3Al cermet on abuse tests at 

Cummins Fuel Systems. 

Accomplishments 

• Found that new source of powder has not compared favorably with the previous source. More work is in 

progress to define the powder and what steps are needed to improve the strength. 

• Did not stabilize prototype component designs during the year, so little progress was made in defining a 

molding tool for the prototype components. 

Future Direction 

• Finalize process development for powder using the new source of NiAl powder. 

• Evaluate powder injection mold prototype components on abuse test rigs at Cummins. 

Introduction injection systems to become more advanced. Injection pressures need to increase (helps with emis- Increasingly stricter diesel engine combustion sions), and more precise control of when and how emission standards and the customer’s desire for 

maintained or improved fuel economy requires fuel much fuel is injected (helping both emissions and 

87 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

fuel economy) is required. New materials that 

exhibit excellent wear properties against steel components are needed for application where components are sliding and impacting against each other. 

One material that Cummins is interested in was 

developed under a DOE Cooperative Agreement 

DE-FC05-97OR22538: TiC/Ni3Al. 

TiC/Ni3Al is a composite of TiC powder in a 

Ni3Al matrix that was developed in collaboration 

with Oak Ridge National Laboratory (ORNL) and 

CoorsTek. Its composition was established to match 

the unique properties required for Cummins Fuel 

Systems applications. CoorsTek and ORNL have 

developed the processing methods to make the material, while Cummins has been responsible to finding 

an application for it and testing the material as fuel 

system components and evaluating their performance vs other materials. 

Approach 

The established composition of TiC/Ni3Al cermet will be processed using the new source powder 

and ten powder injection molded components for 

evaluation of properties. The microstructure, density, and shrinkage rate will be evaluated to determine the suitability of the new powder. Coinciding 

with this work, fuel systems will refine and finalize 

the design of some prototype components to be near 

net shape molded. Following confirmation of the 

properties and final design, prototype components 

will be manufactured for rig testing to evaluate the 

performance and wear resistance of the cermet material in a potential application in the next-generation 

fuel system at Cummins. 

Results 

The benchmark strength for the TiC/Ni3Al powder from pressed and sintered bars is shown below. 

The initial powder injection molding with the same 

formulation powder showed a loss of ~10% strength 

(see below). This was not considered as significant 

because the debinding process was not optimized for 

the powder. The powder injection molding with the 

new NiAl powder showed an additional loss of 15%. 

This loss was confirmed with tests of the powder 

pressing and sintering (see below). Evaluation of the 

test bars showed the presence of NiAl. This may 

account for the lower strength. 

The prototype component design was stabilized, 

and a final drawing was created for a mold tool to 

create the powder injection molded components. 

Mean Standard Coefficient 

Material strength deviation of variation 

(ksi) (ksi) (%) 

Baseline 152 20.4 13 

(pressed) 

Injection 137 21.5 16 

molded (old) 

Injection 114 5.6 5 

molded 

(new) 

Pressed (new) 108 7.0 6 

Conclusions 

Powder processing is critical to the properties of 

the material. Additional work is needed to bring the 

powder properties to the same level as the material 

using the old source of NiAl powder. 

Stabilized design of the prototype is necessary to 

manufacture the prototype pieces for testing. 

88
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

D. Processing and Characterization of Structural and Functional Materials for 

Heavy-Vehicle Applications 

J. Sankar, Z. Xu, and S. Yarmolenko 

North Carolina A &T State University 

1601 East Market Street 

Greensboro, NC 27411 

(336) 256-1151, extension 2282; fax: (336) 256-1153; e-mail: sankar@ncat.edu 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: North Carolina A&T State University, Greensboro, North Carolina 

Objectives 

• Produce yttria-stabilized zirconia (YSZ) thin coatings using electrophoretic deposition (EPD) technique for solid 

oxide fuel cells (SOFCs). 

• Study the effect of the processing parameters on the porosity of the NiO/YSZ anode substrates. 

• Study liquid-phase sintering technique of YSZ. 

Approach 

• Study the experimental parameters for the EPD process and deposit YSZ thin film coatings using this 

technique. 

• Characterize the coatings using scanning electron microscopy (SEM) and X-ray diffraction (XRD). 

• Study the batch preparation, pressing, and sintering parameters for the anode substrate fabrication. 

• Examine the anode microstructure by SEM. 

• Study the liquid-phase sintering of YSZ dense electrolyte and porous substrate using a sintering aid. 

Accomplishments 

• Obtained dense and thin YSZ electrolyte coatings on porous cathode substrates. 

• Identified the effects of the processing parameters of the die-pressing and sintering technique on the porosity of the 

anode substrates. 

• Acquired dense and porous YSZ according to the requirements using liquid-phase sintering at lower 

temperatures. 

Future Direction 

• Optimize the processes to prepare electrodes and electrolyte layers. 

• Fabricate single SOFCs and conduct performance tests. 

• Develop materials for electrolyte with higher conductivities than YSZ. 

89 
FY 2005 Progress Report 

Introduction 

Low-cost production of the high-quality 

components for SOFCs by various techniques is of 

great importance and is the focus of this research 

project. Techniques that could help to reduce the 

overall cost of the fuel cell are identifying costeffective, relatively low-purity raw materials; mass 

production of cell components; use of inexpensive 

and simplified fabrication processes; and enhanced 

cell life.1–5 The future of SOFCs greatly depends on 

the outcome of these techniques. Currently, the cost 

of the fuel cell is mainly influenced by the number 

of processing steps during the fabrication process, 

apart from the direct cost of raw materials. 

Generally, the SOFC is fabricated in either one of 

the two most common designs, tubular or planar 

cell. Electrode-supported planar cells are the most 

popular and commercially efficient cell design, due 

to their high power densities.6

 There are various 

steps at each stage of the fuel cell fabrication: 

developing the anode, depositing electrolyte layer, 

sintering the cell, etc. Each step consumes enormous 

time and cost. This presents a major hindrance in the 

commercialization of the fuel cell. The cost and time 

involved in the fabrication can be reduced a lot by 

careful selection of process techniques and 

parameters. In this project, three different costeffective techniques were used for producing SOFC 

components. The results of the experimental studies 

of each technique will be reported. 

Electrophoretic deposition (EPD) technique is 

especially attractive for SOFC application owing to 

simplicity, low-cost equipment, and ease of 

deposition of films of controlled thickness on 

substrates with even complex shape. EPD is a 

process in which the charged particles in an 

appropriate solvent move under the influence of the 

direct electric field toward an oppositely charged 

electrode and coagulate thereon to form a stable 

deposit. In the past research and applications, it was 

used to prepare ceramic laminates7,8

 and components 

in electrochemical cells,9

 etc. It was recently 

introduced to the area of the manufacturing SOFCs, 

since thin film electrolytes are highly desired for the 

fuel cells to work at reduced temperatures and 

higher efficiencies. 

Optimization of the processing parameters by 

having an experimental design called a factorial 

design for obtaining data with a minimum 

expenditure of time and resources and by meticulous 

Heavy Vehicle Propulsion Materials 

analysis of the data, called analysis of variance 

(ANOVA) is the key to the low-cost production of 

the SOFC components. The analysis provides 

detailed insight into the direct effect of process 

variables on part’s porosity as well as into the 

intensity of the numerous interactions between 

variables. This kind of response surface 

methodology (RSM) helps to determine, on the basis 

of one experiment, where to move in the next 

experiment toward the optimal point on the 

underlying response surface10. The optimized results 

will help to reduce cost and time involved in the 

fabrication process. In this research, the factorial 

design was used for the production of anode 

substrates with optimized porosities. 

High energy expense in the sintering of the fuel 

cell components, especially the YSZ electrolyte, is 

one of the major causes for the cost. Usually the 

sintering of the YSZ powders needs a temperature 

higher than 1500°C for many hours to obtain a fully 

densified electrolyte. Reduction of the sintering 

temperature will save the energy for sintering and 

also save the investment on expensive hightemperature furnaces. Previous research has already 

shown that YSZ can be sintered at a substantially 

lower temperature by adding a small fraction of low 

melting-point oxides, for example, bismuth oxide. 

Bismuth oxide itself is a good oxygen ion conductor. 

However, the purpose of the research at this time is 

to achieve desired microstructures with this 

technique for fuel cell components. The effect of the 

sintering aid addition on the conductivity of the 

electrolyte materials will be addressed in the 

ongoing research. 

Approach 

For the EPD technique, the electrolyte coatings 

were produced by depositing the charged YSZ 

particles in a powder suspension under a direct 

current (dc) electric field. The suspensions were 

prepared by dispersing the powders in the mixture of 

two organic solvents, that is, acetylacetone (acac) 

and ethanol at different ratios ranging from pure 

acac to pure ethanol. The powder concentrations 

used were 5, 10, 25, and 40g/L. All the suspensions 

were subjected to ball milling for 48 h and ultrasonic 

agitation for 0.5 h before use. Stainless steel with the 

size of 15 × 15 × 0.3-mm substrates was used for all 

the depositions for parametric studies. The coatings 

were also applied on porous cathode substrates 

90
Heavy Vehicle Propulsion Materials 

which were made of strontium-doped lanthanum 

manganite (LSM) produced in our laboratory by 

slurry casting. The coated specimens were dried at 

room temperature for more than 20 h and then 

sintered to obtain a dense electrolyte layer. 

In this study, NiO and YSZ, the most preferred 

anode material for SOFCs, were used as the raw 

materials along with the pore former graphite. A full 

24

 factorial design was used. The four variables that 

were considered for the experimental design are 

graphite (vol %), ratio of NiO to YSZ, compaction 

pressure, and sintering temperature. The two levels 

for each variable are listed in Table 1. A total of 16 

sets of experiments with two replicates for each set 

(32 total experiments) were conducted. Powders are 

weighed according to the design. They were mixed 

by ball milling along with ethanol for 48 h and then 

dried. The powder mixture was uniaxially pressed. 

The green samples were sintered in air for 1 h. The 

dimensions and weight of the sintered samples were 

measured using the laboratory calipers and digital 

balance, respectively. The measured dimensions 

were used for the porosity calculations. 

For the liquid-phase sintering study, 4.5 wt % 

bismuth oxide powders were added into YSZ 

powder. The green form of the YSZ laminates was 

prepared by slurry casting. The slurry was prepared 

by mixing the powders with polymeric binder, 

dispersant, plasticizer, and ethanol. To investigate 

the effect of the liquid-phase sintering on the 

densification of the YSZ particles with different 

sizes, the green forms with –325 mesh YSZ power 

were cast and dried firstly. A thin layer comprising 

the suspension of the mixed YSZ powders of 

66.7 wt % of –325 mesh, 23.1 wt % of 56 nm and 

13.2 wt % of 20 nm in ethanol solution stabilized 

with poly(dimethyldiallylammonium chloride) 

(PDDA) was sprayed onto the dried green form YSZ 

substrate. The coated specimens were sintered at the 

Table 1. Level design for the experimental parameters 

Independent variable 

Treatment level 

Level 0 Level 1 

Graphite (vol %) 

YSZ/NiO weight ratio 

Compaction pressure (MPa) 

Sintering temperature (°C) 

40 

50:50 

150 

1350 

30 

42.5:57.5 

125 

1200 

FY 2005 Progress Report 

maximum temperature of 1000°C in air. Figure 1 

shows the programmed sintering cycle for the 

liquid-phase sintering. The microstructure of the 

sintered specimens was examined with scanning 

electron microscopy (SEM). 

1000 

800 

o

Temperature ( C) 

600 

400 

200 

0 

0 200 4 0 0 6 00 800 1000 1 2 00 1400 

T im e ( m in ) 

Figure 1. Sintering cycle for liquid-phase sintering of 

YSZ specimens. 

Results 

Effect of the Solvent on the Preparation of YSZ 

Coatings by EPD 

To produce dense, homogeneous green forms of 

the coatings with EPD, a dilute, well-dispersed, and 

stable colloidal suspension is of great importance. In 

this case the stable suspension was obtained by 

absorbing free protons on the surface of the 

particles, while the free protons were produced by 

the reaction between acac and ethanol. Figure 2 

shows the weight gains as a function of the fraction 

of ethanol in its mixture with acac. It can be noticed 

that deposits were obtained in all the different 

mixtures. The weight gain in pure acac is obviously 

much higher than in any other mixtures in constant 

current deposition mode (curve a). However, this 

advantage is much less in constant voltage mode 

(curve b). If we examine the voltage required for 

pure acac in Figure 3, it is ten times more than the 

lowest one. We can suggest that the amount of the 

free proton is much less in pure acac than in any 

other mixtures and even in pure ethanol. High 

electrical energy input led to the high yield. At the 

constant voltage mode, even though the input power 

in pure acac is the lowest, the deposition yield was 

still the highest one. In constant voltage mode, the 

power inputs in all the mixtures are higher than 

either in pure acac and ethanol. Nevertheless, the 

deposition yields were gradually decreased with the 

91 
FY 2005 Progress Report Weight Gain (mg/cm2) 

18

16

14

12

10 

(b)

(a)

8

6

4

2

0

-2 

0 20 40 60 80 100 

Fraction of ethanol (vol%) 

Figure 2. Weight gains as functions of the volume 

fraction of ethanol in its mixture with acac. 

Depositions were carried out at (a) a 

constant current density of 0.20 mA/cm2 

and (b) a constant voltage of 50 V. 

Deposition time was 120 s with a powder 

concentration of 5 g/L. 

Figure 3. Current and voltage as functions of the mixing 

ratios of ethanol and acac during the 120-s 

depositions. (a) Maximum voltage required to 

maintain a constant current density of 

0.20 mA/cm2

, and (b) minimum current 

density when a constant voltage of 50 V was 

applied. 

Heavy Vehicle Propulsion Materials 

increase of the fraction of ethanol in the mixture. It 

is recommended that in these mixtures, there are free 

protons in addition to those absorbed by the YSZ 

particles. The moving of these free protons partially 

contribute to the total electric current; the deposition 

rates were hence lower than the expectations. 

The microstructures were examined with SEM. 

Figure 4 shows the micrographs taken from the 

surfaces of the deposited green forms at the constant 

voltage of 50 V. Figure 4(a) and (c) show poorly 

packed green forms from the suspensions with pure 

acac and pure ethanol. The poor packing resulted 

from the large agglomerates in the suspensions. This 

means that in the solvents of either pure acac or pure 

ethanol, the YSZ particles were not well dispersed 

because of the low free proton levels. Figure 4(b) 

shows a homogeneous packed deposit in the 

suspension with 45 vol % ethanol/55 vol % acac as 

the solvent. The particle sizes are much smaller than 

in the other samples. Therefore, the deposition 

behavior was further studied for suspension in the 

solvent with the mixing ratio of ethanol and 

acac 1:1. 

Effect of the Parameters in the Factorial 

Experiments 

The obtained porosity values after sintering are 

shown in Table 2, which will be used as the inputs 

for the ANOVA subsequently. 

The effects of the processing parameter on the 

porosity of the sintered anode pellets were analyzed 

by the ANOVA method. It was understood that the 

porosity of the anode increased slightly with the 

increasing of the NiO component in the mixture, and 

the increasing rates were in accordance with the 

level of the pressure used to make the pellets. The 

sintering temperature was found to be one of the 

very significant factors in the porosity control. The 

Figure 4. Micrographs of the dried green forms deposited on stainless steel substrates at a constant voltage of 

50 V for 120 s. The depositions were conducted in the suspensions of 5 g/L YSZ in solvents of 

(a) pure acac, (b) 45 vol % ethanol/55 vol % acac, and (c) pure ethanol for 120 s. 

92
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Table 2. Porosity values of sintered NiO/YSZ samples 

Design No. 1 2 3 

Porosity (%) 

Set 1 

Set 2 

6.85 

7.09 

40.62 

41.80 

8.00 

8.39 

Design No. 9 10 11 

Porosity (%) 

Set 1 

Set 2 

9.58 

9.75 

34.38 

33.68 

9.38 

10.23 

porosity of the anode sintered at 1200°C was four 

times more than that sintered at 1350°C. The effect 

of the amount of graphite used as the pore former 

was shown to be significant, too. The porosity of the 

anode increased as the graphite amount increased 

within the experimental range. 

Effect of the Sintering Temperature in LiquidPhase Sintering 

A comparison of the effect of the sintering 

temperatures on the densification of YSZ was made. 

4 

41.41 

41.41 

5 

10.85 

10.47 

6 

42.94 

42.57 

7 

10.68 

10.03 

8 

40.38 

40.92 

12 

33.80 

34.13 

13 

11.16 

9.68 

14 

35.15 

35.08 

15 

9.31 

9.69 

16 

33.83 

35.77 

Figure 5 shows the micrographs of the specimens 

sintering at 1000°C and 1100°C, respectively. From 

images in Figure 5(a) and (b) that were magnified to 

20,000 times, it seems that both specimens that 

sintered at 1000 and 1100°C are fully sintered. From 

the images in Figure 2(c) and (d), which were taken 

at a lower magnification from the same specimens 

mentioned above, it is shown that there are still 

many pores over the samples. This set of images 

demonstrates that the densification at 1000°C and 

1100°C did not make significant difference. 

(a) (b) 

(c) (d) 

Figure 5. Micrographs of liquid-phase sintered YSZ with addition of 4.5 wt % of bismuth oxide sintered at 

different temperatures, (a) and (c) 1000°C; (b) and (d) 1100°C for 4 h. 

93 
FY 2005 Progress Report 

Therefore, a temperature of 1000°C was chosen for 

all the sintering thereafter. On the other hand, this 

experiment revealed a potential to produce porous 

YSZ frame structure at lowered temperatures using 

liquid-phase sintering. 

Effect of the YSZ Particle Size in Liquid-Phase 

Sintering 

As is stated above, sintering experiments were 

conducted for the samples with substrate and coating 

layered structure. The substrate comprises YSZ 

particles of a larger size (–325 mash), while the 

coating consists of YSZ particles of mixed sizes. 

The purpose of doing it this way is to explore the 

possibility of producing a dense YSZ layer on a 

porous YSZ substrate using one liquid-phase sintering process. The knowledge of solid-state sintering 

advises that small particles have higher densification 

rates than the large particles during sintering 

because of two main reasons. First, large particles 

have lower packing density than the small particles. 

The material comprised with large particles needs 

more diffusion energy and time to reach densification than one with small particles. Secondly, small 

particles, especially nano particles, have higher 

chemical potentials than large particles. As a result, 

the material with small particles possesses higher 

diffusion rates than one with large particles. This 

will further enlarge the difference of the densification during sintering between materials with different particles sizes. These factors will also take effect 

in liquid-phase sintering. 

Heavy Vehicle Propulsion Materials 

Figure 6 shows the SEM images taken from the 

top and bottom surfaces of a coated specimen after 

liquid-phase sintering at 1000°C for 4 h. This observation proved that it is effective to produce dense 

coating on porous substrate by using different particle sizes in the two different parts of the structure 

during liquid-phase sintering. This method can be 

employed in production of SOFC sandwich structures, which will reduce the manufacturing cost 

substantially by simplifying the fabrication process 

and saving sintering energy. 

Conclusions 

Three cost-effective techniques used to produce 

SOFC components were investigated. In the EPD for 

preparation of YSZ electrolyte coatings, microstructure suggested that deposits obtained in the suspensions with ratio of ethanol to acac at about 1:1 outcome homogeneously packed green form that forms 

dense coatings after sintering even through depositions can be achieved in any of the mixing compositions. For the anode factorial experiments, the 

effects of these parameters were examined against 

the porosities of the sintered anode samples. The 

sintering temperature and graphite usage were found 

to have very significant effects, while the ratio of 

NiO/YSZ and the pressing pressure were found to 

have less significant effects on the porosity of the 

anodes. The effects are planned to be further analyzed by ANOVA. Experiments demonstrated that 

liquid-phase sintering is an energy-efficient process, 

and Bi2O3 is an effective low melting point material 

to perform liquid-phase sintering of YSZ. 

(a) (b) 

Figure 6. Micrographs of the (a) coating and (b) substrate sintered at 1000°C for 4 h in liquid phase in one process. 

94
Heavy Vehicle Propulsion Materials 

References 

1. S. C. Singhal, “Status of solid oxide fuel cell 

technology,” In Proceedings of the 17th Risø 

International Symposium on Materials Science: 

High Temperature Electrochemistry: Ceramics and 

Metals,” eds. F. W. Poulsen, N. Bonanos, 

S. Linderoth, M. Mogensen, and B. ZachauChristiansen. Risø National Laboratory, Denmark, 

123–138 (1996). 

2. F. Tietz, F. J. Dias, D. Simwonis, and 

D. Stöver, “Evaluation of commercial nickel oxide 

powders for components in solid oxide fuel cells,” 

Journal of the European Ceramic Society, 20(8), 

1023–1034 (2000). 

3. G. Rietveld, P. Nammensma, and J. P. 

Ouweltjes, “Status of SOFC component 

development at ECN,” In Proceedings of the 

Seventh International Symposium on Solid Oxide 

Cells (SOFC-VII), eds. H. Yokokawa and S. C. 

Singhal. The Electrochemical Society Proceedings 

Series (PV 2001-16), Pennington, New Jersey, 

125–130 (2001). 

4. R. N. Basu, C. A. Randall, and M. J. Mayo, 

“Fabrication of dense zirconia electrolyte films for 

tubular solid oxide fuel cells by electrophoretic 

deposition,” Journal of the American Ceramic 

Society, 84, 33–40 (2001). 

5. T. Ishihara, T. Shibayama, H. Nishiguchi, 

and Y. Takita, “Nickel–Gd-doped CeO2 cermet 

anode for intermediate temperature operating solid 

oxide fuel cells using LaGaO3-based perovskite 

electrolyte,” Solid State Ionics, 132, 209–216 

(2000). 

6. J. W. Kim, A. V. Virkar, K. Z. Fung, 

K. Mehta, and S. C. Singhal, “Polarization effects in 

intermediate temperature, anode supported solid 

oxide fuel cells,” Journal of the Electrochemical 

Society, 146, 69–78 (1999). 

FY 2005 Progress Report 

7. P. Sarkar, X. Huang, and P. Nicholson, 

J. Am. Ceram. Soc., 76, 1055 (1993). 

8. P. S. Nicholson, P. Sarkar, and X. Haung, 

J. Mater. Sci., 28, 6274 (1993). 

9. C. E. Baumgartner, V. J. Decarlo, P. G. 

Glugla, and J. Grimaldi, J. Electrochem. Soc., 132, 

57 (1985). 

10. C. R. Hicks and K. V. Turner, Jr., 

“Fundamental concepts in the design of 

experiments,” Oxford University Press, 5th edition 

(1999). 

Publications and Presentations 

“Property Control of Cathodes and Anodes 

Produced by Slip Casting for Planar Solid Oxide 

Fuel Cells,” Ceramic Engineering and Science 

Proceedings, 26(4), 185–90 (2005). 

“Influence of Processing Parameters on Porosity 

of NiO-YSZ Solid Oxide Fuel Cell Anode 

Material,” Ceramic Engineering and Science 

Proceedings, 26(4), 177–83 (2005). 

“Exploration of Combustion CVD Method for 

YSZ Thin Film Electrolyte of Solid Oxide Fuel 

Cells,” accepted for publication in Mixed Ionic 

Electronic Conducting Perovskites for Advanced 

Energy Systems, eds. N. Orlovskaya and 

N. Browning, Kluwer Academic Publication, 2005. 

“Exploration of Electrophoretic Deposition of 

YSZ Electrolyte for Solid Oxide Fuel Cells,” K8.2, 

MRS Proceedings, Volume 835, eds. P. Knauth, 

C. Masquelier, E. Traversa, and E. D. Wachsman, 

December 2004. 

95 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

E. NDE Development for Valve-Train Components for Diesel Engines 

Jiangang Sun 

Argonne National Laboratory 

9700 South Cass Avenue 

Argonne, IL 60439 

630-252-5169; fax: 630-252-4798; e-mail: sun@anl.gov 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: Argonne National Laboratory, Argonne, Illinois 

Objectives 

• Develop laser-scattering and other nondestructive evaluation (NDE) methods for detection and characterization 

of surface and subsurface defect/damage in diesel-engine valve-train components made from advanced 

materials such as ceramics and intermetallics. 

• Develop quantitative NDE analysis tools to identify strength-limiting flaws and failure mechanisms, evaluate 

damage evolution and growth, and predict mechanical properties for advanced materials. 

Approach 

• Develop fast and reliable automated laser-scattering NDE systems for surface/subsurface inspection of ceramic 

and intermetallic valves. 

• Inspect valve surfaces to determine the level of machining damage and accumulated damages from bench or 

engine tests. 

• Correlate NDE data with microstructure and strength of silicon-nitride specimens for prediction of fracture 

initiation defects and mechanical properties. 

Accomplishments 

• Successfully implemented a photomultiply tube (PMT) detector coupled with an optical fiber in a NDE valvescan system. The PMT detector is more sensitive and more than six times faster in scan speed so it allows for 

fast NDE inspection of a large number of valves to meet the engine test schedule. 

• Evaluated 25 finish-machined SN235P silicon nitride valves. Verified significant improvement in machined 

valve surfaces especially at keeper groove where a severe machining damage had caused a premature valve 

failure in a rig test. 

• Inspected machined surfaces of 11 as-processed TiAl valves. 

• Completed analysis and correlation of NDE data with microstructure and strength for cylindrical SN235P 

specimens machined at various conditions. 

97 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Future Direction 

• Establish laser-scatter NDE detection sensitivity and characterization capability for surface damage in TiAl 

intermetallic specimens. 

• Complete NDE evaluation for damage accumulation in prototype SN235P and TiAl valves by periodic

inspections of these valves during planned engine duration tests. 

• Develop quantitative image-processing methods for prediction of strength-limiting subsurface defects/damage 

in silicon nitride ceramics. 

Introduction 

Advanced ceramics and intermetallics are leading candidates for high-temperature engine applications that offer improved fuel efficiency and engine 

performance. Among them, silicon nitrides (Si3N4) 

and titanium aluminide (TiAl) are being evaluated as 

valve-train materials for diesel and natural gas 

engines because of their lighter weight and high 

strength and corrosion resistance at elevated temperatures. However, these materials are brittle so 

microscopic defect/damage near the surface may 

significantly degrade their mechanical properties. 

Material damage in the subsurface may be induced 

by machining or from service due to impact, wear, 

and corrosion. For silicon nitrides, material strength 

may also be limited by inherent defects such as 

voids and porosities within subsurface. To detect 

and characterize these defects, Argonne National 

Laboratory (ANL) developed a laser scattering NDE 

method to measure detailed surface/subsurface 

microstructure for these materials. The objective of 

this research is to demonstrate that this method can 

be used to assess/evaluate the cost-effectiveness and 

reliability of valve-train components manufactured 

from advanced ceramic and intermetallic materials 

for diesel engines. The primary effort in FY 2005 

was to develop a fast NDE valve-scan system and to 

inspect finish-machined SN235P silicon nitride 

valves to determine the improvement in machined 

surface quality from their initial rough machining 

conducted in FY 2004. Another effort was directed 

on analysis and correlation of NDE data with microstructure and strength for cylindrical SN235P 

specimens machined at various conditions. This 

research is collaborated with Caterpillar, Inc. 

Approach 

The critical region for brittle ceramic and intermetallic components in structural applications is 

near surface. The common types of defects in this 

region are mechanical, such as cracks, spalls, inclusions, and voids. The size of the defects that limits 

the component strength is generally small, say 

~100 μm. To detect these defects, a laser-scattering 

method based on cross-polarization detection of 

optical scattering originated from surface and subsurface microstructure can be used for noncontact, 

NDE of these materials. Because ceramics are partially translucent to light, optical penetration can 

reach the subsurface to directly interact with the 

subsurface defects. For metallic or intermetallic 

materials, optical penetration is not possible. 

However, using the cross-polarization detection 

approach, optical scattering from rough surface or 

open cracks can be significantly enhanced against 

the smoother background. By scanning the entire 

surface (flat or curved) of a component and 

constructing a two-dimensional (2-D) scatter image, 

surface/subsurface defects can be readily identified 

as they exhibit excessive scattering over the 

background, and their type and severity may be 

analyzed. To apply this technology for NDE of 

Si3N4 and TiAl valves, an automated laser-scatter 

system was developed by ANL for scanning the 

entire valve surface at high spatial resolutions. This 

system utilizes two-rotation and two-translation 

stages to align and focus the laser beam on the valve 

surface during the scan, and the resulting 2-D 

scattering image data are used to identify the 

location, size, and relative severity of subsurface 

defects/damage. Analysis of the NDE data can 

determine the statistical properties of the component 

surface, which can be correlated with machining or 

service conditions. Development of quantitative 

image-processing methods may lead to automated 

identification of strength-limiting defect/ 

damage for prediction of mechanical properties. 

98
Heavy Vehicle Propulsion Materials 

Results 

Development of a Fast Valve-Scan NDE System 

Prototype silicon nitride and titanium aluminide 

valves have been planned for engine duration tests in 

FY 2006. Because of a large number (24) of valves 

to be used and a tight schedule for intermittent valve 

evaluation, the NDE system must be able to scan 

each valve at a fast speed to meet the engine test 

schedule. The previous valve-scan system was limited by the speed of a silicon detector; a photomultiplier tube (PMT) was therefore selected to replace 

the silicon detector for detection of back-scattered 

signal. The PMT detector has a data acquisition 

speed up to 20 kHz with 16-bit dynamic range, 

which is more than six times faster and more than 

five times higher in detection sensitivity than the 

silicon detector. 

Figure 1 is a photograph of the new automated 

laser-scatter system. The system utilizes a multimode optical fiber to collect and deliver the backscattered light to the PMT detector. A long translation stage was used to allow for scanning of entire 

valve surface (head and stem) within one pass. In 

addition, data acquisition software was modified to 

improve synchronization between data acquisition 

and stage motion along the valve profile. The 

improved system can scan the complex valve-head 

surface in an hour. 

Figure 1. Photograph of automated laser-scattering NDE 

system. 

Evaluation of Finish-Machined SN235P Valves 

Twenty-five finish-machined SN235P silicon 

nitride valves were evaluated by the automated 

laser-scatter system. They are natural-gas intake and 

FY 2005 Progress Report 

exhaust valves to be used by Caterpillar Inc. in an 

engine duration test. These valves were initially 

machined in FY 2004. Because of a process error, 

most of the valve head and keeper-groove surfaces 

were rough machined, resulting in significant 

machining damage on these surfaces.1

 NDE data of 

the finish-machined valves were compared with 

those of initial machined to verify improvement in 

machined surfaces. 

Figures 2 and 3 show, respectively, laser-scatter 

images of SN235P natural-gas intake valve No. 8 at 

before and after the finish machining. Spatial resolution in both scan directions was ~10 μm. In 

Figure 2, significant machining damage can be 

found on the entire valve head surface except the 

stem surface near the bottom of the image. These 

damages are seen as bright lines (of higher scattering 

intensities) along the machining direction, as shown 

in the detailed image in Figure 4(a). The lengths of 

the machining damage are typically near 1 mm. In 

Figure 2. Laser-scatter scan image of SN235P valve 8 

after initial rough machining. 

Figure 3. Laser-scatter scan image of SN235P valve 8 

after finish machining. 

Figure 4. Detailed laser-scatter images in contact 

surface of SN235P valve 8 after (a) initial 

rough machining and (b) finish machining. 

99 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

addition, the overall optical scatter intensity of the 

valve-head surface is high, indicating a higher level 

of overall damage. In comparison, the finishmachined surface shows minimum machiningdamage indications as seen in Figure 3 and the 

detailed image in Figure 4(b). The higher scatterintensity spots in Figures 4(a–b) come from material 

inherent defects such as porosities with sizes typically smaller than 50 μm. Therefore, it is apparent 

that the finish machining essentially removed all 

machining damages induced in the initial rough 

machining of these SN235P valves. 

Laser scattering scan was also conducted in the 

keeper groove region of the SN235P valves. Before 

the finish machining, a severe machining damage 

was found around the entire circumference of the 

keeper-groove edge, which caused premature failure 

of a SN235P valve in a rig test.2

 In the finishmachined valves, this damage was also removed. 

However, because the edges are still sharp, some 

chip damage can be detected along the keepergroove edges. Figure 5 shows detected edge chips in 

a laser-scatter image and photomicrographs. 

Figure 6. Laser-scatter scan image of an as-processed 

TiAl valve. 

scan directions was ~10 μm. Because intermetallic 

surface is highly reflective to light and the surface 

mismatch during the scan can cause gross change in 

measured scatter intensity, especially at regions with 

sharp curvatures as seen in Figure 6. The enlarged 

image in Figure 6 shows machining marks; the 

nature of these marks will be further investigated. 

Laser scatter scans were also performed for the 

joint region between the TiAl stem and the titanium 

alloy (Ti-6V-4Al) stem. Figure 7 presents a typical 

scan image of the joint surface, which shows clearly 

the joint line and no apparent defect. The image also 

shows axial stripes along the machining direction in 

this region; the nature of these stripes will be further 

studied. 

Figure 5. Detailed laser-scatter image and 

photomicrographs of chip flaws at keeper 

groove edge of a SN235P valve. 

Inspection of As-Processed TiAl Valves 

Eleven as-processed TiAl intermetallic valves 

were inspected by the valve-scan NDE system. 

Figure 6 shows a laser-scatter image of a typical 

TiAl valve-head surface. Spatial resolution in both 

Figure 7. Laser-scatter scan image of joint surface of a 

TiAl valve. 

Correlation of NDE Data with Microstructure 

and Strength 

Three sets of 32 cylindrical SN235P silicon 

nitride specimens machined by two vendors at various conditions were initially examined by laser 

scatter in FY 2004. After fracture test, eight specimens were further studied for correlation of the 

NDE data with material microstructure and strength. 

In the microstructure study, fractographic examinations were performed to identify the fracture sites on 

the fracture surfaces. These fracture sites were then 

mapped into the tensile (or cylinder) surface to 

locate the fracture initiation positions. As expected, 

each fracture initiation position corresponds to a 

defect detected in the laser scatter NDE data. The 

results for all eight SN235P specimens are summarized in Figure 8, which shows the fracture strength 

of each specimen and its corresponding fracture 

100
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 8. (a) Fracture strength and (b) corresponding 

fracture origins in scattering images (image 

size is 0.5 mm × 0.5 mm) of eight ground 

rods. 

origin in the scanning image. It is clear that all 

specimens fractured from inherent material flaws 

(high-porosity defects), except specimen 1, which 

fractured from a machining damage (likely a lateral 

machining crack). The lowest fracture strength 

occurred in specimens 2 and 5. For specimen 2, 

there are a group of material flaws along the fracture 

line, which may explain its low fracture strength. As 

for specimen 5, its fracture origin is more intensive 

(likely deeper flaw) and slightly bigger (~35 μm) 

than those of the other specimens (~25 μm); therefore, its strength is lower. 

Statistical parameters of laser scatter data were 

correlated with fracture strength. These statistical 

parameters include standard deviation, skewness, 

skewness/standard deviation, and coefficient of variance Cv (standard deviation/mean). Figure 9 shows 

fracture strength as functions of statistical parameters for the eight specimens. The standard deviation 

and skewness in Figures 9(a–b) generally correlate 

well with the fracture strength. Both show a similar 

trend, that is, strength decreases with the increase of 

standard deviation or skewness. These results are 

physically reasonable. Because standard deviation 

represents the second-order and skewness the thirdorder difference of local scatter intensity from the 

Figure 9. Correlation of fracture strength with statistical 

parameters processed from laser-scatter data 

around fracture regions. 

background, a higher standard deviation or skewness 

indicates the presence of severe subsurface damages/ 

defects that generate higher scattered light intensity 

above the background. Note that the nondimensional 

parameters skewness/standard-deviation and Cv in 

Figures 9(c–d) also show the same trend. These 

nondimensional parameters can be good candidates 

for predicting the fracture strength. 

Conclusions 

A high-speed automated laser-scatter system 

was developed. It utilizes a PMT detector coupled 

with an optical fiber for collecting and delivering 

scatter signal. The PMT detector is more sensitive 

and more than six times faster in data acquisition 

than a silicon detector. With this new system, scanning of a complex valve-head surface will be 

reduced to about an hour, which allows for fast NDE 

inspection of large number of valves to be evaluated 

in engine duration tests scheduled for FY 2006. 

Twenty-five finish-machined SN235P silicon 

nitride valves and eleven as-processed TiAl valves 

were inspected by the valve-scan system. The 

SN235P valves had been rough-machined in 

FY 2004, which induced significant machining 

damage on the valve surface. In particular, a severe 

damage at keeper groove caused premature failure of 

a valve in a rig test. Laser-scatter examination verified that there is a significant improvement in the 

surface quality of these finish-machined valves, that 

is, most rough-machining damages were removed 

although minor chip damages were detected on the 

edges of keeper grooves. For the as-processed TiAl 

valves, because intermetallic surface is highly 

101 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

reflective to light and the surface scatter is weak, it References 

was challenging to align the laser-scatter system 

with the valve profile during the entire scan. Nevertheless, laser scatter images were obtained for the 

valve-head and stem-joint surfaces for these valves, 

and machining marks and material variation can be 

readily observed. However, correlation of NDE data 

with surface microstructure for TiAl materials need 

to be further investigated. 

Laser scatter measurements were conducted for 

three sets of 32 cylindrical SN235P silicon nitride 

specimens that were machined at different conditions. After fracture test, eight of these specimens 

were delivered to ANL for correlation of NDE data 

with microstructure and strength. In microstructure 

study, fractographic examinations were performed to 

identify the fracture sites that were then correlated 

with NDE data to locate fracture initiation defects on 

tensile surfaces of the specimens. The correlation 

allowed for direct determination of the size, severity, 

and type of defect/damage that caused the fracture. 

The fracture strength was also correlated with statistical parameters of laser scatter NDE data. These 

statistical parameters included standard deviation, 

skewness, skewness/standard deviation, and coefficient of variance (standard deviation/mean). It was 

found that strength decreases as these parameters 

increase; which is physically meaningful because 

these parameters become larger when the number or 

severity of defects increases. 

1. J. G. Sun, “NDE Development for Ceramic 

Valves for Diesel Engines,” Annual Report for 

FY 2004, Heavy Vehicle Propulsion Materials 

Program. 

2. J. S. Trethewey, “High-Temperature 

Advanced Materials for Lightweight Valve Train 

Components,” Annual Report for FY 2004, Heavy 

Vehicle Propulsion Materials Program. 

Presentations and Publications 

J. G. Sun, J. M. Zhang, J. S. Trethewey, J. A. 

Grassi, M. P. Fletcher, and M. J. Andrews, 

“Nondestructive Evaluation of Machining and 

Bench-Test Damage in Silicon-Nitride Ceramic 

Valves,” in Ceramic Eng. Sci. Proc., eds. D. Zhu 

and W. M. Kriven, 26(2), 127–132 (2005), and 

presented at the 29th International Cocoa Beach 

Conference and Exposition on Advanced Ceramics 

& Composites, Cocoa Beach, Florida, 

January 23–28, 2005. 

J. M. Zhang, J. G. Sun, M. J. Andrews, 

A. Ramesh, J. S. Trethewey, and D. M. Longanbach, 

“Characterization of Subsurface Defects in Ceramic 

Rods by Laser Scattering and Fractography,” paper 

accepted for publication in the Proceeding of the 

32nd Annual Review of Progress in Quantitative 

Nondestructive Evaluation, Bowdoin College, 

Brunswick, Maine, July 31–August 5, 2005. 

102
________________________ 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

F. Durability of Diesel Engine Component Materials 

Peter J. Blau, Jun Qu, and John J. Truhan, Jr.* 

Oak Ridge National Laboratory 

P.O. Box 2008, Mail Stop 6063 

Oak Ridge, TN 37831-6063 

(865) 574-537; fax: (865) 574-6918; e-mail: blaupj@ornl.gov 

*University of Tennessee, Knoxville 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: University of Tennessee, Knoxville, TN 

Objectives 

• Provide test data, analyses, and models that enable the use of durable, lower-friction moving parts in diesel 

engines in heavy-vehicle propulsion systems. 

• Develop test methods that simulate the environment of scuffing-prone engine parts. 

• Conduct structural and tribological characterizations of promising new materials, surface treatments, 

composites, and coatings. 

Approach 

• Identify and focus on diesel engine components that require durability and low-friction behavior and might 

benefit from advanced materials. These include (a) wastegate bushings for exhaust gas recirculation (EGR) 

components and (b) fuel injector components. 

• Identify materials, coatings, and/or surface treatments that have the potential to increase the durability of the 

selected engine components. 

• Develop test methods to evaluate and quantify the performance of candidate materials under simulated use 

conditions. 

• Develop graphical methods and models to portray the effects of operating parameters, like speed, load, and 

surface finish, on the scuffing response of the materials. 

Accomplishments 

• Designed, built, and used a high-temperature oscillatory scuffing test system that operates at wastegate bushing 

temperatures (~600–700o

C). 

• Published results of tests on a range of metallic alloys, ceramics, and coatings to determine which of these had 

the best durability under high-temperature conditions. 

• Developed a novel, ‘pin-on-twin’ scuffing test to evaluate fuel injector materials in diesel fuel and low-sulfur 

fuel environments. 

103 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

• Developed criteria for the onset of scuffing damage and evaluated traditional steel fuel injector materials as 

well as ceramics, advanced cermets, and hard coatings in diesel fuel and low-sulfur fuel. Represented results in 

terms of ‘scuffing maps’ and transition diagrams. 

• Developed a model for scuffing tendency that considers lubricant characteristics and solid material

characteristics. Prepared a final report that explains the model and its rationale. 

Future Direction 

• The focus on fuel systems and exhaust gas recirculation (EGR) components ended in FY 2005. It will be 

succeeded by a new effort that focuses on valve train durability issues. 

Introduction 

The diesel engine industry continues to face the 

important challenge of improving fuel efficiency 

while meeting increasingly strict emissions regulations. These challenges are being addressed by 

modifying engine designs and control systems and 

by developing exhaust gas after-treatments. Such 

modifications affect the mechanical, thermal, and 

chemical environments in which the engine materials must operate, and currently used materials may 

fall short of meeting these needs. New materials are 

needed to enable new technological goals. 

The objective of this effort is to enable the 

selection and use of durable, lower-friction moving 

parts in diesel engines for heavy-vehicle propulsion 

systems through the systematic evaluation of promising new materials, surface treatments, composites, 

and coating technologies. The current approach 

involves developing test methods, analyzing microstructures of candidate materials, developing design 

maps with variables like surface finish and frictional 

behavior, and modeling the damage process itself. 

The focus on exhaust gas recirculation (EGR) components and fuel injector plungers was based on discussions with diesel engine manufacturers. Before 

developing tests to evaluate materials for improved 

durability, it was necessary to conduct a tribosystem 

analysis to understand the conditions under which 

the surfaces of these components operate in a diesel 

engine. The nature of contact damage on engine 

components was reviewed to ensure that laboratory 

test methods would adequately reproduce that kind 

of damage. Then suitable tests were developed, surface damage and friction data were analyzed, and 

modeling tasks were undertaken. 

Approach 

In FY 2001, and based on the definition of several key durability problems, a test method and 

apparatus were developed to study the hightemperature friction and wear characteristics of candidate EGR system materials. That methodology has 

seen continued use for evaluating metal alloys, 

ceramics, coatings and other experimental materials 

for moving parts in waste gate bushings. In 

FY 2002, the effort was extended to include an 

investigation of the scuffing of fuel injector component materials. Laboratory tests were developed and 

refined to produce and measure the type of finescale surface damage that is observed in fuel system 

parts. In FY 2003, research continued in two areas: 

(1) evaluating the effects of diesel fuel sulfur reductions on scuffing and (2) identifying materials and 

coatings for high-temperature scuffing resistance in 

EGR components. During FY 2004, a threedimensional ‘scuffing map’ was developed to depict 

conditions for scuffing initiation and propagation in 

time and space domains. By the end of FY 2004 a 

new scuffing model, one that integrates boundary 

film characteristics with material properties, had 

been proposed. The model recognizes the complexities of reciprocating sliding in which velocity is 

continually changing and the direction of motion 

reverses periodically. In FY 2005 the model was 

refined and improved to consider optimization of 

surface finish and the role of abrasion in reducing 

the tendency for materials to adhere to one another. 

Results 

The results presented here are an abbreviated 

version of a more detailed final report. A milestone 

104
Heavy Vehicle Propulsion Materials 

report, dated March 2004, described the reciprocating sliding experiments in which results led to the 

conceptual framework for the three-regime model 

for scuffing. That model considers the failure of a 

lubricating film and the increased deformation of the 

contact surfaces to produce progressively severe 

surface damage. It represents an interdisciplinary 

approach and integrates concepts from both lubrication theory and materials science. The model, schematically shown in Figure 1, uses three submodels 

whose applicability depends on whether the contact 

is effectively lubricated (solid surfaces that are not 

touching), boundary-lubricated, or subjected to significant solid contact and plastic deformation. The 

higher the contact pressure, the less likely the lubricant will effectively separate the moving surfaces 

and the other regimes will come into play. 

Figure 1. The scuffing model can be represented in 

three distinct regimes depending on the 

relative roles played by the liquid lubricant (I), 

thin boundary films (II), and the contacting 

solids (III). 

Conceptually, the horizontal axis, labeled 

‘Contact Pressure’ in Figure 1, can also be used to 

represent time. First, there is a period of effective 

operation whose duration depends on the nature of 

the materials, mechanical design and operating 

parameters, surface finish, and the regime of lubrication. Parts can perform very effectively in 

Regime I for millions of cycles if all goes well. 

However, the system can also operate in a boundary 

lubrication regime (II) in which some contact 

occurs. 

As components and lubricants age, or if the 

contact pressure is too large to sustain a full film 

FY 2005 Progress Report 

condition, then Regime I may either be finite in 

length or not observed at all, and the system may 

find itself in Regime II, which is governed by the 

thin boundary films that form from species initially 

present in the lubricant (or fuel). One can indirectly 

monitor the stability of the shear strength of interfacial films by observing the time-dependence of the 

friction coefficient. If the duration of contact is 

expressed in terms of numbers of oscillating cycles, 

as in a fuel injector plunger, then the transition from 

Regime I to Regime II can be expressed as follows: 

μx = μ ff + Δμ⎜

⎛ x ⎞

⎟

n 

, (1) ⎝ L ⎠

where μx = friction coefficient at x cycles, μff = 

friction coefficient for full-film lubrication, Δμ = the 

change in friction during the transition from fullfilm to boundary lubrication, x = current number of 

oscillating cycles, L = number of cycles to reach 

Regime II, and n = a constant that reflects the rate of 

change during the transition period. In Eq. (1), the 

change in friction during the transition is simply

 Δμ = μbl – μff , (2) 

where μbl = the friction coefficient for boundary 

lubrication at the start of Regime II. Assigning typical values of μff = 0.01, μbl = 0.12, and letting L = 

1,000,000 cycles, we can plot Eq. (1) for a finite life 

in Regime I for several values of n (see Figure 2). 

The greater the magnitude of the rate constant (n), 

the longer is the incubation period, but the more 

rapid the transition. 

0.20 

0.15 

μ 

x 0.10 

0.05 

0.00 

0 2 105 4 105 6 105 8 105 1 106 1.2 106 

Regime I 

n = 10 

n = 5 

μff = 0.01 

μbl = 0.12 

CYCLES 

Figure 2. Portrayal of a finite-length Regime I with 

various transition rates. 

105 
FY 2005 Progress Report 

Physically, the exponent n in Eq. (1) reflects 

such processes as a loss in lubricant load carrying 

capacity from oxidation or degradation of the additive package (in closed systems), a change in the 

contamination level in the system due to debris 

buildup or contaminant entrainment, or even phasetransformation-induced swelling of the plunger 

material to reduce bore clearance and increase surface contact. The referenced final report describes a 

second submodel for the localized changes in the 

composite surface roughness of the contact as 

scuffing initiates and propagates from the ends of 

the reciprocating stroke to its midpoint. 

A smooth, hard material may work very well 

sliding against a soft counterface as long as the 

adhesion of the latter can be mitigated by a lubricating boundary film. That can only happen if there 

is sufficient clearance in the plunger bore to enable 

the formation of such a film. 

Light abrasion can prevent the buildup of adhesive patches (transfer layers). In the absence of significant adhesion (reduced by lubrication), a factor 

that indicates the tendency for abrasion (AT) can be 

defined as 

Hh AT = S , (3) 1.2Hs 

where Hh = the scratch hardness of the harder side, 

Hs, = the scratch hardness of the softer side, and S = 

the deviation from the optimal roughness. The factor 

1.2 enters because of classic work by David Tabor 

(ca. 1951) on hardness that suggests for many materials, there must be about 20% difference in hardness 

in order for abrasion to occur. The deviation from 

the optimum surface roughness can be defined as 

m S = Rh − Ropt , (4) 

where Rh = the roughness of the harder surface, 

Ropt = the optimal roughness (determined by experiHeavy Vehicle Propulsion Materials 

The AT parameter is plotted vs the initiation 

period for various materials against 52100 steel 

(including zirconia, TiN hard coatings, Ni3Al 

bonded cermets, and 52100 steel itself) in low-sulfur 

fuel in Figure 3. A linear fit to the data, excluding 

the self-mated steel, was quite good (R = 0.931). 

The implications of the relationship may initially 

seem counterintuitive because the more abrasive 

(i.e., the higher the AT parameter), the longer the 

period for the onset of scuffing. Closer consideration 

of the physical situation, however, suggests that 

scuffing is controlled by adhesive processes, and if 

abrasion of the surface prevents the formation of 

adhesive bonds, then a small amount of abrasion can 

in effect retard the onset of scuffing. Adhesion is 

promoted by self-mated materials, like the case for 

52100 steel, and therefore, the circled data in 

Figure 3 show a markedly reduced resistance to the 

onset of scuffing. Adhesion also promotes the 

transfer of soft material than can fill up valleys in 

the surface that might otherwise serve as lubricant 

reservoirs. While a modicum of abrasion may be 

good for scuffing reduction, too much abrasion 

could result in an unacceptably high degree of 

abrasive wear damage and an unacceptable loss of 

clearance between mating parts. Therefore, there 

must be a balance between effective boundary 

lubrication and the light abrasiveness of optimally 

rough surfaces to avoid adhesive contact. 

40.0 

IP = M0 + M1*AT 35.0 M0 3.654 

M1 148.3 

30.0 R 0.931 

Self-mated 

52100 steel 10.0 

INITIATION PERIOD (m) 

25.0 

20.0 

15.0 

ments), and m = a weighting factor for the importance of this function Reciprocating scuffing test 

data for a series of materials were used to evaluate 

scuffing based on the foregoing relationships. This 

work showed an optimal roughness of about 

0.07 μm for the given tribosystem and indicated 

values for the time to initiate scuffing. 

5.0 

0.0 

0.00 0.05 0.10 0.15 0.20 

ABRASIVE TENDENCY (AT) ( μm) 

Figure 3. Relationship between the AT parameter and 

the initiation period for scuffing in reciprocating sliding in Jet A, low-sulfur fuel. 

106
Heavy Vehicle Propulsion Materials 

Summary and Conclusions 

• The 3-D friction-scuffing maps can be developed to depict scuffing initiation in both time 

and space domains. These maps can be used to 

construct scuffing transition diagrams that correlate scuffing resistance to surface finish and 

sliding velocity. 

• Current and candidate fuel injector material 

combinations seem to exhibit an optimum surface finish (neither too smooth nor too rough), 

which can extend the times to initiate and 

propagate scuffing damage. The optimum surface finish is a function of the geometric conditions, the normal load, the characteristics of the 

fluid film that is trapped in the valleys of the 

surface roughness, and the adhesion characteristics of the mating surfaces. 

• Abrasion can play a subtle role in avoiding 

adhesive junction formation. Therefore, 

harder/softer material combinations (like zirconia on steel) can be successful in scuffing-sensitive applications. That works if the surface 

finish is optimized and abrasion is kept to a 

minimum. 

• A new scuffing model was developed and 

reported. 

Presentations and Publications 

P. J. Blau, “Challenges for the application of 

engineered surfaces to tribosystems with multiple 

contact modes,” invited presentation at the symposium on Integrated Surface Engineering, 

ASME/STLE Tribology Conference, Ponte Vedra, 

Florida, October 28, 2003. 

FY 2005 Progress Report 

P. J. Blau, J. J. Truhan, and J. Qu, “Scuffing 

Transition Diagrams for Fuel Injection System 

Materials, Project Milestone Report, Oak Ridge 

National Laboratory, Oak Ridge, Tennessee, March 

2004. 

P. J. Blau, J. Qu, and J. J. Truhan, On the 

Definition and Mechanisms of Scuffing in Fuel 

System Components—An Integrated Process Model, 

Project Milestone Report, Oak Ridge National 

Laboratory, Oak Ridge, Tennessee, September 2004. 

J. Qu, J. J. Truhan, and P. J. Blau, “Detecting 

the Onset of Localized Scuffing with the Pin-onTwin Fuel-Lubricated Test for Heavy Duty Diesel 

Fuel Injectors,” SAE International Journal of Engine 

Research, 6(1), 1–9 (2005). 

J. Qu, J. J. Truhan, and P. J. Blau, “Evaluating 

Candidate Materials for Heavy Duty Diesel Fuel 

Injectors Using a ‘Pin-On-Twin’ Scuffing Test,” 

Tribology International, 38(4), 381–390 (2005). 

P. J. Blau, J. Qu, and J. J. Truhan, “A MultiStage Model for Scuffing of Reciprocating Components with Special Consideration of Fuel Injector 

Plungers,” Oak Ridge National Laboratory, Oak 

Ridge, Tennessee, September 30, 2005 (in press). 

107 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

G. Life Prediction of Diesel Engine Components 

H. T. Lin, T. P. Kirkland, and A. A. Wereszczak 

Oak Ridge National Laboratory 

P.O. Box 2008, MS-6068 

Oak Ridge, TN 37831-6068 

(865) 576-8857; fax: (865) 475-6098; e-mail: linh@ornl.gov 

Jeremy Trethewey and Craig Habeger 

Caterpillar, Inc. 

Peoria, IL 61656-1875 

(309) 578-0056; fax: (309) 578-2953; e-mail: trethewey_jeremy@cat.com 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: Caterpillar, Inc., Peoria, Illinois 

Objective 

• Generate a database and characterize damage mechanisms of candidate advanced ceramics and intermetallic 

alloys. 

• Apply and verify probabilistic life prediction and component design and verification methods for advanced 

diesel engine components. 

Approach 

• Evaluate the dynamic fatigue, rotary bending fatigue, and high-temperature fatigue and creep performance of 

candidate silicon nitride ceramics and TiAl alloys at elevated temperatures in air before and after long-term 

exposure to simulated engine environments, as well as engine field tests. 

• Characterize the evolution and role of damage mechanisms, and changes in microstructure and chemistry, 

linked to the long-term mechanical performance and reliability of ceramics and intermetallic alloys. 

• Predict the failure probability and reliability of complex-shaped diesel engine components subjected to 

application conditions via the use of life prediction codes. 

Accomplishments 

• Completed development of a dynamic fatigue database at elevated temperatures up to 1000°C in air for a 

commercial-grade silicon nitride, SN147-31E, acquired from Ceradyne, Inc. 

• Completed characterization of the mechanical strength and microstructure for 16 prototype NT551 silicon 

nitride exhaust valves, as well as half-cylindrical valve stems, after a 500-h bench rig test at Caterpillar. 

• Completed dynamic fatigue studies for diesel particulate filter (DPF) material at elevated temperatures in air. 

109 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Future Direction 

• Characterize the retained mechanical properties and microstructure of Kyocera SN235P silicon nitride and TiAl 

exhaust valves designed for a Caterpillar natural gas engine after a field test, and verify the probabilistic 

component design and life prediction. 

• Develop a static fatigue performance database at elevated temperature for specimens extracted from DPF 

substrates for long-term mechanical reliability and life prediction. 

• Develop a high-temperature tensile creep and step-stress rupture database of TiAl alloys for probabilistic 

component design and life prediction for diesel exhaust valve and turbine wheel component application. 

• Develop a mechanical database for biaxial discs extracted from TiAl turbine wheel airfoils to verify

probabilistic component design and life prediction. 

Introduction 

There has been considerable interest in the 

potential for extensive use of advanced ceramics and 

intermetallic alloys in advanced diesel engine 

systems because of their superior thermomechanical 

properties at elevated temperatures. The implementation of components fabricated from these advanced 

materials would lead to significant improvement in 

engine efficiency and long-term durability and 

reduced nitrogen oxides (NOx) and CO exhaust 

emissions as required in the 21st Century Truck 

Program. Interest has focused primarily on research 

into characterization and design methodology 

development (life prediction) for advanced silicon 

nitride ceramics and TiAl alloys to enable the 

manufacture of consistent, reliable complex-shaped 

components for diesel engine. The valid prediction 

of mechanical reliability and service life is a 

prerequisite for successful use of these materials in 

internal combustion engine components. 

This research project has three primary goals: 

the generation of a mechanical engineering database, 

from ambient to high temperatures, of candidate 

advanced materials before and after exposure to 

simulated engine environments; the microstructural 

characterization of failure phenomena in these 

advanced materials and in components fabricated 

from them; and the application and verification of 

probabilistic life prediction methods using diesel 

engine components as test cases. For all three stages, 

results will be provided to both material suppliers 

and component end-users for use in refining and 

optimizing processing parameters to achieve consistent mechanical reliability, and in validating the 

probabilistic design and life prediction of engine 

components made from these advanced materials. 

Approach 

All flexural test specimens were longitudinally 

or transversely machined according to ASTM standard C1161 from production billets purchased from 

material suppliers and diesel engine components 

acquired from end-users.1

 Flexure testing was conducted in ambient air in a 4-point-bending apparatus 

using 20/40-mm, α-SiC, semi-articulating fixtures at 

temperatures ranging from 20 to 1000°C and at 

stressing rates of 30 MPa/s and 0.003 MPa/s. The 

30-MPa/s test condition was chosen to evaluate the 

inert characteristic strength as a function of temperature, while the 0.003-MPa/s test condition was 

chosen to measure the change in slow crack growth 

(SCG) susceptibility at elevated temperatures. The 

dynamic fatigue tests were carried out per ASTM 

C1465 (Ref. 2). Pneumatic actuators were programmed with a PC to produce the desired loading 

rate (and corresponding stressing rate). Load was 

continuously measured as a function of time, and 

flexure strength was calculated using ASTM C1161. 

The accumulated strength data were then further 

analyzed. The strengths for each test set were fit to a 

two-parameter Weibull distribution using the program CERAMIC,3

 which uses maximum likelihood 

estimation as advocated in ASTM C1239 (Ref. 4). 

Reported results are uncensored because fractography analysis was not conducted in detail to identify strength-limiting flaws for all of the bend bars 

tested. Following the dynamic fatigue test, both optical and scanning electron microscope analysis were 

carried out on fracture surfaces and polished crosssections of selected bend bars to characterize the 

fracture and degradation mechanisms. X-ray analysis was also carried out to evaluate the possible 

phase changes resulting from oxidation during 

dynamic fatigue tests or after long-term exposure to 

110
Heavy Vehicle Propulsion Materials 

simulated engine environments, which could possibly cause degradation in mechanical performance 

and reliability. 

Results 

Dynamic Fatigue Response of Ceradyne SN147-E 

Silicon Nitride 

Studies were completed of the dynamic fatigue 

behavior of a commercial-grade silicon nitride— 

SN147-31E, manufactured by Ceradyne Advanced 

Ceramic, Inc., CA—at temperatures of up to 1000°C 

in air. The SN147-31E has been processed with 

oxide sintering additives (i.e., Al2O3 and Y2O3) and 

contains a crystalline secondary phase achieved by 

proprietary post–heat treatment. The SN147-31E 

silicon nitride test specimens were longitudinally 

and transversely machined per the revised ASTM 

C1161 standard with a 600-grit surface finish.1

 The 

dynamic fatigue tests were carried out at 20 and 

850°C and at stressing rates of 30 and 0.003 MPa/s 

in air per ASTM C1465 (Ref. 2). The 30 MPa/s rate 

is used to evaluate the inert characteristic strength as 

a function of temperature, and 0.003 MPa/s is 

applied to measure the slow crack growth (SCG) 

susceptibility at temperatures. Limited samples were 

also tested at 1000°C and at 30 MPa/s to evaluate 

FY 2005 Progress Report 

high-temperature application limits and changes in 

life-limiting processes. The upper temperature limit 

in a diesel engine environment is about 850°C. 

Dynamic fatigue results at 30 MPa/s showed 

that the transversely machined SN147-31E silicon 

nitride exhibited little or no decrease in characteristic strength at temperatures up to 1000°C compared 

with those obtained at 20°C (as shown in Table 1). 

Also, the SN147-31E samples exhibited higher 

Weibull moduli at elevated temperatures than at 

20°C, indicative of a change in strength-limiting 

flaws as a function of test temperature. On the other 

hand, mechanical results at 850°C indicated that 

there is an apparent decrease in characteristic 

strengths (about 19%) when tested at 0.003 MPa/s. 

In addition, the fatigue exponent of transversely 

machined SN147-31E decreased from 338 at 20°C 

to 39 at 850°C, indicative of an increased 

susceptibility to SCG processes at high temperature 

(Figure 1). The decrease in dynamic fatigue exponent value at elevated temperatures could be attributed to the change in material state (e.g., softening 

of secondary phase), as seen in the case for the 

GS44 silicon nitride.5

 The low fatigue exponent 

obtained for SN147-31E at elevated temperatures 

suggests that the secondary phase was not 

Table 1. Summary of uncensored Weibull and strength distributions for Ceradyne SN147-31E silicon nitride 

specimens longitudinally and transversely machined per ASTM C1161. Data for SN147-31N machined 

longitudinally as well as transversely are used for reference 

Material 

Number 

of 

specimens 

tested 

Stressing 

rate 

(MPa/s) 

Temperature 

(°C) 

Uncens. 

Weibull 

modulus 

± 95% 

uncens. 

Weibull 

modulus 

Uncens. 

char. 

strength 

(MPa) 

± 95% 

uncens. 

char. 

strength 

(MPa) 

SN147-31N-Long 15 30 20 21.73 14.07. 31.09 836 814, 858 

SN147-31N-Long 15 30 850 20.35 13.58, 28.20 777 755, 799 

SN147-31N-Long 15 0.003 850 16.19 10.57, 23.02 732 706, 757 

SN147-31N-Trans 15 30 20 13.76 8.96, 19.59 677 649, 705 

SN147-31N-Trans 15 30 850 18.26 11.67, 26.47 639 619, 659 

SN147-31N-Trans 15 0.003 850 19.95 12.83, 28.63 620 602, 638 

SN147-31E-Long 15 30 20 17.30 10.86, 25.58 668 645, 690 

SN147-31E-Long 15 0.003 20 16.19 10.60, 22.98 620 598, 642 

SN147-31E-Long 15 30 850 9.36 6.01, 13.62 604 567, 641 

SN147-31E-Long 15 0.003 850 12.06 7.70, 17.52 509 485, 533 

SN147-31E-Long 15 30 1000 16.49 10.61, 23.66 585 564, 605 

SN147-31E-Trans 15 30 20 7.44 4.89, 10.58 623 575, 671 

SN147-31E-Trans 15 0.003 20 8.99 5.61, 13.45 607 568, 646 

SN147-31E-Trans 15 30 850 43.94 28.37, 62.78 642 634, 651 

SN147-31E-Trans 15 0.003 850 18.02 11.93, 25.18 521 505, 538 

SN147-31E-Trans 15 30 1000 34.66 21.97, 50.70 602 591, 611 

111 
103 

FY 2005 Progress Report 

Dynamic Fatigue of SN147-31E Silicon Nitride

Transversedly Machined

Uncensored Flexure Strength 

Failure Strength (MPa) 

N ~ 338 

N ~ 39 

31E-TM-20°C 

31E-TM-850°C 

102 

0.001 0.01 0.1 1 10 100 

Stressing Rate (MPa/s) 

Figure 1. Failure strength versus stressing rate curve of 

SN147-31E transversely machined and tested 

at 20 and 850°C in air. 

completely crystallized, and thus the post–heat 

treatment procedure needs to be refined to ensure 

high-temperature mechanical performance and 

reliability. 

Mechanical Characterization of NT551 Silicon 

Nitride Exhaust Valves 

Mechanical strength testing was conducted on 

16 NT551 silicon nitride valves (14 with 30° and 4 

Heavy Vehicle Propulsion Materials 

with 45° valve seat angles) after testing on a 500-h 

bending rig at Caterpillar. In general, optical microscopy examinations showed that no surface damage 

or flaws were apparent after the bench rig test. However, there were distinct valve seat markings at the 

valve and stainless steel valve seat contact point. 

Strength testing was carried out at room temperature, 

with a compression text figure specially designed inhouse, to evaluate the effects of a 500-h bench rig 

test on mechanical performance (Figure 2). The test 

results of as-received NT551 valves obtained using a 

hydraulic chamber facility were used as a baseline 

for comparison (Figure 3). Mechanical results for 

the valves tested at 500 h show that the characteristic strength and the Weibull strength distribution 

were similar to those obtained for the as-machined 

valves, suggesting no surface/subsurface damage 

was introduced during the bench rig test (Figure 4). 

Note that nondestructive evaluation carried out at 

Argonne National Laboratory on the valves tested at 

500 h also indicated no apparent damage/flaws, 

consistent with the present mechanical testing results. 

Studies were also conducted of the mechanical 

properties of NT551 silicon nitride valve stems after 

a 500-h bench rig test. The objective was to see 

whether the 500-h bench rig test would introduce 

any surface damage to the valve stem surfaces and 

thus degrade the mechanical strength. The halfcylindrical valve stems were machined and then 

σc 

Si3N4 

valve 

Figure 2. Photos of silicon nitride exhaust valves after testing on a 500-h bend rig at Caterpillar 

and on a mechanical testing fixture. The valve strength test was on a Universal Instron 

machine. 

112
-5 

-4 

-3 

0 

1 

-5 

-4 

-3 

-2 

1 

99.0 

90.0 

50.0 

20.0 

10.0 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

NT551 Silicon Nitride Valves - As-received NT551 Silicon Nitride Halh-Cylindrical 

Uncensored Flexure Strength Distribution Valve Stems - 500h Test 

Uncensored Flexure Strength Distribution 20°C - 1.3 mm/min - Transversely Machined 20°C - 2 mm/min - Transversely Machined 2 99.9 2 99.9 

Max. Lik. Fit 

Strength Data 

95% Confidence 

Bands Shown 

m = 8.32 (5.41, 11.78) 

σθ

 = 688 MPa (641, 736) 

n = 15 specimens 

Max. Lik. Fit 

Strength Data 

95% Confidence 

Bands Shown 

m = 12.74 (9.11, 17.05) σθ

 = 1135 MPa (1098, 1173) 

n = 26 specimens 

99.0 

1 90.0 Probability of Failure, P , (

%) f , (%) f

0 

50.0 

20.0 

10.0

ln ln ( 1 / ( 1 - P ) ) 

f

-1 

-2 

ln ln ( 1 / ( 1 - P ) ) 

f

-1 

-2 

5.0 -3 

, (%)

f

2.0 -4 

1.0 

-5 

0.5 

-6 -6 

100 1000 2000 300 400 500 600 700 800 9001000 2000 

Failure Stress (MPa) Failure Stress (MPa) 

Figure 3. Strength distribution of as-received Norton Figure 5. Strength distribution of Norton NT551 silicon NT551 silicon nitride valve transversely nitride half-cylindrical valve stems machined and tested at 20°C. 

transversely machined and tested for 500 h at 

20°C. 

NT551 Silicon Nitride Valves - 500h Test the surface smoothing effect resulting from the conUncensored Flexure Strength Distribution tact between the silicon nitride valve stem and the 20°C - 1.3 mm/min - Transversely Machined 2 valve guide. 99.9 

0 

-1 

Max. Lik. Fit 

Strength Data 

ln ln ( 1 / ( 1 - Pf ) ) 

95% Confidence 

Bands Shown 

m = 8.73 (5.79, 12.19) 

σθ

 = 776 MPa (728, 825) 

n = 16 specimens 

Particulate Filters 50.0 

99.0 

90.0 Dynamic Fatigue Performance of Diesel 

Studies of the mechanical properties of a DPF 20.0 substrate, especially dynamic fatigue behavior at 10.0 elevated temperatures, were initiated and carried out 5.0 during this fiscal year. The objective is to develop a 2.0 mechanical database for predicting long-term 

reliability and durability. Two different types of 1.0 

0.5 

-6 

100 1000 2000 

Failure Stress (MPa) 

Figure 4. Strength distribution of Norton NT551 silicon 

nitride valve transversely machined and tested 

for 500 h at 20°C. 

tested using a fixture with 30/60-mm spans with a 

2-mm/min crosshead speed at room temperature. 

Figure 5 shows the uncensored Weibull strength 

distribution of these half-cylindrical valve stem 

specimens. The mechanical results show a 

characteristic strength of 1135 MPa and Weibull 

samples, with 3×4 cells and 4×6 cells, were 

employed to evaluate how size influences the 

mechanical response and fracture process. Tests 

were carried out at temperatures of 600, 800, and 

900°C at stressing rates of 5, 0.005, and 

0.00005 MPa/s in air. The dynamic fatigue 

procedures carried out for DPF materials are similar 

to those employed for monolithic ceramic specimens 

(ASTM C1465). Figure 6 shows the preliminary 

dynamic fatigue results for DPF samples with 4 × 6 

cells. The results showed that the DPF material 

exhibited similar low dynamic fatigue exponents 

(N ~ 40–44) in the temperature range employed. The 

modulus of 12.7, comparable to and/or better than low fatigue exponents obtained suggest that the DPF 

the as-machined cylindrical samples. The results material might exhibit susceptibility to SCG 

indicate excellent wear resistance. The slight processes at temperatures. More test specimens and 

increase in characteristic strength might be due to a lower stressing rate need to be 

113 

2.0 

5.0 

1.0 

0.5 
Fracture Strength (MPa) FY 2005 Progress Report 

8 

7 

6 

5 

4 

3 

2 

1 

0 

10-4 10-3 10-2 10-1 100 101 

600°C (6x3) 

800°C (6x3) 

900°C (6x3) 

N ~ 40-44 

Stressing Rate (MPa/s) 

Figure 6. Fracture strength versus stressing rate of DPF 

material as a function of test temperature. 

studied to provide a better understanding of longterm reliability and life-limiting processes. 

References 

1. “Standard Test Method for Flexural Strength 

of Advanced Ceramics at Ambient Temperatures,” 

ASTM C1161, Annual Book of ASTM Standards, 

Vol. 15.01, American Society for Testing and 

Materials, West Conshohocken, Pennsylvania, 1999. 

2. “Standard Test Method for Determination of 

Slow Crack Growth Parameters of Advanced 

Heavy Vehicle Propulsion Materials 

Ceramics by Constant Stress-Rate Flexural Testing 

at Elevated Temperatures,” ASTM C1465, 

Annual Book of ASTM Standards, Vol. 15.01, 

American Society for Testing and Materials, West 

Conshohocken, Pennsylvania, 1999. 

3. Life Prediction Methodology for Ceramic 

Components of Advanced Heat Engines, Phase 1, 

ORNL/Sub/89-SC674/1-2, prepared by AlliedSignal 

Engines, DOE Office of Transportation 

Technologies, 1995. 

4. “Practice for Reporting Uniaxial Strength 

Data and Estimating Weibull Distribution Parameters for Advanced Ceramics,” ASTM C1239, Annual 

Book of ASTM Standards, Vol. 15.01, American 

Society for Testing and Materials, West 

Conshohocken, Pennsylvania, 1999. 

5. A. A. Wereszczak, H. -T. Lin, T. P. Kirkland, 

M. J. Andrews, and S. K. Lee “Strength and 

Dynamic Fatigue of Silicon Nitride at Intermediate 

Temperatures,” J. Mater. Sci., 37, 2669–2684 (2002). 

Presentations and Publications 

H. T. Lin, T. P. Kirkland, A. A. Wereszczak, 

and M. J. Andrews, “Strength Retention of Silicon 

Nitride After Long-term Oil Immersion Exposure,” 

J. Mater. Sci. (in press). 

114
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

H. Low-Cost Manufacturing of Precision Diesel Engine Components 

P. J. Blau and J. Qu 

Oak Ridge National Laboratory 

P.O. Box 2008, MS-6063 

Oak Ridge, TN 37831-6063 

(865) 574-5377; fax: (865) 574-6918; e-mail: blaupj@ornl.gov 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Objectives 

• Enable the use of advanced materials such as ceramic composites, cermets, and titanium alloys in diesel 

engines by understanding their machining characteristics on a basic and applied level. 

• Develop new, ceramic-based cutting tool materials for high-speed titanium machining. 

• Work with a small business (Third Wave Systems, Inc.) in the development of software packages for 

manufacturers to optimize high-speed machining processes. 

Approach 

• Quantify the grinding characteristics of TiC/Ni3Al matrix cermets that are promising candidates for diesel 

engine fuel injector plungers and wear parts. 

• Investigate the possibility of using cermets as cutting tool materials for high-speed titanium machining. 

• Participate with a consortium of manufacturers led by Third Wave Systems, Inc., to develop finite-element 

analysis (FEA)-based machining modeling software and to acquire new machining process modeling 

capabilities. 

Accomplishments 

• Worked with a visiting faculty member from Louisiana State University to study the effect of wheel condition 

and grinding parameters on the machinability of Ni3Al-bonded advanced composites with varying amounts of 

titanium carbide particles. 

• Worked with the University of Michigan and the ceramic processing staff at Oak Ridge National Laboratory 

(ORNL), to prepare, characterize, and test experimental cermet composites as possible tooling inserts for the 

high-speed machining of titanium alloys. 

• Tested and helped to evaluate a new software package designed to optimize high-speed drilling of titanium 

alloys. 

115 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Future Direction 

• The current project ended in FY 2005. A new effort on synergistic process-enhanced grinding of advanced 

materials is planned to begin in FY 2006. 

Introduction 

Manufacturing broadly comprises the conversion of raw materials into useful products. It encompasses a broad range of engineering disciplines and 

unit processes, including machining and finishing. 

Since the industrial revolution of the late 1800s, 

U.S. industry has invested hundreds of billions of 

dollars in the development of manufacturing technology to promote economic health and security. 

However, with the continuing egress of manufacturing facilities to foreign soil, needed improvements 

in domestic manufacturing facilities and the 

knowledge base to develop new machining technology has been eroding. The machine tool industry in 

the United States has largely closed up shop, and 

most new machine tools are made by foreign 

sources. It is vital therefore, to avoid full dependence on foreign sources, to promote applied 

machining science within the United States. 

Driven by the high price of fuel and increasingly 

stringent emissions regulations, the diesel industry is 

revising its engine designs. Some of these designs 

could benefit from the use of high-performance 

materials, but the ability to machine these hard, 

durable materials cost-effectively becomes a key 

enabler. Therefore, the current project is intended to 

enable the introduction and use of advanced materials for diesel engine components by better understanding and optimizing their machining 

characteristics. 

Approach 

Three tasks addressed new machining technology for diesel engine components. The first two are 

experimental studies and the latter an effort in 

modeling with experimental verification. 

• Characterizing the machining characteristics of 

advanced composite materials containing an 

intermetallic-alloy binder with hard carbides. 

• Developing and testing new cutting tool materials for machining lightweight alloys for fuelefficient heavy vehicles. 

• Working with software developers to test new 

process-specific computer modeling packages 

for machining of advanced materials. 

Results 

Grindability of TiC-Ni3Al cermets 

TiC-Ni3Al metal matrix composites containing 

40–60 vol % Ni3Al were developed as candidate 

materials for diesel fuel injector plungers. They 

possess good mechanical properties and have 

thermal expansion coefficients similar to those of the 

steel typically used for injector bores. Previous 

studies have confirmed their high scuffing resistance 

in fuel-lubricated environments. However, costeffective machining of these composites requires the 

achievement of a high material removal rate (MRR). 

That is difficult due to their relatively high toughness and the distribution of hard phases in the 

microstructure. 

Working in collaboration with Prof. W. Liao, 

Louisiana State University, creep feed and conventional surface grinding characteristics of a TiCNi3Al (50 vol %) composite were studied and compared with two baseline ceramic materials, alumina 

and silicon nitride. The investigation was conducted 

using an instrumented grinder with a 220-grit resinbonded diamond wheels using the grinding conditions in Table 1. Selected properties of the workpiece materials and the measured maximum MRRs 

are shown in Table 2. The composite is tougher but 

lower in hardness than the two ceramics, but it had a 

much lower maximum MRR than alumina. 

Table 1. Grinding conditions 

Wheel diameter (mm) 228.6 

Grinding width (mm) 11.43 

Wheel speed (rpm) 4500, 2500 

Depth of cut (mm) 0.051, 0.254, 1.016, 1.27 

DOCa × work speed 164, 254, 328, 645 (mm2

/min) 

aDOC = depth of cut 

116
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Table 2. Hardness, toughness, and maximum MRRs 

Material 

Vickers 

microindentation 

hardness 

(GPa) 

Fracture 

toughness 

(MPa-m1/2) 

Maximum MRR 

(mm3

/s) 

TiC-Ni3Al 

(50 vol %) 

Si3N4

Al2O3

9.5 

19.4 

24.8 

12 

5 

3–4 

62.5 for conventionala 

48.4 for creep feedb 

(Not measured) 

>125 

aReciprocating surface grinding with relatively shallow DOCs. bDeeper DOCs but with a slower feed rate. 

Grinding forces, spindle power, vibration, and 1 

acoustic emission signals were monitored and analyzed. In general, the cermet required higher grind- 0.8 Power (

K

W) Tangential Force (N) 

Roughness Ra (um)

ing forces and spindle power than the two ceramics 

for a given MRR (see Figure 1), but it also exhibited 0.6 

the best surface finish (see Figure 2). Higher fracture 

toughness is likely to be responsible for both results 0.4 

because it requires more power to overcome the Cermet 

higher toughness, but it also leads to less micro- 0.2 Silicon Nitride 

fracturing of surface. Alumina 

0 

250 

Cermet 

p 

0 0.2 0.4 0.6 0.8 1 

Silicon Nitride 200 

Alumina 

150 

100 

50 

0 

0 0.2 0.4 0.6 0.8 1 

Depth of Cut (mm) p 

Cermet 

Silicon Nitride 

Alumina 

0 0.2 0.4 0.6 0.8 1 

Depth of Cut (mm) 

Figure 1. Grinding forces and power of grinding 

Depth of Cut (mm) 

Figure 2. Arithmetic average surface roughness of 

different workpiece material surfaces ground 

under steady-state wheel conditions at a MRR 

of 48.4 mm3

/s. 

Scanning electron microscopy (SEM) revealed 

the response of the cermet surfaces to differing 

degrees of grinding severity. More surface and subsurface damage occurred for creep feed grinding 

than for conventional grinding. However, no excessive cracking or pull-out of TiC particles was 

observed (see Figure 3). This indicates good bonding between the Ni3Al matrix and the hard TiC 

particles that appear darker gray in the image. 

New Cutting Tools and Process Modeling for 

High-Speed Titanium Machining 

Experimental Cutting Tool Materials: Two TiCNi3Al cermets with 5 vol % and 9 vol % matrix 

material, respectively, were prepared with the cooperation of T. Tiegs at Oak Ridge National Labora9 

8 

7 

6 

5 

4 

3 

2 

1 

0 

tory. Their microindentation hardness is compared 

different work materials in steady-state wheel 

condition (MRR = 48.4 mm3

/s). to that for a commercial WC-Co (6 vol %) tool 

117 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 3. SEM image of a ground cermet surface showing one of the deeper cutting grooves. Ductile 

tearing is evident on the ridges of the groove, 

but there is little evidence for TiC particle 

pull-out. 

insert in Table 3. Figure 4 shows a more uniform 

and less porous microstructure for the TiC-Ni3Al 

(9 vol %) composition than for the TiC-Ni3Al 

(5 vol %). 

Tool inserts for turning were prepared in the triangular shape shown in Figure 5. The material was 

cut by four-axis wire electro-discharge tapered sides 

and corners were manually polished by 600 grit SiC 

abrasive paper to remove the EDM recast layer. 

These TiC-Ni3Al tool inserts were used to turn 

Grade 2 (commercially pure) titanium at University 

of Michigan (UM) using 80-, 160-, and 640-ft/min 

cutting speeds, 0.01-in./rev feed rate, and 0.04 in. 

DOC. Compared with commercial WC-Co tools, the 

TiC-Ni3Al tools showed moderately higher cutting 

forces. This could partially be due to the imperfect 

tool geometry and low sharpness of the manually 

polished cutting edges. The tool wear and tool life 

issues were also briefly studied at UM. Table 4 

summarizes these results. Unlike the progressive 

wear on the WC-Co tool, chipping was the dominant 

failure mode of the TiC-Ni3Al tools, especially at 

Table 3. Microindentation hardness 

comparison 

Material HV (GPa, 200 g) 

TiC-Ni3Al(5%) 20.6 

TiC-Ni3Al(9%) 17.3 

WC-Co(6%) 18.5 

(a) TiC-Ni3Al (5 vol %) 

(b) TiC-Ni3Al (9 vol %) 

Figure 4. Microstructures of TiC-Ni3Al cermets. 

Figure 5. Tool inserts made of TiC-Ni3Al cermets. 

118
Heavy Vehicle Propulsion Materials 

Table 4. Tool life in machining of titanium 

Tool life Cutting speed (ft/min) 

(s) 80 160 640 

TiC-Ni3Al 

(5%) 245.3 18.4 15.3 

TiC-Ni3Al 

(9%) Not tested >294.4 9.2 (tip 1) 

33.7 (tip 2) 

WC-Co Not tested Not tested 69 

machining (EDM), the top and bottom surfaces were 

ground by a diamond grinding wheel, and the high 

cutting speeds. Due to the high porosity as shown in 

Figure 4(a), TiC-Ni3Al (5 vol %) tools broke at all 

cutting speeds. The TiC-Ni3Al (9 vol %) tools 

survived at 160 ft/min with progressive wear after 

294.4 s, but broke fairly quickly at 640 ft/min. The 

wide variability in tool life (9.2 and 33.7 s at high 

cutting speed) for TiC-Ni3Al (9 vol %) tools 

suggests either nonuniformity of the microstructure 

or inconsistent cutting edge sharpness, or both. The 

performance of the TiC-Ni3Al (9 vol %) tools might 

be improved by automated production grinding, but 

in light of the current results and the fact that the 

current project is ending, no further studies of this 

material are planned. 

Machining model testing and verification: Work 

continued on the testing and validation of software 

for high-speed machining. Good correlations have 

been obtained between computer modeling using 

Thermal Wave Systems (TWS) AdvantEdgeTM software and lathe turning experiments. 

For example, in titanium machining, cutting 

chips are known to catch fire when their temperature 

rises above about 1200°C. The TWS AdvantEdgeTM 

FY 2005 Progress Report 

was used to simulate chip ignition in titanium 

turning. Machining simulations and verification 

experiments were conducted on a Ti-6Al-4V workpiece using a Micrograin AR6 (WC-Co) tool. The 

feed rate was 0.1 mm/rev, and DOC ranged from 30 

and 75 μm. Table 5 data show very good agreement 

between the simulation and the experiments. The 

simulated temperature profiles of the titanium chips 

at different DOCs are illustrated in Figure 6. 

Characterization of tool coatings: The physical, 

mechanical, and thermal properties of the tool coatings significantly affect the cutting performance and 

tool life. Selected production tool coatings have 

been characterized to provide input for machining 

models. The data are given in Table 6. The coating 

thickness and microstructure were examined by 

SEM as shown in Figure 7. There was a relatively 

large range of coating thickness (2.5–6.0 μm), 

despite the manufacturers’ claims that they were all 

supposed to be 3 μm. Nanoindentation showed 

similar hardness and lower modulus of elasticity of 

these coatings compared to that of the WC substrate. 

The thermal conductivities of these thin coatings 

were measured using a new photoacoustic technique 

at the Microscale Thermophysical Properties Laboratory at Purdue University. All coatings exhibited 

Table 5. Simulation and experimental results for 

titanium chip ignition 

DOC (μm) 30 75 

Simulation No ignition Ignition 

(Chip Tmax < (Chip Tmax > 

1000o

C) 1400oC) 

Experiment No ignition Ignition 

Figure 6. Simulation of titanium chip ignition (left: no ignition left, right: ignition). 

119 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Table 6. Physical, mechanical, and thermal properties of tool coatings 

Coating characteristics 

Grade 

(kennametal) Tool insert material Thickness 

(μm) 

Nanoindentation 

Ha (GPa) Eb (GPa) 

Thermal 

conductivity 

(W/m-K) 

K313 WC-Co (6% Co) — 23.2 545.2 80.99 

KC730 K313+PVD TiN 2.5 23.1 408.5 18.69 

SP39AH K313+PVD TiCN 6.0 25.5 448.7 17.53 

KC5010 K313+PVD TiAlN 3.0 25.0 355.3 1.86 

aH = nanoindentation hardness number. 

bE = elastic modulus obtained from nanoindentation. 

TiN 

TiCN 

TiAlN 

Figure 7. Cross sections of tool 

coatings clearly indicated a 

variation in thickness. 

much lower thermal conductivity than the WC 

substrate. Results imply that the longer tool lives 

obtained from the coated tools were not due to better 

thermo-physical properties (they were actually 

worse than the substrate), but from improved 

tribological characteristics, such as less adhesion to 

reduce edge buildup. The latter is described in the 

next section on tool wear characterization. 

Tool wear characterization: In titanium 

machining, the high tool wear, particularly 

“diffusion wear,” significantly limits the cutting 

speed and tool life. The wear of four turning WC-Co 

based tools with different coatings and/or corner 

radii was measured and characterized. The four tool 

inserts were collected at the point of tool failure 

(dramatic increase in the cutting force) in titanium 

machining, as shown in Table 7. 

The wear patterns on tool tips were mapped by 

laser profilometry (Rodenstock RM600). Results are 

shown in Figure 8. Different “tool failure” mechanisms were observed: the tools with 0.8-mm corner 

radius showed high crater wear [Figures 8(a) and 

8(b)], while the 0.1-mm corner radius tools failed 

because the work material transferred and built up 

on the cutting edge. Energy dispersive X-ray analysis confirmed a high titanium concentration on the 

builtup material. The volumes of the crater wear and 

edge buildup were then quantified using the True 

MapTM 3D image analysis package, and summarized 

in Table 7. The TiB2 coated tool (0.8 mm rc) had 

slightly higher wear rate compared to the noncoated 

tool, but last much longer. That is possibly due to its 

slower rate of edge buildup, which dramatically 

reduces the edge sharpness and causes high adhesion 

to the work material. For a corner radius of 0.1 mm, 

the TiB2 coated tool significantly outperformed the 

TiAlN coated tool due to the much lower edge 

buildup rate, as shown in Table 7. 

120
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Table 7. Tool wear results 

Tool Corner Tool life Crater wear rate Edge buildup rate 

coating radius rc (s) (mm3

/s) (mm3/s) (mm) 

N/A 0.8 69 1.0 × 10–4 0.4 × 10–4

TiB2 0.8 199 1.3 × 10–4 0.1 × 10–4

TiAlN 0.1 38 — 11 × 10–4

TiB2 0.1 115 — 0.9 × 10–4

(a) Noncoated WC-Co, 0.8 mm rc (b) TiB2 coated, 0.8 mm rc

(c) TiAlN coated, 0.1 mm rc

Figure 8. 3-D profiles of the worn tool tips. 

Presentations and Publications 

R. Li, A.J. Shih, and J. Qu, “High Speed 

Machining of Titanium at ORNL/UM,” 2005 Third 

Wave AdvantEdgeTM International Users’ Conference, Dearborn, Michigan, May 4, 2005. 

(d) TiB2 coated, 0.1 mm rc

J. Qu, T. W. Liao, P. J. Blau, J. E. Shelton, and 

T. N. Tiegs, “Grindability of TiC-Ni3Al Metal 

Matrix Composites,” to be presented at The 30th 

International Conference & Exposition on Advanced 

Ceramics & Composites, Cocoa Beach, Florida, 

January 22–27, 2006. 

121 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

I. Cost-Effective Machining of Titanium Materials 

Albert J. Shih 

Mechanical Engineering 

University of Michigan 

Ann Arbor, Michigan 48109-9739 

(734) 647-176; fax: (734) 936-0363; e-mail: shiha@umich.edu 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: University of Michigan, Ann Arbor, Michigan 

Objective 

• Investigate new technologies for cost-effective machining of titanium alloys. 

• Develop the cost-effective, high-speed machining processes. 

• Investigate new tool material, geometry, coating, and process parameters to enhance process capability and 

product quality. 

• Provide accurate prediction of temperature and stress for tool selection and design. 

Approach 

• Conduct turning test using new tool material. 

• Develop high-speed titanium drilling technology. 

• Carry out the finite-element modeling and inverse heat transfer analysis to study titanium turning and drilling 

processes. 

• Validate the modeling with experimental results. 

Accomplishments 

• Compared the performance of new TiC-based tool material with WC-Co for turning titanium. 

• Demonstrated the feasibility of high-speed drilling process over 180-m/min cutting speed for titanium. 

• Achieved accurate predictions of the temperature distribution of the tool in titanium drilling. 

Future Direction 

• Continue the exploration of high-speed, cost-effective machining of titanium alloys. 

• Advance the thermal analysis and management of the cutting tool. 

• Conduct the 3-D finite-element-based thermal-mechanical coupled drilling modeling to advance the knowledge 

in mechanics of material deformation during the titanium chip formation process. 

123 
FY 2005 Progress Report 

Introduction 

Titanium and its alloys are lightweight, 

corrosion-resistant, biocompatible, and hightemperature materials. The poor machinability of 

titanium alloys has been summarized in several 

review articles.1–4

 Due to the inherent properties of 

titanium, particularly the low thermal conductivity, 

the tool temperature in machining of titanium is high 

and concentrated at the tool tip.1

 High temperature 

softens the tool material and promotes rapid diffusion wear and severe tool edge chipping.5,6 

A new tool material, the TiC-Ni3Al cermet, was 

developed at Oak Ridge National Laboratory with 

the goal to increase high-temperature stability and 

reduce the diffusion wear in titanium machining. 

Titanium turning tests are performed, and results are 

compared with that of a conventional WC-Co tool. 

Among various machining processes, drilling 

has a major economical importance because it is 

typically utilized in final steps in production. 

Technical difficulties of cost-effective drilling of 

titanium include (1) high drill temperature, 

(2) severe drill wear, even at moderate feed rate and 

spindle speed, (3) limited cutting speed and 

productivity, (4) difficulty to eject long ribbon titanium chip, and (5) exit burr formation. The main 

objective of this year is to investigate the drilling 

mechanism and develop the new technologies to 

overcome the technical difficulties mentioned above. 

Heavy Vehicle Propulsion Materials 

(a) TiC-Ni3Al (5%) 

Approach 

Tool Wear Test of TiC-Ni3Al Material 

The TiC-Ni3Al materials were made at ORNL. 

The tool blank face was machined by wire electrodischarge machining (EDM) and polished to the 

same geometry as the Kennametal TPG160308 tool. 

A set of turning tests was conducted using this new 

tool material at the University of Michigan to evaluate the wear of TiC-Ni3Al tool. 

Figure 1 shows the optical microscopic image of 

the polished TiC-Ni3Al. The TiC with 5% Ni3Al 

shows relatively high porosity, compared to 

commercial tool material. The high porosity results 

in the brittleness of tool under high stress. Increasing 

the content of Ni3Al to 9% reduces the material 

porosity, but sacrifices the hardness of TiC-Ni3Al, as 

shown in the hardness indention results in Table 1. 

The tool wear tests were conducted at 24.4-, 

48.8-, and 195-m/min cutting speeds in lathe turning 

with 0.25-mm/rev feed and 1-mm depth of cut. Both 

(b) TiC-Ni3Al (9%) 

Figure 1. Microstructures of TiC-Ni3Al cermets. 

Table 1. Micro-indentation hardness 

Material HV (GPa, 200 g) 

TiC-Ni3Al (5%) 20.6 

TiC-Ni3Al (9%) 17.3 

WC-Co (6%) 18.5 

TiC-Ni3Al and WC-Co tool were evaluated, and the 

cutting forces and tool life were compared. 

High-Speed Drilling of Ti-6Al-4V 

High cutting speed is essential to reduce the 

machining cost and increase the productivity. It is 

particularly important, but difficult to achieve, in the 

machining of titanium and titanium alloys. As 

shown in Table 2, the drilling speed for titanium 

alloys is very low in industrial applications. This has 

124
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Table 2. Comparison of drilling speed 

Drill material Al alloys Cast iron Ti-6Al-4V 

HSS7 105 49 11 

WC-Co7 245 90 46 

Other

8

 533 235 — 

hindered the productivity and increased the cost of 

machining titanium. The goal of this research is to 

enable the high-speed drilling of titanium. 

Three types of drills, (1) M2 HSS twist drill 

(Greenfield Industries 44210), denoted as HSS 

Twist, (2) WC-Co web thinned point twist drill 

(Kennametal KWCD00344), denoted as WC-Co 

Twist, and (3) WC-Co spiral point drill (Kennametal 

K285A01563), denoted as WC-Co Spiral, are evaluated. All drills are double-flute with 4 mm in diameter. The workpiece material is 6.35-mm thick Ti6Al-4V plate. 

The titanium drilling tests were conducted at a 

constant feed of 0.051 mm per revolution or 

0.025-mm feed per tooth. By changing the spindle 

speed, various drill feed rates and peripheral cutting 

speeds could be achieved. Cutting fluid plays an 

important role to reduce the tool and workpiece temperatures and enhance the tool life.6

 Two ways to 

supply the cutting fluid, internally from two 

through-the-drill cutting fluid holes and externally 

from outside the drill, were studied. The cutting 

fluid was 5% CIMTECH 500, a synthetic metalworking fluid. The thrust force and torque during 

drilling were measured. 

Four sets of experiments, denoted as Exps. I, II, 

III, and IV, were conducted. Exp. I was conducted 

using the HSS Twist drill at three peripheral cutting 

speeds, 9.1, 13.7, and 18.3 m/min, without cutting 

fluid. This is the baseline test. To study the effect of 

cutting fluid, Exp. II was conducted with external 

cutting fluid supply using HSS Twist drill at higher 

peripheral cutting speeds at 13.7, 18.3, and 27.4 

m/min. Exp. III was conducted to compare the thrust 

force, torque, energy consumption, and burr formation of three types of drill (HSS Twist, WC-Co 

Twist, and WC-Co Spiral) at 18.3-m/min peripheral 

cutting speed without cutting fluid. Exp. IV was 

high-speed drilling of titanium using the WC-Co 

Spiral drill at 183-m/min peripheral cutting speed. 

Drilling under dry as well as with internal and external cutting fluid supply conditions was evaluated. 

Figure 2 illustrates three stages, marked as A, B, 

and C, in drilling. Stage A occurs when the drill has 

Stage A Stage B Stage C 

d 

Drill 

Workpiece 

Figure 2. Illustration of three stages A, B, and C in 

drilling. 

traveled by a distance d, which is called the drill 

point length, from the initial contact. For the HSS 

Twist, WC-Co Twist, and WC-Co Spiral drills used 

in this study, d is 1.24, 1.19, and 0.74 mm, respectively. Neglecting the deformation of the workpiece, 

the drill cutting edge becomes fully engaged with 

the workpiece at this stage. When the drill tip 

reaches the back surface of the plate, as shown in 

Figure 2, it is marked as Stage B. Stage C is defined 

when the drill cutting edge is disengaged with the 

workpiece without considering the deformation of 

the workpiece and burr formation. 

The torque and force changed throughout the 

drilling process. The value of thrust force and torque 

at Stage A, denoted as FA and TA, and maximum 

thrust force and torque during drilling, denoted as Fm 

and Tm, respectively, were compared to distinguish 

different drilling conditions. 

Besides thrust force and torque, another parameter to compare drilling setup is the energy E required 

to drill a hole. 

2πT E = ∫ Fdl + ∫ f 

dl , l l 

where l is the depth of drilling, F is the thrust force, 

T is the torque, and f is the feed per revolution. In 

titanium drilling, the E is dominated by the torque 

term, which typically accounts for more than 95% of 

the E. 

Inverse Heat Transfer Solution of Tool 

Temperature in Drilling of CP Titanium 

Understanding the heat generation rate and tool 

temperature distribution is important in titanium 

drilling. The inverse heat transfer method, which 

utilizes the temperature measured by thermocouples 

embedded on the drill flank surface as the input, is 

applied to predict the heat flux on the drill chisel and 

cutting edges. Thermocouples are routed from the 

125 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

frank surface, via through-the-drill cutting fluid 

holes, to outside the tool.9 During drilling, these 

thermocouples do not contact nor interfere with the 

chip. A validation process can be performed by 

comparing the drill temperature distributions to 

thermocouple measurements. 

The drilling experiment was conducted in a 

Mori-Seiki TV-30 vertical machining center. 

Figure 3 shows the experimental setup with the 

stationary 9.92-mm-diameter spiral point drill and 

the 38-mm-diameter CP titanium bar rotated by 

spindle. The drill was stationary because four 

thermocouples embedded on the flank face could be 

routed through cutting fluid holes in the drill body to 

a data acquisition system during drilling.9

 The drill 

and machine spindle axes were aligned by a test 

indicator to less than 10-μm eccentricity. Under the 

drill holder was a Kistler 9272 dynamometer to measure the thrust force and torque. Figure 4 shows the 

locations of four thermocouples, denoted as TC1, 

TC2, TC3, and TC4, on the drill frank surface. 

Three drilling experiments were conducted at 

three spindle speeds, 780, 1570, and 2350 rpm, 

which corresponded to 24.4-, 48.8-, and 73.2-m/min 

drill peripheral cutting speed, respectively. The feed 

remained fixed at 0.051 mm/rev. All the experiments were conducted dry without cutting fluid. 

Figure 5 shows the finite-element mesh of the 

drill, which is modeled by 68,757 four-node tetrahedral elements. As shown in Figure 5(b), 13 nodes 

TC3 

TC1 

Figure 4. Thermocouple locations in the drill flank face. 

(a) (b) 

Figure 5. Mesh for the 3-D finite-element thermal 

model: (a) side view and (b) top view. 

Tool holder 

Dynamometer 

Thermocouple 

wires 

Figure 3. Experimental setup with workpiece in the 

holder of spindle and drill in a vertical tool 

holder. 

are located on the chisel edge, and 11 nodes are 

placed on each cutting edge to achieve good resolution in the analysis of drill temperature distribution 

in titanium drilling. 

Results 

Tool Wear Test of TiC-Ni3Al Material 

Cutting forces are summarized and compared 

with a WC-Co tool in Figure 6. Consistently, the 

TiC-Ni3Al tools showed higher cutting forces than 

that of the WC-Co tool. This was likely due to the 

imperfect tool geometry with less sharp cutting 

edges of the TiC-Ni3Al tools. 

The tool life was also evaluated. The major failure mode of TiC-Ni3Al tools is the catastrophic tool 

breakage. Table 3 summarizes the tool life of the 

TiC-Ni3Al tools compared to the WC-Co tool. Due 

to the high porosity, TiC-Ni3Al (5%) tools broke at 

all cutting speeds. The TiC-Ni3Al (9%) tools 

126
Heavy Vehicle Propulsion Materials 

TiC+Ni3Al (9%) TiC+Ni3Al (5%) WC-Co 

450 

400 

350 

300 

200 

150 

100 

50 

250 

200 

150 

100 

50 

0 

40 80 120 160 200 

Cutting speed (m/min) 

Figure 6. Cutting forces of TiC-Ni3Al and WC-Co tools 

in lathe turning of titanium. 

Table 3. Tool lives in machining titanium 

Axial force (N) Radial force (N) Tangentail force (N) 

24.4 

Cutting speed (m/min) 

48.8 195 

TiC-Ni3Al 

(5%) 245.3 18.4 15.3 

TiC-Ni3Al 

(9%) 

WC-Co 

Not tested 

Not tested 

>294.4

Not tested 

 9.2/33.7 

69 

survived at 48.8 m/min with progressive wear after 

294 s, but broke early at 195 m/min. Significantly 

different tool lives, 9.2 and 33.7 s, at high cutting 

speed (195 m/min) were obtained on two tips of a 

TiC-Ni3Al (9%) tool insert, indicating the nonuniformity of the microstructure or the nonconsistent condition of the cutting edges. The higher 

Ni3Al content improves the performance of TiCNi3Al tool material, but some efforts like better 

FY 2005 Progress Report 

preparation of tool edge and surfaces are still needed 

to achieve or exceed the performance of WC-Co. 

We are working with Kennametal to grind the TiCNi3Al tool to the right geometry. 

High-Speed Drilling of Ti-6Al-4V 

Table 4 summarizes the Stage A and peak thrust 

force and torque, and energy consumption in 

Exps. I–IV. Performance of each type of drill is discussed as follows. 

The HSS Twist drill has limited cutting speed 

and material removal rate in drilling titanium. The 

drill tip melted in the second hole due to high 

temperature at 18.3- and 27.4-m/min peripheral 

cutting speeds under dry and external cutting fluid 

supply conditions, respectively. 

The WC-Co Spiral drill has the lowest thrust 

force and torque among all three drills. This drill is 

selected for high-speed titanium drilling. The drill 

can perform at 183-m/min peripheral cutting speed 

(14700 rpm) under the dry as well as with external 

and internal cutting fluid supply conditions. As 

listed in Table 5, the tool life is short without the 

supply of cutting fluid. Using flood external cutting 

fluid supply just slightly increased the tool life 

because the high spindle speed generates high 

centrifugal force, which prevents the cutting fluid 

from reaching the cutting region. The internal cutting fluid supplied through machine spindle and drill 

makes significant improvement of the tool life 

because the cutting fluid can reach the tool-chip 

interface and also assist the chip ejection. 

Figure 7 shows the thrust force and torque in 

drilling at 183 m/min with internal cutting fluid 

supply before the drill breakage. Gradual increase of 

the thrust force and torque can be observed. 

Inverse Heat Transfer Solution of Tool 

Temperature in Drilling of CP Titanium 

Temperatures at four thermocouples are shown 

in Figure 8. Good matches between experimental 

and modeling temperature can be recognized at three 

drilling speeds. 

The temperature distribution near the tip of the 

spiral point drill after 12.7-mm depth of drilling the 

127 
II 

FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Table 4. Experiment design, summary of the Stage A and peak thrust force and torque, and energy consumption in 

Exps. I–IV 

Peripheral Cutting Exp. Drill type fluid cutting speed FA (N) TA (N – m) Fm (N) Tm (N – m) E (J) 

(m/min) 

0.57 s 183 m/min 

(14700 rpm) 

200 

100 

9.1 401–413 0.40–0.44 401–413 0.56–0.58 391–400

I HSS twist No 13.7 464–486 0.39–0.40 464–486 0.72–0.79 426–455 

18.3 390 0.36 390 0.80 453 

13.7 417–420 0.40–0.44 417–420 0.5–0.54 367–390

HSS twist External 18.3 430–446 0.39–0.41 430–446 0.70–0.76 387–427 

27.4 400 0.3 410 0.75 412 

HSS twist 390 0.36 390 0.80 453 

III WC-Co twist No 18.3 233–235 0.28–0.32 235–257 0.71–0.82 603–628 

WC-Co spiral 166–169 0.38–0.42 177–178 0.46–0.48 351–375 

No 178–185 0.24–0.27 178–185 0.82–1.47 497–596 

IV WC-Co spiral External 183 175–178 0.28–0.31 175–190 0.53–1.45 504–574 

Internal 162–210 0.25–0.55 182–230 1.07–1.17 586–783 

Table 5. WC-Co tool life (number of holes) at high 

drilling speed (183 m/min) 

Dry External cutting Internal cutting 

fluid fluid 

10 14 101 

Hole Number 1st 50th 100th 

A B C 

300 

Thrust force (N) Torque (N-m) 

0 

1.5

1

0.5

0

0 2 4 6 8 10 

Drilling distance (mm) 

Figure 7. Thrust force and torque of high-speed drilling 

Figure 8. Comparison of the measured and calculated 

temperature at four thermocouple locations in 

titanium drilling. 

titanium is shown in Figure 9. High temperature is 

concentrated along the cutting edge at the drill tip. 

High cutting speed generates high temperature 

in the drill. The peak temperature increases from 

445 to 893°C when the drill peripheral cutting speed 

is increased from 24.4 to 73.2 m/min. The peak temof Ti-6Al-4V at 183-m/min peripheral cutting perature is located on the cutting edge near the drill 

speed with internal cutting fluid supply. 

128 

margin for all three cutting speeds. 
Heavy Vehicle Propulsion Materials 

445 

409 

373 

337 

301 

299 

265 

229 

193 

24.4 m /min 

601 

544 

487 

430 

373 

315 

258 

201 

144 

48.8 /min 

893 

801 

709 

617 

525 

433 

341 

249 

157 

73.2 m/min 

Figure 9. Temperature distributions at the drill tip. 

Conclusions 

This study demonstrated that high-speed, costeffective machining of titanium and titanium alloys 

was possible. Exploratory evaluation of TiC-Ni3Al 

cermet for machining of titanium had been conducted. Cutting forces were measured and compared 

with the results when using WC-Co tool at different 

cutting speeds and feeds. The porosity of the material was improved with higher Ni3Al content, but the 

manufacturing technique and surface and cutting 

edge preparation of the tool inserts had been identified for further improvement. 

This study demonstrated the feasibility of costeffective drilling of Ti-6Al-4V with peripheral cutting speed over 180 m/min (600 sfpm) using 

commercially available advanced drill geometry and 

tool material. The limitation of HSS drill on cutting 

speed and material removal rate was apparent. The 

test of three drills at the same process conditions 

FY 2005 Progress Report 

demonstrated benefits of the spiral point drill design. 

The WC-Co Spiral drill outperformed the HSS 

Twist and WC-Co Twist drills in terms of thrust 

force, torque, energy consumption, and burr size. 

The importance of cutting fluid supply on drill life 

was identified. The improvement of drill life from 

dry to external cutting fluid supply was limited at 

high-speed drilling. Internal cutting fluid supply was 

proven to be critical on the drill life. 

The inverse heat transfer method was developed 

to analyze the tool temperature in titanium drilling. 

High drill temperature was observed in drilling titanium, particularly at high cutting speed. Temperature measured at various thermocouples provided 

not only the input for inverse heat transfer solution 

but also the validation of the finite-element drill 

temperature model. The proposed method has been 

validated by reasonably good agreement with 

experimental measurements. Detailed finite-element 

analysis of the drill temperature distribution was 

performed to gain insight in titanium drilling. 

References 

1. A. R. Machado and J. Wallbank, “Machining 

of Titanium and Its Alloys—A Review,” Proceedings of Institution of Mechanical Engineers, Part B: 

Management and Engineering Manufacture, 204(1), 

53–60 (1990). 

2. E. O. Ezugwu and Z. M. Wang, “Titanium 

alloys and their machinability—a review,” Journal 

of Materials Processing Technology, 68(3), 262–274 

(1997). 

3. X. Yang and C. R. Liu, “Machining Titanium 

and Its Alloys,” Machining Science and Technology, 

3(1), 107–139 (1999). 

4. M. Rahman, Y. S. Wong, and A. R. Zareena, 

“Machininability of Titanium Alloys,” JSME 

International Journal, Series C: Mechanical Systems, 

Machine Elements and Manufacturing, 46(1), 

107–115 (2003). 

5. M. C. Shaw, Metal Cutting Principles, 

Oxford University Press, New York, 1984. 

6. E. M. Trent and P. K. Wright, Metal Cutting, 

Butterworth-Heinemann, Boston, 2000. 

7. Machining Data Handbook, Metcut Research 

Associates, Inc., 1980 

8. X. Lin and J. Ni, “Drill Design For High 

Throughput Hole Making,” International Conference on Advanced Manufacturing Technology, 

June 16–18, 1999, Xi’an, China, 162–166 (1999). 

129 
FY 2005 Progress Report 

9. J. S. Agapiou and D. A. Stephenson, 

“Analytical and Experimental Studies of Drill 

Temperatures,” Journal of Engineering for Industry, 

116(1), 54–60 (1994). 

Presentations and Publications 

R. Li and A. J. Shih, “Finite Element Modeling 

of 3D Turning of Titanium,” presented at the 2004 

ASME IMECE Anaheim, California, November 

2004. 

Heavy Vehicle Propulsion Materials 

R. Li and A. J. Shih, “Finite Element Modeling 

of 3D Turning of Titanium,” International Journal 

of Advanced Manufacturing Technology (in print). 

R. Li and A. J. Shih, “Inverse Heat Transfer 

Solution of Tool Temperature in Titanium Drilling,” 

Journal of Manufacturing Science and Engineering 

(submitted). 

130
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

J. Advanced Cast Austenitic Stainless Steels for High-Temperature Components 

P. J. Maziasz, J. P. Shingledecker, and N. D. Evans 

Metals and Ceramics Division 

Oak Ridge National Laboratory 

P.O. Box 2008, MS-6115 

Oak Ridge, TN 37831-6115 

(865) 574-5082; fax: (865) 754-7659; e-mail: maziaszpj@ornl.gov 

M. J. Pollard 

Advanced Materials Technology 

Caterpillar Technical Center 

14009 Old Galena Rd. 

Mossville, IL 61552 

(309) 578-6133; fax: (309) 578-2953; e-mail: pollard_michael_j@cat.com 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Objectives 

• Commercialize new CF8C-Plus cast austenitic stainless steel. 

• Complete a high-temperature properties database for test specimens from commercial heats and determine the 

properties of prototype components. 

Approach 

• Support Caterpillar (CAT) evaluation efforts to upgrade turbochargers from cast iron to cast austenitic stainless 

steel. 

• Perform and evaluate direct comparisons between commercial CF8C-Plus steel and cast stainless steels used by 

component suppliers to establish performance benefits at 700–850°C. 

Accomplishments 

• Demonstrated that CF8C-Plus cast stainless steel is far superior to SiMo cast iron at 600°C and above. CF8CPlus also shows strength, creep resistance and aging resistance benefits compared with other commercial cast 

stainless steels being considered for heavy-duty diesel turbocharger housings at 750°C. 

• Successfully cast a good-quality 6700-lb gas-turbine end-cover from CF8C-Plus at MetalTek and cast 

>30,000 lb cast for various applications. 

Future Direction 

• Characterize the microstructure and properties of several different prototype CF8C-Plus steel components to 

facilitate various commercial applications. 

131 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

• Evaluate the properties of commercial heats of new CF8C-Plus Cu/W steel, which boosts tensile and creep 

strength at 750°C and above. 

Introduction 

Advanced heavy truck diesel engines are 

increasingly required to have higher fuel efficiency 

and reduced exhaust emissions without sacrificing 

durability and reliability. The most demanding 

normal duty cycles require exhaust manifolds and 

turbocharger housing materials to withstand temperatures ranging from 70 to nearly 800°C. Such 

materials must withstand prolonged, steady hightemperature exposure (which requires tensile 

strength and creep and oxidation resistance) as well 

as more rapid and severe thermal cycling (which 

requires resistance to aging and thermal fatigue). 

New emissions reduction technology and transient 

power excursions can push temperatures in these 

critical components even higher. 

Higher diesel engine exhaust temperatures push 

components made from SiMo ductile cast iron well 

beyond their current strength and corrosion limits. 

The first Oak Ridge National Laboratory 

(ORNL)/CAT cooperative research and development agreement (CRADA) (3y) developed a new, 

modified CF8C cast austenitic stainless steel (CF8CPlus) as an upgrade alternative to SiMo cast iron and 

performed the initial mechanical properties testing 

on a lab-scale heat. The CF8C-Plus steel was found 

to have outstanding creep resistance at 850°C. The 

second ORNL/CAT CRADA project (3y) capitalized on the initial success and began commercial 

scale-up of the new CF8C-Plus. The new material 

was found to have good castability, as well as much 

better mechanical properties at 600–850°C than 

standard CF8C steel, including resistance to aging, 

creep, and fatigue/thermal fatigue. The new CF8CPlus cast steel won a 2003 R&D 100 Award, and in 

2004 testing continued of different commercial 

casting process effects (static, centrifugal, and larger 

heats). This year, efforts to commercialize CF8CPlus expanded dramatically, and this CRADA project was extended for 2 more years. In 2004–2005, 

successful casting trials of CF8C-Plus were carried 

out for components ranging from exhaust manifolds 

to industrial turbine casing. In addition, the first 

commercial heats of CF8C-Plus Cu/W were made in 

2004–2005 to boost strength and creep resistance at 

750–850°C. Three companies have taken trial 

licenses and, to date, over 30,000 lb of the new steel 

has been produced commercially. 

Approach 

In 2005, three U.S. foundries obtained trial 

licenses for CF8C-Plus steel: MetalTek International, Wollaston, and Stainless Foundry and Engineering. CAT and ORNL both ordered additional 

heats of CF8C-Plus and cast pieces in various sizes, 

using both static casting in sand molds and centrifugal casting of rings of varying thickness, to enable 

more comprehensive mechanical properties testing. 

In addition, a large industrial gas turbine end-cover 

component cast for Solar Turbines (a CAT company) by MetalTek in FY 2004 was evaluated at the 

CAT Technical Center for quality and defects and 

then cut up for mechanical testing. ORNL also had 

MetalTek make a 500-lb commercial heat of the new 

CF8C-Plus Cu/W, selected for testing from screening of lab-scale heats, to evaluate further improvements in strength and creep-resistance above 750o

C. 

Finally, CAT began exploring CF8C-Plus steel and 

other commercial cast stainless steels to replace 

SiMo cast iron for turbocharger casing applications, 

which should enable increased durability at the 

increased engine temperatures needed to allow a 3% 

decrease in fuel consumption for on-highway trucks. 

Both CAT and ORNL do specific mechanical properties testing and microcharacterization in support of 

these various applications or efforts. 

Technical Progress 

Caterpillar 

The large industrial turbine end-cover cast of 

CF8C-Plus steel (8000-lb heat to yield 6700-lb finished component) by MetalTek for Solar Turbines 

was shipped to the CAT Technical Center, as shown 

in Figure 1. Visual and other inspection techniques 

showed no defects. Then sectioning was completed, 

and those initial observations also found no obvious 

cases of hot tearing at the side flange. This is significant because fully austenitic alloys such as 

CF8C-Plus are usually more susceptible to hot tearing than similar steel with a much higher δ-ferrite 

content (standard CF8C steel has 15–25% such 

132
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 1. A-6700-lb gas turbine combustor housing centrifugally cast from CF8C-Plus stainless steel 

by MetalTek. No indications of hot tearing 

have been observed (M. J. Pollard is pictured 

with the casting). 

content). Further nondestructive analysis will be performed to ensure a complete absence of internal hot 

tears. Fatigue and mechanical properties test samples of the bulk material from this housing have 

been produced, and mechanical testing began at the 

end of FY 2005. It is hoped it will show that the 

properties do not change significantly in thicksection components (slow cooling rates) compared 

with smaller test castings (faster cooling rates). 

The 21st Century Truck Partnership technical 

goal of developing emission-compliant engine systems for Class 7–8 highway trucks is an improvement in engine efficiency from 42 to 50% by 2010. 

CAT has determined that replacing SiMo cast iron 

exhaust components with CF8C-Plus cast stainless 

steels allows a 90°C increase in exhaust gas temperatures, which translates directly into a 3% 

increase in fuel economy while also significantly 

improving performance and durability. Efforts to 

produce prototype cast stainless exhaust manifolds 

and turbocharger housings began this year. New 

stainless steel diesel-engine turbocharger housings 

have been ordered for a small quantity off-road 

diesel engine application. Because of current supplier agreements, these first trial turbo-housings are 

being cast from a standard commercial stainless 

alloy, KN2, produced by Diado in Japan. However, 

another part of this effort is side-by-side testing of 

KN2 and CF8C-Plus cast steels at ORNL to directly 

compare properties. This prototyping effort will give 

the designers at CAT confidence with cast stainless 

steel as well as experience with the necessary design 

changes for cast stainless. 

ORNL 

In 2004, ORNL began the creep testing that 

showed an overwhelming advantage of CF8C-Plus 

steel over SiMo cast iron in terms of tensile and 

creep-rupture strength and thermal fatigue resistance 

above 550–600°C. That creep-rupture testing was 

expanded and completed and includes comparison of 

CF8C-Plus with standard CF8C steel and Ni-resist 

austenitic cast-iron (an upgrade for some applications relative to SiMo cast-iron). The results of various creep rupture tests at different temperature and 

stress levels are plotted as a function of the LarsonMiller parameter (LMP) in Figure 2. CF8C-Plus cast 

steel is much stronger than either SiMo or Ni-resist 

cast-irons, and it is almost twice as strong as standard CF8C steel. The latter comparison is important 

to the potential gas- and steam-turbine applications, 

as well as many other chemical/petrochemical, 

energy, or metals processing applications. 

Creep-rupture ductility is shown in Figure 3. 

Clearly, CF8C-Plus has much more creep ductility 

than standard CF8C steel, despite also being 

stronger. This is due to the lack of δ-ferrite in the ascast structure of CF8C-Plus, which makes it resistant 

to the formation of embrittling σ-phase during 

aging. 

Finally, comparison of CF8C-Plus cast stainless 

with the best commercially available creep-resistant 

austenitic stainless steels and alloys on an LMP plot 

(Figure 4) shows that CF8C-Plus is as strong as or 

stronger than NF709 austenitic stainless alloy, and 

comes close to the creep-strength of alloy 617, a NiCr-Co superalloy, particularly at the highest 

temperatures. 

133 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

300 

S

T

R

ESS (

MPa) 100 

80 

60 

40 

20 

CF8C 

CF8C-Plus 

SiMo Cast Iron 

Ni-Resist 

CF8C CF8C-Plus 

SiMo Cast Iron 

Ni-Resist 

18,000 20,000 22,000 24,000 26,000 28,000 

LMP (C=20) 

Figure 2. Creep-rupture stress plotted vs LMP for various ORNL creep-rupture 

tests at 500–850°C in air of various commercial heats of SiMo ductile and 

Ni-resist austenitic cast irons and standard CF8C and CF8C-Plus. CF8CPlus shows significant creep-strength advantage over the standard steel or 

the cast irons. 

CF8C and CF8C-Plus Creep-Rupture Ductility (6-14-2005)

45

Rupture Elongation (%) 

40 

35 

30 

25 

20 

15 

10 

5 

1 10 100 1,000 10,000 105 

CF8C 

CF8C-Plus 

TIME (hrs) 

Figure 3. Comparison of creep rupture ductility data vs rupture time for 

cast CF8C-Plus and standard CF8C (same data as in Figure 2) 

for creep-rupture testing in air at 650–850°C and 35–200 MPa. 

134
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

TEMPERATURE (oC) FOR RUPTURE IN 100,000 HOURS 

600 640 680 720 760 800 840 

300 CF8C-Plus: Solid Diamonds 

CF8C-Plus+Cu/W: Cross 

(Arrows indicate test in progress) 

NF709 

ST

R

ESS (

M

Pa) 100 Alloy 617 

80

60

TP347HFG 

40 

Super 304H 

20 

21,000 22,000 23,000 24,000 25,000 26,000 27,000 28,000 

LMP (C=20) 

Figure 4. Creep-rupture stress plotted vs LMP for various ORNL creep-rupture tests 

at 700–850°C in air of CF8C-Plus and CF8C-Plus Cu/W. CF8C-Plus has 

better creep-strength than most wrought heat-resistant austenitic stainless 

steels and alloys and compares well with alloy 617, a Ni-Cr-Co superalloy. 

In FY 2005, ORNL began testing new commer- comparison. The CF8C-Plus Cu/W has higher yieldcial scale-up heats of CF8C-Plus Cu/W produced by strength than CF8C-Plus at 600–900°C, and the few 

MetalTek. A 500-lb heat of the new CF8C-Plus creep-tests conducted indicate better creep strength 

Cu/W was cast as centrifugal rings and as kiel as well (Figure 4). Comparison of creep-strain 

blocks in sand, together with similar castings of versus time plots for tests at 750°C (Figure 5) 

CF8C-Plus and standard CF8C steels for clearly shows that the new CF8C-Plus Cu/W has a 

30 

Creep Test: 750o

C, 100MPa 

25 Commercial Centrifugal Castings 

CREE

P S

T

RAIN (%) 

CF8C 

20 CF8C-Plus 

CF8C-Plus+Cu/W 

15 

10 

5 

0 

0 500 1,000 1,500 2,000 

TIME (hours) 

Figure 5. Plots of creep-strain vs time for heats of standard CF8C, CF8C-Plus and 

the new CF8C-Plus Cu/W, all tested in air at 750°C. Clearly the CF8CPlus Cu/W has much more creep-resistance, with an extended secondary 

creep-regime and a very low steady-state creep rate. 

135 
FY 2005 Progress Report 

prolonged period of creep at a very low creep rate 

(secondary creep regime) relative to standard CF8C 

or CF8C-Plus steels. 

In 2006, ORNL will continue testing and aging 

studies on the new CF8C-Plus Cu/W steel. ORNL 

will also finish and analyze the side-by-side comparison of CF8C-Plus and the Diado KN2 cast 

stainless steels to support and enable CAT’s development of cast stainless steel turbocharger housings. 

Conclusions 

ORNL and CAT have finished the initial commercialization phase for CF8C-Plus cast stainless 

steel, and have begun the second stage of producing 

more and larger heats of steel. They will expand the 

preliminary testing and include long-term aging, 

creep-rupture and more complex mechanical testing 

(fatigue, thermal-fatigue and creep-fatigue) and 

begin commercial component casting trials. To date, 

more than 30,000 lb of commercial CF8C-Plus 

stainless steel has been cast for various applications, 

and all prototype component castings have been 

successful on the first trial as a result of good castability. Efforts to expand commercial applications 

and facilitate licensing continue. 

For advanced diesel-engine exhaust manifold 

and turbocharger applications, CF8C-Plus steel has 

an overwhelming high-temperature strength and 

creep-resistance advantage over conventional 

exhaust component materials, such as SiMo or Niresist cast irons. CF8C-Plus also has enough of a 

creep-rupture strength advantage over standard 

CF8C steel and other heat-resistant stainless steels 

and alloys that it is being considered and tested for 

gas turbines, advanced steam turbines, and various 

other applications. 

Based on ORNL lab-scale trial heats, a new heat 

of CF8C-Plus Cu/W has been melted commercially 

by MetalTek. Testing began this year and will continue next year. The CF8C-Plus Cu/W has better 

yield strength than CF8C-Plus at 600°C and above, 

and preliminary creep testing at 750°C shows 

improved creep resistance as well. 

Publications/Presentations 

P. J. Maziasz, I. G. Wright, J. P. Shingledecker, 

T. B. Gibbons, and R. R. Romanosky, “Defining 

the Materials Issues and Research Needs for UltraSupercritical Steam Turbines,” pp. 602–622 in Proc. 

4th Internat. Conf. on Advances in Materials 

Heavy Vehicle Propulsion Materials 

Technology for Fossil Power Plants, ASMInternational, Materials Park, Ohio, 2005. 

J. P. Shingledecker, P. J. Maziasz, N. D. Evans, 

and M. J. Pollard, “Creep Behavior of a New Cast 

Austenitic Alloy,” to be published in Proc. ECCC 

Conference on Creep and Fracture in High Temperature Components—Design and Life Assessment 

Issues, London, UK, September 12–14, 2005. 

J. P. Shingledecker, P. J. Maziasz, N. D. Evans, 

and M. J. Pollard, “Alloy Additions for Improved 

Creep-Rupture Properties of a Cast Austenitic 

Alloy,” pp. 129–138 in Proc. Conf. Creep Deformation and Fracture, Design, and Life Extension, The 

Materials Society, Warrendale, PA, 2005. 

P. J. Maziasz, J. P. Shingledecker, N. D. Evans, 

and M. J. Pollard, “Update on ORNL/CAT CRADA 

on CF8C-Plus Cast Stainless Steel: Progress on 

Commercial Scale-Up in 2004/2005,” presented at 

DOE FreedomCAR and Vehicle Technologies 

Review at Oak Ridge National Laboratory, March 2, 

2005. 

M. J. Pollard, P. J. Maziasz, and J. P. 

Shingledecker, “Development of Low-Cost Cast 

Austenitic Stainless Steel for Diesel Engine and Gas 

Turbine Applications,” presented at DOE Project 

Review at Caterpillar, Peoria, Illinois, April 27, 

2005. 

P. J. Maziasz, J. P. Shingledecker, N. D. Evans, 

and M. J. Pollard, “Development, Properties, and 

Applications of CF8C-Plus,” presented during a visit 

to Stainless Foundry and Engineering, Inc., 

Milwaukee, Wisconsin, June 16, 2005. 

P. J. Maziasz, J. P. Shingledecker, N. D. Evans, 

and M. J. Pollard, “Development, Properties, and 

Applications of CF8C-Plus,” presented during a visit 

to MetalTek International, Waukesha, Wisconsin, 

June 27, 2005. 

J. P. Shingledecker, P. J. Maziasz, N. D. Evans, 

and M. J. Pollard, “Creep Behavior of a New Cast 

Austenitic Alloy,” plenary session talk at the 

European Creep Collaborative Committee (ECCC) 

Conference on Creep and Fracture in HighTemperature Components—Design and Life 

Assessment Issues, London, UK, September 12–14, 

2005. 

Special Recognitions and Awards/Patents Issued 

The review panel for the FY 2005 DOE Heavy 

Vehicles Materials Program Merit Review and Peer 

Evaluation Meeting awarded this project the highest 

136
Heavy Vehicle Propulsion Materials 

overall score in the Heavy Vehicles Propulsion 

Materials sub-program. The project also received the 

highest individual evaluation category grades for 

relevance, technical accomplishments, and 

technology transfer. 

The original patent application for both CF8CPlus and CN-12-Plus filed in 2000 by Caterpillar 

was split and continued as two patents: “Heat and 

Corrosion Resistant Cast CN-12 Type Stainless 

Steels With Improved High Temperature Strength 

FY 2005 Progress Report 

and Ductility” (US 2003/0084967) and “Heat and 

Corrosion Resistant Cast CF8C Stainless Steels 

With Improved High Temperature Strength and 

Ductility” (US 2003/0056860), both by P. J. 

Maziasz and R. W. Swindeman (ORNL), 

T. McGreevy (University of Bradley/CAT), M. J. 

Pollard, and C. W. Siebenaler (CAT). Actions to 

obtain final patent approval continued in FY 2005. 

137 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

K. Viscoplastic Twinning in an Ultrafine TiAl/Ti3Al Laminate Composite 

Luke Hsiung 

Chemistry and Materials Science Directorate 

Lawrence Livermore National Laboratory 

L-352, P.O. Box 808 

Livermore, CA 9455-9900 

(925)-424-3125; fax: (925) 424-3815; email: hsiung1@llnl.gov 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: Lawrence Livermore National Laboratory 

Objectives 

• Exploit thermomechanical-processing techniques to fabricate TiAl/Ti3Al nanolaminate composites with lamella 

widths down to nanometer length scales. 

• Characterize microstructure and elevated-temperature creep resistance of the nanolaminate composites. 

• Investigate the fundamental interrelationships among microstructures, alloying additions, and mechanical 

properties of the nanolaminate composites to achieve the desired properties of the composites for hightemperature structural applications. 

Approach 

• Employ in-situ laminate composites with nominal compositions of Ti-47Al-2Cr-2Nb, Ti-46Al-3Nb-1W-0.1B, 

and Ti-46Al-3Nb-2W-0.1B (at. %). The in-situ laminate composites were fabricated at Oak Ridge National 

Laboratory (ORNL) by hot extrusion of cast alloys at 1350°C. 

• Conduct creep tests in a dead-load creep machine with a lever arm ratio of 16:1. Tests were performed in air in 

a split furnace with three zones at 760 and 815°C. 

• Examine the microstructures of creep-deformed samples using a JEOL-200CX transmission electron 

microscope (TEM). 

Introduction 

Two-phase [TiAl (γ-L10) and Ti3Al (α2-DO19)] 

fully lamellar TiAl alloys have recently attracted 

much attention because of their low density (ρ = 

3.9 g/cc), high specific strength, adequate oxidation 

resistance, and good combination of ambienttemperature and elevated-temperature mechanical 

properties, which are of interest for aerospace and 

transportation applications such as high-temperature 

components in turbine and combustion engines. 

Through alloy design and microstructural 

refinement, significant progress has been made 

recently to improve both room-temperature 

ductility/toughness and high-temperature creep 

resistance of the material.1–5 One of the interesting 

deformation substructures of the laminate material 

with an ultrafine microstructure is the formation of 

139 
FY 2005 Progress Report 

high-density twin lamellae with the widths in the 

sub-micrometer or nanometer length scale.5

 The 

most unique characteristic of the twinning 

phenomenon observed from the TiAl/Ti3Al laminate 

material is that it can even take place at elevated 

temperatures under loading conditions such that the 

steady creep rate is as low as 10−7

 s−1 and with a 

stress level much lower than its yield strength. This 

viscoplastic twinning phenomenon is different from 

the twinning phenomenon that commonly occurs in 

conventional metals and alloys deformed under lowtemperature and/or high strain-rate conditions. It has 

been reported previously that the occurrence of 

{111}<112>-type twinning in the TiAl/Ti3Al 

laminate material is intimately related to the 

dissociation of dislocations in laminate interfaces 

and the homogeneous glide of 1/6 < 112 > twinning 

dislocations in the {111} planes of γ lamellae.6 

Results obtained from more rigorous TEM analyses 

are reported here in order to further elucidate the 

underlying mechanisms for the occurrence of 

viscoplastic twinning in the TiAl/Ti3Al laminate 

material. 

Approach 

The ultrafine two-phase [TiAl-(γ)/Ti3Al-(α2)] 

lamellar alloy was fabricated by a powder 

metallurgy process, which involves hot extrusion of 

gas-atomized titanium aluminide powder with a 

nominal composition of Ti-47Al-2Cr-2Nb (at. %) at 

1400°C. After extrusion, the material was stress 

relieved at 900°C in a vacuum (∼10−4

 Pa) for 2 h. 

Creep tests were conducted at 760°C in a dead-load 

creep machine with a lever arm ratio of 16:1. 

Detailed information regarding the creep experiment 

and experimental data were reported elsewhere.5

 For 

the current study, the deformation substructure of a 

specimen tested at 138 MPa with a steady creep-rate 

of 2.0 × 10−9

 s−1, and that of a specimen tested at 

518 MPa with a steady creep-rate of 2.5 × 10−7

 s−1 

were examined. TEM foils were prepared by twinjet electropolishing in a solution of 60 vol % 

methanol, 35 vol % butyl alcohol and 5 vol % 

perchloric acid at ∼15 V and –30°C. The 

microstructures of both as-fabricated material and 

tested samples were examined using a JEOL-200CX 

TEM equipped with a double-tilt goniometer stage. 

Images of microstructures were recorded using highresolution lattice imaging techniques and weakHeavy Vehicle Propulsion Materials 

beam dark field (WBDF) imaging techniques under 

g (3g) two-beam diffraction conditions with the 

deviation factor ω (= ξg s) > ∼1, where ξg is the 

extinction distance and s is the deviation distance 

from the exact Bragg position. The g⋅b invisibility 

criteria used for determining the Burgers vector of 

Shockley partials are described as follows:7 

(a) invisible if g⋅b = 0 or ±1/3; (b) invisible if g⋅b = 

−2/3 but visible if g⋅b = +2/3 provided the deviation 

factor ω > ∼1; (c) invisible if g⋅b = +4/3 but visible 

if g⋅b = −4/3, provided the deviation factor ω > ∼1. 

Results 

Movement and Pileup of Interfacial Dislocations 

The results of an in-situ TEM experiment that 

directly observed cooperative movement of 

interfacial dislocations in a TiAl/Ti3Al laminate 

sample under room-temperature straining conditions 

have been reported previously.

8

 Figure 1 is a typical 

example showing that the interfacial dislocations can 

be blocked by impinged lattice dislocations during 

their movement along a laminate interface. Here, a 

dislocation network formed by the reaction between 

impinged lattice dislocations and interfacial 

dislocations can be readily seen in a γ/α2 interface. 

Figure 2 shows the results of an in-situ TEM 

experiment, which evidence the motion and pile-up 

Figure 1. A WBDF TEM image shows that the 

movement of interfacial dislocations is 

hampered because of the formation of 

dislocation networks as a result of the 

reaction between impinged lattice 

dislocations (LD) and interfacial 

dislocations (ID). 

140
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

(a) (b) 

Figure 2. Consecutive in-situ TEM images show (a) the movement of an interfacial dislocation array in a 

laminate interface (time span: 30 s), and (b) the movement and pile-up of interfacial dislocations in 

a laminate interface adjacent to the tip of an α2 fragment (time span: 30 s). 

of interfacial dislocations during electron-beam 

heating of a thin foil (prepared from a sample precrept at 760°C and 138 MPa) that contains residual 

stress. The original idea of this experiment assumes 

that the relaxation of residual stress can result in the 

motion of interfacial dislocations. Since the purpose 

of the study is only to demonstrate the interfacial 

dislocation motion, the amount of temperature 

increase by beam heating is not a major concern. It 

is also noted that local heating can be achieved when 

an electron beam is focused on a small spot the size 

of several micron meters. Figure 2(a) shows the 

cooperative motion of a dislocation array of eight 

interfacial dislocations on a laminate interface. It is 

interesting to note that the motion of each dislocation was in a viscous drifting or pinning/unpinning 

fashion, and each of them has a different drifting 

velocity. Here, after beam focusing (heating) for 

30 s, the No. 1 leading dislocation of the array 

moved about 375 nm, and the No. 8 trailing 

dislocation moved about 425 nm. This suggests that 

each dislocation has different mobility as a result of 

the solute-dragging effect. Figure 2(b) shows the 

motion of interfacial dislocations within a 

dislocation pileup on a laminate interface. Here, the 

head of the dislocation pileup is at the tip of a 

faulted α2-lamella adjacent to the No. 1 dislocation. 

Notice that the dislocation spacing increases with 

increasing distance from the head of the pile-up, and 

those dislocations close to the head of the pile-up 

moved slower than those distant from the head. For 

instance, the No. 10 and No. 20 dislocations moved 

about 170 nm, and the No. 35 and No. 45 

dislocations moved about 250 nm after beam focusing (heating) for 30 s. That is, during beam heating, 

dislocations in the pile-up start to squeeze, which 

causes the increase of dislocation density within the 

pile-up. 

These in-situ observations of interfacial dislocation motion and pile-up are rationalized below. 

Since the thin foil contains residual stress, each dislocation starts to move under a heating condition 

with a velocity v = MF, where v is the (drift) velocity of dislocation, M is the mobility of dislocation, 

and F is the effective force acting on the dislocation. 

An explicit expression for the dislocation mobility 

M limited by solute drag can be found in Ref. 9: 

DsΩ M = , βb2

CokT 

141 
FY 2005 Progress Report 

where β is a constant, b is the length of Burgers 

vector, Co is the solute concentration, Ds is the 

solute diffusion coefficient, and Ω is the atomic 

volume of the solute. The effective force (F) acting 

on the i-th dislocation in a dislocation array or pileup can be expressed as 

F = τb − μb2 

∑

n 1 , 10 

2π(1 − ν) j=0 xi − xj 

j≠i 

where the term 

n μb2 1 ∑ 2π(1 − ν) j=0 xi − xj 

j≠i 

represents the sum of internal forces acting on the 

i-th dislocation by the other dislocations of the pileup; xi – xj represents the distance between the i-th 

and j-th dislocations. Accordingly, v increases as a 

result of the increase of dislocation mobility as 

temperature (T) increases during beam heating, and 

v diminishes as temperature decreases to an ambient 

temperature. The velocity of dislocations close to the 

head of the pile-up becomes slower, indicating that 

the internal stress acting on the dislocations 

increases as the dislocations move closer to the 

head, and therefore reducing the effective force (F) 

acting on the dislocations. 

Viscoplastic Deformation Twinning 

When an alloy sample was creep-deformed at 

760°C with a constant applied stress of 518 MPa 

under a strain rate of 3 × 10−7

 s−1, a deformation substructure associated with deformation twins was 

developed within γ lamellae. Typical ( 111 ) [211]­

type twin lamellae formed within γ lamellae are 

shown in Figure 3. Notice that one of the twin 

lamellae in Figure 3(a) (labeled by an arrow) was 

still growing toward another interface. Similar features were also observed elsewhere in the same 

sample. This suggests that the laminate interface is 

the major source for the occurrence of viscoplastic 

deformation twinning. The viscoplastic twinning 

phenomenon can be viewed as a relaxation process 

to relieve the local stress concentration caused by 

the pile-up of interfacial dislocations. The effective 

stress (τe) at the head of the pile-up of n dislocations 

Heavy Vehicle Propulsion Materials 

can be evaluated by τe = nτI (Ref. 11), where τi is 

the resolved shear stress acting on the interface. To 

relieve the stress concentration, deformation twinning in γ layers is therefore taking place by a dislocation reaction based upon a stair-rod cross-slip 

mechanism, as schematically illustrated in 

Figure 3(b). For the formation of a ( 111 )-type twin 

lamella, the corresponding dislocation reaction is 

1/6[ 121 ] (111) [b1] → 1/6[011] (100) [b2] + 

1/6[ 112 ] ( 111 ) [b3]. The ( 111 )-type twin is 

accordingly formed within γ lamella by successively 

dissociating the [b1] interfacial dislocations in the 

pile-up and emitting the [b3] twinning dislocations 

into the ( 111 ) γ plane and leaving the [b2] stair-rod 

dislocations in the (100)γ plane. Figures 3(c) and 

3(d) show the edge-on views of the ( 111 )-type twin 

lamellae, in which their widths were measured as 

varying between 3 and 20 nm. 

The formation of [b2] stair-rod dislocations at 

the intersections between twin and α2 lamellae is 

evidenced in Figures 4(a) and 4(b), in which the 

array of 1/6[011] stair-rod dislocations becomes 

invisible [Figure 4(a)] or visible [Figure 4(b)] when 

a reflection vector (g) 200 or 021 is used for imaging. It is noted that the individual stair-rod dislocation is not resolvable because of a narrow distance 

(0.25 nm) between two stair-rod dislocations. The 

formation of [b3] twinning dislocations for the 

( 111 )-type twin lamella is evidenced in Figures 4(c) 

and 4(d), in which the array of 1/6[112] twinning 

dislocations becomes visible [Figure 4(c)] or invisible [Figure 4(d)] when a reflection vector of 111 or 

202 is used for imaging. The significance of the 

proposed mechanism is to reveal that there are 

several barriers to be overcome in order to activate 

the twinning reaction. These barriers include (1) the 

repulsive force between the interfacial (Shockley) 

and stair-rod dislocations, (2) the increase of line 

energy due to the dislocation dissociation, and 

(3) the increase of interfacial energy due to the 

formation of twin faults. Among which the repulsive 

force (F) between the interfacial (Shockley) and 

stair-rod dislocations is considered to be the ratelimiting process. That is, a critical stress (τc) is 

required to activate for the twinning reaction, i.e., 

to activate the dissociation reaction for twinning. 

Assuming an isotropic elasticity, the critical stress 

(τc) can be approximately evaluated as 

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Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

(a) (b) 

(c) (d) 

Figure 3. A dark-field TEM image shows several ( 111 ) type twin lamellae formed 

within a γ lamella (a); a schematic illustration of the nucleation of a 

( 111 ) type twin lamella (γt) from a γ/α2 interface (b), where b1, b2, and b3 

denote the interfacial, stair-rod, and twinning dislocations, respectively; 

and edge-on views of twin lamellae show the widths of the ( 111 )-type 

twin lamellae (c) and (d). 

143 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

(a) (b) 

 (c) (d) 

Figure 4. WBDF TEM images show the g⋅b contrast visibility analyses 7 for the arrays of [b2] 

1/6[011] stair-rod dislocations formed at the intersections (indicated by arrows) between 

the ( 111 )-type deformation twins (DT) and α2 lamellae: (a) Z (zone axis) ≈ [ 012 ], g = 

200, g⋅b = 0 (invisible), (b) Z ≈ [ 012 ], g = 021, g⋅b =1/2 (visible); and the analyses for 

the arrays of [b3] 1/6[112] twinning dislocations of the (111 )-type deformation twins: 

(c) Z ≈ [ 1 2 1 ], g = 111, g⋅b = 2/3, (visible), and (d) Z ≈ [ 1 2 1 ], g = 202 , g⋅b = 1/3 

(invisible). 

τcb1 = F ≈ ημb1b2 cos θ/2πr, where η = 1 for screw between ( 111 ) and (100) planes; and r (0.25 nm) is 

dislocations and η = 1/(1 – υ) for edge dislocations; the distance between the two dislocations. The 

υ is Poisson’s ratio (∼0.3); μ is shear modulus critical stress (τc) required for twinning is thus 

(∼56 GPa at 760°C) (Ref. 12); b1 = 0.163 nm; and approximately 1.95 GPa for screw dislocations and 

b2 = 0.094 nm for Shockley and stair-rod 2.79 GPa for edge dislocations. 

dislocations, respectively; θ (= 54.4°) is the angle 

144
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

In the case of a low applied stress (138 MPa), 

the resolved shear stress (τi) acting on the interface 

plane is about 69 MPa (assuming Schmid factor = 

0.5), and the number of pile-up dislocations required 

to create a sufficient stress concentration (i.e., τe = 

nτi ≥ τc) for twinning to occur is n ≥ 29. In the case 

of a high applied stress (518 MPa), the resolved 

shear stress acting on the same interface plane is 

about 259 MPa, and the number of pile-up dislocations required for twinning to occur is n ≥ 8. Thus 

the accumulation creep strain required to bring on 

the viscoplastic twinning is smaller and therefore 

easier to reach. Besides, the average velocity of dislocations is relatively faster when tested under a 

high applied stress, since the strain rate is higher, 

which gives a shorter time to develop a pile-up configuration when tested under high applied stresses. 

That is, it is more feasible that the viscoplastic twinning will take place when the laminate is creepdeformed under high applied stresses. A recent 

report on deformation twinning of the same material 

creep-deformed under different loading conditions 

has evidenced that the twinning occurred shortly 

after the creep strain reached 0.67% under a high 

applied stress (518 MPa), whereas the twinning took 

place when the creep strain reached ∼1.5% under a 

low applied stress (138 MPa).13 

Conclusion 

The viscoplastic twinning phenomenon, which 

takes place within an ultrafine TiAl/Ti3Al laminate 

composite subjected to an elevated-temperature 

creep deformation, has been studied. The viscoplastic twinning phenomenon is found to be intimately 

related to the movement, pile-up, and dissociation of 

the pre-existing dislocation arrays in laminate interfaces. When the arrays of interfacial dislocation 

(e.g., 1/6<112> Shockley partials) move along the 

interfaces, obstacles such as impinged lattice dislocations, and interfacial ledges can impede their 

movement. The dislocation arrays start piling up 

behind an obstacle because an individual Shockley 

partial is energetically unfavorable to undergo crossslip or climb that otherwise will generate stacking 

fault on its wake if it moves away from the interface. 

Thus the arrays of interfacial dislocations can easily 

be piling up even at elevated temperatures. The 

viscoplastic twinning within the creep-deformed 

laminate can accordingly be rationalized as a relaxation process to dissipate the internal stress resulting 

from dislocation pile-up. The twinning process 

involves the dissociation of interfacial dislocation 

into stair-rod (or residual) and twinning dislocations, 

i.e., 1/6[ 1 2 1 ] (111) → 1/6[011] (100) + 1/6[ 1 12 ] 

( 111 ) for ( 111 ) [ 1 12 ]–type twinning. It is concluded that the dislocation pile-up configuration is a 

prerequisite to begin the process of deformation 

twinning, which is more likely to occur under lowtemperature and high strain-rate deformation conditions for conventional metals and alloys. The occurrence of time-dependent viscoplastic twinning in the 

TiAl/Ti3Al laminate reveals that the dislocation pileup configuration can be readily formed along laminate interfaces even at elevated temperatures. 

References 

1. Y-W. Kim and D.M. Dimiduk, JOM, 43(8), 

40 (1991). 

2. Y–W. Kim, Acta Metall. Mater., 40, 1121 

(1992). 

3. C. T. Liu, J. H. Schneibel, P. J. Maziasz, J. L. 

Wright, and D. S. Easton, Intermetallics, 4, 429 

(1996). 

4. C. T. Liu, P. J. Maziasz, and J. L. Wright, 

Mat. Res. Soc. Symp. Proc., 460, 83 (1997). 

5. L. M. Hsiung and T. G. Nieh, Matls. Sci. & 

Engrg., A364, 1 (2004). 

6. L. M. Hsiung and T.G. Nieh, Matls. Sci. & 

Engrg., A239-240, 438 (1997). 

7. J. W. Edington, Practical Electron 

Microscopy in Materials Science, Van Nostrand 

Reinhold, New York, 1976. 

8. L. M. Hsiung, A. J. Schwartz, and T. G. 

Nieh, Intermetallics, 12, 727 (2004). 

9. E. Arzt, M. F. Ashby, and R. A. Verrall, Acta 

Metall,. 31, 1977 (1983). 

10. J. Weertman, J. Appl. Phys., 28, 1185 

(1957). 

11. J. D. Eshelby, F. C. Frank, and F. R. N. 

Nabarro, Phil. Mag., 42, 351 (1951). 

12. S. C. Huang and J. C. Chesnutt, 

Intermetallic Compounds: Principles and Practice 

2, eds. J. H. Westbrook and R. L. Fleischer, John 

Wiley & Sons, p. 73, 1995. 

13. J. G. Wang, L. M. Hsiung, and T. G. Nieh, 

Scripta Mater., 39, 957 (1998). 

Presentations and Publications 

L. M. Hsiung and T. G. Nieh, “In-Situ TEM 

Study of Interface Sliding and Migration in an 

145 
FY 2005 Progress Report 

Ultrafine Lamellar Structure,” presented at the 2005 

MRS spring meeting. 

L. M. Hsiung, “TEM Study of Microstructures 

in a Creep-Deformed TiAl/Ti3Al Nanocomposite,” 

Heavy Propulsion Materials Program Quarterly 

Progress Report for October through December 

Heavy Vehicle Propulsion Materials 

2004, U.S. Department of Energy, Washington, 

D.C., 2004. 

L. G. Zhou, L. M. Hsiung, and Hanchen Huang, 

“Nucleation and Propagation of Deformation Twin 

in Polysynthetically Twinned TiAl,” Computer 

Modeling in Engineering and Sciences, 6(3) (2004) 

245–252. 

146
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

L. Ti-6Al-4V Billet Feedstock Manufacture and Evaluation 

S. M. Abkowitz, S. Abkowitz, and H. Fisher 

Dynamet Technology, Inc. 

Eight A Street 

Burlington, MA 01803 

(781) 272-5967; fax: (781) 229-4879; e-mail: sabkowitz@dynamettechnology.com 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: Dynamet Technology, Inc., Burlington, Massachusetts 

Objectives 

• Develop low-cost Ti-6Al-4V billet feedstock using a blend of titanium (Ti) and alloy powders and inexpensive 

Ti-6Al-4V machine turnings, or using 100% turnings. 

• Evaluate this low-cost Ti alloy feedstock as a starting billet material for casting, forging, and extrusion 

operations. 

Approach 

• Develop a procedure for producing low-cost Ti billet using a combination of inexpensive Ti alloy machine 

turnings and Ti and alloy powders, or 100% machine turnings. Evaluate the billet for density and 

microstructure. 

• Subject the feedstock billets to casting, forging, and extrusion operations. The three operations will yield Ti6Al-4V test bars. 

• Evaluate the quality of the test bars produced through chemistry and mechanical testing. 

Accomplishments 

• Produced a high-density billet containing 60% turnings/40% Ti alloy powder that was produced using a core of 

100% turnings, surrounded by a “can” made from Ti alloy powder. 

• Produced a billet using 100% turnings, achieving full density, using cold+hot isostatic pressing (CHIP) 

technology with conventional canning in mild steel for the HIP operation. 

• Extruded a billet, composed of 100% turnings surrounded by a “can” made from Ti alloy powder, to high 

density barstock. 

• Cast plates from a billet of 100% Ti-6Al-4V machine turnings. The tensile strength and ductility of the resultant 

cast plate exceeded the specified properties for Ti-6Al-4V castings. 

• Determined that contaminants in the turnings caused low ductility in the extruded barstock and characterized 

these contaminants by metallographic analysis and analysis of tensile fracture surfaces using SEM/EDS. 

• Demonstrated that the contaminants can be eliminated by melting the low cost feed stock. 

147 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

• Preformed economic analyses that showed that extruded bar can be produced from billet cast from low cost 

feedstock at a savings of 48 to 58% over conventional extrusions. 

Future Direction 

• Demonstrate that bar meeting the typical mechanical properties of conventional Ti-6Al-4V can be extruded 

from cast extrusion billet produced by the low cost feedstock process in order to realize the commercial 

potential of this technology. 

Summary 

Dynamet Technology has developed a novel 

method of producing low-cost Ti alloy billet from Ti 

scrap turnings by powder metallurgical processing. 

Dynamet’s process consists of consolidating the 

turnings using a combination of cold isostatic 

pressing and vacuum sintering, followed by HIP, if 

necessary (Figure 1). 

The result is feedstock that can be used to 

produce cast extrusion billet that can then be used to 

produce extrusions. This process promises to 

produce extrusions that are a low-cost alternative to 

extrusions produced from conventional ingot. 

Dynamet’s economic analysis estimates the cost 

savings in producing extrusions at 48 to 58%. The 

use of low-cost billet made by Dynamet’s process in 

manufacturing Ti alloy parts for heavy duty vehicles 

will save energy, increase payloads, and reduce 

emissions. Indeed, the availability of low-cost Ti 

would make Ti attractive to any industry that can 

(a) (b) 

(c) (d) 

Figure 1. The dynamet powder metallurigical process (a) Ti-6Al-4V turnings, (b) LCFP billet, (c) cast test 

plates, and (d) extrusion produced from LCFP billet. 

148
Heavy Vehicle Propulsion Materials 

benefit from the metal’s advantages as a structural 

material. Dynamet’s process would also result in 

significant energy reduction by recycling Ti turnings 

and thus reducing the industry’s dependence on 

energy-intensive Kroll process sponge and the 

subsequent more extensive metalworking. 

The next step is to demonstrate that bar meeting 

the typical mechanical properties of conventional 

Ti-6Al-4V can be extruded from cast extrusion billet 

FY 2005 Progress Report 

produced by the low cost feedstock process in order 

to realize the commercial potential of this 

technology (Figure 2). 

Presentations and Publications 

Dynamet Technology, Inc., Ti-6Al-4V Billet 

Feedstock Manufacture and Evaluation—Final 

Report, Burlington, Massachusetts. 

Figure 2. Bar with the properties of Ti-6 Al-4V can be developed from cast extrusion 

billets produced by the low-cost feedstock process. 

149 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

M. High-Density Infrared Surface Treatment of Materials for Heavy-Duty 

Vehicles 

T. N. Tiegs, F. C. Montgomery, D. C. Harper, J. O. Kiggans, and C. A. Blue 

Oak Ridge National Laboratory 

P.O. Box 2008, MS-6087 

Oak Ridge, TN 37831-6087 

(865) 574-5173; fax: (865) 574-4357; e-mail: tiegstn@ornl.gov 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Prime Contract No: DE-AC05-00OR22725 

Objectives 

• Use high-density infrared (HDI) technology to produce corrosion and/or wear resistance coatings on metal 

substrates. 

• Use lighter or more cost-effective bulk materials with coatings applied to surfaces where improved properties 

required. 

Approach 

• Examine a couple of approaches on surface modification that would be of interest to materials for heavy-duty 

vehicles. 

• Examine the application and formation of adherent, wear/corrosion-resistant coatings. 

• Base the initial tests on hard metal compositions applied onto iron-based parts that are currently used in diesel 

engines. 

Accomplishments 

• Determined that plain hard metal coatings without binder additions have a very narrow window of HDI process 

conditions and are not practical on a large scale. 

• Correlated coating process parameters (compositions, HDI parameters) with coating microstructural 

development. 

Future Direction 

• Investigate the coating of complex shaped parts by HDI technology. 

• Determine potential for producing wear-resistance diffusion coatings for heavy-duty vehicle applications. 

Introduction The HDI processing facility at Oak Ridge National 

Laboratory utilizes a unique technology to produce High-density infrared (HDI) technology is relaextremely high-power densities of up to 3.5 kW/cm2 

tively new to the materials processing area and is 

gradually being exploited in materials processing. with a single lamp. Instead of using an electrically 

151 
FY 2005 Progress Report 

heated resistive element to produce radiant energy, a 

controlled and contained plasma is utilized. 

Because the technology is relatively new, its 

utility to the surface treatment of materials for applications in heavy-duty vehicles is being explored. In 

most cases, the need for wear resistance, or corrosion resistance, or high strength is only necessary in 

selected areas of the part that is exposed to the 

working environment or under high stress. Therefore, it would be desirable to use materials that are 

lighter or less expensive for the bulk of the part and 

only have the appropriate surface properties where 

required. In addition, the HDI approach would be 

more cost-effective than other competitive processes 

such as physical vapor deposition. 

Non-line-of-sight coatings are of interest for 

application on heavy-duty vehicle parts, such as 

gears. Pack cementation is an industrial process 

currently used to apply diffusion coatings onto metal 

parts. Other research has shown that microwave 

heating can be used to produce superior diffusion 

coatings with alloying additives such as aluminum, 

silicon, and boron. The project will examine the 

potential to produce wear-resistance diffusion coatings by both HDI technology and microwave pack 

cementation. 

Results 

Earlier work had shown that adherent coatings 

of hard metal compositions could be applied to a 

variety of metal substrates (D2 tool steel, 4140 alloy 

steel, and cast iron). However, a problem with the 

earlier coatings was that a preheat step was required 

to minimize sintering of the slurry coating prior to 

bonding with the substrate. When the slurry coating 

was heated too fast, it tended to peel off from the 

substrate and crack. The preheating step allowed the 

substrate to get hot enough so bonding occurred. 

Samples were fabricated that used either WC or 

TiC particles only for the coating. The sintering 

temperature for these plain carbides is significantly 

higher than the melting point of the substrates. Consequently, because no binder additions (i.e., Ni or 

Ni3Al) will be present, sintering and cracking will be 

inhibited. This will allow the coating to heat up to a 

sufficiently high temperature so that the underlying 

substrate will melt and be wicked into the coating by 

capillary action. Previous work has shown that 

Heavy Vehicle Propulsion Materials 

WC-cast iron composites have excellent wear 

behavior, so these will be the initial compositions. 

Samples were fabricated with varying coating 

thickness (50 and 100 μm) and particle type (WC, 

TiC, W, or WC/W). The samples were infrared (IR) 

treated at numerous exposure conditions. Cracking 

of the coatings was observed in a similar behavior to 

previous tests where the binder was included. A preheat step helped to minimize the cracking. However, 

getting the coatings hot enough to melt the underlying substrate was extremely sensitive. The conditions to go from little bonding to where extensive 

melting of the entire sample occurred were very 

narrow. Control of the coating process without the 

incorporation of binder particles in the coating is 

very difficult. 

Additional examinations of coatings with binders was done to correlate coating process parameters 

with coating microstructural development. Some of 

the parameters included substrate composition, 

binder chemistry, ceramic particle type, and HDI 

heating cycles. Previous results had shown excellent 

wear resistance for coatings having Ni3Al as the 

binder in the hard metal coating. Consequently, most 

of the correlations were done with these coating 

types. 

Examples of WC-Ni3Al coatings on substrates 

of D2 tool steel, 4140 alloy steel, and cast iron are 

shown in Figures 1–3, respectively. The coatings are 

much thicker than the intended ~100 μm, indicating 

significant mixing with the underlying substrate. A 

penetration zone of Ni3Al into the D2 tool steel was 

significant (≤100 μm), whereas only a minor reaction layer (~10 μm) was observed at the interface on 

the 4140 alloy steel. No reaction layer was evident 

with the cast iron substrate; however, there appeared 

to be a diffuse mixing zone within the coating. In 

addition, the cast iron had graphite precipitates 

throughout the coating which were more pronounced 

adjacent to the interface. Some grain growth of the 

WC was observed at the interface between the coating and the D2 tool steel. Conversely, more uniform 

grain sizes were observed across the coatings on the 

4140 alloy steel and the cast iron. Note that, in general, the grain sizes were on the order of the initial 

starting particle sizes, indicating little to no grain 

growth during the HDI thermal treatments. 

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Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 1. Cross-section of WC-Ni3Al coating on D2 tool 

steel. The coating was initially intended to be 

100 μm thick. Coating bonded by HDI thermal 

treatment using a 4-cm defocus setting and 

lamp parameters of 0 to 300 amps in 5 min 

(preheat), followed by 900 amps for 3.5 s for 

bonding. 

Figure 2. Cross-section of WC-Ni3Al coating on 4140 

alloy steel. The coating was initially intended 

to be 100 μm thick. Coating bonded by HDI 

thermal treatment using a 4-cm defocus setting 

and lamp parameters of 0 to 300 amps in 

5 min (preheat), followed by 900 amps for 

3.5 s for bonding. 

Figure 3. Cross-section of WC-Ni3Al coating on cast 

iron. The coating was initially intended to be 

100 μm thick. Coating bonded by HDI thermal 

treatment using a 4-cm defocus setting and 

lamp parameters of 0 to 300 amps in 5 min 

(preheat), followed by 900 amps for 3.5 s for 

bonding. 

A cross-section of a WC-Co coating on D2 tool 

steel is shown in Figure 4. Less penetration of the 

Co into the D2 tool steel (~25 μm) compared to the 

Ni3Al was evident. Less mixing between the coating 

and the substrate resulted in a uniform grain size 

across the coating. 

When TiC (with Ni3Al) was used for the coating 

there appeared to be less penetration of the Ni3Al 

into the D2 tool steel substrate (Figure 5). This 

needs to be confirmed with additional characterization. Normally, TiC grains are quite faceted; however, the TiC grains in the coating were rounded, 

indicating alloying of the Ni3Al binder with iron 

from the substrate. In addition, no core-rim structures, which are typically observed with TiC grains 

in Ni3Al, were evident. This indicates only minimal 

dissolution and reprecipitation of the TiC during 

melting of the coating. Like the WC-based materials, the final grain size was similar to the starting 

TiC particle size. 

The effect of the use of multiple HDI treatments 

on the microstructure of a WC-Ni3Al coating on D2 

tool steel is shown in Figure 6. The major effect of 

the double treatment was to minimize bubbles in the 

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FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 4. Cross-section of WC-Co coating on D2 tool 

steel. The coating was initially intended to be 

100 μm thick. Coating bonded by HDI thermal 

treatment using a 4-cm defocus setting and 

lamp parameters of 0 to 300 amps in 5 min 

(preheat), followed by 1000 amps for 3.5 s for 

bonding. 

Figure 5. Cross-section of TiC-Ni3Al coating on D2 tool 

steel. The coating was initially intended to be 

100 μm thick. Coating bonded by HDI thermal 

treatment using a 4-cm defocus setting and 

lamp parameters of 0 to 400 amps in 5 min 

(preheat), followed by 900 amps for 2 s and 

900 amps for 1 s for bonding. There was 

approximately 15 s between high-energy 

exposures. 

Figure 6. Cross-section of WC-Ni3Al coating on D2 tool 

steel. The coating was initially intended to be 

100 μm thick. Coating bonded by HDI thermal 

treatment using a 4-cm defocus setting and 

lamp parameters of 0 to 400 amps in 5 min 

(preheat), followed by 900 amps for 2 s and 

900 amps for 1 s for bonding. There was 

approximately 15 s between high-energy 

exposures. 

coating. It is hypothesized that during the second 

HDI exposure, the coating is remelted and any 

trapped bubbles can be released. However, the 

double treatment increases penetration of the Ni3Al 

into the D2 tool steel from ~100 μm for a single 

exposure to ~300 μm for the multiple heating. The 

WC grain size was also increased by the second HDI 

heat treatment. 

Conclusions 

It was determined that plain hard metal coatings 

without binder additions have a very narrow window 

of HDI process conditions and are not practical on a 

large scale. Significant interactions and mixing 

occur between the coatings and the underlying substrates. In most cases, little to no grain growth by 

WC or TiC was observed during HDI heat 

treatments. 

Presentations and Publications 

T. N. Tiegs, “High Density Infrared Surface 

Treatment,” at the Heavy Vehicles Materials Peer 

Review, Oak Ridge, Tennessee, September 13–15, 

2005. 

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Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

N. High-Temperature Aluminum Alloys 

Yong-Ching Chen 

Materials Engineering 

Cummins, Inc. 

1900 McKinley Ave. 

Columbus, IN 47201-3005 

(812) 377-8349; fax: (812) 377-7226; e-mail: yong-ching.c.chen@cummins.com 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: Cummins, Inc., Columbus, Indiana 

Objectives 

• Develop high-temperature aluminum alloys with adequate properties and shape capability for turbocharger 

compressor wheel and housing application. 

Approach 

• Develop ternary-phase compositions by means of modeling of the equilibrium-phase diagram, castings, and 

evaluating the properties of alloys identified. 

• Perform physical and mechanical property measurements on specimens provided by the National Aeronautics 

and Space Administration (NASA) and by Eck Industries to characterize the material and determine the optimal 

properties possible by this method. 

• Evaluate a nanophase particulate-reinforced aluminum alloy patented by Chesapeake Composites, Inc. 

Accomplishments 

• Predicted the binary phase diagrams of Al-Y and Al-Yb accurately and verified them using ThermoCalc. 

• Found that the improvement in high-temperature strength of the NASA 388 alloy was not significant enough 

for further pursuing. 

• Observed that the dispersion-strengthened composite (DSC) material exhibited significantly lower thermal 

conductivity and lower coefficient of thermal expansion than those of current production aluminum alloys. 

• Found that the tensile strength of the DSC sample produced in 2005 was much lower than that of the billets 

tested in 2002. The inferior strength of the current DSC billets was most likely due to the presence of aluminum 

veins and defects inside the aluminum matrix. Improvement in quality is needed. 

• Completed machinability studies of DSC material. Optimal operation parameters for cutting, turning, and 

milling were determined. The tool wear was minimal, similar to the machining of 6061-T6 aluminum. 

Future Direction 

• Use ThermoCalc to generate the Al-Y-Yb ternary system and identify optimal compositions for casting trial and 

material property characterization. 

155 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

• Improve billet quality, and complete notch sensitivity and creep tests of DSC material. 

• Use DSC material for rig testing machine turbocharger compressor wheel. 

Introduction 

The new emission requirement in internal 

combustion of diesel engines mandates that the 

turbocharger compressor be operated at significantly 

higher temperature and pressure. This condition 

makes standard castable aluminum alloys unusable 

as new-generation turbocharger material, because 

their strength deteriorates at elevated temperatures. 

Therefore, there is a need to develop high-strength, 

high-temperature aluminum alloys, which will 

replace these standard alloys. Work at Cummins has 

identified three potential paths to improve the hightemperature strength and fatigue resistance of aluminum casting alloys. These paths have been partially 

investigated under cooperative agreement DE-FC05­

97OR22582. Each path showed some promise, and 

further work is needed to determine the useful 

applications. 

Approach 

The scope of work in the program covers 

investigation of three different paths of providing 

strength retention in aluminum alloys at high 

temperatures. Path 1 uses a ternary-phase aluminum 

alloy utilizing rare earth metals to provide precipitate size control and stability. Limited information 

on these alloys indicates high-temperature strength 

and stability; however, the predicted cost of the 

alloy is high. Additional work is performed to determine if additional rare earth (or other metal) elements produce beneficial properties at a reasonable 

cost. The main goal of this path is to generate the 

phase diagram of Al-Y-Yb, using ThermoCalc software. Because there are few published data and 

assessments on these rare-earth elements, a prudent 

approach demands that a systematic investigation be 

carried out. Path 2 uses a conventional aluminum 

alloy, which has been chemically modified by a process developed at NASA-Huntsville. The elevated 

temperature properties reported by NASA are attractive, but it is preferred for Cummins to use alloy 

with lower silicon content. Eck Industries purchased 

the license for the NASA-developed technology on 

high-silicon casting alloy and has expanded the 

range to include conventional low-silicon casting 

alloys. Limited testing of these modified conventional low-silicon alloys at Cummins has not shown 

the property improvement anticipated. Currently 

NASA 388-T5 alloy is being evaluated to characterize the material and to determine possible process 

improvement. Path 3 uses a particulate-loaded 

aluminum alloy patented by Chesapeake Composites 

Inc. The nanophase particulates at 50 vol % provide 

adequate high-temperature strength in early experiments; however, the particulate-loaded alloy could 

only be forged or squeeze cast so complex shape 

capability is limited. The mechanical behavior of 

this DSC material is being fully characterized. 

Casting modifications will be investigated to 

determine shape capability for the alloy. 

Results 

Path 1: Ternary-Phase Aluminum Alloy 

Development 

In a previous report, the ternary-phase diagram 

of the Al-Fe-Mn was developed by using 

ThermoCalc software. The versatility of the 

ThermoCalc had been demonstrated. A binary Al-Y 

phase diagram was then generated by using the 

ThermoCalc, and it was validated against literature 

data. The work on thermodynamic assessment of 

Al-Yb binary by using ThermoCalc was also carried 

out. In general, it is in good agreement with the 

standard phase diagram published in the literature. 

However, it is important to note that unlike most 

elements, the rare-earth elements yttrium and 

ytterbium are not widely used. For the ternary 

Al-Y-Yb system to be predicted, the minimum 

information needed is the assessment of Y-Yb 

system. Without any experimental data, it is not 

possible to know if such ternary predictions are 

accurate or not. So the minimum step we need is to 

get Y-Yb system; otherwise, one axis of our ternary 

diagram may not be correct. We need data in terms 

of the actual thermodynamics of the system, for 

example, heat of mixing and activity data. We have 

to estimate the actual Gibbs energy for the 

compounds and then make the fit to the polynomial 

terms. Much effort was focused on experimentally 

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Heavy Vehicle Propulsion Materials 

generating the Y-Yb binary phase diagram. This will 

allow us to predict the Al-Y-Yb ternary system. 

Path 2: NASA 388-T5 Aluminum Alloy 

Extensive mechanical tests were conducted on 

the NASA 388-T5 and T6 alloys. Tests included elevated-temperature tensile tests and rotating-beam 

fatigue tests of samples with or without long-term 

thermal soaking. Test results were compared to 

those of the current production materials. Slight 

improvement in the tensile strength was only realized at test temperatures above 300oC for the NASA 

388-T5 alloy as compared to the current materials. 

The degree of improvement was deemed to be 

insignificant as compared to the further development 

needed for the new casting process and prototype 

trials. The decision was made to discontinue the 

evaluation work for the NASA 388-T5 alloy for 

intended application. 

The NASA 388-T5 alloy was originally developed for automotive engine components such as pistons to utilize less material, which can lead to reducing part weights, material cost, as well as improving 

gas mileage and engine durability. The potential 

application of the NASA 388-T5 alloy in diesel 

engine pistons was communicated to the piston 

suppliers. One of the piston suppliers conducted 

fatigue tests of the NASA 388-T5 alloy and their 

production aluminum alloy at 350oC. The NASA 

388-T5 alloy exhibited slightly better fatigue 

strength at 350oC than the production aluminum 

alloy. However, the NASA 388-T5 alloy was found 

to be difficult to cast in diesel piston size without 

defects. The piston supplier also decided not to 

pursue the development of NASA 388-T5 alloy in 

the piston application. 

Path 3: Particulate-Reinforced Alloy 

This task involves a nanophase particulateloaded aluminum alloy patented by Chesapeake 

Composites Inc. A patented, low-cost, liquid-metal 

infiltration process is utilized to produce a billet 

form ready for secondary operations. This composite 

material combines the enhanced elevatedtemperature strength, toughness, and ductility of 

dispersion-strengthened alloys with the stiffness and 

low coefficient of thermal expansion (CTE) of metal 

matrix composites (MMCs). It was claimed that this 

composite can be readily turned using tungsten 

carbide (WC) tooling and drilled and tapped using 

FY 2005 Progress Report 

high-speed steel tools. Potential applications include 

pistons, compressor wheel, and engine components. 

Initial tensile and fatigue tests of 1090 Al and 

2024 Al alloy reinforced with nano-scale Al2O3 

particles at elevated temperatures showed impressive results. Almost 100% improvement in tensile 

strength and yield strength was obtained on the DSC 

material as compared to those of the current 

production materials. No degradation in the tensile 

strength and fatigue strength was observed in the 

DSC material after 500 h of thermal soaking at 

400oC. 

1. Microstructural Characterization 

Extensive microstructural characterization of 

2024Al-DSC material has been carried out. Particular attention was focused on developing an understanding of microstructural development in a 

2024Al-Al2O3 composite during heat treatment. 

Figure 1 shows the variation of microhardness as a 

function of aging time for 2024Al-Al2O3 composite 

aged at 130o

C and 190oC. Generally speaking, the 

aging time at peak hardness for the unreinforced all 

2024 alloys is about 10 h (Ref. 1). However, the 

composite material reaches peak hardness after only 

1 h of aging, which confirms the accelerated aging 

phenomenon we have discussed previously. Overaging and drop in hardness were observable at 

190°C. Solution treatment and water quench greatly 

improve the microhardness of the composite. This is 

particularly because the dislocation generation due 

to thermal expansion mismatch between aluminum 

matrix and Al2O3 reinforcement. 

An accelerated aging phenomenon has been 

reported, which is attributed to the increase in 

dislocation density due to Al2O3 reinforcement. 

Transmission electron microscope (TEM) analysis 

of the aged samples revealed the presence of θ’ and 

S’ precipitates, which were generally quite stable at 

130°C, but underwent some coarsening with a concomitant drop in hardness following aging at 190°C. 

The dislocation density in room-temperature tensiletested samples was quite low, which was consistent 

with the low ductility at this temperature. On the 

other hand, the dislocation density in samples thermally soaked and tested at higher temperatures (500 

and 700°F) was high. This is thought to be associated with the combined effect of ease of plastic 

deformation at high temperatures, coupled with 

differential thermal contraction of Al and Al2O3 on 

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FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 1. Variation of microhardness as a function of aging time for the 2024AlAl2O3 composite. 

cooling from the elevated temperatures of testing. 

Figure 2 shows TEM micrographs of the dislocation 

states of the 2024Al-Al2O3 as solutionized and 

water-quenched material. 

2. Physical Properties 

To further verify that the DSC material exhibits 

adequate mechanical strength to meet the operation 

requirements, mechanical and physical properties of 

DSC materials will be measured. Chesapeake 

Composites Corporations and PCC-AFT manufactured billets of DSC material with 1.7-in. diameter 

and 7.0 in. long. The DSC billets contain 40 vol % 

of Al2O3 nona-particles and were infiltrated with 

1090 aluminum alloy. 

Samples of a DSC material were machined 

for thermophysical property testing from room 

temperature to 400oC. The density of the DSC 

material was measured to be 2.991 g/cm3

, which is 

about 12% heavier than that of the A354-T6 alloy 

(2.677 g/cm3

). The higher density needs to be taken 

into consideration during design because the 

Figure 2. TEM micrographs showing the dislocation states of the 2024Al-Al2O3 as solutionized + water quenched 

material: (a) TEM DF (dark field) image; (b) TEM BF (bright field) image (precipitates). 

158
Heavy Vehicle Propulsion Materials 

compressor wheels made of DSC material will have 

higher reciprocal weight compared to those made of 

traditional aluminum alloy such as A354 alloy. 

Figure 3 displays the thermal conductivity of the 

DSC material as a function of temperature. Note that 

the thermal conductivity of the DSC material 

decreases as the temperature increases. This is most 

likely due to the presence of a high percentage of 

Al2O3 particles inside the aluminum matrix. The 

thermal conductivity of the DSC material was found 

to be much lower than that of a traditional aluminum 

alloy. The lower thermal conductivity in the DSC 

material may result in higher operating temperature. 

The mean CTE of the DSC material increases as the 

temperature increases, Figure 4. Note that the mean 

CTE of DSC material was significantly lower than 

that of a traditional aluminum alloy. The lower mean 

CTE was also most likely due to the presence of a 

high percentage of Al2O3 particles inside the 

aluminum matrix of the DSC material. The lower 

CTE in the DSC material may indicate smaller 

dimensional changes during operation. 

3. Tensile Testing 

Tensile samples with dimensions of 4.2-mm 

diameter by 18-mm gage length were prepared from 

the DSC billet for tensile testing at room temperature, 200oC, 300oC, and 400oC without supplemental 

thermal soaking. Figure 5 shows the tensile strength 

of the DSC material as a function of temperature. 

FY 2005 Progress Report 

The tensile strength of the DSC material tested in 

2002 was also plotted for comparison. Note that the 

tensile strength of the DSC sample produced lately 

was much lower than that tested in 2002. The percent elongation and reduction of area at room 

temperature were almost zero. Some samples broke 

quickly before yielding such that no yield strength 

was available. The much lower tensile strength of 

the DSC samples manufactured in 2005 was 

unexpected. 

4. Microstructural Characterization 

To compare the microstructure of the current 

billets and that of the billets which exhibited much 

higher tensile strength, five tested tensile bars prepared from the DSC billets manufactured by 

Chesapeake Composites in 2002 were submitted for 

microstructure analysis to determine if defects were 

present in the material. The tensile bar samples were 

tested previously at varied temperatures and after a 

500-h soak. Figure 6 shows the comparison between 

the tensile sample tested in 2002 and the billet 

samples manufactured in 2005. Note that the tensile 

sample tested in 2002 exhibited less cellular 

structure and much finer grain size as compared to 

the billet sample manufactured in 2005. Very few or 

no aluminum veins were observed in the tensile 

sample tested in 2002. The difference in the microstructure had resulted in significant difference in the 

tensile strength of the DSC billets. 

Figure 3. Thermal conductivity of the DSC material, 354-T6, and 355-T6 aluminum 

alloys. 

159 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 4. Mean CTE of the DSC material, 354-T6, and 355-T6. 

Figure 5. Elevated-temperature tensile strength of the DSC material manufactured in 

2005 and DSC material manufactured in 2002. 

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Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 6. Comparison of the microstructure of tensile sample (left) to billet sample of different batch (right) at 100×. 

The microstructure of billet contains a larger number of large veins and defects. 

According to the supplier, a different 

manufacturing process was used in making fine 

aluminum powders needed for making preforms. 

The difference in the quality of the powders had 

resulted in the dramatic difference in the final tensile 

strength. The quality of the current DSC billets was 

considered unsatisfactory, and further process 

development is needed by the supplier to get material close to past strength and quality. 

5. Machinability Study 

The machinability study of DSC aluminum was 

performed in the areas of drilling, sawing, singlepoint turning, and five-axis milling. The milling and 

drilling tools tested were selected by the Advanced 

Integrated Manufacturing Center, which consisted of 

carbide end mills from OSG Tooling and high-speed 

steel drills and taps. 

A Daewoo Puma 240MS turret lathe, 

Figure 7(a), was used for the turning test. All turning 

parts were processed to machine the face then rough 

turn the outside to form two different diameters. All 

parts produced broken chips. There was no formation of long “springy” chips. Figure 7(b) shows 

excellent surface finish on both diameters. A Deckel 

Maho DMU50, Figure 8(a), was the five-axis 

machining center used for the test. The nonorthogonal fifth axis provides excellent machine stability. 

The part was 2-D roughed in 0.050-in. increments 

with the 0.375-in. two-flute end mill. The contour 

was finished utilizing all five axes of the machining 

center with the 0.250-in. ball end mill. The blade 

areas were finished, and the corners were removed 

Figure 7. (a) A Daewoo Puma 240MS turret lathe was used for the turning test. (b) Excellent finish on both 

diameters was obtained with optimal cutting parameters. 

161 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 8. (a) A Deckel Maho DMU50 was the five-axis machining center used for the milling test. (b) One 

demonstration part was completed by the five-axis machining center. 

with the 0.125-6in. ball end mill. Figure 8(b) shows 

one demonstration part completed by the five-axis 

machining center. A microscope was used to measure the tool wear. The tool wear for the study was 

found to be minimal, similar to the machining of 

6061-T6 aluminum. No significant difference in 

machining the DSC was noted when compared to 

the machining of traditional wrought aluminum 

alloys. 

Conclusions 

Three different paths are adopted to develop and 

evaluate high-temperature aluminum alloys for 

turbocharger compressor wheel and housing applications. The versatility of the ThermoCalc software 

was demonstrated in the plotting of the phase diagrams. The Al-Y and Al-Yb binary-phase diagrams 

were generated and validated. The Al-Y-Yb ternary 

system will be modeled. Slight improvement in the 

tensile strength was only realized at test temperatures above 300oC for the NASA 388-T5 alloy as 

compared to the current materials. The degree of 

improvement was deemed to be insignificant as 

compared to the further development needed for the 

new casting process and prototype trials. The decision was made to discontinue the evaluation work 

for the NASA 388-T5 alloy for turbocharger 

application. The DSC material exhibited very 

impressive high-temperature tensile and fatigue 

strength. Excellent thermal stability was confirmed 

with SEM and transmission electron microscope 

(TEM) studies. Physical property measurements of 

the DSC material were completed. The thermal conductivity of the DSC material was found to be much 

lower than that of a traditional aluminum alloy. The 

lower thermal conductivity in the DSC material may 

result in higher operating temperature. The mean 

CTE of DSC material was significantly lower than a 

traditional aluminum alloy. The lower CTE of the 

DSC material may indicate smaller dimensional 

changes during operation. The tensile strength of the 

DSC sample produced in 2005 was much lower than 

that of the billets tested in 2002. The inferior 

strength of the current DSC billets was most likely 

due to the presence of aluminum veins and defects 

inside the aluminum matrix. The quality of the 

current DSC billets is not satisfactory. Further 

process improvement is needed by the supplier. 

Machinability study of the DSC material was completed. No significant difference in machining the 

DSC was noted when compared to the machining of 

traditional wrought aluminum alloys. 

References 

1. R. J. Arsenault and R. M. Fisher, Scripta 

Metall., 17, 67 (1983). 

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Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

O. Application of Titanium Alloys for Heavy-Duty Vehicles 

P. C. Becker 

Oak Ridge National Laboratory 

P.O. Box 2008, MS 6066 

Oak Ridge, Tennessee 37831-6068 

(865) 576-4844; fax (865) 574-6098; e-mail: beckerpc@ornl.gov 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

 (865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Prime Contract No.: DE-AC05-00OR22725 

Objectives 

• Complete the final details of the finite element analysis (FEA) modeling of a diesel engine head and block 

where an aluminum (Al) casting alloy is used to replace gray cast iron. 

• Complete the final details of the FEA modeling of a diesel engine head and block where a magnesium (Mg) 

casting alloy is used to replace gray cast iron. 

• Try to seek funding from either a major diesel engine manufacturer or a major customer to construct and test 

engines with heads and blocks made from titanium (Ti), Al, or Mg alloys. 

• Generate fatigue data that could indicate whether the sub-liquidus casting (SLC) process is attractive for 

making Al turbocharger compressor wheels with improved fatigue life. 

Approach 

• Use FEA modeling of a heavy-duty diesel engine head and block to determine whether it is feasible to 

substitute parts from Ti, Al, and Mg instead of gray cast iron. 

Accomplishments 

• Confirmed through FEA modeling that it is feasible to substitute Al and Mg for cast iron in the head and block 

of a heavy-duty diesel engine. 

• Conducted a preliminary evaluation of the SLC process. Mechanical property data from A354-T6 samples were 

generated and compared with data of the same alloy that was permanent mold cast. The data showed that ultimate tensile strength, yield strength, and extrapolated fatigue life for the SLC cast bars was significantly higher 

than for the permanent mold cast bars. 

Future Direction 

• Make a more definitive evaluation of the SLC sample after THT Presses delivers higher-quality cast bars. 

163 
FY 2005 Progress Report 

Project Description and Activities 

Modeling Lightweight Alloys 

In FY 2004, modeling provided by Ricardo, 

Inc., using FEA of temperatures and stresses showed 

that it was feasible to make a heavy-duty diesel 

engine head and block of cast Ti-6Al 4V instead of 

gray cast iron. The engine that was modeled was a 

Cummins 5.9-L B-Series engine (Figure 1). The 

analysis showed that the only problem was unacceptable temperatures generated in the head valve 

bridges; therefore, a novel means of mitigating this 

Figure 1. FEA model of the Cummins engine. 

Heavy Vehicle Propulsion Materials 

effect was proposed and modeled. The novel 

approach was to substitute for the Ti in the valve 

bridges a high-thermal-conductivity alloy insert of 

Be-Cu. The engine was again modeled, but the 

normal 305-hp rating used for modeling was 

increased by 50% to 450 hp. The model showed that 

this approach was viable and resulted in a 10% 

weight savings as well as a higher horse power. It 

was felt that these advantages could be attractive to 

the military, specifically TACOM. In FY 2005, 

numerous presentations were made to the diesel 

engine companies, to Oshkosh Truck, and to the 

Defense Advanced Research Projects Agency 

(DARPA) and TACOM to determine their level of 

interest. 

After the engine was modeled in Ti, two other 

lightweight materials, high-strength Al and highstrength Mg casting alloys, were modeled as 

replacements for cast iron in a diesel engine head 

and block at a normal maximum rating (305 hp). 

The alloys were A354-T6 AL and We-43-T6 MG. 

The stresses and fatigue safety factors in the block 

were found to be acceptable. The same was true for 

the head. However, the temperatures in the valve 

bridges were of concern based on the results of the 

Ti modeling. The results of the thermal FEA of the 

engine head are shown in Figure 2. 

Figure 2. Thermal results comparing the valve bridge region in the head for 

production cast iron, Al, and Mg. 

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Heavy Vehicle Propulsion Materials 

The FEA results show that the temperatures are 

low enough to be considered acceptable, keeping in 

mind that the analysis assumed that the engine was 

running at maximum power throughout its life, 

which does not mirror reality. The temperature of 

310°C in the Mg head was somewhat high, but for 

this study it was not high enough to make the head 

unfeasible in this material. 

The potential overall engine weight saving was 

20% for Al and 33% for Mg. Most of the FEA work 

was done in FY 2004, but some details were completed in FY 2005. These involved calculating 

weight benefits and presenting the work to potential 

customers and manufacturers. 

Since the commercial Cummins engine weighs 

about 1100 lb, 220 lb could saved if the head and 

block were made of Al and 363 lb could be saved if 

made from Mg. All of these data and analyses were 

presented to Cummins and Caterpillar as well as to 

TACOM in FY 2005. 

SLC of Al for Turbocharger Compressor Wheels 

SLC is a lower-cost approach to semi-solid 

forming (SSF) of Al castings. The advantage of the 

SSF process is that a nearly porosity-free casting can 

be made that results in high fatigue strengths. The 

FY 2005 Progress Report 

process is attractive for making high-performance 

diesel engine turbocharger compressor wheels. 

Using the SLC approach solves two major problems 

inherent in traditional SSF processing. First, since 

the semi-solid charge is produced in the heating 

chamber of the die casting tooling, the cost to produce a part is nearly the same as for a traditional die 

casting. This contrasts with a 25% cost premium for 

parts made by traditional SSF. Second, SLC requires 

only about 5000 psi to force the slurry into the mold; 

traditional SSF requires 20–25,000 psi. Therefore, a 

plaster mold or some other destructible material 

could possibly be used to make the compressor 

wheels, which is a requirement. The severe curvature of the compressor wheel’s vanes makes it 

impossible to extract a casting from hard tooling. 

Figure 3 is a schematic drawing of the SLC concept. 

SLC cast test bars of A354 were made at THT 

Presses and then heat-treated to T-6. These bars 

were sent to Metcut Research for machining and 

testing. The data were then to be compared with 

semi-permanent mold castings made from the same 

alloy. Tensile test data are shown in Table 1. The 

table compares data from test bars cast by the SLC 

process to data from permanent mold test samples 

cast with the same alloy and heat treatment process. 

SLC® CONCEPT

Figure 3. Schematic drawing of the SLC casting chamber. 

165 
FY 2005 Progress Report 

Table 1. Comparison of permanent mold cast to 

SLC A354-T6 

Test temp. UTS 0.2% YS Elongation 

(°C) (Ksi) Ksi (%) 

Permanent mold (no chill) 

Room temp 38.5 35 0.57 

150 31 30 0.37 

200 29 29 0.33 

260 21 20 1.1 

Sub-liquidus 

Room temp. 46 40 1.8 

150 42 35 3.9 

200 36 34 1.7 

250 26 24 3.8 

UTS = ultimate tensile strength 

YS = yield strength 

Heavy Vehicle Propulsion Materials 

The fatigue properties at room temperature of 

samples made by the SLC process were measured at 

Metcut Research using the rotating beam technique, 

and the data were compared with permanent mold 

data produced for Cummins in an earlier DOEfunded program. Cummins did the data analyses 

using its proprietary fatigue analysis software 

FATIGUE. The comparison is shown in Figure 4. 

As can be seen from these two S/N curves, the 

SLC process provides a slightly higher fatigue life at 

cycles greater than 107

, which is more pronounced 

at 5 × 107 cycles. Unfortunately, the SLC samples 

had higher than expected porosity, which would 

explain the high scatter in the data. Optical microscopy of the SLC cast bars conducted at ORNL 

(Figure 5) shows the porosity. 

Figure 4. Comparison of rotating beam curves of permanent mold and SLC 

castings of A354-T6. 

Figure 5. Micrographs of SLC samples showing the porosity from the casting process. The micrographs 

are from the infrared heat-treated sample. 

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Heavy Vehicle Propulsion Materials 

Figure 6 shows the microstructure of the SLC 

cast bars. The micrograph on the left shows the conventional T-6 heat-treated microstructure, and the 

one on the right shows the infra-red heat-treated 

microstructure. In both cases, the microstructure is 

finer than for the permanent mold castings of the 

same alloy, but using infra-red provides an even 

finer structure and hence even higher mechanical 

properties. The fine microstructure explains why the 

cycle fatigue life of the SLC samples is higher than 

that of the permanent mold samples even though the 

SLC castings do not have optimum structure 

because of the high porosity. 

THT Presses has promised to replace these samples with higher-quality cast bars so that a more 

definitive evaluation can be made. 

Conclusions 

• FEA modeling has confirmed that it is feasible 

to substitute Al and Mg for cast iron in the head 

and block of a heavy-duty diesel engine. This 

substitution can result in an overall engine 

FY 2005 Progress Report 

weight saving of 20% for Al and 33% for Mg. 

Although this work and the previous work on Ti 

were presented numerous times to TACOM, 

DARPA, Cummins, and Caterpillar, none of 

these organizations has yet committed funding 

to continue the program. 

• The SLC process was given a preliminary 

evaluation. Mechanical property data from 

A354-T6 samples were generated and compared 

with data of the same alloy that was permanent 

mold cast. The ultimate tensile strength at room 

temperature for the SLC cast bars was about 

19% higher than for the permanent mold cast 

bars, and the yield strength was about 14% 

higher. The extrapolated fatigue life of the SLC 

process produced a 17% improvement over that 

of the permanent mold cast samples at 5 × 107 

cycles. 

If the SLC samples contained the very low porosity 

that the process is capable of, even higher fatigue 

properties should be achievable. 

Figure 6. Microstructures of SLC samples that have undergone either conventional T-6 

heat treatment or infra-red heat treatment. 

167 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Presentations 

Puja B. Kadolkar, Craig A. Blue, Qingyou Han, 

and Paul C. Becker, “Rapid Infrared Heat Treatment 

of Al Cast Structures,” presented to the Materials 

Science and Technology 2005 Conference and 

Exhibition, September 25–28, 2005. 

P. C. Becker, presentation on Ti, Al, and Mg in 

diesel engines to Caterpillar, April 27, 2005. 

P. C. Becker, presentations on Ti, Al, and Mg in 

diesel engines to Cummins, April 25, 2005. 

P. C. Becker, presentation on Ti, Al, and Mg in 

diesel engines to TACOM, April 15, 2005. 

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Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

P. Mechanical Behavior of Ceramic Materials for Heavy-Duty Diesel Engines 

Andrew A. Wereszczak and Hong Wang 

Ceramic Science and Technology Group (CerSaT) 

Oak Ridge National Laboratory 

P.O. Box 2008, MS 6068 

Oak Ridge, TN 37831-6068 

(865) 576-1169; fax: (865) 574-6098; e-mail: wereszczakaa@ornl.gov 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Objectives 

• Enable the development of more wear-resistant and mechanically robust ceramic components through the 

systematic evaluation of contact damage in ceramics. 

• Characterize rate effects on deformation processes in Si3N4. 

• Develop supplemental models to assist in the interpretation of deformation. 

Approach 

• Characterize damage processes in Si3N4 ceramics through instrumented static and dynamic indentation and 

instrumented scratch testing. 

• Develop criterion that can be used to rank resistance to damage. 

• Evaluate relevant Si3N4 grades (e.g., existing commercial roller element and ball bearing compositions, 

standard reference materials). 

• Evaluate grain size effects on damage processes through static and dynamic indentation (model Al2O3 ceramics 

used). 

Accomplishments 

• Established dynamic indentation and instrumented scratch testing facilities. 

• Developed a new brittleness parameter that represents static and dynamic indentation responses. 

• Developed a generic finite element method (FEM) that predicts the onsets of quasi-plasticity and 

fracture/cracking in any target material as a function of indentation load, indenter material, friction, and 

confinement. 

• Co-developed and copyrighted μ-FEA software. 

Future Direction 

• Conduct a complete study of grain size effects on rate and deformation processes. 

• Begin a new project involving reliability analysis and design optimization of piezostack fuel injectors for diesel 

engines. 

169 
FY 2005 Progress Report 

Introduction 

The application of more ceramic components in 

engines and transportation systems would be 

enabled if they could be confidently used in more 

wear-related applications and if they could be 

manufactured and machined faster (i.e., more costeffectively). This would occur if contact-induced 

damage mechanisms (that limit their mechanical 

performance or dictate their machinability) were 

understood, predictable, and controlled. 

This project quantifies the loads and stresses 

associated with the initiation of “quasi-plasticity” 

and fracture processes in ceramics using 

instrumented indentation and scratch testings. 

Understanding their competition and interaction is a 

critical part of this project as well. Traditional 

structural ceramics (e.g., Si3N4) under consideration 

for diesel engine components are of primary focus. 

Through the quantification of the loads or 

stresses that initiate quasi-plasticity and fracture, 

contact-induced damage mechanisms in ceramics are 

characterized and modeled and linked to microstructure, and ultimately to wear performance, the 

optimization of machining, and rolling contact 

fatigue. Additionally, the influence that independent 

parameters such as grain size, loading rate, confinement or residual stresses, temperature, field effects 

have on the evolution of contact-induced damage are 

also examined. The thorough understanding of the 

competition of fracture and quasi-plasticity in 

ceramics will enable improved and faster means of 

ceramic component manufacturing and surface 

engineering (e.g., ductile regime machining) and 

maximize mechanical performance when surface 

condition (e.g., bending) or when surface-located 

events (e.g., wear, impact) are service-life-limiters in 

engine and transportation system components. 

Approach 

There were four primary components to the 

FY 2005 work. Instrumented indentation and static 

indentation and instrumented scratch testing were 

used to characterize contact-induced damage 

processes in Si3N4. A criterion was developed that 

can be used to rank resistance to contact-induced 

damage. Several grades of Si3N4 were evaluated 

(e.g., existing roller element and ball bearing 

compositions as well as standard reference 

materials). Lastly, grain size effects (in model Al2O3 

Heavy Vehicle Propulsion Materials 

ceramics) on damage processes from indentation and 

scratch testings were examined. 

Results 

Static and dynamic indentation hardnesses of 

Ceralloy 147-31N, NBD200, SN101C, TSN-03NH, 

and NC132 silicon nitrides were measured as a 

function of load and their results are respectively 

shown in Figures 1–5. The rationale for testing each 

grade: the 147-31N grade is presently used as a cam 

Figure 1. Vickers Hardness for Ceralloy 147-31N Si3N4 

as a function of load for quasi-static and 

dynamic indentation. 

Figure 2. Vickers Hardness for NBD200 Si3N4 as a 

function of load for quasi-static and dynamic 

indentation. 

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Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 3. Vickers Hardness for SN101C Si3N4 as a 

function of load for quasi-static and dynamic 

indentation. 

Figure 4. Vickers Hardness for TSN-03NH Si3N4 as a 

function of load for quasi-static and dynamic 

indentation. 

Figure 5. Vickers Hardness for NC132 Si3N4 as a function 

of load for quasi-static and dynamic indentation. 

roller follower material, so interest exists to quantify 

its indentation and scratch responses, NBD200, 

SN101C, and TSN-03NH are ball bearing grades so 

their testing is relevant to their examination of 

rolling contact fatigue behavior, NBD200 is a NIST 

SRM for Knoop hardness, and NC132 is a NIST 

SRM for fracture toughness so they serve as 

references upon performance comparison among all 

the silicon nitrides. An indentation size effect (ISE) 

and a rate effect are exhibited by all (increased H 

and greater ISE with rate). 

A good fit to the experimental data is often 

obtained using a formula that Frohlich et al.,1 and Li 

and Bradt2 attribute to Bernhardt3 

P = a1d + a2d2 , (1) 

where P is indentation load, d is indent diagonal 

size, and a1 and a2 are constants. The a1 term is 

relatable to the energy consumed in creating new 

surfaces,1–2 whereas the a2 term is relatable to the 

volume energy of deformation.4 The determination 

of those terms for each of the silicon nitrides as a 

function of rate is shown in Figure 6. Additionally, a 

new brittleness parameter was developed that 

involves the use of the ratio of a2/a1 because it can 

provide insight into the competition of those two 

mechanisms. The brittleness of the five silicon 

nitrides is shown in Figure 7, and the relation shows 

171 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 7. Comparison of a new Brittleness parameter 

(a2/a1) for the five evaluated silicon nitrides 

for static and dynamic loadings. The trends 

show that all the grades act less brittle when 

they were dynamically loaded. 

Figure 6. Comparison of a1 and a2 terms [see Eq. (1)] for the five silicon nitrides when indentation was performed 

statically and dynamically. 

that all five silicon nitrides respond in a less brittle 

fashion when they were dynamically indented. 

The exploration of the effect of indentation on 

residual stress generation was initiated using optical 

fluorescence. An illustration of this is shown in 

Figure 8 for a 99.9% alumina that was Knoop 

indented at 2 kg. A maximum tensile stress of 

~100 MPa is observed at the tip of the cracks that 

come from the Knoop indent diagonals. 

Quantification of these effects will enable 

engineering control of ceramic machining or 

allowable wear conditions to inhibit or avoid 

deleterious effects that may be more strongly linked 

to fracture or quasi-plastic processes or both. 

An ANSYS Hertzian indentation model that was 

developed in FY 2004 continued to be refined in 

FY 2005. A three-dimensional (3-D) model was 

developed in FY 2005 using “concrete” elements 

that will enable the prescription of strength 

asymmetry in tension and compression. A tensile 

strength value will be imputable that will be 

associated with fracture processes, whereas a 

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Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 8. Hydrostatic stress map generated with optical 

fluorescence about a 2-kg Knoop indent in a 

99.9% alumina having a ~2- to 3-μm average 

grain size. 

compressive strength value will be imputable that 

will be associated with contact-induced quasi-plastic 

deformation mechanisms. 

A substantial amount of instrumented scratch 

characterization occurred during the latter half of 

FY 2005. Testing of silicon nitride focused on the 

evaluation of Ceradyne’s Ceralloy 147-31N, which 

is a reaction-bonded material used in cam roller 

followers. Evaluation of three 99.9% aluminas 

having mean sizes of 2, 15, and 25 μm (see 

Figure 9) also proceeded; the rationale behind the 

alumina testing is to evaluate scratch response as a 

function of controlled grain size with the same 

chemical composition. Additionally, the testing of 

alumina affords the opportunity to map out residual 

stresses about the contact or scratch damage using 

optical fluorescence. Scratch damage evolution and 

its quantification are being related to machining and 

wear performance. 

The interpretation of the load histories during 

scratch testing with respect to microstructural 

examination is ongoing. Normal and tangential 

forces are continuously monitored during the 

(pendulum) scratch test. Examples of these with 

associated microstructural damage are shown in 

Figures 10–11 for Ceralloy 147-31N, and 

Figures 12–17, for the 2-, 15-, and 25-μm alumina, 

respectively. The effects of lateral crack formation 

on normal and tangential load histories and depth 

changes are illustrated in Figure 18. SEM 

observations show that there is an apparent grain 

size effect on the scratch responses; in particular, 

Figure 9. Microstructure of 99.9% 

aluminas with mean average 

grain size of (top) 2 μm, (middle) 

15 μm, and (bottom) 25 μm. 

173 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 10. Example of normal (blue) and tangential 

(red) load histories during a Ceralloy 147­

31N silicon nitride scratch test. 

Figure 12. Example of normal (blue) and tangential 

(red) load histories during a 99.9% 2-μm 

alumina scratch test. 

Figure 11. Scratch on Ceralloy 147-31N associated 

with boxed region in Figure 10. 

Figure 13. Scratch on 99.9% 2-μm alumina associated 

with boxed region in Figure 12. 

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Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 14. Example of normal (blue) and tangential 

(red) load histories during a 99.9% 15-μm 

alumina scratch test. 

Figure 16. Example of normal (blue) and tangential 

(red) load histories during a 99.9% 25-μm 

alumina scratch test. 

Figure 15. Scratch on 99.9% 15-μm alumina associated 

with boxed region in Figure 14. 

Figure 17. Scratch on 99.9% 25-μm alumina associated 

with boxed region in Figure 16. 

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FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 18. Evidence of lateral crack formation evident in the normal and tangential load histories (top 

graph) and in depth changes too (profilometry–bottom graph). 

scratch forces oscillate more significantly during the 

scratching of larger-grained alumina corresponding 

to a greater amount of lateral crack formation. 

μ-FEA, a “microstructure-based finite-element 

analysis” computer program, was copyrighted in 

FY 2005. μ-FEA is a graphics-user-interfaced 

LabVIEW executable program that serves as a preand post-processor to the ANSYS finite-element 

analysis solver and enables stress analyses of both 

real (e.g., imaged with SEM or OM) and simulated 

material microstructures. For one of its applications, 

μ-FEA is being used to study stress concentrators, 

deformation, and microcracking around pores and 

large grains and in secondary phases in materials 

subjected to a Hertzian contact stress gradient. 

Conclusions 

Static and dynamic indentation and instrumented 

scratch testing facilities were established to quantify 

the loads and stresses associated with the initiation 

of quasi-plasticity and fracture processes. Numerous 

silicon nitride ceramics were investigated, and rate 

effects on hardness (i.e., quasi-plasticity) were 

studied. A new brittleness parameter was developed 

that enables static and dynamic indentation 

responses to be ranked for different ceramics. A 

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Heavy Vehicle Propulsion Materials 

generic FEM was developed that can be used to 

predict the onsets of quasi-plasticity and fracture/ 

cracking in any target material as a function of 

indentation load, indenter material, friction, and 

confinement. Lastly, microstructure-based finiteelement analysis software (μ-FEA) was copyrighted. 

The software enables stress analysis characterization 

at the microstructural scale and to help with the 

interpretation of quasi-plastic and fracture processes. 

References 

1. F. Fröhlich, P. Grau, and W. Grellmann, 

Phys. Stat. Sol. (a), 42, 79–89 (1977). 

2. H. Li and R. C. Bradt, J. Non-Crystalline 

Solids, 146, 197–212 (1992). 

3. V. E. O. Bernhardt, Z. Metallkde, 33, 

135–144 (1941). 

4. H. Li, A. Ghosh, Y. H. Han, and R. C. Bradt, 

J. Mat. Res., 8, 1028–1032 (1993). 

Presentations and Publications 

A. A. Wereszczak and T. P. Kirkland, 

“Exclusivity of Strength-Limiting Extrinsic Surface 

and Hybrid Flaws,” in press, Ceramic Engineering 

and Science Proceedings (2006). 

FY 2005 Progress Report 

H. Wang and A. A. Wereszczak, “Mechanical 

Responses of Silicon Nitrides Under Dynamic 

Indentation,” Ceramic Engineering and Science 

Proceedings, 26, 275–283 (2005). 

G. Subhash, M. A. Marszalek, A. A. 

Wereszczak, and M. J. Lance, “Evaluation of 

Scratch Resistance and Residual Stresses in 

Machined Silicon Nitride,” Ceramic Engineering 

and Science Proceedings, 26, 285-292 (2005). 

H. Wang, A. A. Wereszczak, and M. K. Ferber, 

“Characterization of Dynamic Indentation of Silicon 

Nitrides by Using Vickers Indentation,” in 

preparation. 

“Research Overview of ORNL’s Ceramic 

Science and Technology Group,” University of 

Alabama, Tuscaloosa, Alabama, April 20, 2005. 

Special Recognitions and Awards/Patents Issued 

μ-FEA Software, M. K. Ferber and A. A. 

Wereszczak, U.S. Copyright Office, Certificate of 

Registration No. TXu1-234–195, 25 April 2005. 

177 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Q. Powder Processing of Nanostructured Alloys Produced by Machining 

S. Chandrasekar, K. P. Trumble, and W. D. Compton 

College of Engineering 

Purdue University, IE, GRIS 

315 North Grant Street 

West Lafayette, IN 47907-2023 

(765) 494-3623; fax: (765) 494-5448; e-mail: chandy@purdue.edu 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Prime DOE Contract No.: DE-AC05-00OR22725 

Subcontractor: Purdue University, West Lafayette, Indiana 

Objective 

• Develop manufacturing technologies for creating advanced nanostructured alloys with enhanced mechanical 

properties for transportation applications. 

Approach 

• Create nanostructured chips by machining alloys of aluminum, titanium, and iron under controlled machining 

conditions. 

• Characterize microstructure and determine relationships between structure and mechanical properties of the 

nanostructured chips. 

• Develop methods for converting the nanostructured chips into particulates suitable for processing bulk forms. 

• Develop low-temperature consolidation routes for creating monolithic and composite materials without 

significant coarsening of the nanocrystalline microstructures. 

• Characterize microstructure, mechanical properties, and performance of the bulk forms. 

Accomplishments 

• Produced nanostructured chips of a variety of alloys and metals—including Al6061-T6, copper, iron, 

commercially pure titanium, stainless steel, and Inconel—by machining. The nanostructured chips are up to 

three times as hard as the bulk material. 

• Developed low-temperature consolidation routes that limit coarsening in Al6061-T6, and steels. Achieved a 

high degree of interparticle bonding in room temperature powder extrusion of nanostructured Al6061-T6. 

• Developed highly stable nanostructured (HSN) materials, for example, Inconel 718, Waspaloy, of high-strength 

through a combination of grain-size refinement by machining and creation of fine-scale precipitates by heat 

treatment. 

• Developed a new process, extrusion-machining, to engineer bulk samples, for example, foils, filaments, with 

ultra-fine grained (UFG) microstructure and enhanced mechanical properties. 

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FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Future Direction 

• Continue development of low-temperature consolidation routes to produce bulk monolithic and/or composite 

alloys with nanocrystalline microstructure. 

• Characterize microstructure and properties of the bulk forms to demonstrate retention of nanostructures and 

enhanced mechanical properties. 

• Determine tensile and wear properties of bulk forms to demonstrate integrity of the consolidated alloy samples 

and guide application. 

• Develop (identify 1–2) applications in the transportation sector for prototype development. 

Introduction 

This project builds on the recent discovery of a 

low-cost means of producing nanostructured 

materials in essentially any alloy via normal 

machining. Comminution of the nanostructured 

chips enables large-scale production of nanocrystalline particulate, which can be converted into 

nanostructured bulk forms using powder processing 

methods. Particulate having controlled size and 

shape can also be produced by a modulated 

machining process. The chips and particulates also 

may be used as continuous or discontinuous 

reinforcements in metal or polymer matrices to 

create advanced composites. 

Progress from FY 2005 in two main areas is 

highlighted in this report: (1) low-temperature 

consolidation routes, primarily of ultrafine-grained 

Al6061-T6 and (2) microstructure development in 

machining of solution-treated alloys. Modulated 

machining could be in there as a third, but not 

necessary or enough space. 

Approach 

Three routes for consolidation and densification 

of the nanostructured particulate to produce 

monolithic and composite materials were pursued: 

(1) shear-based deformation processing (e.g., 

powder extrusion); (2) rapid infiltration processing 

using low-melting metal alloys; and (3) polymer diepress bonding. 

These routes are all geared toward densification 

and bonding of the particulate while minimizing 

thermal coarsening. For the powder extrusion, 

nanostructured Al-6061-T6 chips were converted 

into particulate using attrition milling. The resulting 

particulate was mixed with varying weight 

percentages of commercial gas-atomized aluminum 

powder; this mixture was powder extruded through 

steel dies at extrusion ratios of 10–60. The extruded 

samples were prepared in the form of cylinders up to 

10 mm in diameter and 5–25 mm in length. 

The polymer bonding is a new approach in 

which a thermosetting polymer resin is used, 

together with high-pressure compression molding in 

a steel die to yield a structure consisting of lowvolume fractions of polymer bonding the 

nanostructured metal particles. The focus so far has 

been on Al-6061-T6 chip powders bonded with 

different epoxies. This approach is based on the 

observation that the cure conditions of the resin are 

well below the known time-temperature conditions 

for coarsening (and loss of hardness) of the 

nanostructures in the chip particles. Application for 

lightweight, wear-resistant components is 

envisioned. 

Another avenue of investigation has focused on 

creation of high-strength, high-temperature 

nanostructured materials such as titanium and 

nickel-based alloys, with exceptional thermal 

stability of the microstructure. In this context, we 

have demonstrated a means of extending the 

temperature range for structural applications of 

nanomaterials through the introduction of a 

dispersion of second-phase precipitates, whose 

thermal stability determines the temperature at 

which the strength deteriorates. Precipitatestabilized nanostructured materials synthesized from 

solution-treated Inconel 718 are shown to be 

extremely stable even after prolonged heat treatment 

for 500 h at temperatures as high as one-half of the 

melting point; we designate them as “highly-stable 

nanostructured (HSN) materials.” The enhanced 

thermal stability of the material’s strength results 

from retention of both the fine microstructure and 

the fine precipitate phase coexisting at the high 

temperatures. This novel microstructure was created 

by first deforming solution-treated Inconel 718 to a 

180
Heavy Vehicle Propulsion Materials 

large shear strain by machining, followed by thermal 

aging of the chip to create fine-scale precipitates. 

Results 

Room Temperature Powder Extrusion of 

Al-6061-T6 Chip Particulate 

Extrusion of nanostructured Al6061-T6 chip 

particulate results in poor densification and bonding. 

However, by blending the chip particulate with pure 

aluminum powder, near complete densification is 

possible. The effect of chip powder particle size 

relative to the pure aluminum (Ampal, 99.9%, gasatomized, ~60 μm) and extrusion ratio are under 

investigation. Figure 1 shows the extrudate and 

corresponding microstructure for a ~180-μm nanochip particulate blended 50–50 with pure aluminum 

and extruded at room temperature with an extrusion 

ratio of 42. Immersion density measurements indicated a relative density of 98.7%, consistent with 

the small amount of porosity observed in optical 

Figure 1. (a) Macrograph showing Al-6061-T6 

nanochip/pure aluminum (1:1 by volume) 

extrudate; extrusion ratio = 42. (b) Optical 

micrograph of corresponding microstructure in 

cross-section showing small amount of 

residual porosity on the order of the nanochip 

particle size but otherwise well-bonded 

structure. 

FY 2005 Progress Report 

microscopy. The chip particles are clearly distinguished in the micrographs and appear well— 

bonding with the continuous pure aluminum matrix. 

Tensile testing (3-mm-diam gage sections) completed late this year confirmed a significant degree 

of bonding. A fully sintered pure aluminum powder 

specimen was tested in the same manner to serve as 

a control sample. The detailed test results are as 

follows: 

Properties 

Material 

(YS/UTS/%El) 

Fully sintered pure 60 MPa/75 MPa/5% 

aluminum 

6061 nanochips + pure 170 MPa/170 MPa/nil 

aluminum 

6061 nanochips + 601AB 200 MPa/200 MPa/nil 

The measured strength values indicate 

significant interparticle bonding in the ambient 

temperature extruded materials based on machining 

chips. The tensile strength values furthermore 

support the apparent interparticle bonding revealed 

by metallography, as we have reported before. For 

comparison, sintered 601AB exhibits tensile 

property sets from 50 MPa/125MPa/8% to 

240 MPa/250 MPa/2%, depending on the density 

and temper.1 Work is ongoing to quantify the degree 

of bonding in terms of tensile strength, which, with 

such limited ductility, is also strongly dependent on 

the size of pores and other flaws. 

Larger scale extrusions using the ORNL 

extrusion facilities are in progress. 

Polymer-Bonded Chip Particulate Composites 

Die press bonding has been used to produce 

bulk Al-6061-T6 samples from nanostructured chip 

particulate using a small amount (<10%) of epoxy as 

binder. Figure 2 shows a typical example. Small 

amounts of porosity persist, especially where larger 

particles were not packed uniformly, but most 

particle boundaries show a thin bond layer of epoxy. 

Epoxy and porosity volumes as low as 6% have 

been achieved so far, and preliminary nanohardness 

measurements show little or no loss of hardening 

imparted in producing the chips, consistent with 

separate measurements of coarsening-softening 

kinetics in this alloy. 

181 
FY 2005 Progress Report 

Figure 2. Optical micrograph showing epoxy-bonded 

Al-6061-T6 chip particulate (15–100 μm) 

pressed at 1000 MPa and cured at 100°C. 

Work in the next year will proceed to tensile 

testing in order to evaluate the degree of bonding. 

Although high strength and ductility are not 

predicted, the high hardness should impart improved 

wear resistance together with the shaping advantages 

of powder processing. 

Highly Stable Nanostructured Materials 

Figure 3 illustrates the nanostructured material 

composed of grains typically 100 nm in size 

produced from solution-treated microcrystalline 

Inconel 718 (grain size ~50 μm in the initial state) 

Figure 3. TEM micrograph of nanostructured Inconel 

718 produced by large strain machining (shear 

strain ~6) from solution-treated microcrystalline Inconel 718. Average grain size is 

~100 nm. 

Heavy Vehicle Propulsion Materials 

due to plastic shear strains of ~6. The microstructure 

is composed primarily of equi-axed grains, along 

with some elongated structures. The diffraction 

pattern inset in Figure 3 indicates the presence of 

large angle grain boundaries in this nanostructured 

material. The Vickers microhardness value of this 

microstructure was 550 kg/mm2

 (~5.4 GPa), nearly 

twice that of the undeformed bulk microcrystalline 

solution-treated material (300 kg/mm2

). 

When this nanostructured material was heat 

treated at 600°C (~ one-half the melting point), a 

rapid hardening was found that can be associated 

with precipitation in the nanostructured solutiontreated material. Figure 4 illustrates the variation in 

microhardness of the nanostructured Inconel 718 

with temperature and the attainment of a peak 

hardness value of 630 kg/mm2 (~6.2 GPa) after 

about 6 h of annealing. This hardness value 

corresponds to a tensile yield strength of ~2 GPa for 

this material. Figure 4 confirms the extraordinary 

stability of material strength (hardness) of this HSN 

Inconel 718 even after 30 h at 600°C (~ one-half of 

the melting temperature). The peak hardness of 

Figure 4. Variation of Vickers hardness with annealing 

time at 600°C for high-stability nanostructured 

(HSN) Inconel 718 and bulk microcrystalline 

Inconel 718 (undeformed bulk, grain size 

~50 μm). The bulk peak hardness of 

450 kg/mm2

 is that of microcrystalline Inconel 

718 after prolonged double-ageing treatment 

and is shown for comparison. HSN Inconel 

718 was also heat treated at 600°C for 500 h, 

and its hardness was found to remain stable at 

~630 kg/mm2

. 

182
Heavy Vehicle Propulsion Materials 

630 kg/mm2 is unchanged throughout this heat 

treatment, and this hardness value is still ~40% 

greater than that of peak-aged Inconel 718 

(~450kg/mm2

) in the conventional microcrystalline 

form. We also performed annealing experiments at 

600°C for 500 h on the HSN Inconel 718, and even 

such prolonged heat treatment was found to have a 

negligible effect on the stable hardness value 

(~630 kg/mm2

). 

The materials design principle propounded 

above can be generalized to other precipitationtreatable materials and offer a new line of 

development for this important class of structural 

materials. Other material systems that form stable, 

fine precipitate phases, for example, P-series 

stainless steels, Waspaloy, can be expected to 

behave in a similar manner in the nanostructured 

state. Exceptional thermal stability and retention of 

the material strength are ensured by the stable, fine 

precipitate phase and the fine microstructure 

coexisting at high temperatures for significant 

lengths of time, resulting in a class of HSN materials 

with enormous potential for high-strength and/or 

high-temperature applications. These avenues will 

be investigated in detail in the coming year. 

Nanostructured Pure Titanium by Machining 

Figure 5 shows a TEM micrograph and selected 

area diffraction pattern of nanostructured titanium 

formed in machining chips. The starting grain size 

of the commercially pure titanium was ~50 μm, and 

TEM revealed a grain size range of 50 to 200 nm in 

the chip specimens. Correspondingly, Vickers 

Figure 5. TEM micrograph. 

FY 2005 Progress Report 

hardness increased from 144 ± 5 to 230 ± 9 kg/mm2, 

consistent with the machining response of a wide 

range of other pure metals and alloys. 

References 

1. Metals Handbook, 9th ed., Vol. 2 (1979). 

Presentations and Publications 

S. Swaminathan, S. Chandrasekar, W. D. 

Compton, K. P. Trumble, and A. H. King, 

“Microstructural Refinement in Single-Phase 

Copper Solid Solutions by Machining,” in 

Nanoscale Materials and Modeling-Relations 

Among Processing, Microstructure and Mechanical 

Properties, Materials Research Society (MRS) 

Proceedings, 821, p. 8.27 (2004). 

S. Swaminathan, C. Swanson, T. L. Brown, 

R. F. Kezar, S. Chandrasekar, W. D. Compton, and 

K. P. Trumble, “Microstructure Refinement in Steels 

by Machining,” in Nanoscale Materials and 

Modeling—Relations Among Processing, 

Microstructure and Mechanical Properties, 

Materials Research Society (MRS) Proceedings, 

821, p. 9.5 (2004). 

S. Swaminathan, M. Ravi Shankar, S. Lee, 

J. Hwang, A. H. King, R. F. Kezar, B. C. Rao, T. L. 

Brown, S. Chandrasekar, W. D. Compton, and K. P. 

Trumble, “Large Strain Deformation and Ultra-Fine 

Grained Materials by Machining,” Materials Science 

and Engineering A, 410–411, 358–363, 2005. 

M. R. Shankar, S. Chandrasekar, A. H. King, 

and W. D. Compton, “Microstructure and Stability 

of Nanocrystalline Aluminum Created by Large 

Strain Machining,” Acta Metallurgica, 53, 4781– 

4793 (2005). 

M. Ravi Shankar, S. Chandrasekar, W. D. 

Compton, and A. H. King, “Characteristics of 

Aluminum 6061-T6 Deformed to Large Strains by 

Machining,” Materials Science and Engineering A, 

410–411, 364–368 (2005). 

M. Sevier, S. Lee, M. Ravi Shankar, H. T. Y. 

Yang, S. Chandrasekar, and W. D. Compton, 

“Deformation Mechanics Associated with Formation 

of Ultra-Fine Grained Chips in Machining,” 

Materials Science Forum, 503–504, 379–384 

(2005). 

Y. Uluca, B. C. Rao, M. Ravi Shankar, T. L. 

Brown, J. B. Mann, S. Chandrasekar, and W. D. 

Compton, “Low-Cost Ultra-Fine Grained Materials 

from Aerospace Machining Chips,” Proceedings 

183 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

SAE Aerotech Congress & Exhibition, Dallas, 

Texas, in press. 

S. Swaminathan, S. Chandrasekar, W. D. 

Compton, A. H. King, and K. P. Trumble, “Large 

Strain Deformation of Single-Phase Solid Solutions 

by Plane Strain Machining,” Materials Science 

Forum, 503–504, 651–656 (2005). 

S. Swaminathan, M. Ravi Shankar, B. C. Rao, 

S. Chandrasekar, W. D. Compton, A. H. King, and 

K. P. Trumble, “Nanostructured Materials by 

Machining,” in Proceedings of ASME IMECE, 2005. 

Patents 

J. B. Mann, M. Ravi Shankar, S. Chandrasekar, 

W.D. Compton, and W. Moscoso, “Direct 

Production of Nanocrystalline Particulate by 

Modulated Machining,” U.S. Patent Pending Ser. 

No. 60/677,247, May 3, 2005. 

184
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

R. Deformation Processes for the Next-Generation Ceramics 

P. F. Becher 

Oak Ridge National Laboratory 

P.O. 2008, MS 6068 

Oak Ridge, Tennessee 

(865) 574-5157; fax: (865) 574-6098; e-mail: becherpf@ornl.gov 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Objectives 

• Assess the plastic deformation behavior in nanocrystalline oxide ceramics and applicability to fabrication of 

complex shapes. 

• Characterize thermal and mechanical properties to assess influence of nanocrystalline microstructures for 

gaining improved performance for wear and other components. 

Approach 

• Collaborate with members of the Pennsylvania State University task on Synthesis of Nanocrystalline Ceramics, 

develop advanced colloidal technology-based processing to fabricate monolithic bodies and sintering 

technology to be able to fabricate dense monolithic nanocrystalline (i.e., grain sizes <100-nm) bodies. 

• Characterize the potential for enhanced plasticity in ceramics resulting from reducing grain size below 100 nm 

and with the application of applied electric fields with regard to mechanical behavior improvements and 

applicability to forming and surface finishing technologies. 

Accomplishments 

• Expanded the fabrication of powder compacts to include granulation technology, which can readily produce 

larger samples and more complex geometries. The recent sintering studies of granulated compacts were sintered 

to >92% of theoretical density and exhibited grain sizes of 6 to 140 nm. However, these bodies contained 

mainly monoclinic phase, which was attributed to a loss of yttria during powder synthesis. This is now being 

addressed in the Pennsylvania State University task. 

• Observed electric field effects on deformation to exhibit a temperature dependence that apparently is altered by 

the mechanism dictating the plastic flow. Above 1200°C, the reduction in yield stress increases with decrease in 

temperature. At 1200°C, the field effect becomes more sensitive to the strain rate, which suggests a shift from 

grain boundary sliding to a dislocation-related process. 

Future Direction 

• Develop a two-step sintering profile to be able to optimize the microstructure of large compacts produced by 

the granulation process. Room temperature fracture and wear properties will be examined to seek approaches 

for developing ceramics with greater wear resistance. 

185 
FY 2005 Progress Report 

Introduction 

In addition to lower specific weight, ceramics 

have a number of attractive properties for application as components in various heavy-vehicle 

systems. However, many heavy-vehicle components 

require ceramics with complex shapes as well as 

function. Furthermore the demanding service conditions necessitate improved mechanical reliability 

together with fracture and/or wear resistance. 

Underlying the performance and use of ceramics is 

their tendency to fail in a brittle fashion except at 

temperatures generally well above 1000°C. Complex 

shape forming processes for ceramics are relegated 

to approaches based on powder processing as a 

result of the excessive temperatures for hot forming. 

On the other hand, metallic alloys can be hot worked 

at temperatures below 1000°C to not only form the 

desired shape but also to enhance their properties. 

The ductility (>20% plastic strain) of metallic alloys 

at temperatures below 500°C is the source of their 

10- to 100-fold greater fracture toughness as compared to that of ceramics. 

In ceramic systems, one needs to combine processing to tailor microstructures with novel concepts 

that offer the potential to enhance deformation. One 

such concept that has seen considerable interest is 

that of superplastic deformation of submicron-grainsized ceramics above 1000°C.1

 Recent studies 

suggest that reducing the microstructure to achieve 

grain sizes <100 nm could lead to permanent/ 

plastic deformation in ceramics at quite modest 

temperatures.

2,3

 Recent studies have shown that the 

application of electric fields can reduce the yield 

stress or increase the creep rates of ceramic 

systems.4

 To date, this has not been exploited to 

determine the effects on the deformation of 

nanocrystalline oxides. The present study seeks to 

examine the influence of true nanocrystalline microstructures on the mechanical properties of ceramics, 

as well as approaches to economically fabricate 

nanocrystalline ceramics monoliths. 

Hardness values generally obey the Petch-Hall 

behavior with the hardness rising as the grain size is 

reduced. However, results show that when the grain 

size is ~20 nm, the room temperature hardness then 

decreases as the grain size continues to be 

decreased.1,5 This behavior has two important 

implications for applications of interest here. One is 

that some permanent deformation might be achieved 

in ceramics with grain sizes <30 nm, and this might 

Heavy Vehicle Propulsion Materials 

offer improved reliability. More likely of interest is 

the fact that ceramics with grain sizes below 100 nm 

but above the softening point will be extremely hard, 

and that offers substantial advantages in overcoming 

wear problems in rolling stock. In addition, surface 

roughness of as-sintered ceramics decreases as the 

grain size decreases. In fact, the as-sintered surfaces 

of alumina substrates with grain sizes <500 nm 

exhibit a mirror-like finish without any surface 

preparation. This feature should be applicable to the 

fabrication of low-friction components. 

Approach 

This project involves two major thrusts. The 

first is a collaborative effort with researchers in the 

Synthesis of Nanocrystalline Ceramics task at 

Pennsylvania State University in the processing and 

sintering of nanocrystalline ceramics. The goal 

being to develop the technology to form “green” 

(unfired) bodies of nanocrystalline particles with 

densities well in excess of 50% of the theoretical 

density to allow one to devise pressureless sintering 

profiles to produce dense monolithic bodies with 

grain sizes <100 nm (see Ref. 6 for a review of 

processing of nanocrystalline ceramics). This is 

aimed at finding approaches to produce sizeable 

monolithic bodies for subsequent evaluation. The 

second thrust focuses on the mechanical and friction 

and wear behavior of nanocrystalline ceramics and 

the influence of external parameters [e.g., stressing 

rate, temperature, environment (e.g., air, vacuum)], 

and electric fields on the plastic deformation 

behavior. 

Synthesis, Evolution, and Stability of 

Nanostructures in Bulk Materials 

Outstanding challenges in particle processing to 

form bulk, dense nanograin-sized materials include 

the dispersion of the nanosize particles and control 

of gas evolution from the very high surface area 

compacts during sintering. Typically, nanometer 

size particulates suffer from a high degree of 

agglomeration, which results in poor particle packing (green densities <40% of theoretical), which 

inhibits sintering and promotes undesirable grain 

growth. Two routes to obtaining well-dispersed 

powders include comminution and colloidal 

processing.7,8

 By optimizing colloidal processing 

approaches, one should be able to produce “green” 

bodies with densities in excess of 60% of theoretical 

186
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

by control of interparticle forces through appropriate 

chemical additives to suspensions of nanoparticulates combined with filter pressing or pressure 

casting of the high-solids-content slurries. By controlling the packing of nanoparticulates one should 

be able to control the porosity to achieve the narrow 

size distribution necessary, which then will enhance 

densification. 

Densification of nanocrystalline ceramics has 

typically involved the use of applied pressure to 

achieve dense sub-100-nm grain sized samples.9

 In 

those few cases where pressureless sintering has 

been employed successfully, long hold times at the 

sintering temperatures were required to achieve densities ≥95% of theoretical, which allows undesired 

grain growth to proceed.10 Heating to a temperature 

sufficient to achieve densities of ≥75% of theoretical 

followed by rapidly lowering the temperature to a 

level to obtain sufficient boundary diffusion to 

remove pores trapped on boundaries while minimizing grain growth have been only partially 

successful because long hold times at the lower temperature were still required.11 Here we seek to optimize the processing of the “green” bodies and utilize 

modified sintering cycles to achieve dense 

sub-100-nm grain size oxides. 

Deformation Behavior in Nanostructured 

Materials: Baseline deformation data will be established through studies of high purity, dense submicron grain size zirconia ceramics in order to 

evaluate the influence of nanocrystalline microstructures. These studies will then be extended to 

<100-nm grain sized ceramics. This will include 

(a) nanostructured ceramics to explore for the 

softening effects observed in nanocrystalline metals 

by evaluating both the creep response over a range 

of temperatures and applied stresses, and the yield 

stress as a function of strain rate and temperature, 

and (b) the influence of external factors (e.g., 

electric fields) that have been shown to enhance 

deformation processes in ceramics with micrometer 

grain sizes. Modeling studies will be conducted to 

describe the deformation behavior and to address 

such phenomenon as space charge/solute 

segregation coupled with electric field effects. 

Additionally, the fracture, hardness, and wear 

response of nanocrystalline oxides will be evaluated 

to assess their application in wear components. 

Results 

Sintering 

The Synthesis of Nanocrystalline Ceramics task 

recently expanded the powder processing method to 

include granulation to more readily produce large 

powder compacts. In the granulation step, the 

chemically derived particulates are dried and treated 

to form granules, which are basically large, soft 

agglomerations. These then could be pressed to form 

bodies of various dimensions to densities of 50 to 

55% of theoretical density! The process is very 

amenable to commercial manufacturing. The ORNL 

studies have focused on exploring the densification 

of compacts based on the granulation process 

approach. Sintering trials at ORNL showed that 

these granulated compacts could then be densified to 

greater than 90% of theoretical density, while 

achieving grain sizes <200 nm. Subsequent characterization revealed that the sintered materials consisted of primarily monolithic zirconia rather than 

the tetragonal phase. The monoclinic phase is a 

result of the martensitic transformation from the 

tetragonal phase upon cooling from the sintering 

temperature. Transformation to the monoclinic 

phase involves a volume expansion of ~4%, which is 

quite large and the associated stress state results in 

extensive microcrack formation. Such microcracking leads not only to a decrease in density but also to 

damage and loss of structural intensity. 

Earlier work demonstrated that colloidal processed ZrO2-1.7 mol % Y2O3 bodies could be sintered to densities >95% of theoretical density (TD) 

when held at 1150°C for only 5 min. These were 

fully tetragonal phase and had grain sizes <100 nm. 

Even with 200- to 400-nm grain sizes, ZrO2­

1.7 mol % Y2O3 bodies were fully tetragonal. On the 

other hand, granulated compacts sintered at 1125°C 

for 5 min had densities > 90%TD but were composed of 75% monoclinic and 25% tetragonal phases 

even though the grain size was below 100 nm. 

Because the grain size is well below 200 nm, the 

formation of monoclinic phase must be attributed to 

an insufficient yttria solute content, which is added 

to stabilize the tetragonal phase. Current synthesis 

efforts are addressing steps to control and maintain 

the yttria levels in the as-synthesized powders. 

187 
FY 2005 Progress Report 

Deformation studies 

Collaborations with Prof. Hans Conrad and 

Dr. Kang Jung of North Carolina State University 

are assessing the reduction in flow stress in zirconia 

ceramics due to the application of an applied direct 

current (dc) field. In this case, the chemical/ 

electrochemical potential associated with vacancy 

flux can be altered by both an applied stress and an 

applied electric field: 

μ* = μo + Zeφi − σΩ − Zeφa , (1) 

where μ is the potential (superscript o for the standard state, superscript * with electric field); Z is the 

valence of the ion; e is the charge in coulombs; σ is 

the stress; Ω is the atomic volume; Φa is the applied 

electric potential (= EΧSC); E is the applied electric 

field; ΧSC is the width of the space charge zone at 

the grain boundaries; and Φi is the local potential. 

Thus, one can enhance the flow of vacancies needed 

for material transport and dislocation motion by 

applying stress and an electric field. The direction of 

the vacancy flux is a function of the sign of the 

applied stress (i.e., vacancies flow to regions under 

compression). This applies locally to vacancy flow 

in the deformation of a polycrystalline ceramic 

where vacancies move from grain boundaries under 

tension to those in compression while ions/atoms 

move in the reverse direction. As a result, one can 

then reduce the applied stress required to deform the 

material by applying an electric field: 

ΔσE = ZeExSC Ω , (2) 

where ΔσE represents the reduction in the flow stress 

due to the applied field. 

The width of the space charge zone is related to 

the Debye length of the material: 

XSC = βλ , (3) 

where it is a proportionality factor in the range of 2 

to 5. Theoretical treatments show that the Debye 

length (δ) in ionic solids increases as both the bulk 

solute content and temperature are decreased.12,13

 In 

addition, δ must be much less than the distance 

between opposite space charge zones (i.e., << grain 

size). The influence of temperature and solute content on the Debye length in alkali halides is shown in 

Heavy Vehicle Propulsion Materials 

Figure 1. The increases in the Debye length due to 

decreases in temperature and solute content will be 

reflected as increases in the space charge width 

[Eq. (3)]. One should then expect that the decrease 

in flow/yield stress brought about by application of 

an electric field would increase as the temperature 

decreases [Eq. (2)]. This is indeed the case found for 

polycrystalline MgO and ZrO2-3 mol % Y2O3 

ceramics (Figure 2). 

Initial studies on the influence of an electric 

field on the yield stress in compression of tetragonal 

zirconia ceramics containing 2 and 2.5 mol % yttria 

with grain sizes of 350 to 400 nm were conducted at 

1200°C to determine a lower temperature limit for 

such effects. At 1200°C, the electric field effects 

apparently are much more sensitive to strain rates as 

analysis of the results indicated the field effects 

increase with decrease in strain rates. Analysis of the 

data reveals the stress exponent at 1200°C is ~4 in 

the zirconia with 2 and 2.5 mol % yttria. On the 

other hand at ≥1400°C the stress exponent is 2 both 

for the zirconia containing 2 and 2.5 mol % yttria 

ceramics without an applied field and the zirconia-3 

mol % yttria with an applied field. As a result, the 

field effect just begins to be indicated at strain rates 

of 10–4/s at 1200°C, which is the limit of the test 

facility available for the field effect tests. Studies are 

now under way to examine the effects of the electric 

field on the yield stress of the zirconia with 2 and 

2.5 mol % yttria at 1450°C. This will allow us to test 

the effects of solute content on the predicted space 

Figure 1. The Debye length is predicted to increase as 

either the temperature or the solute content 

decrease as shown for the case of potassium 

chloride.12 

188
Heavy Vehicle Propulsion Materials 

(a) 

15 

FY 2005 Progress Report 

140 nm. However, these bodies contained mainly 

monoclinic phase, which was attributed to a loss of 

yttria during powder synthesis. This is now being 

addressed in the Pennsylvania State University task. 

Electric field effects on deformation were 

observed to exhibit a temperature dependence that 

apparently is altered by the mechanism dictating the 

plastic flow. Above 1200°C, the reduction in yield 

stress increases with decrease in temperature. At 

1200°C, the field effect becomes more sensitive to 

the strain rate, which suggests a shift from grain 

boundary sliding to a dislocation-related process. 

References 

1. T. G. Nieh and J. Wadsworth, “Superplastic 

Behavior of a Fine-Grained Yttria-Stabilized, 

Tetragonal Zirconia Polycrystal (Y-TZP),” Acta 

Metall et Mater., 38, 1121–1133 (1990). 

2. H. Conrad and J. Narayan, “On the Grain 

Size Softening in Nanocrystalline Materials,” Script. 

Mater., 42, 1025–30 (2000). 

3. M. Yoshida, Y. Shinoda, T. Akatsu, and 

F. Wakai, “Deformation of Monoclinic ZrO2 Polycrystals and Y2O3-Stabilized Tetragonal ZrO2 Polycrystals below the Monoclinic-Tetragonal Transition 

Temperature,” J. Am. Ceram. Soc., 85(11), 2834–36 

(2002). Ibid, “Superplasticity-Like Deformation of 

Nanocrystalline Monoclinic Zirconia at Elevated 

Temperatures,” J. Am. Ceram. Soc., 87(6), 1122–25 

(2004). 

4. H. Conrad, “Space Charge and the Dependence of the Flow Stress of Ceramics on an Applied 

Electric Field,” Scripta Mater., 44, 211–16 (2001). 

5. T. G. Nieh and J. Wadsworth, “Hall-Petch 

10 

5 

1400 1450 1500 1550 1600 1650 

ZrO -3 mol.% Y O 2 2 3 

G.S. = 440 nm 

ε = 1.3 x 10-4/s . 

E = 0 kV/cm 

E = 1 kV/cm 

0 

Temperature, °C 

(b) 

Figure 2. The effectiveness of the application of a dc Yield Stress, MPa 

electric field in decreasing the tensile yield 

stress increases as the temperature in the range 

of 1650°C to 1450°C in the case of both cubic 

magnesia (a)

14 and cubic zirconia containing 3 

mol % yttria15 (b) ceramics. 

charge zone size and the role of the space charge 

zone on the yield behavior. 

Conclusions 

The fabrication of powder compacts was 

expanded to include granulation technology, which 

can readily produce larger samples and more complex geometries. The recent sintering studies of 

granulated compacts were sintered to >92% of theoretical density and exhibited grain sizes of 6 to 

Relation in Nanocrystalline Solids,” Scripta Metall., 

25, 955–58 (1991). 

6. M. J. Mayo, “Processing of Nanocrystalline 

Ceramics from Ultrafine Powders,” Int’l Mater. 

Rev., 41(3), 85–115 (1996). 

7. J. H. Adair et al., “Dynamic Control and 

Enhanced Chemical Milling of Ceramics to Submicron Particle Sizes,” U.S. Patent 6,415,966 B1, 

July 9, 2002. 

8. A. Bleier, P. F. Becher, K. B. Alexander, and 

C. G. Westmoreland, “The Effect of Aqueous Processing Conditions on the Microstructure and Transformation Behavior of Al2O3-ZrO2(CeO2) 

Composites,” J. Am. Ceram. Soc., 75(10), 2649–58 

(1992). 

189 
FY 2005 Progress Report 

9. R. Chaim and M. Hefetz, “Fabrication of 

Dense Nanocrystalline ZrO2-3 wt. % Y2O3 by HotIsostatic Pressing,” J. Mater. Res., 13(7), 1875–80 

(1998). 

10. A. Bravo-Leon, Y. Morikawa, 

M. Kawahara, and M. J. Mayo, “Fracture Toughness 

of Nanocrystalline Tetragonal Zirconia with Low 

Yttria Content,” Act. Mater., 50, 4555–62 (2002). 

11. I.-W. Chen and X.-H. Wang, “Sintering 

Dense Nanocrystalline Ceramics Without FinalStage Grain Growth,” Nature, 404, 168–71 (2000). 

12. K. L. Kliewer and J. S. Koehler, “Space 

Charge in Ionic Crystals. I. General Approach with 

Application to NaCl,” Phys. Rev. A, 140(4A), 

1226–40 (1965). 

Heavy Vehicle Propulsion Materials 

13. M. F. Yan, R. M. Cannon, H. K. Bowen, 

and R. L. Coble, “Space-Charge Contribution to 

Grain-Boundary Diffusion,” J. Am. Ceram. Soc., 

60(3–4), 120–27 (1977). 

14. H. Conrad and D. Yang, “Influence of an 

Electric Field on the Plastic Deformation of FineGrained MgO at High Homologous Temperatures,” 

Acta Mater., 48, 4045–52 (2000). 

15. K. Jung and H. Conrad, personal communication, 2005. 

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Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

S. Development of Titanium Component Applications in Heavy-Duty Diesel 

Engines 

Nan Yang, Jeremy Trethewey, Michael Pollard, Jesus Chapa-Cabrera, and John Grassi 

Advanced Materials Technology, Caterpillar, Inc. 

Tech Center 854-80 TC-E 

14009 Old Galena Rd. 

Mossville, IL 61552 

(309) 578-0056; fax: (309) 578-2953; e-mail: Trethewey_Jeremy_S@cat.com 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: Caterpillar, Inc., Peoria, Illinois 

Objectives 

• Develop and demonstrate a high-strength, lightweight titanium aluminide (TiAl) turbine for a heavy-duty diesel 

engine. The unique material properties of TiAl will enable the next generation of turbocharger designs to have 

⎯ improved thermal efficiency, and 

⎯ increased transient response. 

Approach 

• Select the appropriate TiAl compounds and procure high-quality TiAl castings/forgings. 

• Develop robust TiAl joining technology to deliver prototype TiAl turbine assemblies for engine testing. 

• Demonstrate improved thermal efficiency and transient response in engine test. 

• Establish a TiAl material property database to be used as a design tool and for improving next-generation TiAl 

turbine design. 

Accomplishments 

• Demonstrated joining of TiAl turbine wheel to Ti-6Al-4V shaft. 

• Demonstrated joining of TiAl specimens to steel shaft material. 

• Completed foreign object damage (FOD) testing. 

• Evaluated physical and mechanical properties of candidate TiAl compounds. 

• Assessed the commercial viability of using titanium alloys for turbocharger applications for heavy-duty diesel 

engines. 

Future Direction 

• Contract end September 30, 2005 

191 
FY 2005 Progress Report 

Introduction 

Turbochargers on diesel engines play an integral 

role in enhancing engine performance, controlling 

fuel economy, and meeting emission regulations. 

After a half century of development, the performance of a turbocharger is now greatly limited by the 

properties of turbine materials. Turbines in diesel 

engines, subjected to long-term exposure to hightemperature exhaust gases above 750ο

C, are traditionally made of nickel-based superalloys. These 

superalloys have a relatively low specific strength 

and large inertia due to their high density, approximately 8.0 g/cm3

. Therefore, turbines made of 

superalloys usually utilize a low air efficiency 

design and inherently have a slow transient 

response. A lightweight, high-strength, heat-resistant 

material to replace the current turbine materials is 

critical for further improving turbocharger 

performance. 

Caterpillar has selected candidate TiAl intermetallic compounds to enable next-generation turbocharger designs. The candidate TiAl compounds 

possess a favorable combination of low-density and 

attractive high-temperature capabilities. Turbines 

made of TiAl compounds are anticipated to provide 

faster response (with reduced transient emissions) 

and better fuel economy. One of these turbine 

designs has already been chosen as the prime path 

for select Caterpillar off-highway applications in 

2008. Experience with low-volume product under 

harsh conditions will also encourage high-volume 

applications aligned to Caterpillar ACERTTM engines 

for on-highway trucks in 2010. 

Approach 

The primary focus of the TiAl turbocharger 

project in the year 2005 is to develop technology to 

join TiAl turbine wheel to shaft materials. The difference between a TiAl turbocharger and a conventional one is the material used for the turbine wheel. 

In the 2004 budget year, Caterpillar identified two 

potential suppliers along with the candidate TiAl 

alloys and procured good quality cast TiAl wheels. 

Subsequently, the most important technology in 

manufacturing TiAl turbochargers is now the joining 

of the turbine wheel to its shaft. 

Similar to other intermetallic compounds, limited ductility at ambient temperature of TiAl compounds intrinsically presents a great technical 

challenge for joining TiAl turbine to its shaft. FricHeavy Vehicle Propulsion Materials 

tion welding (FW) was selected as a prime path for 

the joining because of previous success in joining 

TiAl valves to Ti6Al4V shafts in a Department of 

Energy (DOE)-funded project. However, weak and 

cracked joints with strengths of approximately 

50 MPa were produced when the welding was scaled 

up from valve to turbine. A parametric study to 

conduct FW at different conditions could not avoid 

cracking, which typically occurred at the later stage 

of FW process. Therefore, two approaches have 

been pursued for the joining in this year. One is to 

continue the effort in friction welding to solve the 

cracking problem and achieve acceptable joint 

strength. Another is to examine brazing methods. 

Once the turbine wheel is successfully joined to 

the turbine assembly with acceptable strength, a gas 

stand test will be completed for preliminary turbo 

efficiency and reliability. Subsequently, engine and 

machine tests will be conducted for demonstrating 

improved thermal efficiency and transient response. 

Material property characterizations have been conducted in parallel to the turbine rig tests to establish 

a TiAl materials database to be used as a design tool 

for improving next-generation TiAl turbine designs. 

Results 

Successful FW was accomplished to join TiAl 

turbine wheels to Ti6Al4V shafts. The successful 

welding was achieved through extensive finiteelement analysis (FEA) simulations, aimed at 

understanding the correlation among the welding 

parameters, and intensive welding trials, designed 

according to the input acquired from the FEA simulations and microstructure characterization. The 

ultimate tensile strength of the joints between TiAl 

turbines and Ti6Al4V shafts has achieved 570 MPa. 

The FEA simulations results confirmed the 

extremely high thermal stresses inside the TiAl turbine near the welding interface. The FEA simulations were conducted using commercial simulation 

tools and customized user subroutines based on a 

thermal mechanical model. Welding parts geometry 

and welding parameters were input to the simulation. Temperature profile, stress components, and 

strain components (deformation) were output from 

the simulation (an example is shown in Figure 1). 

Simulations using the conditions for the early FW 

trials showed that the hoop stresses near the welding 

line were more than 1 GPa. The simulation result 

agreed with the failure analysis work conducted at 

192 
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 1. A temperature profile at the beginning of the forging stage obtained through simulations 

using a thermal mechanical model. FEA of the joining process is used to understand the 

sensitivity of the thermal stresses to FW control parameters. The simulations were carried 

out for the friction and forging stages, respectively, for the welding process. 

cracked joints. Excessive thermal stresses were 

therefore identified as the primary cause for the joint 

cracking. 

The critical factors determining the thermal 

stresses were then identified through FEA sensitivity 

studies. These simulations were performed using a 

systematic combination of the factors affecting FW 

process, including part geometry, rotational speed, 

friction and forge loads, and friction and forge times. 

Analyses were then conducted on the simulation 

results to provide a general understanding of the 

effect of these factors on the resulting temperature 

profile and thermal stresses. Among these factors, 

welding part geometries were relatively independent 

to other factors and were therefore optimized separately to minimize their effect on thermal stresses 

(Figure 2). Optimizing other critical factors was 

completed through intensive experimental trials. 

These experiments have been designed based on the 

understanding gained from the FEA simulations. 

Modification of the friction welding facilities has 

been implemented to conduct the design of 

experiments. 

Figure 2. FEA used to study the effect of geometry 

constraint on the joint quality. 

With the above efforts, the thermal stress 

cracking issue has been resolved and crack-free 

welding has been achieved. Figure 3 is a 

TiAl 

Ti6Al4V 

193
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 3. TiAl turbine wheels joined to Ti-6Al-4V shafts by FW. 

representative photograph of these joined TiAl tur- several representative NDE scans obtained with 

bines (~1 kg). Nondestructive evaluation (NDE) Eddy current and ultrasonic inspection techniques. 

techniques have been used in examining the weld These scans could not capture any surface/ 

joints. No crack or defect could be identified in the subsurface cracks or radial cracks, but rather demonwelding zone or heat affected zone. Figure 4 shows strate the uniform weld interface. Compared to the 

Figure 4. NDE results at the weld and heat affected zone. No surface or near surface cracks could be detected using 

(a) Eddy current technique, (b) ultrasonic inspections also indicate that the welding interface is uniform, and 

(c) free of radial cracks. 

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Heavy Vehicle Propulsion Materials 

NDE results in the previous year annual report, the 

current NDE results indicated a significant improvement in joint quality. Tensile bars prepared from 

these welded joints demonstrated that the highest 

tensile strength was about 570 MPa. These tensile 

test bars usually failed several millimeters away 

from the welding interface in the TiAl turbine wheel 

side. In a full-scale tensile test on the turbine wheel, 

the fracture surface is about 20 mm away from the 

joint interface as shown in Figure 5. The joint 

strength has reached the strength of the TiAl alloy. 

The FW process has been further scaled up to larger 

size TiAl turbine wheels (~2 kg). Joints have been 

successfully produced free of any visual cracks. 

NDE examinations and tensile tests are yet to be 

completed on these joints. 

Eighteen FW trials were conducted to join 

50-mm-diam. Ti6Al4V bars to 37.5-mm-diam. 4140 

steel bars with a goal to use steel instead of Ti6Al4V 

as a shaft material. Fifteen out of eighteen trials 

were successfully joined without any visual cracks. 

NDE examinations could not detect any cracks or 

defects in these joints. Mechanical tensile test have 

been conducted on tensile bars prepared from the 

FY 2005 Progress Report 

joints. The highest tensile strength is slightly more 

than 300 MPa, but the joints are rather brittle. 

Materials characterizations have been carried 

out for the TiAl materials obtained from the two 

suppliers to compare the properties and establish a 

materials database to be used as a design tool for 

next-generation turbine applications. Mechanical 

tensile tests were performed at room and elevated 

temperatures for the TiAl materials from both suppliers. Figure 6 shows normalized results of ultimate 

tensile strength and elongation at failure for the two 

TiAl materials. Material from Supplier B has lower 

strength at low temperatures than Supplier A. 

However the curves cross over at about 700ο

C, 

above which the TiAl material from Supplier B 

becomes stronger than Supplier A. In addition, the 

ductility of the TiAl material from Supplier B is 

higher than that from Supplier A. Physical properties 

have also been measured, and the results show that 

the density of the TiAl materials from Supplier A is 

about 2% higher than that of Supplier B. Thermophysical properties of the two TiAl materials are 

apparently different. As shown in Figure 7, thermal 

conductivity of the TiAl material from Supplier B is 

Figure 5. Friction-welded TiAl turbine wheel failed about 20 mm away from the FW interface in 

mechanical tensile test. The hub of TiAl turbine wheel is completely pulled out of TiAl 

wheel. 

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FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Tensile Strength 

0.0 

0.2 

0.4 

0.6 

0.8 

1.0 

0 250 500 750 1000 

Temperature (C) 

Ultimate Tensile Strength 

(N

ormalized)

Nickel Superalloy

TiAl Alloy A

TiAl Alloy B 

Ductility 

0.0 

0.2 

0.4 

0.6 

0.8 

1.0 

0 250 500 750 1000 

Temperature (C) 

Elongation at failur

(Nor

malized) 

Nickel Superalloy 

TiAl Alloy A 

TiAl Alloy B 

Figure 6. Comparison of tensile strength and ductility of TiAl compounds and nickel-based alloy. 

Thermal Conductivity 

0.00 

0.25 

0.50 

0.75 

1.00 

0 250 500 750 1000 

Temperature (C) 

Thermal conductivity

(Normalized)

TiAl compound B 

TiAl compound A 

Ni-Superalloy 

Mean CTE 

0.00 

0.20 

0.40 

0.60 

0.80 

1.00 

0 250 500 750 1000 

Temperature (C) 

Mean CTE

(Normalized) 

TiAl compound B 

TiAl Compound A 

Nickel Superalloy 

Figure 7. Comparison of thermal conductivity and coefficient of thermal expansion for TiAl compounds and nickelbased alloy. 

more than 30% higher than that from Supplier A, but 

the coefficient of thermal expansion of the TiAl 

material from Supplier B is about 5% less than that 

of the material from Supplier A. 

FEA simulations have been conducted to evaluate the effect of material property differences of the 

two TiAl alloys in the turbocharger applications. 

This involved the analysis of the strain/stress differences in the turbine at engine operating conditions 

based on the measured properties data. Consequently, no significant effect could be found despite 

the discrepancies in properties of the two TiAl materials. Creep tests are being conducted at Oak Ridge 

National Laboratory (ORNL) as well as an external 

test facility. These tests are scheduled to be 

completed by the end of 2005. Low-cycle fatigue 

(LCF) tests will start in the next quarter and be 

followed with full-scale thermal mechanical fatigue 

(TMF) tests. 

Conclusions 

Successful FW has been achieved in joining 

~1-kg TiAl turbines to Ti6Al4V shaft. The joint 

tensile strength has been increased more than tenfold 

from ~50 MPa to 570 MPa. The FW process is yet 

to be fully scaled up to a ~2-kg TiAl turbine wheel. 

Sufficient confidence has already been gained in FW 

TiAl turbine wheels to Ti6Al4V shafts. Validating 

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Heavy Vehicle Propulsion Materials 

the joint quality needs to be completed before a gas 

stand or engine test on these turbines to demonstrate 

improved transient response and thermal efficiency. 

The option of using steel shaft through FW joining 

is still dubious. The tensile strength of the joints 

produced between 4140 steel and Ti6Al4V shafts is 

about 300 MPa, lower than the tensile strength of 

TiAl material. Microstructure characterization needs 

to be conducted to investigate the brittleness of these 

joints. 

Mechanical tensile tests, physical properties 

characterizations and FEA simulations have been 

carried out for the TiAl materials from the two suppliers. No significant differences in terms of turbine 

FY 2005 Progress Report 

applications have been identified for the two TiAl 

materials despite some discrepancy in their mechanical and thermophysical properties. Characterizations 

for creep properties and thermomechanical fatigue 

properties of the two TiAl materials are yet to be 

completed. 

Presentations and Publications 

Presented for a DOE annual program review on 

April 27, 2005. 

Poster for DOE Heavy Vehicle Propulsion 

Materials Merit Review, September 12–16, 2005. 

197

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

T. High-Speed Machining of Titanium 

Rahul Aphale 

Third Wave Systems, Inc. 

Minneapolis, MN 

(952) 832-5515 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: Third Wave Systems, Inc., Minneapolis, MN 

Objectives 

• To improve the material removal rate in machining of Ti-6Al-4V by developing modeling capabilities for 

various machining processes. 

Approach 

• Develop drilling modeling capability in the finite element modeling–based tool, AdvantEdge. 

Accomplishments 

• Delivered and validated drilling capability for titanium workpieces. 

• Validated predictions from the AdvantEdge modeling software using drilling test data generated by the 

University of Michigan. 

Future Direction 

• Continue to work on reducing computational times and improving accuracy. 

Summary 

Delivery and validation of drilling capability for 

titanium workpieces has been successfully accomplished. Third Wave Systems developed the drilling 

modeling capability in a finite element modeling– 

based tool, AdvantEdge. Drilling test data generated 

by the University of Michigan were used to validate 

predictions from the AdvantEdge modeling software. The agreement between predicted and experimental data was within 15%. Third Wave Systems 

continues to concentrate its efforts on reducing 

computational times and improving accuracy. 

Task Details 

Development of the drilling model was performed for two fluted twist drills via modifications 

to adaptive meshing and contact algorithms for 

improvement in overall performance and robustness. 

The graphical user interface (GUI) development 

was performed with the implementation of 

parametric definition of drill geometry. Twist drills 

can be defined with either drill design parameters or 

drill grind parameters (Figure 1). The fundamental 

information used by the initial mesh generator are 

the grind parameters, since this is the information 

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FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 1. Inputs parameters are shown for drill geometry specification. 

used in fabrication. Consequently, a nonlinear set of 

equations needs to be solved to determine the grind 

parameters if the user defines the drill geometry via 

design parameters. The solution of these equations is 

very fast, taking only a few seconds. Figure 2 shows 

an example of drilling modeling setup, with chips 

curling up along the flutes and heat being conducted 

into chips and deposited on the workpiece. 

Additional modifications to contact and mesh 

adaptation algorithms allow even more chip 

Figure 2. Chip formation for drilling of aerospace 

alloys. Temperature plots are shown for each 

chip formed by the two-fluted twist drill. 

formation development and drill break-through, . 

The drill is removed from the analysis, revealing the 

chip and workpiece with a break-through hole 

(Figure 3). The adaptation and contact algorithms 

are able to model both the cutting action of the drill 

flutes as well as the plowing action of the chisel 

edge. 

In order to model the entire hole drilling 

process, it is necessary to analyze the entrance, full 

engagement, and exit stages of the cut. To model 

full engagement it is necessary to have a pre-drilled 

hole in the workpiece. This hole has the same 

geometry as the drill would make in one complete 

revolution. The drill is then inserted into the predrilled hole and programmed speeds and feeds are 

applied. As the drill rotates it progresses deeper in 

the hole until it finally breaks out, Figure 4. 

Modeling and simulation of indexable drills was 

performed. Automated initial mesh generation was 

performed for a number of ISO standard inserts. An 

example of indexable drilling modeling with 

rectangular inserts is shown in Figure 5. Inserts are 

offset from center axis of drill, so each insert 

experiences a different chip load and cutting speed. 

Experimental thrust force and drilling torque 

data were generated by the University of Michigan 

(Figures 6–9). The drills used were 3.97 mm in 

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Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 3. AdvantEdge model of drill breakout shown with 

drill geometry removed. Burr formation on 

breakout (top) and chip formation with 

temperature contours (bottom). 

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FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 4. Different stages of drilling with pre-drilled holes are shown. (a) The initial predrilled hole in the workpiece; (b) the drill and workpiece combined (translucent 

mesh); (c) drill and workpiece after several rotations of the drill; (d) and drill break 

out on the bottom side of the workpiece. 

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Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 5. Modeling of indexable drilling is shown with finite element mesh and temperature contours 

overlaid. Chip formation is seen from the inserts which are offset from center, top view, left; side 

view, right. 

Figure 6. Experimental thrust force data from University of Michigan, for drill diameter 

of 3.97 mm, 734 rev/min, and 0.050 mm/rev. 

203 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 7. Experimental torque data from University of Michigan, for drill diameter of 

3.97 mm, 734 rev/min, and 0.050 mm/rev. 

Figure 8. Model predicted torque and thrust force. 

204
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 9. Validation data for drilling with 3.97-mm diameter drill, 734 rev/min, 

and 0.05-mm/rev feed. 

diameter and tests were performed at 0.05 mm/rev concentrate its efforts on significant improvements 

feed, and 734 rev/min and 1468 rev/min spindle in the computational times and accuracy of the 

speeds. AdvantEdge simulations were performed by predictions. 

Third Wave Systems. Simulation results were 

analyzed and compared with the experimental test Future Work 

data, F. It was observed that the model prediction for Third Wave Systems continues to work with thrust force was within 15% of experimental thrust University of Michigan and ORNL for completion force at 734 rev/min speed and within 25% at of remaining tasks per milestone. 

1468 rev/min, F. Third Wave Systems continues to 

205 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

5. TEST AND MATERIALS STANDARDS 

A. Rolling Contact Fatigue 

Andrew A. Wereszczak 

Ceramic Science and Technology Group (CerSaT) 

Oak Ridge National Laboratory 

PO Box 2008, MS 6068 

Oak Ridge, TN 37831-6068 

(865) 576-1169; fax: (865) 574-6098; e-mail: wereszczakaa@ornl.gov 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

Field Technical Manager: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Prime DOE Contract Number: DE-AC05-00OR22725 

Objectives 

• Enable greater use of next-generation ceramic and coated-metal roller elements for diesel engines through 

rolling contact fatigue (RCF) and microstructural characterizations. 

• Link surface and subsurface flaws to RCF response to enable microstructural engineering that will promote 

microstructure engineering and improved RCF performance. 

• Link RCF performances measured by different internationally used RCF test techniques. 

Approach 

• Establish three-ball-on-rod (3BOR) RCF test facilities and collaborations with domestic and international 

institutions that employ different RCF test techniques and correlate test results. 

• Vary machining-induced subsurface damage in a Si3N4 and correlate that to measured RCF performance and 

model hypothetical microstructures that will promote improved RCF performance. 

• Develop supporting mechanical test methods that enable the study of surface and subsurface flaws in roller 

elements. 

Accomplishments 

• Established 3BOR RCF test facility at Oak Ridge National Laboratory. 

• Completed summary report on RCF test methods and result interpretations used in Germany, Japan, United 

Kingdom, and the United States. 

• Developed method to evaluate and discriminate the elastic properties of ball bearings in situ using resonance 

ultrasound spectroscopy (RUS). 

• Conceived “c-sphere” flexure strength specimen to enable surface flaw examination. 

207 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Future Direction 

• Conduct International interlaboratory study to compare 3BOR performances of a selected Si3N4. 

• Compare RCF damage on rods and balls of a selected Si3N4 with the same coarse and fine grinding. 

• Refine the geometry and testing of the “c-sphere” flexure strength specimen and the ability to characterize 

surface and subsurface damage in finished ceramic balls and link to RCF performance. 

• Link surface and subsurface flaws to RCF performance. 

Introduction 

Technical ceramics used as materials for rolling 

contact bearing components show several practical 

advantages over traditional bearing steels. The properties of ceramics, specifically low density, high 

stiffness, high hardness, low coefficient of thermal 

expansion, and high-temperature capability are 

properties suited to rolling element materials. Silicon nitride (Si3N4) has been found to have a good 

combination of properties suitable for these applications. Research over the past three decades on 

structure, quality control, and manufacturing techniques has produced ceramic materials that are suitable alternatives to steel for rolling contact bearing 

design. This is especially true for hybrid ball bearings (i.e., precision angular contact ball bearings 

using ceramic rolling elements) that are now offered 

as standard components within the ball bearing 

manufacturing industry. 

An expansion in the use of ceramic bearings 

necessitates guidelines for the choice of bearings or 

rolling elements. In light of these requirements, 

ASTM standard F2094-01, “Standard Specification 

of Silicon Nitride Bearing Balls,” has been issued 

recently. The dimensional specifications of balls are 

based on those of rolling element bearing steel balls. 

The most remarkable thing in the standard is that the 

ranking of Si3N4’s mechanical properties is provided 

depending on various applications of bearings. 

However, nothing is described about rolling contact 

fatigue (RCF) behavior, which is fundamental 

information to assess the lifetime of the material. 

Due to the lack of such information, some concern 

for the reliability of ceramic bearings still remains. 

RCF is the surface damage process due to the 

repeated application of stresses when the surfaces of 

two bodies roll on each other. RCF is encountered 

most often in rolling element bearings and gears. 

The failure process of RCF involves fatigue crack 

initiation and propagation, which is related to the 

characteristics of the surface and subsurface quality 

or flaw population, stress distributions, and lubrication condition. Presently, surface crack defects are 

considered as a main factor that causes RCF failure. It 

is difficult to detect these surface cracks during 

high-volume production processes, and hence it is 

crucially important to understand their influence and 

the fundamental mechanism of the failures they 

cause. 

To address these issues, this project characterizes the RCF performance of Si3N4 compositions 

that are under consideration for use, or that are presently used, in rolling element components (e.g., 

hybrid bearings—see Figure 1, cam roller followers, 

etc.). The study of the effects that subsurface damage and the stress state from Hertzian contact have 

on RCF performance is a primary component to this 

project. Several different test methods are used 

internationally to measure RCF performance, so this 

project also works toward reconciling their 

Figure 1. Schematic of a hybrid bearing system that 

includes silicon nitride balls. Image from 

Cerbec Web site (www.cerbec.com). 

208
Heavy Vehicle Propulsion Materials 

measured differences so that their results can be 

validly pooled. Based on all these findings, 

hypothetical microstructures that will promote 

improved RCF performance will be identified and 

shared with ceramic manufacturers. 

Approach 

There are three primary aspects to the project: 

(1) examine the effects that subsurface damage and 

microstructure have on RCF performance, 

(2) measure RCF performance and reconcile results 

that were generated using different RCF test techniques, and (3) iteratively work with ceramic manufacturers and communicate RCF results and interpretations that those manufacturers can then use to 

develop lower cost Si3N4 compositions or improve 

Si3N4 machinability or improve RCF performance or 

combinations thereof. 

Results 

Effects of Surface and Subsurface Damage 

The “C-Sphere” flexure strength specimen was 

conceived and preliminarily developed to measure 

the strength of bearing-grade Si3N4 balls and to 

relate that property to surface-located strengthlimiting flaws (e.g., machining damage) and to 

ultimately link those flaw populations to RCF performance. A slot is machined into the balls to a set 

depth to produce the “C-sphere” geometry. C-sphere 

specimens are then diametrally compressed to produce a monotonically increasing hoop tensile stress 

at their surface that ultimately caused their fracture 

(see Figure 2). Strength is determined using the 

combination of failure load, C-sphere geometry, and 

FEA. 

The stress field was used to determine C-sphere 

effective areas and effective volumes as a function 

of Weibull modulus. A quarter symmetry model for 

the C-sphere specimen is shown in Figure 3 with the 

mesh distribution that was ultimately used to evaluate its effective volume and area. Solid95 tetrahedral 

elements were used, and a 100-N point load was 

applied to the sphere’s apex. The mesh contained 

52703 elements and 76682 nodes. As can be seen 

from Figure 3, finer mesh density was used in the 

region where high tensile stresses would develop, 

gradually getting coarser toward the area where the 

load is applied and the overhang section. The ball 

from which the specimen was simulated had the 

FY 2005 Progress Report 

dimensions of a 12.7-mm diameter and a slot width 

of 6.35 mm. 

The issue of artificially high stress concentration 

within the narrow zone under the applied point load 

and its effect on the computed effective sizes was 

considered. In the virtual world of finite element 

analysis (FEA), the load is applied as a perfectly 

concentrated point load, which causes stress singularity in that region. This is an idealization of what 

physically takes place where distributed load acting 

over a narrow area is actually applied. In other 

words, as the mesh gets finer, the stress under the 

load will increase indefinitely making the model and 

hence the effective sizes mesh dependent. 

To understand the issue of stress singularity 

under the point load, both nodal and element first 

principal stress distributions are shown in Figures 4 

and 5 for the intermediate and fine mesh models, 

respectively. A nodal plot smoothes the stress distribution by averaging the stresses at that point in all 

the elements having that common node. This 

smoothing is masking the very high stresses taking 

place within the individual elements, which if not 

removed would yield erroneous effective size calculations. This is because when the effective size is 

computed, the stress distribution is normalized with 

respect to the maximum effective stress in the component. If this maximum effective stress were 

incorrect, due for example, to a localized stress concentration, as is the case for this model, then the 

computed value would be incorrect. To unmask the 

Figure 2. Schematic of the “C-Sphere” test specimen 

and its diametral compressive loading. 

Fracture will initiate at the outer surface and 

propagate radially inward. 

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FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 3. Mesh distribution for one-fourth symmetry model of the C-sphere specimen, 

having 52703 elements and 76682 nodes. Three different mesh densities were used 

to evaluate the specimen, and they were coarse (8497 elements), intermediate 

(16090 elements), and fine mesh (shown—52703 elements). 

Figure 4. Nodal first principal stress distribution (left) and element first principal stress distribution (right) for the 

C-sphere specimen, using the intermediate mesh model. 

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Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 5. Nodal first principal stress distribution (left) and element first principal stress distribution (right) for the 

C-sphere specimen, using the fine mesh model. 

high nodal stresses, element stress plots are used that 

display the actual nodal stresses as computed for 

each element. 

The nodal and element first principal stress distributions for the C-sphere specimen using the 

intermediate mesh model are shown in Figure 4. As 

can be seen from the nodal plot, the maximum 

tensile stress is computed to be 132.1 MPa, taking 

place at the side of the sphere as expected. However, 

when the element stress plot is examined, the high 

tensile stress (306-MPa) region shifts to the area 

under the load, which is obviously an artifact of the 

point load effect. For the fine mesh model 

(Figure 5), both nodal and element stress plots show 

the highest tensile stress region to be under the point 

load, with the maximum first principal nodal and 

element stresses increasing to 244 MPa, and 

644 MPa, respectively. Figures 4 and 5 clearly show 

that the stresses under the load are extremely mesh 

dependent and will keep changing as the mesh 

changes; for the rest of the specimen, the stresses 

have long converged even for the coarse mesh 

specimen model. 

To take out the effect of localized stress concentration on the effective size calculations, the artificially high stressed elements under the load were 

removed. The remaining elements, comprising the 

vast majority of the specimen, were then used to 

assess the specimen’s effective sizes. The nodal and 

element stress distribution plots for the C-sphere 

model are shown in Figure 6, with the highly 

stressed elements carved out. One would then 

observe from these plots how the maximum tensile 

stress of 132 MPa shifts back to the side of the 

specimen where it should be. The fact that both the 

nodal and element stress plots yield the same stress 

distribution indicates that the suspect elements were 

successfully removed and that the remaining model 

can now be used to compute the effective sizes. 

Hence, it is the model in Figure 6 that was utilized to 

compute the effective area and effective volume for 

the C-sphere specimen. 

The effective area (Ae) and effective volume 

(Ve) as function of Weibull modulus for the 

C-sphere specimen are illustrated in Figures 7–8, 

respectively. The effective sizes were computed 

using the following equations: 

Ae = ⎜

⎜

⎛

⎝

σ

σ

0 

e ⎟

⎟

⎞

⎠

m 

ln⎜

⎛

⎜

⎝1 −

1 

Pf 

⎟

⎞

⎟

⎠

, 

Ve = ⎜

⎜

⎛ σ

σ

0 ⎟

⎟

⎞

m 

ln⎜

⎜

⎛

1 −

1 

P ⎟

⎟

⎞

, 

⎝ e ⎠ ⎝ f ⎠

where σ0 is the scale parameter, m is the Weibull 

modulus, σe is the maximum effective stress 

(computed by CARES/Life), and Pf is the probability of failure (computed by CARES/Life). The 

effective sizes are independent of the scale parameter because they only vary with geometry, loading, 

and the Weibull modulus. Hence, the scale parameter was assigned a random value in order to carry out 

the reliability and effective size calculations. 

To achieve familiarity with the new C-sphere 

geometry, approximately 30 finished Si3N4 spheres 

211
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 6. Nodal first principal stress distribution (left) and element first principal stress distribution (right) for the 

C-sphere specimen using the fine mesh model with the elements within the high stress concentration zone 

(under the load) carved out. This is the model used to evaluate the effective sizes whose results are shown in 

Figures 7–8. 

Figure 7. Effective area vs Weibull module for the C-sphere specimen. 

Figure 8. Effective volume vs Weibull module for the C-sphere specimen. 

212
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

were machined according to Figure 2 and were 

strength tested. The failure loads and geometrical 

parameters are now being combined to calculate 

strength and for interpretation. Additional C-sphere 

specimens of NBD200, SN101C, and TSN-03NH 

will be machined to generate more confident 

Weibull statistics, and C-sphere tests from balls 

machined according to Table 1 will be performed 

too. Fractography will then be performed to link 

strength-limiting flaw type to strength and ultimately 

RCF performance. 

All four machined ball sets have been received 

(see Table 1). Bournemouth University is using the 

one larger diameter set to reduce to a 12.7-mm 

diameter using an internally developed lapping procedure that will essentially yield a fourth (lapped) 

finish condition. 

The utility of using resonant ultrasound spectroscopy (RUS) to quantify elastic properties of 

Si3N4 balls, to assess consistency of those properties, 

and (hopefully) be able to nondestructively proof 

test them or identify preexisting flaws is under 

exploration. H. Trivedi of UES (an Air Force 

Research Laboratory contractor) forwarded numerous Si3N4 balls that were damaged from their RCF 

testing for attempted RUS characterization and that 

is under way. Part of those supplied balls included 

pristine Toshiba TSN-03NH balls whose elastic 

properties were measured with the RUS and compared with those of NBD 200 and SN101C silicon 

nitrides. Their comparison is shown in Figure 9, and 

it was observed that the TSN-03NH balls showed 

the least amount of variability in those properties of 

the three Si3N4 compositions. The elastic properties 

Figure 9. Elastic modulus and Poisson’s ratio distributions for three bearing grade Si3N4 materials as 

measured with RUS. 

of the three 12.7-mm-diameter sets listed in Table 1 

were also measured with RUS, and those results are 

shown in Figure 10. Several of the balls showed 

variability, but it may be argued that close clustering 

of data suggest that most of the ball diameters and 

material density were uniform and that machining 

condition did not affect that. 

To better understand RCF damage and its link to 

material microstructure, several supplemental characterization test methods are used. For example, 

instrumented static and dynamic indentation testing, 

and instrumented scratch testing of the Ceralloy 

147-31N is used to explore how contact loading 

damage (as a function of loading rate) is affected by 

subsurface damage. The competition of quasi-plastic 

damage and cracking processes in Ceralloy 147-31N 

Table 1. Grinding conditions for silicon nitride spheres 

Step Diameter & finish Wheel Removal 

(in.) 

Removal per 

pass 

(in.) 

12.7 mm/0.500 in. 

Coarse 

1 (roughing) 

2 (induce damage) 

3 (finishing) 

Accepted practice 

100 grit 

600 grit 

0.004 

0.0005 

0.001 

0.001 

0.0001 

13.2 mm/0.520 in. 

Coarse 

1 (roughing) 

2 (induce damage) 

3 (finishing) 

Accepted practice 

100 grit 

600 grit 

0.004 

0.0005 

0.001 

0.001 

0.0001 

12.7 mm/0.500 in. 

Fine 

1 (roughing) 

2 (induce damage) 

3 (finishing) 

Accepted practice 

180 grit 

600 grit 

0.004 

0.0005 

0.001 

0.001 

0.0001 

12.7 mm/0.500 in. 

RCF-conventional 

“Accepted” practice for RCF test bar finishing 

213
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Figure 10. Elastic modulus and Poisson’s ratio 

distributions for Ceralloy 147-31N Si3N4 

balls as measured with RUS. The three sets 

were machined per the conditions in 

Table 1. 

is being interrogated as a function of the depth of the 

subsurface damage and compared to polished 

material as well. NBD200 Si3N4 (NIST SRM for 

Knoop hardness) and NC132 Si3N4 (NIST SRM for 

fracture toughness) are also being tested; specimens 

of those materials are presently undergoing 

indentation and scratch testings whose results will 

serve as a performance reference for comparisons of 

performances of Ceralloy 147-31N and whatever 

additional ceramics that are ultimately tested in this 

program. SN101C and TSN-03NH (both bearing 

grades of Si3N4) are undergoing instrumented 

indentation and scratch testings as well, and their 

performances will be compared with those of the 

other listed silicon nitride compositions. Raman 

spectroscopy, and its ability to measure residual 

stresses, is being used. Preliminary results suggest 

that residual stresses (manifested by changes in 

wave number peak location in Figure 8) can indeed 

be quantified. 

Reconciliation of RCF Test Results 

A summary report was completed that describes 

RCF test methods used in Germany, Japan, United 

Kingdom, and the United States and the interpretation of their results. The report will be published as 

an ORNL Technical Report. 

Japan’s National Institute of Advanced 

Industrial Science and Technology (AIST) will start 

a 3-year standardization project of RCF testing 

funded by Japan’s Ministry of Economy, Trade, and 

Industry (METI), Japan, in Quarter 4, and the present project will collaborate. AIST’s W. Kanematsu 

invited this project participation for the purpose of 

performing three-ball-on-rod (3BOR) RCF testing. 

The UK National Physical Laboratory (NPL) and 

Germany (BAM) will also participate. AIST will 

provide materials for the testing. 

Interaction with Domestic Manufacturers 

Dialog with ceramic bearing manufacturers is 

critical to relevance being maintained in this program. Frequent communication occurs between the 

principal investigator and chief scientists at several 

domestic manufacturers (Cerbec/Saint-Gobain, 

Ceradyne, Enceratec, Kennametal, and Cercom) 

regarding this project’s test matrix, plans, and 

progress. 

Conclusions 

A 3BOR RCF test facility was established in 

ORNL, and the evaluation of ceramic RCF 

performance of Si3N4 roller element materials is now 

well under way. A formal collaboration is under way 

with BU for three ball-on-ball (3BOB) RCF testing 

of ceramics, and this enables the eventual 

reconciliation of RCF performances measured with 

those two techniques. The effects of subsurface 

damage on RCF performance are being scrutinized 

with the C-sphere flexure strength specimen. A 

summary report was completed that describes RCF 

testing and result interpretations used in Germany, 

Japan, the United Kingdom, and the United States. 

A method was developed to evaluate and 

discriminate the elastic properties of ball bearings in 

situ using resonance ultrasound spectroscopy. 

Ultimately, the results from this project will link 

surface and subsurface flaws to RCF performance 

and enable greater use of next-generation ceramic 

and coated-metal roller elements for diesel engines 

through RCF and microstructural characterizations. 

Presentations and Publications 

A. A. Wereszczak, W. Wang, O. M. Jadaan, and 

M. J. Lance, “Strength of a C-Sphere Specimen,” 

abstract submitted to the January 2006 Cocoa Beach 

conference. 

Y. Wang, M. Hadfield, W. Wang, and A. A. 

Wereszczak, “Rolling Contact Fatigue of Ceramics,” 

in internal review, to be published as a DOE/ORNL 

Technical Memorandum. 

214
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

B. Implementing Agreement for a Programme of Research and Development on 

Advanced Materials for Transportation Applications 

M. K. Ferber 

Ceramic Science and Technology Group 

Oak Ridge National Laboratory 

P.O. Box 2008, MS: 6068 

Oak Ridge, TN 37831-6068 

(865) 576-0818; fax: (865) 574-6098; e-Mail: ferbermk@ornl.gov 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Prime Contract No.: DE-AC05-00OR22725 

Objectives 

• Facilitate the integration of new technologies into the transportation sector by implementing research that 

validates the applicability of this technology to improve material properties while maintaining acceptable lifecycle costs. 

• Promote commercialization of new materials technologies by developing standard testing and characterization 

methods in conjunction with national and international standards communities. 

Approach 

• Define and implement research under the International Energy Agency (IEA) Implementing Agreement (IA) 

entitled Implementing Agreement For A Programme Of Research And Development On Advanced Materials 

For Transportation Applications (IA-AMT). 

• Conduct major research themes as annexes under the current IA: 

— Annex II: Co-Operative Program on Ceramics for Advanced Engines and Other Conservation Applications 

— Annex III: Co-operative Program on Contact Reliability of Advanced Engine Materials 

— Annex IV: A Cooperative Program on Integrated Engineered Surface Technology 

Accomplishments 

• Completed final work on project Annex II, “Subtask 13: Burner Rig Round Robin.” The final report can be 

found at http://ia-amt.ornl.gov/index.html. 

• Completed a final report summarizing all activities conducted under Annex II. The final report can be found at 

http://ia-amt.ornl.gov/index.html. 

• Completed preliminary assessment of the laser spallation test for the measurement of coating adherence. The 

technique was capable of generating controlled spallation of both metallic and ceramic coatings. The remaining 

challenge is to quantify stress levels associated with the spallation event. 

215 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Future Direction 

• Develop a plan for the lightweighting materials annex and present it to the Executive Committee for approval. 

• Expand the scope of the implementing agreement by adding an annex on characterization/qualification of 

ceramic coatings for wear, thermal, and environmental protection. 

Introduction 

The current mission of the IA-AMT is to investigate promising new technologies for evaluating 

and ultimately improving the performance of materials for transportation systems. The primary motivation for this activity is the fact that new material 

technologies are required to increase efficiency and 

reduce harmful emissions in these systems. Examples of these technologies include (1) lightweighting 

to improve fuel efficiency; (2) surface engineering to 

improve the resistance to wear and contact damage; 

(3) development of durable coating systems for 

thermal, wear, and environmental management; and 

(4) development of revolutionary materials (structural ceramics and ceramic matrix composites) for 

operation at much higher temperatures and pressures. As discussed in this report, the research 

activities within the IA-AMT focus specifically on 

(1) identifying promising new technologies for 

improving materials performance and (2) developing 

specialized characterization techniques for validating the applicability of this technology to improve 

material properties while maintaining acceptable 

life-cycle costs. 

Approach 

In performance improvement, the current 

emphasis is on integrated engineered surface technology (IEST) and lightweighting of materials. IEST 

encompasses the synthesis, processing, characterization, and application of technologies that enhance 

the functionality of surfaces in contact with the environment or with the surfaces of other solids. Activities on lightweighting of materials focus on aluminum, high-strength steels, magnesium, metal and 

polymer composites, titanium, intermetallic alloys 

and other advanced materials. Current topics under 

consideration include (1) data on production and 

resource availability; (2) life cycle data on environmental impacts associated with production, processing and use of lightweight materials; 

(3) recycling information, including regulatory 

frameworks; (4) data on crashworthiness, design, 

and testing methodologies, (5) data on base material 

costs; (6) data on energy impacts of lightweight 

materials; and (7) shared information on research 

programs on lightweight materials. 

In terms of performance evaluation, the primary 

focus is on techniques for (1) assessing environmental degradation of structural (non-oxide) ceramics; (2) evaluating time-dependent degradation of 

mechanical performance of structural ceramics; 

(3) quantifying key properties of coatings for wear, 

thermal, and environmental protection of current 

transportation materials; and (4) developing techniques for the measurement of key properties 

(topography, chemistry, subsurface damage) of 

engineered surfaces. Items 1 and 2 are motivated by 

the need to address key barriers to the use of this 

important class of material. For example, given the 

recent concern over environmental degradation of 

non-oxide ceramics in combustion environments, 

cost-effective techniques are required to simulate 

these effects as well as to assess the effectiveness of 

environmental barrier coatings. The IA-AMT is 

currently evaluating a variety of techniques ranging 

from complex high-pressure burner rig tests to a 

simple cost-effective steam injection system. 

Ceramic coatings hold considerable promise for 

(1) improving wear resistance, (2) providing thermal 

protection, and (3) reducing environmental degradation of critical metallic components used in internal 

combustion engines. Unfortunately techniques for 

assessing key properties particularly with respect the 

interface are unproven. Item 3 addresses this limitation. In a similar fashion, as surface modification 

technologies mature, proven characterization techniques will also be required to validate their performance (Item 4). 

Results 

A major result this year was the completion of 

Annex II. The Annex II subtasks primarily focused 

on addressing key technology barriers to the use of 

structural ceramics in both automotive and landbased gas turbines. The motivating factors for 

216
Heavy Vehicle Propulsion Materials 

ceramic gas turbine development were the (1) desire 

to dramatically improve efficiency through higher 

turbine inlet temperatures, (2) potential for reducing 

the dependency on strategic metals, and (3) lowering 

of harmful emissions. Structural ceramics were considered for use as combustors, transition sections, hot 

gas ducts, turbine vanes and shrouds, turbine blades, 

and integrally bladed rotors (blisks). Highperformance SiN and SiC ceramics were considered 

as leading candidates for use in these applications.1–10 

These materials offered the following advantages 

over their metallic counterparts: (1) low density, 

(2) refractoriness, (3) adequate short-term strength, 

and (4) good oxidation resistance at ambient 

pressures. The technical barriers to the 

FY 2005 Progress Report 

implementation of ceramic technology in gas turbines 

included maturation of powder processing methodologies for component fabrication; establishment of 

proven methodologies to assess key mechanical 

properties (required for the development of design 

database); and optimization of the machining process 

used in the final stages of component manufacturing 

(related to cost optimization). Given these barriers, 

Annex II focused on (1) characterization of powder 

properties, (2) quantification of the green-state 

characteristics, (3) evaluation of mechanical 

performance (strength and thermal shock), 

(4) machining, and (5) evaluation of environmental 

degradation of non-oxide ceramics (see Table 1). An 

extensive set of results was generated from the work 

Table 1. Summary of all subtasks implemented under Annex II 

Subtask Title (final report) Started Completed Countries 

1 Information Exchange 1985 2004 Germany, Japan, 

Sweden, United States 

2 Characterization of Ceramic Powders: Data and 

Analyses 

1985 Final Report—March 1990 Germany, Sweden, 

United States 

3 Characterization of Sintered Silicon Nitride and 

Silicon Carbide Structural Ceramics 

1985 Final Report—October 

1989

 Germany, Sweden, 

United States 

4 (1) Fractography Analysis of Silicon Nitride 

and Silicon Carbide Structural Ceramics, 

(2) Statistical Analysis of Flexure Strength 

Data, and (3) Analysis of Error Sources in 

Four Point Flexure Strength Measurements of 

Structural Ceramics 

1985 Final Reports—June 1989 

(1 and 2) December 

1989 (3) 

 Germany, Sweden, 

United States 

5 Study of the Flexure and Tensile Strength of a 

United States Silicon Nitride 

1990 Final Report—September 

1993

 Germany, Japan, 

Sweden, United States 

6 Development and Testing of Procedures for 

Characterization of Ceramic Powders 

1990 Final Report—September 

1993

 Germany, Japan, 

Sweden, United States 

7 Effect of Machining Conditions on the Strength 

of Silicon Nitride 

1993 Final Report—March 1998 Germany, Japan, 

Sweden, United States 

8 Development and Testing of Primary and 

Secondary Properties of Ceramic Powders 

1993 Final Report—December 

1996

 Belgium, Germany, 

Japan, Sweden, United 

States 

9 Thermal Shock Testing of Advanced Ceramics 1996 Final Report—March 2000 Belgium, Germany, 

Japan, Sweden, United 

States 

10 Assessment of Powder Characterization Methods 

for Advanced Ceramics 

1996 Final Report—February 

2000

 Belgium, Germany, 

Japan, Sweden, United 

States 

11 Techniques for the Measurement of Thermal and 

Mechanical Fatigue 

2000 Final Report—March 2004 Germany, Japan, 

Sweden, United States 

12 Characterization of Ceramic Powders and Green 

Bodies 

2000 Final Report—September 

2001

 Belgium, Germany, 

Japan, Sweden, United 

States 

13 Burner Rig Round Robin 2003 Final Report—March 2005 Germany, Japan, United 

States 

217 
FY 2005 Progress Report 

conducted in Annex II. The major reports summarizing most of the subtasks can be downloaded from 

http://ia-amt.ornl.gov/index.html. 

The results from Annex II were subsequently 

used in the establishment of standards [via American 

Society for Testing and Materials (ASTM), Japan 

Industrial Standards (JIS), Committee for European 

Normalization (CEN), International Organization for 

Standardization (ISO), and National Institute of 

Standards and Technology (NIST) guidelines (see 

the publications list). Table 2 summarizes some the 

existing standards that benefited from this work. In 

the case of room-temperature flexural strength, the 

standard, ASTM C1161, was revised to reflect 

lessons learned about fixturing and test specimen 

configurations. ISO 14704 evolved from several 

standards (ASTM C1161, CEN EN843-1 and JIS 

R1601) and lessons learned from the IA-AMT work. 

The high-temperature flexural strength standard, 

ASTM C1211, evolved about the time of the IEA 

round robin and included lessons learned from 

C1161 and the IEA work. ISO DIS 17565 (not yet a 

standard) has evolved from several standards 

(ASTM C1211, CEN prEN820-1, JIS R1604) and 

lessons learned from the IA-AMT subtask. The 

room-temperature tensile strength standard, ASTM 

C1273, was developed as the IA-AMT results were 

being reviewed and analyzed and included many 

Table 2. Standards that have benefited from the IA subtasks 

Heavy Vehicle Propulsion Materials 

lessons learned. ISO 15490 evolved from two 

standards (ASTM C1273 and JIS1606), and the ISO 

round robin results were instrumental in establishing 

test specimen configurations, gripping arrangements, 

allowable bending, and test rates. The thermal shock 

standard, ASTM C1525, was developed after the IAAMT subtask on thermal shock was completed. 

Although this standard follows a more conventional 

approach to thermal shock by using water and 

standard MOR bars, insights garnered from the IAAMT work are used in providing guidance to users 

in notes and discussions. 

Thirteen subtasks were ultimately conducted 

under Annex II. Technical experts in the field 

defined the work scope for each subtask, with final 

approval given by the executive committee members 

during the annual meetings. Once the subtask was 

approved, each country was responsible for selecting 

technical task leaders; the executive committee 

appointed an overall subtask coordinator. Working 

group meetings were used to monitor the progress of 

the effort and provide a venue for technical data 

exchange by the task leaders. Once a subtask was 

completed, the coordinator was responsible for 

collecting and analyzing the data and preparing the 

final reports. Many of the reports are available on 

the IA-AMT web site: http://ia-amt.ornl.gov/ 

index.html. 

Property JIS ASTM CEN ISO 

Flexural strength: RT R1601-95 C1161-96 EN 843-1:95 14704:2000 

Flexural strength: HT R1604-95 C1211-98 prEN 820-1 DIS 17565 

Statistical analysis R1625-96 C1239-95 ENV 843-5:97 CD 20501 

Fractography C1322-96 prENV843-6 

Tensile strength R1606-95 C1273-95 15490:2000 

Sample preparation for the determination 

of particle size distribution of ceramic 

powders 

R1619:95 C1282-94 EN 725-5:96 14703:2000 

Surface area R1626-96 C1274-95 EN 725-6:96 DIS 18757 

Particle size distribution of powder by 

laser diffraction method 

R1629:97 TC206 NP02 

Thermal shock R1615-93 C1100 prEN 820-3 

Flowability R1639-4:99 TC206 WI93 

Size distribution of granules R1639-1:99 

Binder content of granules 

Drying loss of granules R1639-3:99 

Bending fatigue: RT R1621:95 TC206 PWI 07 

Bending fatigue: HT R16xx:01 

RT = room temperature; HT = high temperature. 

218 
Heavy Vehicle Propulsion Materials 

As shown in Table 1, the 13 subtasks included 

an ongoing information exchange (subtask 1), which 

occurred in conjunction with the annual executive 

committee meeting, and specific research activities 

(subtasks 2–13), which consisted of task sharing 

collaborations between the participating national 

organizations. Most of these subtasks included 

extensive industry participation. Because the more 

recent subtasks involved fairly specialized hardware 

(thermal shock facilities for the centrally heated disk 

specimen—subtask 9, burner rig facilities— 

subtask 13), industrial participation was limited. 

Throughout Annex II, the industrial participants 

provided (1) expertise in the development of 

experimental research plans, (2) materials and 

facilities required for the implementation of those 

plans, and (3) personnel to actively participate in the 

research activities. In this approach, each country 

benefited from the experience and work of the other 

countries by gaining access to a larger quantity of 

data required to substantiate the standard methods 

FY 2005 Progress Report 

under development. This approach also minimized 

the effort required to transition standards from the 

national to international level. As a result the 

structure of the IA led to considerable cost savings. 

The motivation for subtask 13, the last task 

implemented under Annex II, was the need to 

develop a reliable method to assess the 

environmental attack on non-oxide ceramics arising 

from the global removal of the normally protective 

silica scale via reaction with water vapor present in 

the combustion gases. In stagnant environments, 

oxidation of both SiN and SiC is increased by (1) the 

replacement of oxygen by water vapor and (2) an 

increase in the pressure of the oxidant. In addition, 

the high velocities and presence of water in the 

environment can lead to the volatilization of the 

normally protective silica layer. The resulting 

recession of non-oxide ceramics such as SiC and 

SiN ultimately leads to excessive material loss, 

which increases the stress levels arising from both 

thermal and mechanical loading (Figure 1). 

Figure 1. Photograph of uncoated AS800 prior to field testing (a). The 

dotted lines show the sections that were machined for subsequent 

polishing and evaluation by scanning electron microscopy. 

Typical sections for an as-received vane and a vane exposed for 

815 h are compared in (b) trail edge and (c) leading edge. The 

dotted lines are predicted recession profiles. 

219 
FY 2005 Progress Report 

Because of this recession problem, considerable 

emphasis has been placed on the development of 

environmental barrier coatings (EBCs) to act as 

diffusion barriers to water vapor. The effectiveness 

of EBC systems is defined by their (1) ability to 

isolate the non-oxide ceramic from the water vapor 

and (2) phase stability in the gas turbine 

environment. Both aspects can be evaluated from 

burner rig studies. Unfortunately, the results 

obtained from a given burner rig system depend 

upon a number of factors, including specimen 

geometries, gas flow/specimen configuration, 

temperature uniformity, and burner stability. The 

goal of this subtask was to implement a round robin 

in which a selected SiN material was evaluated in 

burner rig facilities located in Japan and the United 

States. 

In subtask 13, three techniques were used to 

evaluate the recession behavior of the SN282 

material. The first method consisted of exposing 

rectangular specimens for 400, 600, and 1000 hours 

in a high-pressure, high-velocity burner rig located 

at General Electric Global Research Center. This rig, 

which was used extensively used in the Continuous 

Fiber Ceramic Composite program,

11 runs largely 

unattended with extensive computer control, data 

acquisition, and safety monitoring. The second 

technique considered involved the development of a 

Heavy Vehicle Propulsion Materials 

system for injecting steam onto a small region of the 

SiN specimen using an alumina or SiC tube. As 

shown in Figure 2, this method was quite effective at 

inducing localized recession in the injection area. 

The final technique involved exposing prismatic 

bars in an atmospheric burner rig. A detailed 

comparison of the resulting recession data12 

indicated that all three techniques gave comparable 

values of the recession rates (Figure 3). 

Conclusions 

The IA-AMT has made significant progress in 

expanding its scope, as evidenced by the addition of 

two new annexes, IV and V. Both annexes are 

expected to generate new members. Given this 

progress, a formal request has made to the IEA to 

extend the IA-AMT for a period of 3–5 years. 

Presentations and Publications 

M. K. Ferber, Burner Rig Round Robin— 

Subtask 13: Final Report, March 2005. Available at 

http://ia-amt.ornl.gov/index.html. 

M. K. Ferber, Annex II: Co-Operative 

Program on Ceramics for Advanced Engines 

and Other Conservation Applications— 

Final Report, September 2005, available at 

http://ia-amt.ornl.gov/index.html. 

Figure 2. Sintered alpha SiC tensile specimen after 500 h exposure to a 

saturated water vapor environment at 1200°C. The white dotted line 

on the right shows the position of the original surface prior to testing. 

220
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 3. Comparison of normalized recession rate versus 1/T data obtained in 

present study with similar data reported in the literature. 

References 

1. D. Anson, K. S. Ramesh, and M. DeCorso, 

Application of Ceramics To Industrial Gas Turbines, 

DOE/CE/40878-1&2, Volume 1—Technology 

Background and Volume 2—Economics, Market 

Analysis and R&D Plan, Battelle Columbus, Ohio, 

March 1991. 

2. H. E. Helms, R. A. Johnson, and L. E. 

Groseclose, “AGT 100-Advanced Gas Turbine 

Technology Development Project,” pp. 137–155 in 

Proceedings of the Twenty-Third Automotive 

Technology Development Contractor’ Coordination 

Meeting, P-165, Society of Automotive Engineers, 

Warrendale, Pennsylvania, March 1986. 

3. M. van Roode, “Ceramic Retrofit Program,” 

pp. 77–93 in Proceedings of the Joint Contractors 

Meeting: FE/EE Advanced Turbine Systems 

Conference—FE Fuel Cells and Coal-Fired Heat 

Engines Conference, DOE/METC-93/6132, August 

1993. 

4. D. Carruthers and L. Lindberg, “Critical 

Issues for Ceramics for Gas Turbines,” pp. 1258–72 

in Third International Symposium on Ceramic 

Components and Materials for Engines, V. J. 

Tennery, ed., American Ceramic Society, 

Westerville, Ohio, 1988. 

5. P. J. Haley, R. L. Holtman, L. E. Groseclose, 

S. J. Hilpisch, and A. H. Bell III, “Advanced 

Turbine Technology Applications Project 

(ATTAP),” pp. 19–29 in Proceedings of the TwentyNinth Automotive Technology Development 

Contractors’ Coordination Meeting, Society of 

Automotive Engineers, Warrendale, Pennsylvania, 

June 1992. 

6. Ceramic Technology for Advanced Heat 

Engines Program Plan 1983–1993, ORNL/M-1248, 

Oak Ridge National Laboratory, July 1990. 

7. A. Nagahiro, Present Status of the National 

Project of R and D of Fine Ceramics, FC Annual 

Report, 1986, pp. 61–66. 

8. J. I. Mueller, “Handicapping the World’s 

Derby for Advanced Ceramics,” Am. Ceram. Soc. 

Bull., 61(5), 588–591 (1982). 

9. S. O. Kronogård and L. Malmrup, “Ceramic 

Engine Research and Development in Sweden,” in 

pp. 51–80 in Ceramics for High-Performance 

Applications III Reliability, E. M. Lenoe, R. N. 

Katz, and J. J. Burke, eds., Plenum Press, New 

York, 1983. 

10. W. Bunk, E. Gugel, and P. Walzer, 

“Overview of the German Ceramic Gas Turbine 

Program,” pp. 29–50 in Ceramics for HighPerformance Applications III Reliability, E. M. 

Lenoe, R. N. Katz, and J. J. Burke, eds., Plenum 

Press, New York, 1983. 

11. G. S. Corman, A. J. Dean, S. Brabetz, M. K. 

Brun, and K. L. Luthra, “Rig and Engine Testing of 

Melt Infiltrated Ceramic Composites for Combustor 

and Shroud Applications,” in Proceedings of ASME 

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FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Turbo Expo 2000, May 8–11, 2000, Munich, 

Germany, Journal of Engineering for Gas Turbines 

and Power, 124(3) (July 2002), pp. 459–464. 

12. M. K. Ferber, Burner Rig Round Robin— 

Subtask 13: Final Report, March 2005. 

222
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

C. Mechanical Property Test Development 

George D. Quinn 

National Institute for Standards and Technology 

Ceramics Division, Stop 8529 

Gaithersburg, MD 20899 

(201) 975-5765; fax: (301) 975-5334; e-mail: geoq@nist.gov 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: National Institute for Standards and Technology, Gaithersburg, Maryland 

Objective 

• Develop mechanical test method standards in support of the Propulsion Systems Materials Program. New 

methods and sound engineering data will facilitate adoption of new materials in heavy vehicle propulsion 

systems. 

Approach 

• Conduct pre-standardization research on test methods that need refinements, or develop new methods. 

• Develop draft recommendations for practices or procedures based upon the needs identified by the research. 

• Conduct round robins as necessary. 

• Standardize procedures in the American Society for Testing and Materials (ASTM) and/or the International 

Organization for Standards (ISO). 

Accomplishments 

• Conducted new experiments on zirconia split cylinder flexural strength test specimens using the latest 

generation material and compared results with full rod flexural strength test results. Conducted extensive 

fractographic analysis to find the strength-limiting flaws in this commercially important material. 

• Neared completion of the comprehensive Guide to Practical Fractography for Ceramics and Glasses. 

• Completed a set of papers and made presentations. 

• Revised, refined, and improved several current ASTM and ISO standards. 

Future Direction 

• Finish the split cylinder flexural strength test work and write a paper on the method. Show how a simple 

modification of ordinary bend fixtures for rectangular bars can do double duty and also work with split 

cylinders. 

• Finish the comprehensive Guide to Fractography. 

• Complete prestandardization work on flexure testing of rods and write the first draft of an ASTM standard. 

223 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

• Resume pre-standardization evaluation of diametral compression strength testing for small round-shaped 

specimens 

Introduction 

This project creates new test methods that will 

facilitate the use of advanced materials in heavyweight propulsion systems. Much of the work is for 

brittle materials such as ceramics, for which classical metal mechanical test methods are not suitable. 

For example, tension strength test specimens of 

many ceramic materials made in short, stubby cylindrical shapes (e.g., diesel engine fuel injector pins, 

timing plungers, valves) and classical dog bone 

shapes are impractical (Figure 1). Our goal is to 

adapt or refine existing methods or invent new ones 

that will allow engineers and researchers to measure 

mechanical properties with good accuracy and precision. They then will be able to construct their own 

databases with greater confidence. Formal test 

method standards are our primary objective. Sound 

test methods and high-quality databases will 

enhance the credibility of new materials and encourage engineers to use them in advanced heat engines. 

This purpose of this project is to develop the test 

methods. 

Figure 1. Round ceramic engine components. 

Approach 

Over the course of this program, we have formulated or contributed to the development of 17 

ASTM and ISO standards. We currently are working 

aggressively on three more: 

• Flexural strength of cylindrical rods 

• Flexural strength of split cylinders 

• Fractographic measurement of fracture mirror 

sizes 

We also plan to tackle the diametral compression strength test method when those three standards 

are completed in FY 2006. 

The standards adopted so far include methods 

for ceramics such as flexural strength, elastic 

modulus, Weibull statistical analysis, fractographic 

analysis, and fracture toughness. As a direct result of 

this work, there has been a dramatic improvement in 

test data quality and reliability in the structural 

ceramics field. We have also contributed to a 

ceramic material specification for silicon nitride ball 

bearings. The specification uses several of the test 

method standards developed in this program. The 

test method standards have been sufficiently generic 

that they have even been used in to the biomedical 

field for materials specifications for ceramic surgical 

implants. 

Considerable energy was put into finishing a 

Guide to Practice for Fractography of Ceramics and 

Glasses this year. This is a large document that has 

been written with a strong practical slant. It will 

complement the ASTM standard practice for fractography, and it is intended to help engineers and 

scientists find flaws and fracture origins and help 

them do their jobs better. The goal is to make fractographic analysis less an art and more an engineering 

practice. 

Results 

Flexural Strength of Split Cylinders 

Splitting a rod and testing the halves is one way 

of evaluating the strength of short, stubby cylindrical 

parts that otherwise might be difficult to test correctly. The break forces for solid full rods are very 

high and fixtures may be damaged, or high contact 

stresses could cause Hertzian cracking in the specimen. Splitting the rods makes them much easier to 

test, as shown in Figure 2. Our goal is to refine this 

simple procedure and make it available to engineers 

and scientists as an optional test configuration. 

We did a thorough fractographic analysis of the 

split Coors zirconia rods that we had previously 

broken. Our goal was to identify the strengthlimiting flaws and ascertain whether the new fixture 

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Heavy Vehicle Propulsion Materials 

Figure 2. Split zirconia fuel injector pin in the NIST 

bend fixture for split rods. 

scheme produced legitimate test results and breakage patterns. Fractography was very difficult in this 

material because of its coarse microstructure, but we 

found that flaw pockets (local clusters of tiny 

defects), as shown in Figure 3, and grain boundary 

faults were the origins. A paper on this work is in 

preparation. 

We did additional split cylinder testing to fill in 

a gap in our results. We previously had tested split 

rods made from the Coors zirconia, but we did not 

have any solid rod strengths with which to directly 

compare our results. To get a better sense of whether 

the split cylinders were giving data comparable to or 

different from those for solid rods, we obtained a 

new set of solid zirconia pins from Cummins. These 

were from a set of Carpenter zirconia rods that had 

been set aside. Some of these were split into two 

halves and each half was tested, as shown in 

Figure 2. The 4-point testing was with effective 

spans of 20 × 40 mm. The 3-point testing was with a 

FY 2005 Progress Report 

40-mm span. The strength outcomes are shown in 

Figure 4. Three-point strength numbers for the same 

batch of material that were obtained by Cummins 

are also shown for comparison. The good news is 

that the 3-point strengths of the split rods are nearly 

identical to those of the full rods. The 4-point 

strengths are less, and the shift was in perfect 

accordance with Weibull size effects. 

The fracture origins these specimens were also 

thoroughly analyzed. An interesting outcome was 

that most specimens failed from a flaw type not 

detected in the Coors batch. The Carpenter 

specimens had lovely fracture surfaces with distinct 

mirrors and broke from transformed grains right on 

the outer ground surfaces, as shown in Figure 5. 

Evidently improved processing in the Carpenter 

material has eliminated many of the grain boundary 

and flaw pocket faults that controlled properties in 

the earlier-generation Coors material. 

A remarkable finding was that although the split 

cylinders had the usual strength variability associated with ceramics, the strengths of the two halves 

of a particular rod were remarkably consistent. That 

suggests that the strength-limiting flaws are very 

consistent within any given rod. 

Our work is converging on a very simple solution 

for this test method. An ordinary bend fixture 

designed for rectangular specimens may be easily 

adapted by using simple cradles with alignment 

parts that enable it to instantly convert to test split 

cylinders. Engineering drawings of the new singlepiece cradles are being prepared. 

Flexural Strength of Solid Cylindrical Rods 

Work resumed in conjunction with the split 

cylinder testing described in the next paragraph. As 

Figure 3. Fracture surface of a Coors zirconia slit cylinder specimen showing the very rough surface. The origin was a 

region of locally greater porosity coupled with grain boundary faults. The schematic suggests how these 

pockets could control strength. 

225 
FY 2005 Progress Report 

Probability (Pf) % 99.9 

99 

95 

90 

80 

63.2 

50 

20 

10 

5 

2 

1 

0.5 

3 point flexure - NIST 

split cylinder 

4 point flexure - NIST 

split cylinder 

m = 19.5 

σθ = 756 MPa 

(MLE) 

n = 20 

m = 15.6 

σθ = 666.5 MPa 

(MLE) 

n = 20 

Cummins solid rod 

3 point data in black 

n = 10 

400 450 500 600 700 800 900 1000 1100 

Stress (MPa) 

Figure 4. Flexural strengths of split zirconia rods tested 

measured at NIST. Cummins 3-point data for 

solid rods is shown in black. Weibull 

parameter estimates are noted. The 3- and 

4-point distributions are simply shifted in 

accordance with the customary Weibull size 

effect. 

Figure 5. Fracture surface and origin of a Carpenter 

zirconia split rod specimen. The 

microstructure is less rough than in Figure 3, 

and there is a fracture mirror. The second 

photo shows a typical origin: a transformed 

grain on the ground surface. 

Heavy Vehicle Propulsion Materials 

part of our fixture design work, we tested one oversized steel dowel rod in our flexure fixture that has 

special cradles to hold round specimens (Figure 6). 

We did this to ascertain whether minor yielding in 

the steel cradles (which had been detected after 

doing some ceramic tests) was due to fixture overload or an artifact of a transient load pulse when the 

ceramic specimens fractured. The yielding was only 

on a tiny portion of an edge of the cradles, but we 

decided to investigate the source of the deformation 

to possibly refine our new fixture design. The fixture 

and hardened steel dowel were loaded to static loads 

well above those used in the ceramic strength tests 

and then unloaded for inspection. No deformation 

whatever was detected in the steel fixture parts or 

the hardened steel dowel rod. This tends to confirm 

the notion that a shock–pulse in the fixture when the 

ceramic test pieces broke was the cause of the 

observed small deformations. 

We have resumed our error analyses work for 

flexure testing of round specimens. This will be the 

cornerstone of the standard. Once we have analytical 

estimates of errors for different specimen and fixture 

geometries, then we will be able to intelligently 

select an optimum configuration and set ranges of 

acceptable configurations and close practical engineering tolerances. The error analysis builds upon 

similar work for rectangular specimens done by 

Francis Baratta and Quinn in the early 1980s. 

Figure 6. Steel dummy specimen in the same 4-point 

bend fixtures that were used to test split 

cylinder specimens. Special cradles hold the 

round specimens in the fixtures that are 

normally used to test rectangular bend 

specimens. Deformations on the cradle edges 

were observed after ceramic strength tests. 

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Heavy Vehicle Propulsion Materials 

Nuances pertaining to rod shapes must be considered, however, before the analysis can be directly 

applied to the rod. Most errors are similar or identical, but contact point tangency shift with the new 

cradles, errors from severe contact loadings from 

crossed rollers, and wedging stresses are different 

for rods and rectangular bars. So, for example, in 

October 2005 we completed a new, simpler analysis 

for contact point tangency shift for beams in flexure. 

This error occurs when the specimen deflects and its 

point of contact with the loading rollers shifts. This 

in turn affects the moment that is applied to the 

specimen and alters the stress in the beam. The error 

is minimal (< 1%) in the standardized rectangular 

bend bars so long as loading roller sizes are kept 

small. A similar effect is expected for rods, except 

that if cradles are used, the error could get as large 

as 5% or more. The original Baratta analysis of 1982 

was cumbersome and required an iterative analysis 

to solve for a particular geometry. The newer, 

simpler analysis we derived gives virtually identical 

error estimates and is far simpler to use. Our goal is 

to finish the complete error analysis by December 

2005, so that we can use the analysis to set reasonable testing limits for an ASTM draft standard. 

Guide to Fractography 

Work continued on a NIST user-friendly Guide 

to Best Practice for fractographic analysis. This 

pamphlet-style document will be about 400 pages 

long with hundreds of figures illustrating key fractographic markings in ceramics and brittle materials. It 

will complement ASTM standard C 1322, Standard 

Practice for Fractography and Characterization of 

Fracture Origins in Advanced Ceramics. 

Small revisions to five flaw definitions in 

ASTM C 1322, Standard Practice for Fractography, 

were balloted. Two negatives were obtained. These 

negatives were discussed at the Reno, Nevada, 

meeting of ASTM Committee C-28, Advanced 

Ceramics, on May 16, 2005, and were found to be 

unpersuasive. The final revised version 2005b of 

C1322 was reviewed and has been sent to press. 

This concludes all our planned revisions to this 

standard for the time being. 

One important byproduct of the Guide to Best 

Practice is a new set of guidelines for how to measure fracture mirrors. The guidelines are one step 

away from a formal standard. They will be shown to 

FY 2005 Progress Report 

ASTM Committee C-28 in January 2006 for review 

and consideration. Fracture mirrors are relatively 

smooth surrounding a fracture origin in a brittle 

material, such as is shown in Figure 7. Mirrors are 

formed when the crack radiates outward from the 

origin and initially generates a smooth surface, but 

then creates a rough surface once the crack reaches 

terminal velocity. A remarkable characteristic of 

fracture mirrors is that their size is inversely related 

to the square of stress in the part at the instant of 

fracture. Thus mirrors may be used to estimate the 

breakage stress in failed components, even if the 

engineer has no idea how the parts were loaded or 

stressed. Mirrors are discussed in the ASTM fractography standard C 1322, which focused on how to 

characterize fracture origins, but there was no guidance on how to measure them. There are wide variations in how these fractures are currently measured. 

We are ready to rectify this now with a new standalone standard if the ASTM Committee concurs. 

Figure 7. Fracture mirror in a ceramic. The fracture 

stress can be estimated fairly accurately from 

the size of the relatively smooth circular 

fracture mirror. 

Other 

A review paper, “Design and Reliability of 

Ceramics: Do Modelers, Designers, and 

Fractographers See the Same World?” was presented 

at the American Ceramic Society conference in 

January. One of the case studies cited in the paper 

was the excellent work done by the Ford Motor 

Company in the early 1980s on a model silicon 

227 
FY 2005 Progress Report 

nitride gas turbine rotor that was spun to failure in a 

hot test rig (Figure 8). 

Heavy Vehicle Propulsion Materials 

and furnished a summary report to the current award 

subcommittee leader. Quinn also reviewed the 

history of the development of two ASTM powder 

characterization standards, C 1282 and C 1274, and 

two nondestructive evaluation standards, C 1212 and 

C 1336, and sent a summary report to Steve Gonczy, 

the C-28 chairman. Gonczy wanted to know the 

backgrounds so that plans could be made to upgrade, 

Figure 8. The Ford model turbine rotor was a featured 

example in a review paper on design and 

reliability. 

Although a high-quality ASTM standard for 

fracture toughness evaluation has been on the books 

for 8 years, and although NIST has a Standard 

Reference Material (SRM) 2100 to support it, people 

are still using the defective Vickers indentation 

crack length method to evaluate fracture toughness. 

Therefore, we have begun some Vickers indentation 

crack length measurements to estimate the “indentation fracture resistance” of the NIST SRM 2100 in 

order to demonstrate the shortcomings of the method 

and the errors that can result. Our goal is to get these 

findings into the literature as soon as possible before 

the adoption of indentation crack length methodologies gets codified too far. We advocate the use of 

genuine fracture mechanics tests for evaluating 

fracture toughness. The Vickers indentation crack 

length method evaluates a particular measure of 

fracture resistance, which crudely approximates 

fracture toughness. The indentation crack length 

method may have some utility, but data from it 

ought to be deemed “indentation fracture resistance” 

and not “fracture toughness.” 

We helped to update and refine some existing 

ASTM and ISO standards. We are continuing to 

work with W. Mandler of Enceratec/Cummins and 

B. Mikijelj of Ceradyne as the silicon nitride ball 

bearing material specification F2094 evolves. Some 

corrections to data specifications were reviewed and 

balloted this year. The ASTM specification is also 

being converted into an ISO international standard. 

There is some controversy about some of the procedures, and we are working with the stakeholders to 

ensure that a technically sound specification is prepared. Another example is the new static fatigue test 

method draft standard that we helped review. Quinn 

reviewed the history of Committee C-28’s awards 

refine, or drop these standards as appropriate. 

Conclusions 

Step by step, we are building a national and 

international standards infrastructure to facilitate the 

commercial utilization of new advanced materials in 

engine applications. The generic test method standards developed to date have proved to be so practical, reliable, and versatile that they are now being 

used to support a wide range of applications, 

including surgical implants in humans and even 

ceramic military body armor. 

Publications/Presentations 

G. D. Quinn, L. K, Ives, and S. Jahanmir, “On 

the Nature of Machining Cracks in Ground 

Ceramics: Part I: SRBSN Strengths and 

Fractographic Analysis,” Machining Science and 

Technology, 9, 169–210 (2005). 

G. D. Quinn, L. K, Ives, and S. Jahanmir, “On 

the Nature of Machining Cracks in Ground 

Ceramics: Part II: Comparison to Other Silicon 

Nitrides and Damage Maps,” Machining Science 

and Technology, 9 211–237 (2005). 

G. D. Quinn, L. K. Ives, and S. Jahanmir, 

“Machining Cracks in Finished Ceramics,” pp. 1–14 

in Fractography of Advanced Ceramics, II, ed. J. 

Dusza, R. Danzer, and R. Morrell, TransTech Publ., 

Zurich, 2005. 

G. D. Quinn, “Design and Reliability of 

Ceramics: Do Modelers, Designers and 

Fractographers See the Same World?” presented at 

the Cocoa Beach Ceramic Society Conference, 

January 2005, to be published in Ceramic 

Engineering and Science Proceedings. 

G. D. Quinn, “Indentation Size Effects for 

Glasses: Yes, There is a Fracture Contribution,” 

pp. 149–172 in Fracture Mechanics of Glasses and 

Ceramics, Vol. 14, eds. R. C. Bradt, D. Munz, 

M. Sakai, and K. W. White, Klewer/Plenum, New 

York, 2005. 

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Heavy Vehicle Propulsion Materials 

G. D. Quinn, K. Xu, J. A, Salem, and J. J. Swab, 

“SRM 2100: the World’s First Fracture Toughness 

Reference Material,” pp. 499–530 in Fracture 

Mechanics of Glasses and Ceramics, Vol. 14, eds. 

R. C. Bradt, D. Munz, M. Sakai, and K. W. White, 

Klewer/Plenum, New York, 2005. 

J. A. Salem, G. D. Quinn, and M. G. Jenkins, 

“Measuring the Real Fracture Toughness of 

Ceramics—ASTM C 1421,” pp. 531–554 in 

Fracture Mechanics of Glasses and Ceramics, 

FY 2005 Progress Report 

Vol. 14, eds. R. C. Bradt, D. Munz, M. Sakai, and 

K. W. White, Klewer/Plenum, New York, 2005. 

G. D. Quinn, R. J. Gettings, and L. K. Ives, 

“A Standard Reference Materials for Vickers 

Hardness of Ceramics and Hardmetals,” 

pp. 121–128 in Proceedings of HARDMEKO 2004, 

Hardness Measurements Theory and Application in 

Laboratories and Industries, IMEKO Technical 

Committee 5 Hardness Conference, Washington, 

D.C., November 11–12, 2004, ed., S. Low. 

229 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

D. An Integrated Surface Modification of Engineering Materials for Heavy 

Vehicle Applications 

Stephen Hsu 

National Institute of Standards and Technology 

100 Bureau Drive, Gaithersburg, MD 20899 

301-975-6120; fax: 301-975-5995; e-mail: stephen.hsu@nist.gov 

Xaiolei Wang, Lewis K. Ives, Huan Zhang, Yanan Liang, Charles Ying 

National Institute of Standards and Technology 

DOE Technology Development Area Specialist: James J. Eberhardt 

(202) 586-9837; fax: (202) 586-1600; e-mail: james.eberhardt@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(865) 576-6832; fax: (865) 574-6098; e-mail: johnsondr@ornl.gov 

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: National Institute of Standards and Technology, Gaithersburg, Maryland 

Objectives 

• Organize an international cooperative research program on an integrated surface modification technology under 

the auspices of the International Energy Agency (IEA). 

• Design and identify surface features and patterns that can achieve friction reduction and enhanced durability for 

heavy-duty diesel engine components. 

• Develop understanding and appropriate models to explain the effects of texturing on frictional characteristics. 

Develop appropriate thin films and coatings to achieve synergistic and complementary relationship with 

texturing to enhance performance. 

• Discover and develop surface chemistry for protecting the films and coatings that work in synergy with the 

coatings. 

Approach 

• Determine the effects of the size, shape, pitch, and patterns of surface textural features on friction under (1) a 

high-speed, low-load regime; (2) a high-load, high-speed regime; and (3) a high-load, low-speed regime. 

• Develop cost-effective fabrication technologies for creating surface textural features on various surfaces, 

including metals, ceramics, and coatings. 

• Develop a test methodology to measure the effects of the textures on friction. 

• Conduct research to develop an integrated system approach to combine the best practices in thin films, coatings, 

and surface chemistry for performances unrealizable by individual approaches alone. 

• Concurrently, organize an international cooperative research program under the auspices of the IEA to pool 

resources and share this energy conservation technology worldwide. 

231 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

Accomplishments 

• Attended the IEA End Use Working Party meeting in Paris 2005 and presented the Implementing Agreement 

on Advanced Materials for Transportation (IA-AMT) end-of-term report. Based on input from the meeting, 

numerous revisions were made to incorporate new activities and future emphases. 

• Won a 3-year extension of the IEA work from the Committee on Energy Research and Technology (CERT). 

• Contacted potential participants in China, Australia, Finland, Japan, the United Kingdom, and Belgium to ask 

them to participate and to attend the Executive Committee meeting to be held in Porto, Portugal. 

• Participated in a Special Symposium on Surface Texture in Porto in conjunction with the COST 532 

conference. 

• Demonstrated the geometric effects of various surface features under high-speed, low-load conditions and 

achieved significant friction reduction. 

• Developed a new design principle incorporating a hydrodynamic wedge geometry inside the surface feature and 

achieved success in friction reduction, first on steel on soft metals (because of limitations of fabrication 

techniques). 

• Overcame the fabrication limitation and demonstrated for the first time bearing steel on bearing steel under 

high-load, low-speed conditions (boundary lubrication conditions). Such surface texturing would yield a 30% 

reduction in friction in bench testing. 

• Completed the initial development effort on an enhanced selective electrochemical etching technique in 

combination with photolithography on steel surfaces. 

• Successfully demonstrated that cavitation played a role under high-speed, high-load conditions in a single 

dimple observation apparatus. 

• Initiated an elasto-plastic lubrication model development effort to describe the friction reduction mechanism of 

surface textures under boundary lubrication conditions. 

Future Direction 

• Continue parametric study of the size, pattern, density, and L/D ratio in this uncharted territory. 

Introduction 

Frictional losses are inherent in most practical 

mechanical systems. The ability to control friction 

offers many opportunities to achieve energy conservation. Over the years, materials, lubricants, and surface modifications have been used to reduce friction 

in automotive and diesel engine applications to promote energy efficiency. However, in recent years, 

progress in friction reduction technology has slowed 

because much obvious inefficiency has been eliminated. A new avenue is needed. 

Recently, laser-ablated dimples on surfaces have 

shown friction reduction properties and have been 

demonstrated successfully in conformal contacts, 

such as seals, where the speed is high and the load is 

low. The friction reduction mechanism in this 

regime appears to depend on the size, patterns, and 

density of dimples in the contact. The success of 

dimples in reducing friction has opened a new 

avenue for friction control for engine applications. 

The objective of this project, therefore, is to 

explore the feasibility of surface texture designs to 

reduce frictional losses in energy transmission 

components, cylinder liner ring contacts, cam and 

lifters, and overhead bearing components. To ensure 

that the durability criteria are met, the texture 

patterns must be protected by thin films/coatings and 

appropriate lubricating chemistry over a broad range 

of operating conditions. The final outcome of the 

project will be to develop design principles and 

guidelines for various engine components under a 

wide range of operating conditions. To achieve 

maximum energy conservation, this technology 

should be implemented worldwide through international avenues for cooperation such as the IEA IAAMT. Therefore, concomitant international cooperative research under IEA auspices is being 

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Heavy Vehicle Propulsion Materials 

conducted to pool resources from various countries 

to accelerate this technology development. Toward 

this end, the United Kingdom, China, Australia, 

Finland, Sweden, Israel, and Japan have agreed to 

participate under IEA annex IV on surface 

technology. 

Approach 

Previously, experiments were conducted to 

examine various surface textures on steel surfaces 

using photolithography and chemical etching. Using 

the same area coverage (percentage of area occupied 

by the surface textural features), surface features 

such as grooves, triangles, ellipses, and circles were 

compared under high-speed, low-load conditions 

similar to the operating conditions for seals. Results 

suggested that (1) surface texture size and shape had 

significant influence on friction, (2) orientation of 

the surface features with respect to the sliding direction changed the friction; and (3) the primary effect 

of surface textures is to accelerate the transition into 

a hydrodynamic lubrication regime. What is the fundamental friction reduction mechanism in this 

regime is still not clearly understood. 

From an energy conservation point of view, 

most parasitic losses take place under high-load, 

medium- to low-speed conditions. Therefore, we 

need to extend this concept to increasingly severe 

contact conditions. However, when we ran the same 

surface texture patterns under boundary lubrication 

conditions, friction increased. A collaborative effort 

with Northwestern University using a sophisticated 

elastohydrodynamic friction model confirmed our 

experimental observations. Basically, under boundary lubricated conditions (high-load, low-speed), the 

edge stresses around the dimple increase frictional 

losses as if the surface is much rougher. So there is a 

trade-off in this regime. 

Technical Highlights 

During the past year, we developed a new surface 

textural feature with an inclined plane at the bottom 

of the feature to artificially generate a hydrodynamic 

wedge effect under plastic deformation loading; we 

called this the elasto-plastic lubrication model. 

Experiments conducted demonstrated significant 

friction reductions for steel on brass, aluminum, 

copper, and steel under boundary lubricated 

conditions. 

FY 2005 Progress Report 

When we used hard metals such as 52100 steel, 

the initial high contact pressure caused severe wear, 

damaging the surface textures; and the friction 

reduction results were erratic. To overcome this 

issue, we developed a wear-in procedure to create a 

smooth crater surface approximately 0.5 mm in 

diameter and then carved the texture onto the surface 

inside the crater with a diamond tip on the triboindenter. This way, a controlled interface was created 

to avoid severe wear-in and destruction of the surface features. We demonstrated that friction in a 

52100 steel/52100 steel contact could be reduced 

significantly from 0.09 to 0.06. 

We also focused on developing a low-cost, rapid 

fabrication method and conducted tests to verify the 

friction reduction characteristics of the new fabrication technique. We used a microlithography/ 

chemical etching process previously to fabricate 

dimples with depths ranging from 3 to 10 μm. For 

shallow dimples, we needed to control the etching 

process precisely. Numerous factors could affect the 

electrochemical etching process, including the kind 

of electrolyte, the voltage, and the composition of 

the material. Since electricity always passed through 

the high-conductivity points, an inhomogeneous 

oxidation rate on a heterogeneous work piece will 

make the etching process selective. This was critical 

when a shallow depth was needed. A special electrolyte composed with acids was used for this purpose. Acid was effective in removing the oxidant on 

the surface so that etching could start quickly and 

uniformly. A circular dimple pattern was produced 

by this technique. A four-ball wear tester was used 

to carry out the friction tests using a ball-on-three 

flat configuration. The results showed that friction 

reduction of up to 32% has been achieved by the 

texture generated by this fabrication technique. 

Compared with previous textures generated by 

nanoscratching with sloped bottoms, the friction 

reduction was not as high (40–80%). 

L/D Ratio and Size Effects 

One of the critical issues in fabricating the new 

surface features was the effect of the slope at the 

bottom of the feature. In bearing applications, this 

typically was called the L/D ratio (bearing length 

over radius of the bearing) to estimate the lift force. 

We borrowed this concept and applied it to our 

design in terms of the length over the depth of the 

slope fabricated at the bottom of the surface features. 

233 
FY 2005 Progress Report 

Another important factor was the size of the surface 

feature over the contact area. In boundary lubrication conditions, the edges of the surface features 

served as rough spots, increasing the friction; therefore, the optimum number of the features might be 

much smaller than for the high-speed, low-load 

regime. So we decided to conduct a series of 

experiments to investigate this effect. 

We fabricated several surface features by varying the L/D ratios and the sizes of the surface features. We ended with three patterns, A3, B2, and C4. 

The micrographs are shown in Figure 1. The 

densities of the three patterns are kept at 5%. 

Pattern A has a reversed wedge; pattern B has a 

Heavy Vehicle Propulsion Materials 

symmetrical wedge, and pattern C has the classic 

wedge design orientation. 

Figure 2 shows that the classical wedge orientation has the highest friction reduction capacity, 

reducing the average friction coefficient of the 

untextured pattern (UNT) from 0.09 to about 0.05, a 

44% decrease in a 52100 bearing steel on 52100 

bearing steel contact. Another significant point is 

that the friction starts at 0.09, which is at the low 

end of the boundary lubrication regime; the surface 

texturing features lower it to the hydrodynamic 

lubrication regime even though the contact is still 

operating at the boundary lubrication regime, as 

evidenced by the wear of the contact. 

A3 

B2 

C4 

Figure. 1. Three wedge orientation 

patterns for testing. 

Comparison, 20kg 

0 

0.02 

0.04 

0.06 

0.08 

0.1 

0.12 

0 1000 2000 3000 4000 

Sliding speed, rpm 

C

oefficient of fricti

on

UNT 

B2 

A3 

C4 

Figure 2. Comparison of friction coefficients of the 

three patterns under boundary lubrication 

conditions. 

IEA Highlights 

We attended the IEA End-Use Working Party 

(EUWP) meeting in Paris on April 18–19, 2005, and 

presented the IA-AMT end-of-term report. Based on 

input from the meeting, numerous revisions were 

made to incorporate new activities and future 

emphases. The EUWP voted to recommend a 4-year 

extension of the IEA work to CERT. CERT voted in 

June to give tentative approval for a 3-year extension, but it was noted that CERT wishes to see rapid 

progress on implementing the changes. 

Potential participants in China, Australia, 

Finland, Japan, the United Kingdom, and Belgium 

were contacted to ask them to participate and to 

attend the IEA Executive Committee meeting to be 

held in Porto, Portugal, in conjunction with the 

COST 532 (European Cooperative Research Consortium on Triboscience and Tribotechnology) 

234
Heavy Vehicle Propulsion Materials 

technical conference. The Executive Committee was 

successfully convened in Porto with delegates from 

China, the United Kingdom, Australia, Finland, 

Japan, Belgium, Portugal, and the United States. 

Annex 2 was successfully concluded; annexes 3 and 

4 are continuing. We discussed two new activities in 

coatings and nanomaterials. The delegates will take 

the information and seek their governments’ 

approval for formal participation in IA-AMT. The 

next meeting was scheduled for Calgary, Canada, 

for May 7–12, 2006, in conjunction with the Society 

of Tribologists and Lubrication Engineers (STLE)) 

meeting. 

A Special Symposium on Surface Texture was 

held in Porto in conjunction with the COST 532 

conference on October 12–15, 2005. The symposium featured speakers from eight countries— 

representing academia, industry, and government 

laboratories—who described their latest results in 

surface texturing in controlling friction. 

Future Direction 

We will continue our parametric study of size, 

pattern, density, and L/D ratio in this uncharted territory. Until we understand the basic operating 

mechanism and the effects of the operating parameters on the mechanism, we will not be able to 

develop design guidelines. 

Presentations and Publications 

S. M. Hsu, “Surface Texturing under Boundary 

Lubrication for Friction Control,” Society of 

Tribologists and Lubrication Engineers annual 

meeting, Toronto, Canada, May 17–20, 2004. 

Jorn Larson Basse, X. Wang, L. Ives, and S. M. 

Hsu, “Some Friction Experiments with Dimpled 

Surface Texture,” Fourth China International 

Symposium on Tribology, Xian, China, November 

8–11, 2004. 

S. M. Hsu, “An Integrated Surface Texture 

Technology for Friction Control in Engine 

Components,” DOE DEER Conference, San Diego, 

August 30–September 2, 2004. 

S. M. Hsu, “Surface Texturing to Reduce 

Friction in Engine Components and IEA Program,” 

COST 532 management meeting, Ghent, Belgium, 

October 17, 2004. 

S. M. Hsu, “An Integrated Surface Modification 

Technique to Control Friction: A New Paradigm,” 

Keynote speech at the Fourth China International 

FY 2005 Progress Report 

Symposium on Tribology, Xian, China, 

November 8–11, 2004. 

Y. Chae, X. Wang, and S. M. Hsu, “The Size 

Effect of Surface Texture on Lubricated Friction,” 

First International Conference on Advanced 

Tribology, Singapore, December 1–3, 2004. 

Jorn Larson Basse, Lew Ives, and S. M. Hsu, 

“Lubricated Friction Experiments with Coarse 

Groove Texture,” Society of Tribologists and 

Lubrication Engineers annual meeting, Las Vegas, 

May 15–19, 2005. 

S. M. Hsu and Xiaoli Wang,” Surface Texturing 

for Friction Control,” Society of Tribologists and 

Lubrication Engineers annual meeting, Las Vegas, 

May 15–19, 2005. 

S. M. Hsu, “Surface Texturing: A New Design 

Principle for Friction Control under Boundary 

Lubricated Conditions,” International Tribology 

Congress, Kobe, Japan, May 29–June 2, 2005. 

S. M. Hsu, “Surface Texturing: Comparison of 

Various Geometric Shapes on Friction under Highspeed Low-load Conditions,” International 

Tribology Congress, Kobe, Japan, May 29–June 2, 

2005. 

S. M. Hsu, “An Integrated Surface Modification 

Technology: Project Review,” DOE project review, 

Oak Ridge National Laboratory, Oak Ridge, 

Tennessee, September 13–15, 2005. 

S. M. Hsu, “Integrated Surface Modification 

Technology: A Progress Report,” COST 532 

Symposium on Surface Textures, Porto, Portugal, 

October 12, 2005. 

Jorn Larson Basse, X. Wang, L. Ives, and S. M. 

Hsu, “Some Friction Experiments with Textured 

Surfaces,” in Proceedings of Nordic Symposium on 

Tribology, Troms, Norway, June 2004. 

X. Wang, S. M. Hsu, “An Integrated Surface 

Modification Technique to Control Friction: A New 

Paradigm,” in Proceedings of the Fourth China 

International Symposium on Tribology, Xian, 

China, November 8–11, 2004. 

X. Wang and S. M. Hsu, “Surface Texturing: 

Comparison of Various Geometric Shapes on 

Friction under High-Speed Low-Load Conditions,” 

in Proceedings of the International Tribology 

Conference, Kobe, Japan, May 29–June 2, 2005 

X. Wang and S. M. Hsu, “Surface Texturing: A 

New Design Principle for Friction Reduction Under 

Boundary Lubrication Conditions,” in Proceedings 

of the International Tribology Conference, Kobe, 

Japan, May 29–June 2, 2005. 

235 

Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Appendix A. ACRONYMS AND ABBREVIATIONS 

2-D two-dimensional 

3BOB three ball-on-ball 

3BOR three ball-on-rod 

3-D three-dimensional 

ACEM aberration-corrected electron microscope

 ACERT Advanced Combustion Emissions Reduction Technology 

ADF annular dark-field 

AE acoustic emission 

AFM atomic force microscope 

AFRL Air Force Research Laboratory (Wright-Patterson Air Force Base, Dayton, Ohio) 

AIST (National Institute of) Advanced Industrial Science and Technology (Nagoya, 

Japan) 

Al aluminum 

AMCL Advanced Materials Characterization Laboratory 

ANL Argonne National Laboratory 

ANSYS commercial finite-element analysis software 

ASM American Society for Metals 

ASME American Society of Mechanical Engineers 

ASTM American Society for Testing and Materials 

Ba barium 

BAM Bundesanstalt für Materilforschung und - prüfung (Berlin, Germany)

 BET Brunauer-Emmett-Teller theory

 BMOs Base metal oxides 

BU Bournemouth University (United Kingdom) 

CARES Ceramics Analysis and Reliability Evaluation of Structures 

CC current collector 

CCVD combustion chemical vapor deposition 

CDF cumulative distribution function 

CEN Committee for European Normalization 

CerSaT Ceramic Science and Technology 

CERT Committee on Energy Research and Technology 

CHIP cold + hot isostatic pressing

 CIM chemically interactive material 

CMM coordinate measuring machine 

CP commercially pure 

CRADA cooperative research and development agreement 

CRT continuously regenerating technology 

CTE coefficient of thermal expansion 

CVD chemical vapor deposition 

CWRU Case Western Reserve University 

DDC Detroit Diesel Corporation 

DISI direct-injection spark-ignition 

DITS dynamic indentation test system 

DOE U.S. Department of Energy 

237 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

DPF diesel particulate filter 

DRIFT diffuse reflectance infrared Fourier transform spectroscopy

 DSC dispersion-strengthened composite 

DT deformation twin 

DTA differential thermal analyzer 

EDM electro-discharge machining 

EDS energy-dispersive spectroscopy 

EERE Office of Energy Efficiency and Renewable Energy 

EGR engine gas recirculation 

EPA U.S. Environmental Protection Agency

 EPMA electron probe microanalysis 

EVD electrochemical vapor deposition 

FCVT FreedomCAR and Vehicle Technologies 

FEA finite element analysis 

FEM finite element method 

FEM finite element method/finite element modeling 

FTIR Fourier transform infrared spectroscopy 

FY fiscal year 

GPa gigapascals 

HA-ADF high-angle annular dark-field 

HA-ADG high-angle annular dark field 

HC hydrocarbons 

HDI high-density infrared 

HFIR High Flux Isotope Reactor 

HIP hot isostatic pressing 

HIPed hot isostatically pressed 

HS high stress 

HSWR High-Strength Weight-Reduction Materials Program 

HTML High Temperature Materials Laboratory 

HVOF high velocity oxygen-fuel 

IA implementing agreement 

IA-AMT Implementing Agreement for a Programme of Research and Development on 

Advanced Materials for Transportation Applications 

IEA International Energy Agency 

IEST integrated engineered surface technology 

IITS instrumented indentation test system 

INL Idaho National Engineering and Environmental Laboratory 

ISO International Organization for Standards 

JIS Japan Industrial Standards 

238 
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

LabVIEW Programming software made by National Instruments 

LANL Los Alamos National Laboratory

LBNL Lawrence Berkeley National Laboratory 

LNT lean NOx trap 

LS low stress 

LST laser surface texturing 

METI Ministry of Economy, Trade and Industry (Japan; formerly MITI) 

Mg magnesium

 MMC metal matrix composite 

NASA National Aeronautics and Space Administration 

NDE nondestructive evaluation 

NiAl nickel aluminide 

NiCr nickel chromium

NiCrFe nickel chromium iron 

NIST National Institute for Standards and Technology

NIST National Institute of Standards and Technology 

NOx oxides of nitrogen; nitrites/nitrates species 

NPD neutron powder diffraction 

NPL National Physical Laboratory (UK) 

NSF National Science Foundation 

NSR NOx stoichiometric ratio 

NTN NTN Bower 

NTRC National Transportation Research Center

 OBD on-board diagnostic 

OFCVT Office of FreedomCAR and Vehicle Technologies 

OFHC oxygen-free high-conductivity

 OM optical microscopy

ORNL Oak Ridge National Laboratory

 PBC phosphate-bonded composite 

PCA principal component analysis 

PDS probabilistic design system 

PI principal investigator 

PM particulate matter 

PMT photomultiplier tube 

PNNL Pacific Northwest National Laboratory

PoT pin-on-twin oscillating wear test 

Pt platinum

PZT lead zirconate titanate 

RBS rotary bend strength 

RCF rolling contact fatigue 

RE reference electrode 

239 
FY 2005 Progress Report Heavy Vehicle Propulsion Materials 

SAcM scanning acoustic microscopy

 SAM scanning auger microprobe 

SAXS small angle x-ray scattering 

SCG slow crack growth 

SCR selective catalytic reduction 

SE sensing electrode 

SEM scanning electron microscopy

 Si3N4 silicon nitride

 SiC silicon carbide 

SIDI spark-ignition direct-injection 

SIUC Southern Illinois University–Carbondale

SMPS scanning mobility particle sizer 

SOFC solid oxide fuel cell 

SOx sulfur and oxygen containing compounds 

SpaciMS spacially resolved mass spectrometer 

SRBSN sintered reaction-bonded silicon nitride 

SRM standard reference material

 SS stainless steel

STEM scanning transmission electron microscopy

TBCs thermal barrier coatings 

TEM transmission electron microscopy

TFS through focal series 

TGA thermal gravimetric analysis (or analyzer) 

Ti titanium

 TiAl titanium aluminide 

TiC titanium carbide 

TTBCs thick thermal barrier coatings 

TTZ transformation-toughened zirconia 

TWC three-way catalyst 

TWS Third-Wave Systems 

UHMW-PE ultra-high molecular weight polyethylene 

UHR-TEM ultra-high resolution TEM 

UTBM University of Technology at Belfort-Montbeliard 

WBDF weak beam dark field 

WC tungsten carbide 

XPS X-ray photoelectron spectroscopy

XRD X-ray diffraction/X-ray powder diffraction 

YSZ yttria-stabilized zirconia 

240