A Strong Energy Portfolio for a Strong America 

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

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

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

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 2004 

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 

Sidney Diamond Technology Development Area Specialist 

May 2005 

Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

iii 

CONTENTS 

1. INTRODUCTION ······························································································································· 1 

2. MATERIALS FOR FUEL SYSTEMS······························································································· 7 

 A. Low-Cost, High-Toughness Ceramics·························································································· 7 

 B. Low-Cost Manufacturing Processes for Ceramic and Cermet Diesel Engine Components········· 11 

 C. Cermet Composites for Wear Applications ·················································································· 15 

3. MATERIALS FOR EXHAUST AFTERTREATMENT ··································································· 19 

 A. Materials for Exhaust Aftertreatment ··························································································· 19 

 B. Development of Materials Analysis Tools for Studying NOx Adsorber Catalysts······················· 25 

 C. Development of NOx Sensors for Heavy Vehicle Applications ·················································· 33 

 D. Ultra-High Resolution Electron Microscopy for Characterization of 

 Catalyst Microstructures and Deactivation Mechanisms···························································· 39 

 E. Microstructural Changes In NOx Trap Materials Under Lean and 

 Rich Conditions at High Temperatures ······················································································· 45 

 F. Aftertreatment Catalysts Materials Research················································································ 53 

 G. Durability of Diesel Particulate Filters ························································································· 59 

4. MATERIALS FOR AIR HANDLING, HOT SECTION AND 

 STRUCTURAL COMPONENTS ····································································································· 65 

 A. High-Temperature Advanced Materials for Lightweight Valve Train 

Components ·································································································································· 65 

 B. Engineered Surfaces for Diesel Engine Components ··································································· 71 

 C. Processing and Characterization of Structural and Functional Materials 

for Heavy Vehicle Applications ··································································································· 77 

 D. NDE Development for Ceramic Valves for Diesel Engines························································· 83 

 E. Durability of Diesel Engine Component Materials ······································································ 89 

 F. Life Prediction of Diesel Engine Components ············································································· 95

 G. Low-Cost Manufacturing of Precision Diesel Engine Components············································· 103 

 H. Cost-Effective Machining of Titanium Materials········································································· 109

 I. Advanced Cast Austenitic Stainless Steels for High-Temperature Components·························· 115 

 J. Interfacial Control of Deformation Twinning in Creep-Deformed 

 TiAl/Ti3Al Nanolaminate ············································································································· 119 

 K. Laser Surface Texturing of Lubricated Ceramic Parts ································································ 123 

 L. Ti-6Al-4V Billet Feedstock Manufacture and Evaluation···························································· 129 

 M. High-Density Infrared Surface Treatment of Materials for Heavy-Duty Vehicles······················· 131 

 N. High-Temperature Aluminum Alloys··························································································· 135 

 O. Applications of Titanium Alloys for Heavy-Duty Vehicles ························································· 141 

 P. Mechanical Behavior of Ceramic Materials for Heavy-Duty Diesel Engines······························ 159 

 Q. Powder Processing of Nanostructured Alloys Produced by Machining ······································· 165 

 R. Deformation Processes for Next-Generation Ceramics································································ 171 

 S. Synthesis of Nanocrystalline Ceramics ························································································ 177 

 T. Development of Titanium Component Applications in Heavy Duty Diesel Engines ·················· 181 

5. TEST AND MATERIALS STANDARDS ························································································ 187 

 A. Rolling Contact Fatigue················································································································ 187 

 B. Implementing Agreement for a Programme of Research And Development 

on Advanced Materials for Transportation Applications······························································ 193 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

iv

 C. Testing Standards·························································································································· 199 

 D. Surface Modification of Engineering Materials for Heavy Vehicle Applications························ 205 

Appendix A: ACRONYMS and ABBREVIATIONS··········································································· 211 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

1

1. 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 

The fuel systems for diesel engines are complex, expensive systems that are critically important to 

meeting the efficiency and emissions targets of the Office of FreedomCAR and Vehicle Technologies 

(OFCVT). Virtually every path to improving control of the combustion process in order to improve efficiency 

and reduce emissions depends on improvements in the fuel injection system. 

Fuel injectors are highly precise. The individual components must fit together with clearances that are 

sometimes less than 1 × 10−6 meters. Control of the combustion process requires precise control of the size, 

shape, and surface finish of the injector components. 

The current systems for reliably controlling multiple injections are limited by the ability of mechanical 

and electronic systems to respond precisely and quickly enough to provide additional injection control. Smart 

materials, such as piezoelectric materials, offer the potential for better control of fuel injection and have been 

recently introduced for automobiles. However, a number of improvements in materials and manufacturing 

methods for the materials are yet required for heavy vehicles. 

Presently, 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), Southern Illinois 

University–Carbondale (SIUC), CoorsTek, Inc. and Cummins, Inc. ORNL prepared a large batch of 

processed powder, which was provided to CoorsTek for injection molding of test components. The injectionmolded components were then sintered at ORNL and provided to Cummins for testing in fuel injection 

systems. In a parallel effort, SIUC evaluated a low-cost, high-throughput process, continuous belt sintering, 

for manufacturing the cermet components. SIUC determined that rapid heating, at rates up to 750°C/minute, 

resulted in higher densities and more uniform microstructures and offers the potential to reduce sintering 

costs by as much as 50% by increasing part throughput. 

Materials for Exhaust Aftertreatment 

The reduction of nitrogen oxides (NOx) and particulate emissions is critically important to OFCVT’s 

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. 

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 development of more durable systems. The development of advanced NOx sensors is being 

conducted to facilitate optimal engine and aftertreatment control strategies. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

2

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 upon 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 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. 

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 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. Accomplishments included the 

characterization via high-resolution annular dark-field scanning TEM (STEM) that imaged the structure of a 

model NOx trap material of near-atomic clusters and rafts of platinum atoms 1–3 atomic layers thick; 

characterization via TEM and STEM of core samples taken from catalyst monoliths removed from vehicles 

after up to 82,000 km of driving; and taking delivery of the new aberration-corrected electron microscope 

(ACEM), which will enable imaging of ultra-fine clusters of catalyst species on real catalyst samples of 

interest. 

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 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, via cooperative research and development agreements, in characterizing lab and engine-tested 

catalysts via X-ray diffraction, spectroscopy, and microscopy; and in characterizing diesel particulate filters 

and developing probabilistic design tools to predict the useful lifetime 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 demonstrating NO2 sensing elements 

with extremely high sensitivity; demonstrating biased NO-selective sensing elements; and constructing both 

types of elements to consist of co-planar electrodes, one of oxide and the other of a noble metal. 

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. 

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 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

3

made from a titanium alloy. Accomplishments this year included procuring cast TiAl turbine wheels, 

developing friction welding technology for joining the turbine wheels to shafts and for realizing a higher joint 

strength than in the base TiAl wheel, and conducting preliminary turbocharger bench tests. 

In a related project, Dynamet Technologies is conducting a research and development project to develop 

low-cost Ti-6Al-4V billet feedstock using a blend of titanium and alloy powders and inexpensive Ti-6Al-4V 

machine turnings. Dynamet has evaluated this low-cost titanium alloy feedstock as starting billet material for 

casting, forging, and extrusion operations and has produced high-density material in each case. 

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. Thermal-mechanical fatigue (TMF) of the initial commercial heats was 

completed this year, together with aging of various specimens. There was a significant advantage for CF8CPlus relative to CN-12 or high-SiMo cast iron in TMF testing to 760°C. This year, MetalTek International 

produced additional static sand-cast and centrifugally-cast heats of CF8C-Plus for testing at Caterpillar and 

ORNL. That testing began this year and will be completed next year. 

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

fabricate prototype engine valves from silicon nitride and titanium aluminide 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 was evaluated and indicated that 

valves with good surface finish performed well in bench tests. 

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. Specific objectives are (1) to 

develop laser technology for surface dimpling, cleaning, and laser-assisted spraying to enhance adherence and 

increase coating strength; (2) to develop phosphate-bonded composites for thermal management coatings; and 

(3) to evaluate quasicrystalline materials as potential thermal barrier and wear coatings. This year, a cost 

analysis of laser pre- and post-treatments was conducted, and post-laser tacking was selected as the method 

for treating coatings to increase adherence. Phosphate-bonded composite coatings survived thermal cycling to 

650°C with good residual adherence to the substrate. Quasicrystalline coatings were sprayed using the highvelocity oxygen-fueled technique, and evaluation began. In a collaborative effort, ORNL demonstrated the 

use of high-density infrared surface treatment as a means of producing hard-metal coatings with low porosity. 

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. 

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 investigated surface modification of lubricated 

ceramic parts for applications in fuel systems and related applications. Purdue University is investigating the 

consolidation of low-cost machining chips to produce nanocrystalline components with high strength and 

hardness. 

In a new effort this year, ORNL is investigating the plastic deformation behavior of industrial ceramics 

and the potential for low-temperature forming of components. In collaboration with Pennsylvania State 

University, a novel sintering method was developed to fabricate nanocrystalline (less than 100-nm grain size) 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

4

zirconia ceramics. Enhancement of the plastic deformation of sub-micron zirconia in applied electric fields 

was observed in collaborative work with North Carolina State University. ORNL also established dynamic 

indentation and instrumented scratch testing facilities and associated finite-element analysis models to 

characterize and model contact-induced damage and its links to wear performance and to the optimization of 

machining and rolling contact fatigue. 

Materials and Testing Standards 

OFCVT has an International Energy Agency (IEA) “Implementing Agreement for a Programme of 

Research and Development on Advanced Materials for Transportation Applications” (IA-AMT). The 

objectives of the IA-AMT are as follows: 

1. to identify and evaluate promising new processing and surface engineering technologies capable of 

improving materials performance in transportation systems and 

2. to promote and implement pre-competitive development and verification of advanced characterization 

methods appropriate for advanced materials for transportation applications. 

Annex III, which was approved in July 2002, consists of two subtasks on contact reliability of advanced 

engine materials including structural ceramics, composites, and nanostructured friction/wear coatings. 

Subtask 1 is an information exchange, and Subtask 2 focuses on the development of standard test methods 

and procedures for determining the rolling contact fatigue resistance of advanced materials. At present, the 

active contracting parties for the IA-AMT are Germany (Bundesanstalt für Materialforschung und–prüfung), 

Canada (CANMET), United Kingdom (Department of Trade and Industry), and the United States (DOE). 

A new activity to be conducted under the IEA program was initiated last year by the National Institute of 

Standards and Technology (NIST). The objectives of this new effort are as follows: 

1. organize an international cooperative research program on an integrated surface modification technology 

under the auspices of the IEA 

2. design and identify surface features and patterns that can achieve friction reduction and enhanced 

durability for heavy-duty diesel engine components 

3. develop understanding and appropriate models to explain texturing effects on frictional characteristics 

4. develop appropriate thin films and coatings to achieve a synergistic and complementary relationship with 

texturing to enhance performance 

5. discover and develop surface chemistry for protecting the films and coatings that work in synergy with 

the textures 

Accomplishments this year include the development 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. 

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. In FY 

2004, the pre-standardization work on the flexural strength of split cylinders was completed, a NIST Guide to 

Practice on Fractographic Analysis of Brittle Materials was 50% completed, and several standards from the 

American Society for Testing and Materials and the International Organization for Standards were revised. 

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. In FY 2004, a three-ball-on-rod 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

5

test facility was established and informal collaboration was initiated with Bournemouth University in the 

United Kingdom. 

Rogelio Sullivan 

Team Leader, Materials Technologies 

Office of FreedomCAR and Vehicle 

 Technologies 

Energy Efficiency and Renewable Energy

Sidney Diamond 

Technology Development Area Specialist 

Office of FreedomCAR and Vehicle 

 Technologies 

Energy Efficiency and Renewable Energy
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

7 

2. MATERIALS FOR FUEL SYSTEMS 

A. Low-Cost, High-Toughness Ceramics 

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

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: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-2476; e-mail: sid.diamond@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 

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

Approach 

• Develop TiC-based composites with 40–60 vol % Ni3Al to take advantage of their expansion characteristics, 

which are very close to those for steel. 

• Conduct a development effort in collaboration with CoorsTek, Inc., and Cummins Engine on processing scaleup and engine testing. 

Accomplishments 

• Supplied a large batch of processed powder mixtures to CoorsTek for injection molding of test components. 

• Sintered injection-molded components for engine testing in fuel injection systems. 

Future Direction 

• Complete work with CoorsTek (a parts supplier) to scale up the processing and to supply CoorsTek with pilotplant–scale quantities of powder mixtures for injection molding trials. 

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

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 content is on the order of 

30 to 50 vol % in order 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 liquidphase sintering, and most of the early work used 

gas-atomized Ni3Al particles with fine TiC powders. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

8 

Later work was done using Ni 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 were 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. 

Results

Large-Batch Processing of TiC-Ni3Al 

Composites 

Several large batches (>3 kg each) of a 50% 

TiC-Ni3Al (with 2% molybdenum) composition 

have been milled and blended together to produce 

~13 kg of powder. Approximately 10 kg was sent to 

CoorsTek for injection molding testing. A large 

batch (>3 kg) was also produced that used a cheaper 

NiAl powder from the same manufacturer that was 

coarser than that previously used. Much of the earlier work used powder that had particle sizes ≤20 

μm. The batch produced recently used NiAl powder 

with particle sizes ≤150 μm (−100 mesh). 

The sintering furnace used for processing studies was out of service for many months because of a 

bad controller and was repaired during the last quarter of the fiscal year. However, sintering runs were 

completed on injection-molded test components 

from CoorsTek. These parts were sintered and sent 

to Coors-Tek and Cummins Engine Company for 

final component finishing and engine testing. 

Alternate Precursors for Ni3Al Formation 

All previous work used a combination of Ni and 

NiAl for an in-situ reaction to form the Ni3Al. 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. Several new batches were milled and 

pressed into discs and billets. The compositions are 

described in Table 1. 

As shown, several different compositions are being examined. The new batch of Al-30 Ni from the 

catalyst manufacturer will determine if these materials can be made without the impurities, as observed 

previously. Earlier work demonstrated that these 

precursors produced high-density parts but had low 

strengths because of impurity contamination. One of 

the other samples will re-examine the use of prealloyed Ni3Al (made by gas atomization). Earlier 

work at high TiC contents (70–90 vol %) showed 

these Ni3Al powders produced non-homogeneous 

microstructures with large “pools” of Ni3Al surrounded by the finer TiC grains. However, with the 

TiC content of interest now at 50 vol %, the presence of the Ni3Al pools may not be significant to the 

performance of the cermets. Samples were also fabricated using a commercial Hastelloy X powder. 

This is a Ni-Cr-Fe-Mo alloy with good oxidation 

and corrosion resistance. 

An attempt was made to make a pre-reacted AlNi precursor that could be milled to fine sizes during 

the normal processing of the cermets (Sample 

DC-3-10). Thermodynamic calculations indicated 

that the lowest heat of formation per gram was associated with making Al3Ni. Appropriate amounts of 

aluminum and nickel were dry-milled and then fired 

(at a slow heating rate) to 700°C, which is sufficient 

to melt the aluminum. The resulting reacted product 

was very hard and non-friable. Consequently, this 

route to producing an Al-Ni precursor did not appear 

to be one with significant potential. 

Table 1. Summary of TiC-Ni3Al compositions fabricated

Batch no. TiC content ( vol %) Ni3Al source 

DC-3-7 50 NiAl + Ni (milling done with larger WC media) 

DC-3-8 50 Al-30 Ni Catalyst (new batch from manufacturer) 

DC-3-9 50 Pre-alloyed Ni3Al (-325 mesh) 

DC-3-10 50 Reacted Al3Ni 

DC-3-11 50 Hastelloy X 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

9 

Conclusions

Pilot scale batches of TiC-Ni3Al composite mixtures were produced. In collaboration with CoorsTek, injection-molded components were fabricated 

and sintered to high density. The parts will be used 

for engine testing by Cummins Engine Company. 

Publications/Presentations

T. N. Tiegs, F. C. Montgomery, and D. A. 

Menchhofer, “Effect of NiAl Precursor Type on 

Fabrication and Properties of Ni3Al-Bonded Carbide 

Composites,” presented at the American Ceramic 

Society Conference on Ceramics and Composites, 

Cocoa Beach, FL, Jan. 25–28, 2004. 

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

Menchhofer, “Effect of Ni-Al Precursor Type on 

Fabrication and Properties of TiC-Ni3Al Composites,” to be published in Ceram. Eng. Sci. Proc. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

10 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

11 

B. Low-Cost Manufacturing Processes for Ceramic and Cermet Diesel Engine 

Components 

Dale E. Wittmer, Professor 

Southern Illinois University, 

Carbondale, IL 62901 

(618) 453-7006; fax: (618) 453-5847; e-mail: wittmer@engr.siu.edu 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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: Southern Illinois University, Carbondale, Illinois 

Objective 

• Investigate higher continuous sintering rates as a more cost-effective manufacturing process for ceramic and 

cermet diesel engine components.

Approach 

• Process cermet compositions containing titanium carbide (TiC) and nickel aluminide intermetallic matrix phase 

into cylindrical shapes by standard dry pressing techniques. 

• Sinter the cermets by continuous sintering at much higher sintering rates than previously investigated. 

• Measure physical properties and compared them with previously processed materials. 

Accomplishments 

• Modified the continuous furnace at Southern Illinois University–Carbondale (SIUC) to allow for the investigation of higher sintering rates, which appear to produce more uniform microstructures in nickel aluminide 

bonded TiC cermets. Higher sintering rates could lead to improved economics of production, while more uniform microstructures give inherently more uniform properties and performance. 

• Continued interaction with Oak Ridge National Laboratory (ORNL) and CoorsTek to promote commercialization of cermet diesel engine components. 

Future Direction 

• The past year was the final year of funding for this project. However, exploration of the use of higher sintering 

rates for other potential cermet formulations should be continued. 

Introduction

All manufactured parts can potentially benefit 

from improved manufacturing processes and the 

incorporation of advanced materials. Injection molding has been previously shown to be a very costeffective method for producing precision parts with 

minimal labor cost, and continuous sintering has 

been proven industrially to provide an economic 

advantage for sintering many advanced ceramics. 

More recently, continuous sintering has been used as 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

12 

a means of rapid sintering of cermets composed of 

TiC in an intermetallic matrix. 

To reduce costs, it is essential to minimize material waste while maximizing the yield of finished 

parts that are within specifications. A potential 

means of maximizing the furnace yield of parts that 

are in-specification while minimizing material losses 

due to furnace-related problems, thereby reducing 

part cost, is the use of continuous sintering. The continuous furnace at SIUC has been used to sinter a 

wide range of pre-alloyed intermetallic-TiC formulations and similar formulations where the intermetallic is formed by reaction sintering of the individual 

elements. Based on this work, the most promising 

intermetallics contained 30–50 vol % of NiA, NiC, 

or NiCrFe added to a fine-grained commercial TiC. 

These formulations were found to have high 

strength, hardness, toughness, and corrosion resistance. In addition, their thermal expansion can be 

engineered to be very close to that of cast iron and 

steel, which will reduce thermal expansion mismatch in several key diesel engine applications. 

The present work was focused on modifying the 

continuous furnace to allow for higher transport 

rates through the furnace without causing damage to 

the furnace structure. Once modified, cermets were 

sintered at much higher sintering rates than previously investigated. Also, interactions were continued 

with ORNL and CoorsTek to promote commercialization of these cermets for diesel engine components. 

Approach

Continuous Furnace Modification 

The continuous furnace was constructed using 

an ultra-high molecular weight polyethylene 

(UHMW-PE) gear on the exit end of the furnace. 

This was done to give some flexibility to the ceramic link belt and to avoid stresses that might cause 

the ceramic pins in the belt to fracture. Attempting 

to run the continuous furnace at high transport 

speeds, in excess of about 1.5 in./min (4.31 cm/min), 

when the furnace temperature exceeded 1400°C, 

caused overheating of the UHMW-PE gear. It was 

believed that running for extended periods of time at 

these rates or higher would eventually cause damage 

to the ceramic link belt. Figure 1 shows the continuous furnace and Figure 2 shows the original 

Figure 1. Continuous furnace at SIUC showing location 

of original polymer gear. 

Figure 2. Original ultra-high molecular weight polyethylene gear. 

polymer gear. To eliminate this problem, an aluminum alloy, 6061-T6, gear was designed and machined at SIUC. Figure 3 shows the new aluminum 

alloy gear installed on the continuous furnace exit 

drive. As seen, several holes were machined in the 

gear to improve heat transfer, which allows the gear 

to run cooler. 

Investigation of High Heating Rates

Previous continuous sintering of cermets at 

SIUC was accomplished at sintering rates less than 

or equal to 125°C/min. This task investigated the use 

of continuous sintering as a means of obtaining high 

heating rates (>125°C/min) in the liquid-phase sintering of Ni3Al-TiC cermets. Since the continuous 

furnace uses a ceramic belt to transport the load 

through the hot zone, the heating rate is dependent 

on the belt speed. The modification of the furnace 

allowed the investigation of much higher heating 

rates than previously investigated. The cermets investigated are sintered by liquid-phase sintering; 

therefore, employing high heating rates could enhance particle rearrangement and solution precipita- 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

13 

Figure 3. Aluminum alloy gear installed. 

tion during the formation of the liquid. In the Ni3AlTiC system, TiC is known to be soluble in the intermetallic, with about 4% Ti remaining in solution 

in the Ni3Al. It was hoped that the high heating rates 

would lead to enhanced solution and precipitation of 

the TiC, thereby producing finer TiC in the microstructure. Ni3Al-TiC cermets containing 50 vol % 

intermetallic were sintered in flowing argon at 

1450°C, using four different belt speeds that corresponded to heating rates of 125, 250, 500 and 

750°C/min. Soak time at peak temperature was held 

constant at 30 min. Following sintering, density and 

microhardness were determined and the microstructure investigated by scanning electron microscopy 

(SEM). 

Industrial Collaboration

Collaboration between SIUC, ORNL, and 

CoorsTek has continued since the start of this 

project. Cermets processed by low-pressure injection 

molding and continuous sintering were submitted to 

ORNL for wear testing, and materials have been 

submitted to CoorsTek for evaluation. During this 

project period, CoorsTek submitted materials to 

SIUC for sintering and evaluation. 

Results

Furnace Modification 

The furnace was modified as planned by the addition of the aluminum alloy gear, as shown in Figure 3. This modification allowed the belt speed to be 

increased from 0.75 to 4.5 in./min (1.9 to 

11.43 cm/min), which is equivalent to 125 and 

750°C/min, respectively. This is assuming a maximum hot-zone temperature of 1450°C and all three 

zones being at temperature. Obviously, higher hotzone temperatures would lead to higher heating 

rates. This is significant in that traditional production furnaces usually do not exceed heating rates of 

over 50°C/min, and usually operate at much lower 

temperatures.

Investigation of High-Heating Rates

As reported, four different heating rates were investigated. Figure 4 shows the equivalent heating 

rate as a function of belt speed. 

Figure 4. Heating rate as a function of belt speed. 

The resulting density results are shown in Figure 5. The relative density achieved was higher at 

the higher heating rates, with the belt speed of 

3.0 in./min (7.62 cm/min) or heating rate of 

500°C/min giving the highest density results. 

As seen in Figure 6, the lower sintering rates 

produced microstructures in which the intermetallic 

appears as large pools. The bright phase is the in 

termetallic, and the dark phase is the TiC or, in some 

cases, porosity. These intermetallic pools were 

found to be reduced in size at the higher heating 

rates. At the highest heating rate, 750°C/min, the 

pools appear to have been eliminated. In addition, 

there appears to be more rounding of the TiC particles at the higher heating rates. 

Some potential advantages of the higher heating 

rates appear to be obvious. The use of the continuous furnace has been found previously to offer a reduction of as much as 50% in the sintering costs for 

advanced ceramics. A similar saving could be anticipated for cermets if the properties obtained are 

acceptable. The higher heating rates obviously 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

14 

Figure 5. Relative density as a function of belt speed. 

would allow for the production of more parts per 

shift, which would further reduce the overall manufacturing cost. An additional advantage of the high 

heating rate appears to be more uniform microstructures, which should in turn produce more uniform 

properties for the cermets. 

Industrial Collaboration

SIUC, ORNL, CoorsTek, and Cummins have 

been involved in a confidential collaboration, and all 

data related to this project have been the responsibility of the individual contributors. During this reporting period, SIUC received four test bars from 

CoorsTek for continuous sintering and evaluation. 

None of the test bars was sintered without warping, 

and the densities obtained were well below the 

benchmark set by CoorsTek 

The results of wear testing at ORNL have been 

very favorable for many of the cermets investigated. 

The hardness, toughness, and strength of the cermets 

investigated are responsible for the very high wear 

resistance. The effects of wear resistance of the intermetallic and of particle size of the TiC in the cermet are not very well understood. Several of the 

simple cylindrical shapes produced at SIUC were 

previously submitted to Peter Blau at ORNL for 

sliding wear evaluation. The results continue to be 

encouraging. 

Figure 6. SEM microstructures for the four heating rates 

investigated. 

Conclusions

The continuous furnace was modified to allow 

for higher heating rates. Ni3Al-TiC cermets containing 50 vol % intermetallic were sintered in flowing 

argon at 1450°C using four different belt speeds corresponding to heating rates of 125, 250, 500 and 

750°C/min. Higher heating rates produced higher 

densities and more uniform microstructures than the 

lower heating rates. Higher heating rates could lead 

to improved production costs by increasing part 

throughput. 

Presentation

Dale E. Wittmer and Joshua W. Steffen, “Effect 

of High Sintering Rates on Microstructure of Ni3alTic Cermets,” (poster presentation at PM2TEC 

2004, Chicago, IL, June 13–17, 2004), pp. 38–43 in 

Advances in Powder Metallurgy and Particulate 

Materials, Part 9, 2004. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

15

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) 377-7678; fax: (812) 377-2041; e-mail: matt.anderson@cummins.com 

*CoorsTek, Golden, CO 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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 

determination of the shrink rate during sintering for prototype tool construction and demonstration. 

• Test prototype components in an intended application on Cummins Fuel System test rigs to show the viability 

of the material in an application. 

Approach 

• Make a prototype injection molding tool of a proposed Cummins Fuel Systems component and produce test 

pieces of the TiC/Ni3Al cermet to near-net shape for evaluation of shrinkage, microstructure uniformity, and 

density. 

• Conduct abuse tests of prototypes machined from isostatically pressed rods of the TiC/Ni3Al cermet at Cummins Fuel Systems. 

Accomplishments 

• Achieved a 99+% dense structure with the injection molding process. The TiC distribution within the Ni3Al matrix was uniform, and there was good wetting of the TiC particles. The shrink rate was also determined. 

• Successfully tested prototypes at Cummins Fuel Systems. The components are required to withstand both sliding and impact wear. No adverse wear was observed after the testing. 

Future Direction 

• Finalize the design and test the injection-molded materials on Cummins Fuel Systems test rigs. 

• Determine and/or develop a supply base for raw material constituents. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

16

Introduction

Increasingly stricter diesel engine combustion 

emission standards and the desire of buyers for 

maintained or improved fuel economy require that 

fuel injection systems become more advanced. 

Higher fuel injection pressures (to aid emissions 

control) and more precise control of when and how 

much fuel is injected (for both emissions control and 

fuel economy) are required. New materials that exhibit excellent wear properties against steel components are needed for applications where components 

slide and impact against each other. One material 

that Cummins is interested in was developed under a 

DOE cooperative agreement: 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 for finding an application for it, testing the material in fuel 

system components, and evaluating its performance 

against that of other materials. 

Approach

The established composition of TiC/Ni3Al cermet will be processed using the powder injection 

molding process to create a near-net shape component. Its microstructure, density, and shrinkage rate 

will be evaluated to determine the effectiveness of 

the manufacturing process. Coinciding with this 

work, fuel system rig testing of some prototype 

components will be conducted 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 first injection-molded batch was successfully 

completed. It achieved over 99% density when sintered in a non-overpressure argon atmosphere. The 

microstructure (see Figure 1) shows the porosity and 

TiC particles are well distributed throughout the 

structure. The high- magnification scanning electron 

microscope (SEM) image in Figure 2 shows good 

wetting of the TiC particles. This first batch had a 

higher shrinkage rate than expected; therefore, 

Figure 1. SEM backscattered image of polished crosssection of injection-molded TiC/Ni3Al cermet. The black 

spots are porosity, the dark gray phase is the TiC, and the 

light phase is the Ni3Al matrix. 

Figure 2. SEM backscattered image of the sample shown 

in Figure 1. No delamination between the TiC particles 

and the Ni3Al particles is observed. 

a second sintering run was performed after some 

slight tooling changes to accommodate the new 

shrinkage rate. This run produced similar results to 

the first. 

The application testing at Cummins Fuel Systems commenced with parts machined from isostatically pressed rods of the TiC/Ni3Al material. The 

application experiences both sliding and impact 

loading, and 350 hours of abusive testing was conducted between two tests. No observable damage 

was present on the impact surface, and the sliding 

wear surface was pristine, as was the mating steel 

component. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

17

Conclusions

Process optimization work at CoorsTek has progressed well, and CoorsTek is able to create functional test pieces for component evaluation at 

Cummins. Cummins has shown that the material has 

a viable application in its next-generation fuel system. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

18
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

19 

3. MATERIALS FOR EXHAUST AFTERTREATMENT 

A. Materials for Exhaust Aftertreatment 

Paul W. Park, Herbert DaCosta, Svetlana M. Zemskova, Alex Panov, Jill R. Akers, Christie S. Ragle 

and Craig F. 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: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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: Caterpillar Inc., Peoria, Illinois 

Objectives 

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

• Develop a prototype lean-NOx (nitrogen oxides) catalyst system to demonstrate >50% NOx conversion with 

<5% fuel penalty. 

• Evaluate capabilities of commercial and near-commercial NOx sensor technologies to meet future Environmental Protection Agency (EPA) requirement for on-board NOx monitoring. 

• Develop novel NOx sensor materials and concepts functioning at low NOx concentrations (<100 ppm). 

• Evaluate filtration properties of various sintered metal media for application in diesel particulate filters (DPFs). 

Approach 

• Identify a reformer technology to provide desirable reductants for selected lean-NOx catalyst materials, which 

will maximize NOx reduction performance. 

• Develop catalyst durability test protocols for sulfur and phosphorus poisoning. 

• Develop test protocols and evaluate various prototype NOx sensor technologies. 

• Synthesize perovskite-type materials for NOx sensor application. 

• Develop novel NOx sensor materials based on semiconductors functionalized with organic sensitizers. 

• Use our modified test bench and new test protocols to evaluate the capability of various particulate matter trap 

materials. 

• Identify the critical parameters for soot oxidation in catalyzed particulate filters. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

20

Accomplishments 

• Identified a promising reformer technology, “reformer-assisted lean-NOx catalyst system” to achieve high NOx

conversion. 

• Performed steady state engine tests and achieved 44% NOx conversion with 5.6% fuel penalty. 

• Identified the challenges of NOx sensor application, including poor accuracy at low NOx concentrations 

(<200 ppm) and cross-sensitivity to other gas components. 

• Among prototype NOx sensors tested, determined that a mixed-potential NOx sensor showed better performance at low NO concentration below 100 ppm. 

• Fabricated perovskite materials in-house that showed the best sensitivity to NO reported in the literature among 

electrochemical-type NOx sensor materials. 

• Identified semiconductor materials functionalized with organic sensitizers to show pronounced response to low 

NO concentrations (10 ppm). 

• Found a combination of sintered metal materials that provided competitive filtration efficiency and backpressure with substantially lower volume compared with a conventional wall-flow filter material. 

• Developed a test protocol that provides reliable soot loading (amount and type of soot) on cordierite, SiC, and 

sintered metal DPF samples, as well as regenerative capabilities of cordierite filters. 

Future Direction 

• Assess sulfur and phosphorus poisoning behavior on the lean-NOx catalysts. 

• Document intellectual properties resulting from the lean-NOx project and transfer the invention to a strategic 

partner. 

• Study the effect of a “rich” exhaust environment and cross-sensitivity of NH3, H2S, H2 and hydrocarbons on 

NOx sensor technologies. 

• Perform durability and rapid aging tests to assess the lifetime of the selected NOx sensors. 

• Conduct a fundamental study of NOx -sensing mechanism using a computational modeling technique. 

• Catalyze sintered metal materials and evaluate their DPF regeneration features. 

• Investigate the effects of temperature, space velocity, and gas composition on the soot oxidation rates in DPF 

samples. 

• Evaluate particulate-matter regeneration and material durability of various DPF materials. 

Introduction

The objective of this effort is to develop and 

evaluate materials that will be used in aftertreatment 

systems for diesel engine applications to comply 

with future government emission regulations. The 

materials include catalysts for NOx abatement, filtration media for particulate control, and novel materials to improve NOx sensing capabilities in the 

exhaust system. This year’s focus is on 

• development of a prototype reformer-assisted 

lean-NOx catalyst system to improve NOx reduction (>50%) and to minimize fuel penalty (<5%) 

• identification of the critical parameters for soot 

oxidation in a catalyzed particulate filter 

• further evaluation of particulate filtration characteristics of sintered metal materials, evaluation 

of the current state-of-the-art NOx sensor technologies 

• development of novel NOx sensor materials to 

improve sensing capabilities at low NOx concentrations (<100 ppm) 

Approach

Lean NOx. This effort focused on identifying 

new lean- NOx catalyst formulations that effectively 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

21 

utilize selected reductants, including H2, CO, alkane, 

oxygenated hydrocarbons, or mixtures of these. The 

concept of reformer-assisted lean- NOx catalysis has 

been assessed with various reformer technologies to 

identify the right reformer technology that is able to 

provide the desirable reductants to the catalyst materials. A prototype system consisting of a promising 

catalyst and a selected reformer was developed to 

conduct engine tests in order to demonstrate the 

proof of concept. 

NOx sensor. A commercially available NOx sensor does not demonstrate adequate sensitivity and 

selectivity to measure the 0.2 g NOx/bhp*h level 

proposed for 2010 and OBDII regulations. The objective of this project is to develop a promising NOx

sensor technology capable of meeting or exceeding 

EPA’s 2010 standard for OBDII monitoring systems 

for diesel engine exhaust emissions. The approach is 

aimed at thorough evaluation of commercial or 

nearly-commercial electrochemical NOx sensors in 

order to guide sensor developers to further improvements. The development of novel sensor concepts was pursued as an alternative NOx sensor 

technology. 

Particulate matter trap. Our modified test 

bench used to evaluate particulate filter materials 

includes a diesel fuel burner, a chemical reactor, a 

bench of Horiba analyzers, and a scanning mobility 

particle sizer (SMPS). The amount of particulate 

matter captured is determined using the pressure 

change across the sample, while the efficiency of the 

filter is determined using SMPS data. The composition of the exhaust from the burner is measured with 

the Horiba analyzers. The capability of various particulate matter trap materials, including sintered 

metals, was evaluated using the bench system. The 

system can easily be adapted to test other materials, 

such as cordierite and SiC particulate filters. 

Results

Lean NOx

Lean-NOx catalyst material. A synergy effect 

was identified when alumina and zeolite-based materials were combined in a specific order. The combination catalyst significantly improves lean-NOx

performance with a liquid hydrocarbon reductant. 

(Figure 1) The kinetic parameters and the influence 

of fuel components on NOx reduction performance 

over the selected catalyst were obtained using inhouse bench test systems to further understand the 

Temperature (oC)

250 300 350 400 450 500 550 600

NOx Conversion (%)

0

10

20

30

40

50

60

Alumina based catalyst

Zeolite based catalyst

Combination catalyst

Figure 1. Comparison of NO reduction over individual 

catalysts and a combination catalyst tested with a liquid 

hydrocarbon (0.05% NO, C1/ NOx =3, 9% O2, 7% H2O, 

35,000 h−1

). 

efficiency of the catalysts and to provide insight into 

optimization strategies for a reformer technology. 

Engine test. The engine tests were performed 

using two different levels of horsepower (275 and 

500 hp) and two different NOx emission-compliant 

engines (years 1998 and 2004) with catalysts of two 

different sizes (17 and 34 L), respectively. The effects of space velocity, catalyst temperature, reductant species, catalyst formulation, and reductant-toNOx ratios were studied. The results showed 44% 

NOx conversion was achieved with a 5.6% fuel penalty from weighted OICA steady state cycles (nine 

selected modes). The test results showed that the 

reformer significantly enhanced NOx reduction performance and reduced the fuel penalty (Figure 2). 

NOx Sensor

Commercial and nearly-commercial NOx sensor technologies. Evaluation of state-of-the-art NOx 

sensor technology in the presence of 0–1000 ppm 

ammonia showed that the sensor output increases 

with increasing NH3 concentration, while the NOx

analyzer reading (used for control purposes) remained constant (Figure 3). This cross-sensitivity 

substantially limits sensor applications under “fuelrich” environments or in the presence of reductants 

such as NH3 (ammonia/urea selective catalytic reduction), H2 (reformer technology), and CxHy (hydrocarbons selective catalytic reduction). In 

addition, the sensor showed poor resolution at low 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

22

Speed (RPM)

Torque (N.m)

Fuel: 5.7%

NOx: 59.5%

Fuel: 6.4%

NOx: 41.2%

Fuel: 6.0%

NOx: 45.5%

Fuel: 4.9%

NOx: 29.7% Fuel :6.0%

NOx: 50.2%

Fuel: 5.6%

NOx: 57.5%

Fuel: 9.2%

NOx: 40%

Fuel: 6.7%

NOx: 35.2%

Fuel: 8.2%

NOx: 37.3%

Fuel: 7.0%

NOx: 37.8%

Fuel: 7.5%

NOx: 36.2%

Fuel: 6.9%

NOx: 56.8%

Figure 2. NOx reduction performance and fuel penalty 

data obtained from selected OICA steady state cycles 

of a diesel engine equipped with a prototype reformerassisted lean- NOx catalyst system. Black circles: 

without reformer, Shaded circles: with reformer. 

0

200

400

600

800

1000

1200

0 200 400 600 800 1000 NH3 (ppm)

NOx (A.U.)

Figure 3. Output from a commercial sensor (diamonds) and a NOx analyzer (squares) as a function of 

ammonia concentration. 

NOx concentrations. Evaluation of nearlycommercial NOx sensor technologies in low NOx

concentrations has been completed. One sensor has 

been identified as exhibiting a good linear response 

to low NO concentrations, specifically in the range 

of 0–20 ppm. Reproducibility tests showed that the 

sensor response can vary by up to 4.5% and needs 

further improvement 

Perovskite-type NOx sensor materials. Novel 

amperometric-type electrodes on yttria-stabilized 

zirconia (YSZ) as solid electrolyte and La/Fe 

perovskites (ABO3 composition) as chemically interactive material (CIM) were prepared in-house by 

screen-printing and tape-casting techniques 

(Figure 4). Materials were tested at low NO concentrations (<100 ppm) in the presence of 0–20% O2 in 

N2 at 700–800oC. Materials exhibited much higher 

conductivity (up to 400 mA) in comparison with 

other mixed metal oxides described in the literature 

and showed better NOx sensitivity. The CIM cited in 

literature1

 exhibited an amperometric NOx sensor 

output that yielded Δ47 mA for the increment of 

500 ppm NO, while the La/Fe materials prepared in 

our study showed a Δ5–25mA output increase for 

the increment of 15–50 ppm NO at applied voltages 

greater than 0.8V (Figure 5). These are the best results reported in the literature for electrochemical 

type NOx sensor materials. 

Figure 4. Schematic of a novel electrode 

fabrication for a NOx sensor. 

0

50

100

150

200

250

300

350

400

450

0 0.2 0.4 0.6 0.8 1 1.2

Voltage (V)

Current (mA)

10ppm

25ppm

50ppm

100ppm

Figure 5. I(V) response curves of an amperometric 

La/Fe sensor at various NO concentrations (ppm). 

Semiconductors functionalized with organic 

sensitizer NOx sensor materials. Functionalized 

semiconductor materials were evaluated for optical 

detection of NO. Semiconductor-coated quartz sub-
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

23 

strates were provided by a supplier or prepared inhouse by the spray-pyrolysis method using organometallic precursors. The dip-coating method 

was used to prepare sensitized semiconductor materials. The UV-VIS tests at ambient conditions 

showed that pronounced response to 10 and 

100 ppm NO was exhibited. Therefore, a novel NOx

sensor concept employing functionalized semiconductor materials can be developed. This concept offers many advantages over traditional 

electrochemical sensor technologies such as direct 

NO measurement, including no requirement for a 

catalyst or a heater.

Particulate Matter Trap

We have completed and submitted for publication a work in collaboration with Jinyu Zhu and 

Kyeong Lee from Argonne National Laboratory on 

the characterization of particulates produced in our 

diesel fuel burner, focusing on morphology, microstructures, and fractal geometry. This work revealed 

that diesel particulate matter produced using the 

newly developed test protocol shows morphological 

characteristics similar to those of the particulates 

produced in diesel engines. 

The first and second generations of test protocols were developed for soot oxidation studies on 

DPFs that provide consistent soot loading on the 

DPF samples. The protocol was used with our bench 

to evaluate new DPF filtration media based on sintered metal technology. Preliminary results indicated 

that a combination of new sintered metal materials 

yielded competitive filtration efficiency (Figure 6) 

and backpressure (Figure 7) with substantially lower 

volume compared with a commercial wall-flow cordierite DPF tested under the same conditions. Further work is in progress to validate the tests 

described in this report and to apply this approach to 

regenerate particulate matter accumulated on various 

filtration materials.

Conclusions

The research has demonstrated a prototype reformer-assisted lean-NOx system that has the potential to achieve >50% NOx conversion with a <5% 

fuel penalty with the optimization of catalyst materials as well as fuel injection strategy. The comercialization plan has been being developed through 

the collaboration of strategic partners. The current 

0

20

40

60

80

100

120

Filteration Efficiency (%)

Figure 6. Comparison of filtration efficiencies of various 

sintered metal media (SM series) and an un-catalyzed 

wall flow cordierite particulate filter. 

0

1

2

3

4

5

6

7

0 200 400 600

Space Velocity (s-1)

Pressure (kPa)

SM AB

SM B

SM C

SM BB

SM AA

Cordierite

Figure 7. Comparison of backpressure as a function of 

space velocity for various sintered metal media (SM series) and an un-catalyzed wall flow cordierite particulate 

filter. 

electrochemical NOx sensor concept has limitations 

owing to poor accuracy at low NOx concentrations 

and cross-sensitivity to ammonia. The novel NOx

sensor concepts studied in this program showed 

promising results and warrant further development. 

The testing performed with sintered metal filtration 

media indicates that this material has competitive 

filtration efficiency and backpressure compared with 

a traditional wall-flow cordierite filter and provides 

a significant reduction in volume. This material, 

therefore, deserves further testing to confirm these 

results and examine its regenerative properties. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

24

Reference

1. F. Menil et al., “Critical Review of Nitrogen 

Monoxide Sensors for Exhaust Gases of Lean Burn 

Engines,” Sensors and Actuators B67, 1–23, 2000. 

Publications/Presentations

Carrie L. Boyer and Paul W. Park,“Effect of 

SO2 on the Activity over Ag/γ-Al2O3 Catalysts for 

the Reduction of NOx in Lean Conditions,” submitted to Appl. Catal. B. 

Svetlana Zemskova, Paul W. Park, J. S. Lin, J. 

Wen, and I. Petrov, “Study of High Surface Area 

Alumina and Ga-alumina Materials for deNOx Catalyst Applications,” presented at the 28th International Conference and Exposition on Advanced 

Ceramics and Composites, Cocoa Beach, FL, January 25–30, 2004. 

Craig F. Habeger, “Requirements for NOx Sensors in Heavy Duty Diesel Exhaust Environments,” 

presented at the 28th International Conference and 

Exposition on Advanced Ceramics and Composites, 

Cocoa Beach, FL, January 25–30, 2004. 

Jinyu Zhu, Kyeong Lee, Alexander Panov, and 

Jill Akers, “An Investigation of Particulate Morphology, Microstructures, and Fractal Geometry for 

a Diesel Engine-Simulating Combustor,” SAE paper 

2004-01-3044, Society of Automotive Engineers, 

2004.

Patents Issued

Paul W. Park, “Metal/Metal Oxide Doped Oxide 

Catalysts Having High deNOx Selectivity for Lean 

NOx Exhaust Aftertreatment Systems,” U.S. Patent 

6,706,660, March 16, 2004. 

Paul W. Park, “Method of Preparing Doped Oxide Catalysts for Lean-NOx Exhaust,” U.S. Patent 

6,703,343, March 9, 2004. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

25 

B. Development of Materials Analysis Tools for Studying NOx Adsorber 

Catalysts 

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

Chaitanya Narula 

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: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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 process/product interdependence leading to catalyst systems 

with improved final product quality, resulting in diesel emission levels that meet 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 

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

• Obtained rates of desulfation and activation energies for various catalytic compositions after different periods of aging. 

• Continued ex-situ microstructural, microchemical, and crystallographic studies of these new materials in 

simulated engine environments. 

Future Direction 

• Evaluate gradient formation on a macro scale of active elements on a catalyst as a function of catalyst history and operating conditions. 

• Evaluate thermal degradation of a Nox stoichiometric reduction (NSR) catalyst as a function of macroscopic position within the catalyst support brick. 

• Determine the soot and ash distribution as a function of macroscopic position within the NSR catalyst support brick. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

26

Introduction

To meet the 2007 emission requirements for 

diesel exhaust, aftertreatment in diesel engines 

may be necessary. The technology for 2007 will 

need to integrate aftertreatment with engine control systems. Currently, no commercial off-theshelf technologies are available to meet these 

standards. Consequently, Cummins, Inc., is working to understand the basic science necessary to 

effectively use these catalyst systems. ORNL is 

assisting with the materials characterization effort. 

Base-metal oxides (BMOs) are major components in current nitrogen oxides (NOx) adsorber 

catalysts that Cummins seeks to use in aftertreatment systems that contain NOx adsorber catalysts. 

Although the function of these adsorbers is to collect surface nitrite/nitrate (NOx) species, they also 

collect oxy-sulfur (SOx) species. Both species are 

to be released from these surface sites during different regenerations, during which the adsorber 

BMO is either heated to some critical temperature 

and/or exposed to a reducing or reactant atmosphere. Sulfur adsorption is unfortunately a form of 

poisoning of adsorber catalysts and 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, a Cummins catalyst supplier provided 

“core” samples taken from cordierite “bricks” containing the catalyst along the exhaust path for the 

following conditions: fresh (unused), degreened 

(zero desulfation hours), and engine-aged catalysts 

(>zero desulfation hours). Three main variables 

were considered with these samples: platinum (Pt) 

loading (1 or 2 wt %), brick position (1, 2, 3 or 4; 

1 most upstream), and desulfation hours (the number of hours the catalyst system is at the desulfation temperature of ~500°C; actual engine time 

~150% of desulfation time). The X-ray and spectroscopy work on these samples was discussed 

previously.1

 The remainder of the microscopy 

work for these samples is presented in this report. 

New engine-tested brick/core and powder 

samples were examined in FY 2004. These samples consisted of a series of cores with varied 

amounts of Pt loading, a sample that had lost performance after numerous desulfations, and a sample loaded with sulfur. The powder samples 

contained Pt, γ-alumina, and BMOs, such as 

BaCO3. 

Results

X-ray 

XRD and neutron powder diffraction (NPD) 

were employed to aid understanding of the crystalline nature of the samples. For the unfamiliar, the 

following analogy can be applied: As a fingerprint 

identifies a person, so a diffraction pattern identifies a crystalline material. 

Preliminary NPD was attempted with the idea 

that phases containing lighter elements, namely 

the target BMO, could be observed. The washcoat 

components, cordierite (Mg2Al2Si5O18) and possibly Pt, were identified in brick 3, 2 wt % Pt, 1000 

desulfation hours. However, no BMO-containing 

phases were identified/observed. Given the circumstances under which the data were taken, NPD 

should be revisited in the future when a more appropriate and optimized instrument is available. 

Three phases were observed in the E270-3 

samples (composition differs from the sample described in the preceding paragraph) with synchrotron radiation: Pt, γ-alumina (washcoat), and 

cordierite. The samples varied in Pt content and 

were run for months at 500°C, cycled in with hydrogen and no sulfur. XRD indicated that the order 

of decreasing Pt content was E271 (most), E272, 

273, and E270 (least); while the chemisorption 

indicated E270 (most), E271, 272, and E273 (200, 

100, 50, 25 g/ft3

, respectively). The crystallite size 

was effectively ~34 nm for all four samples. Two 

other core samples (”commercial samples”) were 

examined with laboratory XRD: one that lost performance after numerous desulfations (Exp6&7) 

and one that was “loaded” with sulfur (E358-03). 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

27 

Washcoat components, cordierite, and Pt were 

identified in both samples. No sulfur-containing 

compounds were observed. The crystallite size 

was 10 and 4 nm for samples Exp6&7 and E358-

03, respectively, which is in agreement with the 

observed loss in catalytic performance. Comparison with prior data1

 suggests that sample Exp6&7 

has undergone the equivalent of 150 hours of 

desulfation at 500°C. 

A set of FRESH powders was examined with 

high-temperature XRD. Samples E362a, E363a, 

and E364a contain a series of BMOs, respectively, 

which are reported to be stabilizers for Pt. Pt and 

γ-alumina were identified by XRD in all of these 

samples. The high-temperature XRD data for all 

samples showed Pt coalescence/sintering occurring at temperatures in excess of 450°C (e.g., Figure 1), as evidenced by the breadth narrowing and 

intensity increase of the (111) and (200) Pt peaks. 

The high-temperature XRD data for sample E364a 

(FRESH, BaCO3) showed the BaCO3 decomposing/reacting with the Al2O3 to form a spinel, 

BaAl2O3 (see Figure 2). 

Figure 1. The diffraction patterns for E364a as a function of temperature. 

Spectroscopy 

Spectroscopic techniques provide important 

information about the chemical state of and identification of adsorbed surface species. Two complementary spectroscopic techniques are Raman 

spectroscopy and X-ray photoelectron spectroscopy (XPS), which provide molecular and lowfrequency vibrational information, respectively.

Figure 2. The diffraction pattern for E364a after exposure to elevated temperatures. 

Raman 

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. In 

this case, frequency changes of a washcoat component vibration were monitored with desulfation 

time. Previously, we discovered a decrease in peak 

width and the increase in peak position of a washcoat-attributed Raman band, which occurs as a 

result of the growth of the Pt particle size. Although Raman spectroscopy could not directly 

measure the Pt particle size in these samples, the 

changes in the Raman spectra with Pt size were 

calibrated for this specific catalyst formulation so 

that Raman could nondestructively measure Pt 

particle coarsening. The origin of the change in 

Raman spectrum was determined using molecular 

simulations, as shown in Figure 3. When a Pt atom 

is adsorbed onto a washcoat particle, the Pt will 

change the electron charge density by forming a 

metallic bond with one of the washcoat atoms. 

This in turn increases the force constants of the 

washcoat components (see Table 1). Since the 

bonds are getting stronger, they vibrate at a higher 

frequency, which explains why the Raman bands 

increase in frequency with Pt. 

Raman spectroscopy is also very sensitive to 

adsorbed sulfur on the support. The vibration of 

the SO4

−2

 species has a band at ~960 cm−1. Using a 

high-temperature stage, the amount of sulfur present on a BaO support was monitored in situ in a 

reducing atmosphere. By monitoring the relative 

peak heights of BaSO4, BaCO3, and BaO with 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

28

Figure 3. The change in electron charge distribution 

that originates from chemical bonding between Pt and 

the washcoat material. Regions where electrons are 

withdrawn (blue), and where the charge density is increased (red). The predicted change in force constant 

describes the observed shift in the Raman spectrum. 

Table 1. The force constants of a washcoat component 

before and after Pt adsorption. The K value (in 

mdyne/Å) generally increases, causing the Raman 

peaks to shift to higher frequency with Pt concentration 

K1(host) K2(w.Pt) Site 

0.032 0.086 O1 x 

0.093 0.112 O1 z 

0.033 0.038 WC x 

0.077 0.113 WC z 

0.003 0.005 O2

 x 

0.017 0.009 O2

 z 

time at different temperatures, the activation energies for sulfation were directly measured. Table 2 

shows activation energies for sulfation with and 

without Pt present and with and without oxygen 

present. The activation energies for both BaO and 

BaCO3 were measured simultaneously since these 

two phases had peaks in different positions. Three 

main observations were taken from these results. 

First, Pt reduces the activation energy in all cases, 

which is expected. Second, the presence of oxygen 

reduces the activation energy for sulfation. This 

suggests that O2 participates in the formation of 

BaSO4 by donating an oxygen atom by the following equations: 

Table 2. The activation energy of sulfation (kcal/mole) 

in four conditions, with and without Pt and O2. The 

presence of Pt reduces the activation energy, as does an 

oxidizing environment

O2 present? Pt present? BaO BaCO3

No No 17.8 25.4 

No Yes 11.0 17.9 

Yes No 11.6 15.9 

Yes Yes 7.4 13.1 

BaCO3 + SO2 + 1⁄2 O2 → BaSO4 + CO2 

and 

BaO + SO2 + 1⁄2 O2 → BaSO4. 

However, this means that the reaction can proceed only with O2 present. Since we have measured sulfation in an anaerobic environment, we 

believe that the extra oxygen must therefore come 

from the bulk lattice of the BaO and BaSO4. Third, 

the sulfation of BaO is easier than that of BaCO3. 

These results show the power of Raman spectroscopy for monitoring surface reactions in situ on 

these catalyst materials. By changing the gaseous 

environment, we can infer what chemical reactions 

are taking place and then take steps to improve the 

catalyst system based on this new information. 

X-ray Photoelectron 

XPS was used to analyze chemistry along the 

length (~5 cm) of individual “troughs” or channels 

of several “as run” brick/core samples. Figure 4 

shows the S 2p, Cl 2p, and Pt 4f core level intensities as a function of position along the catalyst 

trough of sample E358-03. The sulfur concentration was highest on one end and fell to near-zero 

intensity by the midpoint of the core. As the sulfur 

intensity diminished, the chlorine signal grew, 

reaching a maximum of >1 at. %. Measurements 

of the Pt 4f core level show a slight decrease at the 

high-sulfur end, but the overall signal intensity is 

weak, making the significance of this decrease 

uncertain. In contrast, the results for sample 

Exp6&7 showed significant sulfur at only one end 

of the core, at the position of the first measurement 

(see Figure 5). Figure 5 also shows that the chlorine intensity was higher in this sample overall and 

showed an increase from its lowest value at the 

high-sulfur end of ~2 at. % to a maximum of 5–6 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

29 

Figure 4. S 2p, Cl 2p, and Pt 4f signals as a function of 

position along a single catalyst trough/channel of sample E358-03. These signal intensities, along with other 

minor and major component intensities, were used to 

calculate the at % of each constituent. Measurements 

were made every ~5 mm, as indicated on the photo. 

at. % at the opposite end. As with sample E358-

03, the Pt showed little variation from one end to 

the other. A small but significant amount of iron, 

~8 at. %, was detected on this sample at the sulfurend of the core. This rapidly decreased to ~2 at. % 

by the third measurement location and remained 

near this level for all other measurements. Taken 

together, these measurements demonstrated that 

this measurement technique was helpful in understanding sulfur distribution. The value of this 

technique was that the measurements were done 

on real monolith core samples, not samples that 

were handled or treated specially for ultra-high 

vacuum surface analytical techniques. 

Four brick/core engine-aged samples were 

desulfated at temperatures of 450, 500, 550 and 

600°C and examined using XPS. Figure 6 shows 

the S 2p core level intensity as a function of position along an individual catalyst channel/ trough 

for each. There is a definite trend going from 

450 to 550 to 600°C. All samples show similar 

(1.5–2 at. %) and high sulfur at one end versus the 

other. Each also shows a drop from this initial 

Figure 5. S 2p, Cl 2p, Fe 2p and Pt 4f signals as a function of position along a single trough/channel of sample 

Exp6&7. These signal intensities, along with other minor and major component intensities, were used to calculate the atom % of each constituent. Measurements 

were made every ~5 mm, as indicated on the photo. 

Figure 6. Sulfur concentration, as determined by XPS 

measurements, as a function of position along individual catalyst troughs for samples desulfated at 450, 500, 

550, and 600°C. Measurements of the S 2p signal intensity were made every 3–5 mm. 

value along the first centimeter of the core. After 

the initial decrease, all samples show an increase. 

The length over which the increase occurs varies 

from 450°C (~1.5 cm) to 550°C (~1 cm) to 600°C 

(~0.5 cm). The sulfur concentration then de-
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

30

creases, reaching near-zero levels by the midpoint 

of the core. There also appears to be a second, 

smaller rise for the 450 and 550°C samples at 

~3 cm. The 500°C sample showed a similar initial 

high sulfur level (in fact, the highest of all the 

samples) but then exhibited a general decrease 

over the first half of the core before reaching nearzero levels. Two separate trials on two different 

(but nearly adjacent) troughs showed the 500°C 

was self-consistent. Further studies should be attempted to determine the reason that this sample 

did not follow the trend of the others. No significant trends for other minor components were 

noted. Chlorine levels for the four samples were 

generally low and did not correlate with the sulfur 

signal as noted above for sample E358-03. This 

study showed that tracking the sulfur on real catalyst cores was feasible and showed trends versus 

desulfation temperature. 

Microscopy

Transmission electron microscopy (TEM), 

scanning transmission electron microscopy 

(STEM), and electron probe microanalysis 

(EPMA) techniques were used to continue characterization of the microstructure of the NOx trap 

catalysts. In the previous report, the microstructure 

and chemistry of samples with 1 wt % Pt loading 

at different positions along the reactor were studied; and the changes in morphology of Pt particles, 

as well as the distribution of elements such as the 

trapping component within the washcoat, were 

elucidated. In subsequent work, a series of samples with 2 wt % Pt loading were studied in like 

fashion. Some example results and main conclusions from the microscopy and microprobe experiments are given below. 

In contrast to the 1 wt % Pt loading results, 

high-resolution bright-field TEM images of the 

2 wt % Pt fresh catalyst did in fact clearly reveal 

discrete Pt catalyst particles (see Figure 7). Similar 

results were obtained with high-angle annular 

dark-field (HA-ADF) images taken in the Hitachi 

HD-2000 STEM instrument. The de-greened samples showed discrete fine particles of Pt, on the 

order of 5 nm in diameter, using both TEM and 

STEM methods. HA-ADF images on the STEM 

provided a facile method for imaging the particles 

in high contrast and for measuring particle sizes. 

These measurements were made on all observed 

Figure 7. Bright-field TEM micrograph of fresh 2 wt % 

Pt catalyst sample with discrete Pt particles evident in 

brick 2, FRESH.\ 

samples (except the fresh specimen). Samples 

from the bricks with 2 wt % Pt loading with 150, 

500, and 1000 desulfation hours displayed 8-, 

10.5-, and 11.5-nm particles, respectively. Figure 8 shows the HA-ADF image of brick 2 with 

150 desulfation hours, for example. The results 

were nearly identical, with respect to average particle size, to those obtained from the 1 wt % Pt 

samples, with only a slight increase of about 1 nm 

in the diameters of particles in the 500 and 1000 

desulfation hour samples. 

Figure 8. Z-contrast STEM image illustrating dispersion and size of 2 wt % Pt particles in brick 2 after 150 

hours of deS. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

31 

EPMA was used to determine the bulk distributions and quantitative measurements of the Pt 

trapping component and sulfur species by acquiring X-ray counts from small regions from the 

outer surface of the washcoat layer into the cordierite framework. These measurements mapped Pt 

at the 2 wt % level, as expected. The trapping 

component initially was present also at about the 

0.7–0.8 wt % level, and sulfur in all the 2 wt % Pt 

samples was less than 0.1 wt % in the bulk of the 

washcoat; but it seemed to concentrate in thin layers at the cordierite-washcoat interface. Unlike the 

1 wt % Pt samples, where Pt showed concentrations primarily at free surfaces within and external 

to the two washcoat layers, Pt was very uniformly 

dispersed within the washcoats of the 2 wt % Pt 

samples (see Figure 9). The differences are likely 

due to different rates of drying employed in deposition of the washcoats on the cordierite support. 

In the fresh samples, the trapping component was 

sometimes observed to be strongly concentrated at 

the interface between the inner washcoat and the 

cordierite framework material. In most samples, 

however, the trapping component showed a uniform concentration, with occasional high concentrations in very small, widely scattered regions. 

The trapping component concentration remained very uniform throughout the washcoat 

with engine aging, but the overall concentration 

decreased by 150 desulfation hours to about the 

0.5 wt % level. As with 1 wt % Pt samples, this 

result is important because it shows the trapping 

component that is available for NOx adsorption, 

whereas the XRD data showed only total trapping 

component. Thus the decrease in the total amount 

of trapping component coupled with the increase 

in Pt particle size could suggest a reduction in the 

total number and surface area of Pt-trapping component contact sites upon aging. This would also 

explain the reduced performance of the NOx trap. 

Energy-dispersion spectroscopy spectra in the 

TEM showed small trapping component peaks, but 

no discrete trapping component phase was detected. Platinum particle sizes at brick positions 1 

and 4 in 2 wt % Pt loaded samples were similar to 

those of brick 2, indicating that the particle growth 

behavior was similar regardless of sample position 

in the reactor. 

Figure 9. Two EPMA Pt X-ray maps demonstrating 

differences in dispersion of Pt between the (a) 1 wt % 

Pt and (b) 2 wt % Pt in both bricks after 150 hours of 

deS. Pt is uniformly distributed in the 2 wt % sample; in 

the 1 wt % sample, it is higher at the surface and free 

surfaces throughout the washcoat than in the bulk. 

References

1. Thomas Watkins, Larry Allard, Doug 

Blom, Michael Lance, Chaitanya Narula, and 

Harry Meyer, “Development of Materials Analysis 

Tools for Studying NOx Adsorber Catalysts,” 

Heavy Vehicle Propulsion Materials 2003 Annual 

Progress Report, Energy Efficiency and Renew-
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

32

able Energy Office of FreedomCAR and Vehicle 

Technologies, March 2004. 

Publications

H. L. Fang and M. J. Lance, “Influence of 

Soot Surface Changes on DPF Regeneration,” accepted for publication in J. Soc. Auto. Eng. (2004).

J. Parks, A. Watson, B. Epling, L. Campbell, 

R. England, N. Currier, A. Yezerets, H. Fang, 

T. Yonushonis, T. Gallant, T. Watkins, D. Blom, 

L. Allard, M. Lance, H. Meyer and C. Narula, 

“NOx Absorber Catalyst Durability: Light- and 

Heavy-Duty Perspectives,” pp. 304–311 in Proceedings of 2003 DEER Conference: 9th Diesel 

Engine Emission Reduction Conference, August 24–28, 2003, Newport, RI.

Presentations

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

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

H. Fang, and T. Yonushonis, “Development of 

Materials Analysis Tools for Studying NOx Adsorber Catalysts,” presented at NLCat 2003, 4th 

DOE National Laboratory Catalysis Conference, 

Oak Ridge, Tennessee, October 23, 2003. 

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

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

H. Fang, and T. Yonushonis, “Development of 

Materials Analysis Tools for Studying NOx Adsorber Catalysts,” presented at the 53th Annual 

Denver Conference, Steamboat Springs, Colorado, 

August 2, 2004. 

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

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

H. Fang, and T. Yonushonis, “Development of 

Materials Analysis Tools for Studying NOx Adsorber Catalysts,” presented by D. R. Johnson at 

the 10th Diesel Engine Emission Reduction 

(DEER) Conference, Coronado, California, August 30, 2004. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

33 

C. Development of NOx Sensors for Heavy Vehicle Applications 

David L. West, Fred C. Montgomery, Timothy R. Armstrong, 

Jaco H. Visser* and David J. Kubinski* 

Oak Ridge National Laboratory 

P. O. Box 2008, Mail Stop 6063 

Oak Ridge, TN 37831-6063 

(865) 574-5377; fax (865) 574-6918; blaupj@ornl.gov 

*Scientific Research Laboratory, Ford Motor Co. 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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, TN 

Prime DOE Contract Number DE-AC05-00OR22725; 

Subcontractor: Ford Motor Company, Detroit, Michigan 

Objectives 

• Develop nitrogen oxides (NOx) sensors for remediation and monitoring of diesel engine exhausts. 

⎯ Sensors should have an operating temperature of 500–700°C and be able to measure NOx concentrations 

from ∼1 to 1500 ppm at oxygen levels from 5 to 20 vol %. 

⎯ Since NOx is a mixture of NO and NO2, it would be useful to have sensors selective for either NO or NO2

or able to measure “total NOx” ([NO] + [NO2]).

Approach 

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

[typically yttria stabilized zirconia (YSZ)]. 

• Operate elements either in a non-Nernstian mode (where the output is a voltage) or under dc electrical bias. 

• Characterize elements for NOx response, sensitivity to varying O2, and recovery/response kinetics. 

Accomplishments 

• Demonstrated NO2 sensing elements with extremely high sensitivity (50 mV/decade at 700°C, 7 vol % O2 over 

the concentration range 20–200 ppmV NO2). 

• Demonstrated biased NO-selective sensing elements. 

• Constructed both types of elements to consist of co-planar electrodes, one oxide electrode and the other a noble 

metal (usually platinum). 

Future Direction 

• Determine the effects of H2O on sensing element performance and stability. (Diesel exhausts are rich in H2O, 

typically 4–15 vol %). 

• Characterize cross-sensitivity to other species often found in diesel exhausts (CO, hydrocarbons, CO2). 

• Evaluate long-term stability of these sensing elements in atmospheres simulating diesel exhaust. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

34

• Investigate use of low-frequency ac biasing, especially the effect on drift of the baseline sensor response 

compared with dc-biased elements. 

Introduction

The primary pollutants (excluding the 

greenhouse gas CO2) from the combustion of lowsulfur fuels are CO, hydrocarbons, and NOx. 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 the current absence of a “lean-NOx” catalyst, 

three technologies have been proposed to meet the 

challenge of NOx remediation of diesel exhausts: 

lean-NOx 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). 

This project is developing NOx sensing elements 

that can be used in diesel engine exhausts. The 

minimum requirement for this application is an 

operating temperature of 500–700°C 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 that two of the following three 

concentrations be measured: NO, NO2, or NOx (= 

NO + NO2). This project is a cooperative research 

and development agreement with Ford Motor 

Company. 

Approach

Some sample geometries for prototype sensing 

elements are shown in Figure 1. The YSZ substrate 

is manufactured in-house using Tosoh TZ8YS YSZ 

powder, and the electrodes (oxide and platinum) are 

screen-printed and fired. Inks for the platinum 

electrode are typically obtained commercially; for 

the oxide electrode, inks are formulated in-house 

Figure 1. Schematic semicircular (a) and interdigitated 

(b) sensing element geometries.

using proprietary methods and materials. The 

precise composition of the oxide electrode is 

dictated by the target application of the sensing 

element. For non-Nernstian NO2 sensing elements, a 

binary transition metal oxide electrode is sufficient. 

[The term “non-Nernstian” is used because in the 

presence of 10–100 ppmv of NO2, these elements 

can produce large dc output voltages (on the order of 

0.1 V), which far exceed the predictions of the 

Nernst equation.] For sensing elements that will be 

biased to yield NO-selective behavior, a more 

complex oxide electrode is required. 

The sensing elements depicted schematically in 

Figure 1 can be operated in either non-Nernstian or 

biased mode. In non-Nernstian mode, the voltage 

between the oxide and platinum electrodes is 

monitored. In the biased mode, either a constant dc 

current or voltage is imposed across the electrode 

pair. The measured voltage (current biased) or 

current (voltage biased) is then monitored and is 

useful as a sensing signal for NOx. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

35 

Prototype sensing elements are evaluated both 

microstructurally and for sensing performance. 

Isothermal sensing performance is characterized by 

exposing the sensing elements to varying 

concentrations of NOx at a fixed O2 level or varying 

O2 levels at fixed concentrations of NOx. Because 

NOx is a mixture primarily of NO and NO2, in our 

characterization we typically expose the sensing 

element to either NO or NO2. 

Results

Representative electrode and substrate 

microstructures are shown in Figure 2. The screenprinted and fired (1200°C) ZnO electrode is 

relatively open and porous compared with the YSZ 

substrate. The typical electrode thickness after firing 

is 20–30 μm (determined by non-contact 

profilometry); this compares well with the predicted 

thickness given the mesh size and emulsion 

thickness of the screen used for printing the 

electrodes. 

Figure 2. Plan view of a ZnO electrode on a YSZ 

substrate. Secondary electron image at 5 kV.

As Figure 3 shows, when paired with platinum 

in the geometry shown in Figure 1a, the binary 

oxides ZnO, Cr2O3, and NiO have yielded excellent 

non-Nernstian NO2 sensing elements. For example, 

a ZnO/Pt sensing element, operating at 700°C in 7 

vol % O2, displayed an NO2 sensitivity of ∼50 

mV/decade over the concentration range 20–200 

Figure 3. Response trace of an NiO/Pt sensing element 

as the input NO2 is varied at 700°C in 7 vol % O2 (a) 

and the variation in measured voltage with input NOx

for different binary oxide/Pt sensing elements in 7 vol 

% O2 (b).

ppmV NO2, with 20 ppmV NO2 giving rise to a 

voltage of approximately 40 mV. 

Diesel exhaust varies in O2, typically ranging 

from 5 to 20 vol %.1

 Therefore, it is important to 

characterize the O2 dependence of the sensing 

element response. Figure 4 shows that for these nonNernstian NO2 sensors, the measured voltage is a 

function of O2 only in the presence of NO2. We 

currently believe that this is observed because the 

voltage developed in the presence of NO2 is related 

to the different behavior of the oxide and platinum 

electrodes with respect to the reduction of NO2 and 

oxidation of oxygen ions (O2−

) in the YSZ solid 

electrolyte: 

 NO2(g) + 2e− → NO(g) + O2− , and (1) 

 O2− → 1/2O2(g) + 2e−

 .(2) 

The simultaneous occurrence of these reactions 

on both the oxide and platinum electrodes gives rise 

to a mixed potential on each electrode, the precise 

value of which will depend on the rate of Eqs. (1) 

and (2) on each electrode, the absorptive behavior of 

NO2 and O2 on each electrode, and the catalytic 

behavior of each electrode.2

 In the absence of NO2, 

(Figure 4b), there is no gaseous species readily 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

36

Figure 4. Measured voltage as O2 is varied between 7 

and 20 vol % with 190 ppmV of NO2 (a) and 0 ppmV

NO2 (b) for an NiO/Pt sensing element operated in 

non-Nernstian mode.

available for reduction, and a mixed potential cannot 

develop. 

The dominant equilibrium NOx species above 

500°C is NO.3

 Therefore, although we cannot 

assume that the NO/NO2 ratio in diesel exhaust is at 

equilibrium, it seems likely that NO would be more 

prevalent in the exhaust at the desired operating 

temperatures of 500–700°C. Thus sensor design may 

be simplified if sensing elements sensitive to NO 

could be developed. This can be done with the 

application of bias, as described below. 

Equations (1) and (2) describe electrochemical 

reduction of NO2 and oxidation of O2−

, respectively. 

A similar pair of equations can be written for the 

electrochemical oxidation of NO (to NO2) and 

reduction of O2 (to O2−

). Since these are 

electrochemical phenomena, it may be possible to 

adjust the rates of the various reactions by applying 

electrical bias to the sensing element (as is done in 

non-Faradaic electrochemical modification of 

catalytic activity).4

 To demonstrate this with biased 

sensing elements, a dc current bias was stepped at 

discrete levels (e.g., −16, −8, 0, 8, 16 μA for 5 min 

each), and a brief (2-min) 450-ppmV pulse of either 

NO or NO2 was applied at each bias level. The result 

is shown in Figure 5, where it can be seen that at 

negative biases (oxide electrode biased negatively 

Figure 5. Voltage traces as 2-minute pulses of 450 

ppmV NOx are applied at different current biases. Data 

collected at 600°C in 7 vol % O2.

with respect to the platinum electrode), the change 

in voltage induced by NO2 is much larger than that 

induced by NO. However, as the bias is increased to 

positive values, the changes induced by NO are 

larger relative to those due to NO2, and at a certain 

positive bias (∼+16 μA for the sensing element in 

Figure 5) the introduction of 450 ppmV NO gives a 

large signal while the same amount of NO2 gives 

only a small response. 

This asymmetry (stronger response to NO than 

NO2) to was displayed over a wide concentration 

range (∼50–1500 ppmV NOx) as shown in Figure 6. 

Although the mechanism of the enhanced NO 

response with bias is still under investigation, it 

seems clear that the oxide electrode must promote 

the oxidation of NO in order for the effect to be 

seen. Under conditions of positive bias, for current 

to flow through the YSZ electrolyte, oxygen ions 

must flow through the YSZ to the oxide electrode 

(the anode), and the external electrical circuit must 

supply electrons to the platinum electrode (the 

cathode). Therefore, it would appear that the 

electrochemical reactions occurring in the presence 

of NO are the reverse of Eq. (1) (oxidation of NO) at 

the oxide electrode and the reverse of Eq. (2) 

(reduction of O2) at the platinum electrode. 

Less clear is the mechanism underlying the 

weak response to NO2 seen in Figure 6. One 

possible explanation is that the oxide electrode 

strongly promotes the reduction of NO2 via Eq. (1), 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

37 

Figure 6. Response (600°C, 7 vol % O2) of a currentbiased sensing element to input NO and NO2. The 

concentration variations of the two NOx species were 

nominally identical.

and this interferes with the expected anodic reaction 

[Eq. (2)] at the oxide electrode with positive bias 

and only NO2 and O2 present. Lending credence to 

this is the observation that at sufficiently large 

positive biases, as seen in Figure 5, the introduction 

of either NO or NO2 decreases the measured voltage. 

The O2 dependence of these biased sensing 

elements is quite different from the reaction of the 

non-Nernstian elements discussed earlier. As 

Figure 7 shows, these elements show an O2

dependence that is a decreasing function of NO. In 

the absence of NO, the measured voltage of a 

current-biased element can vary ∼20% as O2 is 

varied from 7 to 20 vol % (Figure 7b). However, in 

the presence of 1500 ppmV NO, variation of O2 over 

this range has little effect (Figure 7a). This behavior 

points to the presence of NO as having an 

anomalously strong effect on the electrode/ 

electrolyte interfacial resistance. 

An important concern as sensor development 

progresses is the role of electrode geometry, as rigid 

space constraints will apply to the final sensor 

design. As an initial step in addressing these 

concerns, semicircular and interdigitated sensing 

Figure 7. Variation of current-biased sensing element 

signal (at 600°C) with O2 with different input NO (a) 

and response trace measured while varying O2 with no 

input NOx (b).

elements (Figure 1) have been fabricated and 

characterized. Figure 8 indicates that the geometry 

variation had little impact on the magnitude and 

functional form of the NO response. It is also seen in 

Figure 8 that for these biased sensing elements, an 

almost linear response is obtained at low input NO, 

and that the extent of this linear regime increased 

with the element operating temperature. This is in 

contrast to the logarithmic behavior exhibited by the 

non-Nernstian elements in Figure 3. The linear 

behavior at low NOx, combined with the observation 

that the NO-selectivity of the biased elements 

seemed to decrease with NOx (Figure 6), indicates 

that careful sensor design will be necessary to use 

these biased sensing elements for low NOx levels. 

Finally, temperature was much more important than 

electrode geometry in governing the recovery time 

from NO exposure, as shown in Figure 9. 

Conclusions

Progress in NOx sensor development has 

proceeded along two main fronts. In the first 

approach, non-Nernstian sensing elements using 

binary oxides and platinum have been developed. 

These elements respond strongly to NO2 and 

produce an output voltage that is proportional to the 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

38

Figure 8. Response to input NO for biased sensing 

elements with the two different geometries 

(semicircular (SC) and interdigitated (ID)) of Figure 1. 

Data collected in 7 vol % O2.

Figure 9. Recovery from NO exposure for biased 

sensing elements with the two different geometries of 

Figure 1, semicircular (SC) and interdigitated (ID).

logarithm of NO2. Use of a more complex oxide 

electrode and a dc electrical bias enables selective 

detection of NO, the stable form of NOx at high 

temperature. Pending evaluation of selectivity and 

stability, these two types of sensing elements are 

candidates for incorporation into NOx sensors for 

heavy vehicle applications. 

Publications

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

Armstrong, “Electrically Biased NOx Sensing 

Elements with Co-Planar Electrodes,” in 205th 

Meeting of the Electrochemical Society, San 

Antonio, TX, May 2004. 

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

Armstrong, “NO-Selective NOx Sensing Elements 

for Combustion Exhausts,” accepted by Eurosensors 

XVIII, Rome, Italy, September 2004. 

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, Long 

Beach, CA, November 2004. 

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

Armstrong, “Electrode Materials for MixedPotential NOx Sensors,” in Proceedings of the 28th 

International Cocoa Beach Conference on Advanced 

Ceramics and Composites, Cocoa Beach, FL, 

January 2004. 

Patent

D. L. West, F. C. Montgomery, C. Maxey and T. 

R. Armstrong, “Sensing Devices Having CoPlanar 

Electrodes,” filed with the U.S. Patent Office. 

References

1. J. Kaspar, P. Fornasiero, and N. Hickey, 

“Automotive Catalytic Converters: Current Status 

and Some Perspectives," Catalysis Today 77, 419–

49 (2003). 

2. A. A. Arias de Velasco, P. T. Moseley, R. 

Peat, and J. G. Peláez, “Atmosphere-dependent 

Potentials at Oxide Interfaces,” Sensors and 

Actuators B, 15–16, 55–62 (1993). 

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

Control,” in Air Pollution Control Technology 

Handbook, CRC Press, Boca Raton, 2001. 

4. C. G. Vayenas, S. Bebekis, I. V. Yentekakis, 

and S. Neophytids, “Non-Faradaic Electrochemical 

Modification of Catalytic Activity: The Work 

Function of Metal Electrodes in Solid Electrolyte 

Cells,” Solid State Ionics 53–56, 97–110 (1992). 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

39

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 

Oak Ridge, TN 37831-6064 

(865) 574-4981; fax: (865) 576-5413; e-mail: allardlfjr@ornl.gov 

*Dept. of Materials Science and Engineering, Lawrence Berkeley National Laboratory, Berkeley, CA 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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 

• 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 reducing emissions of nitrogen oxides (NOx) 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). 

Approach 

• Use new field emission microscope at the High Temperature Materials Laboratory to image catalyst particles at 

near-atomic level to characterize systematic series of model NOx trap catalyst materials (composed of 2% platinum by weight on a mixed CeO2-ZrO2-La2O3-BaO-Al2O3 support material) and to develop techniques applicable for imaging future catalysts at sub-Ångstrom levels with the new Aberration-Corrected Electron Microscope 

(ACEM). 

• Conduct studies of a series of NOx trap catalyst specimens, provided by Ford Research Laboratory colleague 

Dr. G. Graham, that had been field-tested by running in vehicles with spark-ignition direct-injection (SIDI) engines, rather than by simulating exhaust compositions and cycles using bench-top reactors. 

Accomplishments 

• Characterized via high-resolution annular dark-field scanning TEM (STEM) imaging the structure of a model 

NOx trap catalyst composed of 2% platinum by weight on a mixed CeO2-ZrO2-La2O3-BaO-Al2O3 support material. The imaging showed, for example, that initial dispersions of platinum species consisted of near-atomic 

clusters and “rafts” of platinum atoms one to 3 atomic layers thick. 

• Characterized the progression of changes in the microstructure of NOx trap catalysts as a result of vehicle aging, 

by TEM and STEM analysis of core samples taken from the centers of catalyst monoliths removed from the vehicles after 30,000, 53,000 and 82,000 km of driving.

• Successfully factory tested and subsequently accepted delivery of the ACEM, funded by the Office of Energy 

Efficiency and Renewable Energy. Achieved an Oak Ridge National Laboratory (ORNL) PEP milestone for 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

40

beneficial operation of all critical imaging systems and demonstrated the capabilities of the new microscope for 

imaging ultra-fine clusters of catalyst species on real catalyst samples of OFCVT interest. 

Future Direction 

• Use new high-resolution electron microscopy imaging techniques and energy-loss electron spectroscopy on the 

ACEM to characterize diesel particulate NOx –reduction and lean NOx trap samples in collaboration with Ford 

Research Laboratory colleagues. 

• Develop new thrust capability for characterization of catalyst structures at the atomic level using the ACEM, 

using techniques of electron tomography. This will involve constructing a new specimen holder to allow high 

tilts (up to ±75 degrees), and assessing the ability of annular dark-field STEM imaging to provide unambiguous 

information on the shapes of fine catalyst particles and the shape changes that occur on aging. 

Technical Progress

Microstructure of Pt/Mixed-Oxide Model 

Catalysts 

A model NOx trap catalyst composed of 2% 

platinum by weight on a mixed CeO2-ZrO2-La2O3-

BaO-Al2O3 support material (or more formally, 

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

98%BaO.6Al2O3)]) is being studied (in a separate 

task) to determine the effects of exposure to lean–

rich exhaust mixtures on the microstructure of the 

catalyst. In the present task, the model catalyst material was characterized using the high-resolution annular dark-field imaging capability of the JEOL 

2010F field emission TEM on loan as part of the 

new ACEM project. This effort served as a precursor to imaging of the ultrastructure of our catalyst 

samples using similar (but aberration-corrected) 

STEM imaging techniques on the ACEM. 

The STEM/TEM instrument provides both 

bright-field TEM and dark-field STEM imaging capabilities, and typical results at the resolution level 

of the 2010F allowed the imaging of “rafts” of platinum species on the alumina/baria and ceria/zirconia 

phases of the experimental catalysts (e.g., Figure 1). 

However, the significantly higher resolution expected from the ACEM should permit unambiguous 

and more quantitative characterization of the exact 

structure of the as-dispersed precious metal species 

on both experimental and “real” catalyst samples. 

This will provide a better understanding of the initial 

stages of growth of catalyst particles under aging 

conditions, so that advances can be made in controlling heavy metal particle size changes with exposure 

to temperature and exhaust gas conditions. 

Microstructure of SIDI Vehicle-Tested 

Catalysts 

In collaboration with Dr. George Graham of 

Ford Research Laboratory, we characterized a series 

of NOx trap catalyst specimens that were field-tested 

by running in vehicles with SIDI engines. These 

catalysts comprised a double washcoat configuration 

on standard cordierite brick monoliths. The initial 

washcoat is composed of a baria-alumina composition on which platinum is dispersed at the 1 wt % 

level as the heavy-metal catalytic species. The second washcoat is composed of ceria-zirconia with 

rhodium dispersed at about the 0.2 wt % level. Catalyst core samples were taken from the centers of 

catalyst monoliths removed from the vehicles after 

30,000, 53,000, and 82,000 km of driving, and their 

microstructures were compared with the microstructure of a fresh catalyst sample that had seen no vehicle aging. 

The technique of “spectrum imaging” was used 

extensively to obtain quantitative information on the 

microstructural and chemistry changes on aging. 

Figure 1. (a) Alumina phase with platinum rafts (arrows). 

(b) Ce-Zr phase with platinum rafts (arrows). Possible 

edge-on rafts are also shown (arrows labeled “e”). Scale = 

5 nm
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

41

With this technique, an annular dark-field image is 

acquired from a chosen sample area. Then the same 

sample area is scanned point by point over a long 

time (up to hours, with drift correction implemented 

during the scan), with an energy-dispersive spectroscopy (EDS) spectrum acquired and saved at each 

point, to make the so-called spectrum image. In the 

present work, spectrum images of 512 × 384 pixel 

count (about 20 K spectra) were acquired, providing 

the ability to perform “principal component analysis” (PCA) with the data processing capabilities of 

the acquisition software. It is a way of identifying 

patterns in data and expressing the data in such a 

way as to highlight their similarities and differences. 

Patterns in data can be hard to find in data of high 

dimensions (i.e., a large number of points, each with 

an EDS spectrum), and PCA is a powerful tool for 

analyzing such data (details are beyond the scope of 

this report). A straightforward application is to extract “maps” of elements and phases from the spectrum image. 

For example, an annular dark-field image from 

the 30,000-km catalyst sample is shown in Figure 2a. The EDS spectrum represented by the entire 

area of the image is shown in Figure 2c. The majority component is alumina, with cerium and zirconium also clearly present. The platinum component 

is essentially in the noise at this full scale, but the 

platinum L family is clearly present, as indicated in 

Figure 2d, with no overlaps of other elements. When 

a window is applied to show only the counts from 

the platinum L family, as shown in Figure 2b, the 

one-to-one correspondence between the brightcontrast particles in the annular dark-field image and 

the high-density platinum concentrations in the 

platinum image (as marked by arrows) clearly confirm the location and size distribution of platinum in 

the specimen area. 

A similar result from the 82,000 km sample is 

shown in Figure 3a, which shows platinum particle 

growth as a result of aging phenomena during the 

extended driving cycle. A benefit of the spectrum 

imaging process is that the presence of any minor 

components in an image area can be determined. In 

the case of these SIDI catalysts, rhodium is supposed to be present at about 10% of the level of the 

platinum (i.e., at the 0.1 wt % level), but in the images acquired to date, no rhodium was identified (all 

the bright contrast particles in any image acquired 

corresponded to platinum). It is likely that either 

Figure 3. (a) Annular dark-field image of catalyst area in 

82,000-km SIDI sample. (b) Platinum map extracted from 

spectrum image. Note the ability to locate the platinum 

particles unambiguously, even in the presence of a large 

bright-contrast aggregate of Ce-Zr in the center of the 

image (the increased contrast in this area in the platinum 

map is due to an increase in the background from the 

high-atomic-number material).

rhodium is atomically dispersed even after significant aging in the test vehicles, or it is present in 

widely isolated particles that will require further 

analysis to locate. Data such as these are providing 

a significant input to the understanding of the effects 

of vehicle aging of NOx trap catalysts. 

Ultra-High Resolution Imaging of Catalyst Clusters Using the ACEM

ORNL’s ACEM (Figure 4) was delivered in 

May 2004 and installed in the newly constructed 

Advanced Materials Characterization Laboratory. 

The instrument is one of the first two combination 

STEM/TEM instruments in the country that are 

Figure 2. (a) Annular dark-field image of 30,000-km SIDI 

catalyst. ( b) Platinum map extracted from sum spectrum 

of Fig. 2c. (c) Sum spectrum from spectrum image collection. (d) Expanded sum spectrum showing clear platinum 

peaks. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

42

equipped with a special corrector element on the 

illumination to enable a probe size of 0.7Å (compared with a 1.4-Å probe in the uncorrected instrument). This increase in probe resolution results in a 

10-fold improvement in contrast in STEM dark-field 

imaging mode and should enable resolution of heavy 

metal species such as platinum on light oxide support materials at the single-atom level. An early result from the ACEM is the annular dark-field image 

shown in Figure 5, which is from a sample of asdeposited platinum on alumina. The platinum particles show in bright contrast; the two large particles 

are in the 2.5–4 nm size range (and are the only two 

particles visible unambiguously in standard brightfield images). The finest spots may be single atoms 

of platinum, or at most clusters of two or three atoms. The image illustrates the advanced capability 

the ACEM will provide for future catalyst characterization. 

Analysis of catalyst materials at ultra-high resolution provides the potential to characterize directly 

the shapes of nanometer-sized particles and their 

changes with aging. The ultimate goal is to obtain a 

full understanding of the mechanisms of catalyst 

activity, selectivity, and aging. In parallel work to 

the development of advanced imaging techniques in 

the laboratory, we are collaborating with Lawrence 

Berkeley National Laboratory in a computational 

component of the work that involves modeling cataFigure 5. Example of an early very high-resolution 

annular dark-field image of platinum species on alumina support. The finest bright spots represent platinum clusters of a few atoms, with perhaps single atoms 

also seen.

lyst particles and calculating the images expected 

from them (conventional TEM images have been 

simulated so far, with STEM images to be done in 

the future). It is of great interest to determine, with 

very small particles comprising only a relatively few 

atoms, what crystal facets are exposed at the surface 

from the initial fine particle dispersions to larger 

particles as the catalyst ages. A typical model is 

shown in Figure 6, along with the corresponding 

computed TEM image and a chosen line profile that 

illustrates that, in the ideal case, there is a one-toone correspondence between image intensity and the 

number of atoms in a column. This suggests that real 

atomic structure images of a particle taken from different directions (i.e., in a tomography application) 

might allow unambiguous characterization of the 

exact particle shape, with the specific crystal planes 

exposed on the surface thereby analyzed. A new research thrust area is being developed to implement 

the capability and techniques for nanoparticle shape 

determination with sub-Ångstrom imaging in the 

ACEM. 

Conclusions

The acquisition and imminent beneficial operation of the ACEM will provide OFCVT programs 

with the top capability in the nation for characterization of catalyst materials and other materials of inFigure 4. ORNL’s ACEM, the JEOL 2200FS-AC, during 

final assembly and test-out. The instrument will ultimately be operated from the adjacent control room (seen 

through window behind the instrument).
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

43

Figure 6. Example of model catalyst particle imaging 

computation. A cuboctohedron particle geometry is 

shown, with the image of the particle looking down the 

plane of a-a’. A line trace thru the image shows a direct 

relationship between atom column intensity and the number of atoms in each column.

terest to the program. Imaging at the single-atom 

level of heavy catalyst species on oxide substrates—

coupled with electron spectroscopy to obtain, for 

example, chemical binding information—will add a 

new dimension to our understanding of the behavior 

of these complex materials. The present work using 

ORNL’s current TEM and STEM instruments, with 

techniques such as annular dark-field imaging and 

especially spectrum imaging, has paved the way for 

rapid implementation of the next level of imaging 

and analysis capabilities. Both model NOx trap catalysts and catalysts run in vehicles for known times 

have been characterized. Changes in platinum particle sizes and the relationship between the various 

oxide support species upon aging have supplied information on the mechanisms of catalyst degradation. Finally, initial work on modeling of catalyst 

particle shapes and the resulting images has shown 

promise for allowing the determination of particle 

morphologies (exposed crystal facets) with highresolution electron tomography using the ACEM. 

Publications/Presentations

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

Kiely, M. Watanabe, M. Kawasaki, T. Kaneyama, 

and H. Sawada, “First Results from the AberrationCorrected JEOL 2200FS-AC STEM/TEM, pp. 110–

111 in Proceeding M&M 2004, San Francisco Press, 

San Francisco. 

M. A. O’Keefe and L. F. Allard, “A Standard 

for Sub-Ångstrom Metrology of Resolution in Aberration-Corrected Electron Microscopes.” Microscopy & Microanalysis 10 1002–1003 (2004). 

M. A. O’Keefe and L. F. Allard, “Sub-Ångstrom 

Microscopy for Materials Science,” abstract for 

Nanometrology Workshop, National Institute of 

Standards and Technology, Gaithersburg, MD, May 

2004. 

M. A. O’Keefe and L. F. Allard, “Sub-Ångstrom 

Electron Microscopy for Sub-Ångstrom NanoMetrology,” pp. 52–53 in Proceedings of National 

Nanotechnology Initiative Grand Challenge Workshop on Instrumentation and Metrology for 

Nanotechnology, National Institute of Standards and 

Technology, Gaithersburg, MD, 2004. 

L. F. Allard and M. A. O‚Keefe, “Sub-Angstrom 

Metrology of Resolution in Aberration-Corrected 

Transmission Electron Microscopes using the A-OK 

Standard Test,” p. 9 in Proceedings of the 1st Aberration-Corrected MSA Meeting, 2004. location of 

meeting? 

M. A. O‚Keefe, L. F. Allard, and K. H. Downing, “Atomic-Resolution Tomography,” presented at 

the Center for Nanophase Materials Sciences Workshop on Microscopy, Metrology and Manipulations 

using Electrons, Ions, and Photons, Oak Ridge, Tennessee, September 15–16, 2004. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

44
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

45 

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

Conditions at High Temperatures 

C. K. Narula, K. Lester, H. Meyer, L. Allard 

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.gov 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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 

• Facilitate deployment of a nitrogen oxides (NOx) trap for lean diesel or gasoline exhaust by 

⎯ investigating materials issues related to deterioration of the performance of NOx traps upon aging as a result of thermal and sulfation-desulfation cycles 

⎯ investigating materials that are robust under the lean NOx trap operating conditions 

Approach 

• Synthesize new model catalyst systems containing platinum (Pt): Pt/BaO.6Al2O3, [10%CeO2-ZrO2,

90%(2%La2O3, 98%BaO.6Al2O3)], and 2%Pt, 5%MnO2, 93%[10%CeO2-ZrO2, 90%(2%La2O3, 98% 

BaO.6Al2O3)]. 

• Treat them under thermal, lean-diesel, and rich-diesel conditions at 500oC for 4 h and monitor changes in microstructure. The next step involves evaluation under lean-rich cycles and aging under laboratory aging cycles. 

• Evaluate the performance of the selected systems on a bench-top flow reactor. 

Accomplishments 

• Completed the microstructural evaluation of model catalysts Pt/BaO.6Al2O3, [10%CeO2-ZrO2, 90%(2%La2O3, 

98%BaO.6Al2O3)] and 2%Pt, 5%MnO2, 93%[10%CeO2-ZrO2, 90%(2%La2O3, 98% BaO.6Al2O3)]. The latter is 

a better representative of this class of NOx trap systems. 

⎯ Completed and reproduced the study of microstructural changes on aging under diesel lean and rich conditions at 500oC. 

⎯ Completed microstructural changes under thermal aging studies of these samples for comparison. 

⎯ Completed microstructural changes in thermally aged samples after extended exposure to lean and rich 

conditions. 

• Began updating the ex-situ reactor to enhance its capabilities to enable treatment of transmission electron microscopy (TEM) samples under lean, rich, or, stoichiometric conditions, as well as lean-rich cycles. This reactor 

will be capable of rapid screening of the model catalyst powders for their durability in exhaust under the operating conditions for diesel and gasoline engines. 

• Equipped the synthesis laboratory to enable preparation of NOx trap materials. The bench-top-flow reactor is 

being set up. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

46

Future Direction 

• Evaluate model catalyst systems on a bench-top flow reactor for their efficiency and durability as NOx traps under simulated diesel and gasoline conditions [with or without sulfur oxides (SOx) in the simulated exhaust]. 

These results will provide a benchmark for evaluation of new materials. 

• Carry out microstructural characterization of model NOx trap powders after aging under lean, rich, or 

stoichiometric conditions on the bench-top flow reactor and TEM samples on an ex-situ microreactor. 

• Investigate new materials that can withstand NOx trap operating conditions without detrimental structural 

changes. 

Introduction

NOx traps are at the forefront of various strategies under investigation to treat NOx from diesel and 

lean-gasoline engines.1

 NOx traps collect engine-out 

NOx during lean operation and treat it during short 

rich-operation cycles.2

 Fresh NOx traps work very 

well but cannot sustain their high efficiency over the 

lifetime of a vehicle. The efficiency loss is believed 

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

by the SOx in the emissions from the oxidation of 

sulfur in fuel. 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. 

NOx traps are derived from commercial threeway catalysts installed to treat emissions from engines operating at stoichiometric air-fuel ratios. 

Therefore, the basic components of NOx traps are 

identical to three-way catalysts. The advanced version of a three-way catalyst is a two-layer system on 

a honeycomb substrate with the inner layer based on 

Pt-alumina and the outer layer on rhodium-ceriazirconia. The NOx traps derived from advanced 

three-way catalysts are identical to the catalysts except for a high baria content (the upper limit being 

close to 20%) in the alumina layer. Thus aging can 

lead to intermixing of layers, crystallization of bariacontaining phases that are not good NOx absorbers, 

and sintering of precious metals. The first goal of 

this project is to determine if one or all of the microstructural changes take place and if these changes 

occur 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 lean 

conditions to determine the species formed during lean cycles. 

• Complete microstructural characterization of 

fresh NOx trap materials after exposure to rich 

conditions to determine the species formed during rich cycles. 

The second goal of the project is to investigate 

and design new 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 synthesizing new model 

catalyst systems; treating them under thermal, lean, 

and rich conditions; and monitoring microstructural 

changes to understand the deterioration mechanism. 

In order to rapidly screen model catalyst systems, 

we have designed an ex-situ reactor system that is 

being further upgraded to enhance its capabilities. 

The updated system will operate in the 25–1000°C 

range and will have the ability to introduce periodic 

rich pulses while the catalyst is being treated under 

lean or stoichiometric conditions. 

The information from this study will be used to 

design thermally durable catalyst systems and will 

be tested on a bench-top flow reactor system. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

47 

Results

Over the last 2 years, we determined the microstructural changes that occurred in a supplier lean 

NOx trap system (based on Pt/BaO-Al2O3 and CeO2-

ZrO2 materials) upon being aged on (1) a pulsator at 

Ford Motor Company, (2) a dynamometer at Ford, 

and (3) vehicles with gasoline direct-injection, 

spark-ignition (DISI) engines in Europe. 

• After pulsator aging, lean and rich aged samples 

showed that the sintering of Pt particles occurs 

during aging, and barium migrates into the 

ceria-zirconia layer. Both of these factors reduce 

the Pt-barium oxide surface area where NOx adsorption and reduction take place during lean 

and rich cycles, respectively. The stoichiometric 

aging also leads to the migration of barium into 

the ceria-zirconia layer, but the sintering of Pt is 

less severe. 

• The dyno-aged samples showed extensive sintering of Pt and its migration into the ceriazirconia layer. The sintering of rhodium was 

also observed. The migration of barium into 

ceria-zirconia and the sintering of the precious 

metal component could explain the deterioration 

of performance 

• The analysis of the samples evaluated on-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 the cause of performance deterioration in NOx trap materials in 

gasoline engine exhaust conditions. While the laboratory aging protocols for diesel engines are in development, we have compiled data showing earlystage changes in model NOx trap materials under 

diesel conditions. We assumed that lowertemperature simulated diesel exhaust would slow 

down the aging compared with gasoline engine exhaust. Thus we completed a series of thermal aging, 

diesel lean aging, and diesel rich aging cycles of a 

model NOx trap and a model NOx trap modified with 

manganese oxide, as well as Pt/Al2O3, for comparison purposes. The results are summarized in the following paragraphs. 

Synthesis of New Model Catalysts

Pt/Al2O3: The Pt/alumina was prepared by the 

impregnation method using H2PtCl6 salt. X-ray 

powder diffraction (XRD) shows broad diffraction 

peaks for alumina and no peaks due to Pt, as expected because of the small particle size of Pt in the 

freshly prepared sample (Figure 1). 

Figure 1. X-ray diffraction patterns of Pt/Al2O3.

2%Pt, 5%MnO2, 93%[10%CeO2-ZrO2, 

90%(2%La2O3, 98% BaO.6Al2O3)]: The model 

catalyst was synthesized by impregnating 2%Pt, 

97%[10%CeO2-ZrO2, 90%(2%La2O3, 98% 

BaO.6Al2O3)] with manganese nitrate and subsequent decomposition. 

Microstructural Changes in Model

Catalysts 

The microstructural changes in model catalysts 

(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)]) after thermal treatment at 500, 600, 

700, 800, and 900oC; lean treatment at 500oC for 

4 h; and rich treatment at 500oC were followed with 

XRD and transmission electron microscopy (TEM). 

Here lean and rich treatment refers to exposure of 

sample in an ex-situ reactor (Figure 7a, described 

previously) to diesel lean and diesel rich conditions. 

A summary of changes in the precious metal (Pt) is 

presented in Table 1. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

48 

Table 1. Summary of Pt particles size changes under various treatment conditions in model catalysts I, II, 

and III 

Sample Treatment conditions 

 Fresh Thermal treatment 

(XRD of powders) 

Thermal 

treatment 

at 500°C 

on ex-situ 

reactor 

Treatment 

under lean 

conditions 

on ex-situ 

reactor 

Treatment 

under rich 

conditions on 

ex-situ reactor 

Pt/Al2O3 1–1.2 nm 600oC, 3.4 nm 

700oC, 17.1nm 

800oC, 26.1nm 

900oC, 39.5nm 

2–4 nm 8–10nm 2–4 nm 

Model NOx trap catalyst 

I: 2%Pt-98%[10%CeO2-

ZrO2-90%(2%La2O3-

98%BaO.6Al2O3)] 

1–2nm 600oC, 2.6 nm 

700oC, 21.3nm 

800oC, 37.2nm 

900oC, 48.4nm 

1–2 nm 1–2nm 1.5–3.5 nm 

Model NOx trap catalyst 

II: 2%Pt, 5%MnO2–

93%[10%CeO2-ZrO2-

90%(2%La2O3-98% 

BaO.6Al2O3)] 

1–2 nm 700oC, 20.7nm 

800oC, 27.0nm 

900oC, 34.0nm 

1–2 nm 2–3nm 1–2nm 

Thermal Treatment

• Thermal treatment of samples was carried out at 

500, 600, 700, 800, and 900oC in a bench-top 

furnace in air. 

• Pt/Al2O3: In XRD patterns (Figure 1) of 

Pt/Al2O3, diffraction peaks due to Pt particle size 

start to grow in 500–900oC range. The calculated Pt particle sizes (using the Scherrer formula, assumes spherical particles) are 3.4, 17.1, 

26.1, and 39.5 nm at 600, 700, 800, and 900oC, 

respectively. At 500oC, the alumina is a mixture 

of γ-alumina and hydrated γ–alumina. At 800oC, 

a broad peak assignable to the alpha phase can 

be seen. At 900oC, the diffraction peaks due to 

α– and γ – alumina are observed. TEM of the 

fresh sample shows that the Pt particle size is in 

the 1–1.2 nm range (Figure 2). This particle size 

grows to 2–4 nm on thermal treatment in a furnace at 500°C for 4 h. 

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

98%BaO.6Al2O3)]: In XRD patterns (Figure 3), 

diffraction peaks due to Pt particle size start to 

grow in the 500–900oC range. The calculated Pt 

particle sizes (using the Scherrer formula, assumes spherical particles) are 2.6, 21.3, 37.2, 

and 48.4 nm at 600, 700, 800, and 900oC, respectively. TEM of the fresh sample shows that 

the Pt particle size is ~1 nm (Figure 4). This 

Figure 2. TEM of fresh and thermally aged (500oC, 4 h) 

Pt/Al2O3. 

Figure 3. XRD patterns of 2%Pt-98%[10%CeO2-ZrO2-

90%(2%La2O3-98%BaO.6Al2O3)].
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

49 

Figure 4. TEM of fresh and thermally aged (500oC, 4 h) 

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

98%BaO.6Al2O3)].

particle size grows to 1–2 nm upon thermal 

treatment in a furnace at 500oC for 4 h. 

• 2%Pt, 5%MnO2, 93%[10%CeO2-ZrO2, 

90%(2%La2O3, 98% BaO.6Al2O3)]: In XRD 

patterns (Figure 5), diffraction peaks due to Pt 

particle size start to grow in the 500–900oC 

range. The calculated Pt particle sizes (using the 

Scherrer formula, assumes spherical particles) 

are 20.7, 27, and 34 nm at 700, 800, and 900oC 

respectively. TEM of the fresh sample shows 

that the Pt particle size is ~1 nm (Figure 6). This 

particle size grows to 1–2 nm upon thermal 

treatment in a furnace at 500oC for 4 h. 

Figure 5. XRD patterns of 2%Pt, 5%MnO2–93% 

[10%CeO2-ZrO2-90%(2%La2O3-98% BaO.6Al2O3)].

Lean and Rich Treatment in Ex-situ

Reactor

The aging of model catalysts was carried out in 

our ex-situ reactor system (schematics in Figure 7) 

under lean or rich diesel exhaust conditions at 500°C 

for 4 h [CO, CO2, H2, HC, NOx, H2O, flow rate 

100cc/min] showed some sintering of Pt particles, 

Figure 6. TEM of fresh and thermally aged (500oC, 4 h) 

2%Pt, 5%MnO2–93%[10%CeO2-ZrO2-90%(2%La2O3-

98% BaO.6Al2O3)]. 

Figure 7. a. Schematics of current ex-situ reactor; 

b. Schematics of modified ex-situ reactor. 

and the size grew to ~2.5nm. However, the sintering 

was significantly less than that observed upon thermal aging (~30–40nm) at 500°C in air. 

It is important to note that the experiments in the 

ex-situ reactor system do not precisely duplicate the 

aging of NOx traps during normal operation. The 

alternating lean/rich conditions are not present in our 

experiments; instead, our samples are exposed to 

lean conditions, and a lack of control of space velocity due to the very small amount of catalyst leads to 

NOx absorber saturation conditions. The saturation 

condition persists throughout our treatment. 

Lean Treatment 

• Pt/Al2O3: TEM of the fresh sample shows that 

the Pt particle size is in the 1–1.2 nm range (Figure 8). This particle size grows to 8–10 nm in 

diesel lean simulated exhaust in the ex-situ reactor at 500oC for 4 h. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

50 

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

98%BaO.6Al2O3)]: TEM of the fresh sample 

shows that the Pt particle size is ~1 nm (Figure 9). This particle size grows to 1–2 nm upon 

lean treatment in the ex-sit reactor at 500oC for 

4 h. 

• 2%Pt, 5%MnO2, 93%[10%CeO2-ZrO2, 

90%(2%La2O3, 98% BaO.6Al2O3)]: TEM of the 

fresh sample shows that the Pt particle size is 

~1 nm (Figure 10). This particle size grows to 

2–3 nm in diesel lean simulated exhaust in the 

ex-situ reactor at 500oC for 4 h. 

Figure 8. TEM of fresh and lean aged (500oC, 4 h) 

Pt/Al2O3.

Figure 9. TEM of fresh and lean aged (500oC, 4 h) 2%Pt98%[10%CeO2-ZrO2-90%(2%La2O3-98%BaO.6Al2O3)]. 

Figure 10. TEM of fresh and lean aged (500oC, 4 h) 

2%Pt, 5%MnO2–93%[10%CeO2-ZrO2-90%(2%La2O3-

98% BaO.6Al2O3)]. 

Rich Treatment 

• Pt/Al2O3: TEM of the fresh sample shows that 

the Pt particle size is in the 1–1.2 nm range 

(Figure 11). This particle size grows to 2–4 nm 

in diesel rich simulated exhaust in the ex-situ reactor at 500oC for 4 h. 

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

98%BaO.6Al2O3)]: TEM of the fresh sample 

shows that the Pt particle size is ~1–1.2 nm 

(Figure 12). This particle size grows to 1.5–

3.5 nm in diesel rich simulated exhaust in the 

ex-situ reactor at 500oC for 4 h. 

• 2%Pt, 5%MnO2, 93%[10%CeO2-ZrO2, 

90%(2%La2O3, 98% BaO.6Al2O3)]: TEM of the 

fresh sample shows that the Pt particle size is 

~1 nm (Figure 13). This particle size grows to 

1–2 nm in diesel rich simulated exhaust in the 

ex-situ reactor at 500oC for 4 h. 

Figure 11. TEM off fresh and rich aged (500oC, 4 h) 

Pt/Al2O3. 

Figure 12. TEM of fresh and rich aged (500oC, 4 h) 

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

98%BaO.6Al2O3)]. 

Figure 13. TEM of fresh and lean aged (500oC, 4 h) 

2%Pt, 5%MnO2–93%[10%CeO2-ZrO2-90%(2%La2O3-

98% BaO.6Al2O3)].

Modifications in Ex-situ Reactor

The ex-situ reaction is under modification to include the ability to cycle through lean and rich conditions. This involves the addition of a gas-flow 

controller and installation of solenoid valves that can 

be controlled through a computer using the Labview 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

51 

software system. The installation of plumbing is 

complete, and the new ex-situ reactor should be in 

operation shortly. The schematics of the modified 

reactor are shown in Figure 7b. 

Surface Studies of Fresh and Aged Model 

NOx Trap Catalyst

The results of X-ray photoelectron spectroscopy 

analysis (XPS) of 2%Pt, 98%[10%CeO2-ZrO2, 

90%(2%La2O3, 98%BaO.6Al2O3)] are presented in 

Table 2. Powder samples of the fresh, thermally 

aged (500oC in air for 4h), and lean simulated exhaust-exposed (after thermal aging) catalysts were 

pressed onto an indium foil and then analyzed. The 

high concentration of Pt and lathanum in fresh samples is due to the migration of these elements to the 

surface. Thermally aged samples probably achieve a 

redistribution of elements, leading to a reduced concentration of these elements on the surface. When a 

thermally aged sample is exposed to simulated lean 

NOx conditions for about 4 h at 500oC, the surface 

composition does not change significantly. The lack 

of a significant increase in nitrogen or carbon after 

extended exposure to lean diesel exhaust is puzzling. 

Further experiments are in progress to determine if it 

is just a surface phenomenon, or if the bulk material 

loses nitrate during our cooling protocol. 

Conclusions

Our results from the study of microstructural 

changes in a supplier lean NOx trap system (based 

on Pt/BaO-Al2O3 and CeO2-ZrO2 materials) upon 

aging on (1) a pulsator at Ford, (2) a dynamometer 

at Ford, and (3) vehicles in gasoline DISI engines in 

Europe show that precious metal sintering and barium migration occurs. This could be a cause of performance deterioration in NOx trap materials. 

We have successfully employed an ex-situ reactor to study aging in model NOx trap catalysts. This 

system has allowed us to study changes in samples 

in a selected area under thermal, lean diesel, or rich 

diesel conditions. We have updated the ex-situ reactor to a lean–rich cycle NOx trap system. This protocol will allow us to rapidly screen catalyst systems 

without having to do start-and-stop experiments to 

understand aging. 

Our next goal is to complete the bench-flow reactor and start evaluating new model catalyst systems for their performance and relate 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 LenaBurn Engines,” SAE paper 95-0809, Society of 

Automotive Engineers.. 

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 Catalyst for 

Automotive Lean-Burn Engine: NOx Storage and 

Reduction Catalyst,” Catalysis Today 27, 63–69 

(1996). 

Table 2. XPS analysis of 2%Pt-98%[10%CeO2-ZrO2-90%(2%La2O3-98%BaO.6Al2O3)] 

Element Fresh Thermally aged at 500oC in 

air for 4 hours 

Aging under simulated lean 

diesel exhaust at 500oC for 4 

hours 

 Atomic % Wt % Atomic % Wt % Atomic % Wt % 

O 57.8 33.5 59.3 34.9 61.7 37.7 

Al 26.6 26.0 27.6 27.4 26 26.8 

C 6.3 2.7 3.7 1.6 4.1 1.9 

Ba 3.9 19.4 5.9 29.8 5.1 26.7 

La 2.2 11.0 0.5 2.5 0.6 3.2 

N 0.3 0.2 0.2 0.1 0 0 

Pt 1 7.1 0.5 3.6 0.5 3.7 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

52 

Presentations

C. K. Narula, “Recent Advances in Lean NOx

Treatment in Automotive Exhaust,” invited seminar 

at Purdue University, West Lafayette, Indiana, April 

22, 2004. 

C. K. Narula, “Applied Catalysis Program,” a 

presentation of results to the Science and Technology Technical Review Panel, National Transportation Research Center, Oak Ridge, TN, 2004. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

53 

F. Aftertreatment Catalysts Materials Research 

Matthew Henrichsen and Randy Stafford 

Cummins, Inc. 

1900 McKinley Ave. 

Columbus, IN 47203 

(812) 377-418-0428; e-mail: matthew.henrichsen@cummins.com 

ORNL Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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 the 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 nitrogen oxides (NOx) reduction catalysts 

through reactor studies. 

Accomplishments 

• Evaluated the kinetics of desulfation on a commercial NOx reduction catalyst. 

• Demonstrated that microwave heating results in catalyzed soot oxidation rates substantially similar to those 

from conventional heating. 

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 mandate 

significantly lower emissions than the current regulatory standards. These regulations are seen as aftertreatment-forcing 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 in capturing and destroying 

NOx emitted from diesel engines. Even though sulfur levels will be reduced in future diesel fuels, any 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

54

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 an extended 

time at high temperature. If the desulfation process 

requires 10 minutes 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 minutes. 

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 catalysts more than conventional 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

An Emerachem NOx adsorber catalyst was selected for evaluation of sulfation and desulfation 

behavior. The impacts of desulfation temperature, 

reductant concentration, duration of exposure to reductant, and sulfur loading on desulfation rates and 

desorbed sulfur species were measured. The behavior of other NOx adsorber catalysts will be similar to 

the behavior of this catalyst, although the temperature required for desulfation will likely be higher. 

Using the microwave thermal gravimetric analysis (TGA) method developed under this program at 

Notre Dame University, microwave and conventional TGA was applied to several catalysts. Catalysts were selected based on other literature reports 

of accelerated activity in microwave heating and the 

principal investigators’ knowledge of microwave 

heating and soot oxidation. 

Results 

Sulfation and Desulfation 

A proprietary Emerachem catalyst was applied 

to a 3-in.-long 400-cpsi NGK extruded cordierite 

catalyst substrate. A core drill was used to separate 

0.6-in.-diameter cores from the catalyzed substrate 

for use in a bench flow reactor. All cores were prepared for desulfation experiments using the degreening cycle detailed in Table 1. Before sulfation/ 

desulfation testing, the performance of each catalyst 

core was evaluated at the conditions described in 

Table 2. 

Table 1. Experimental conditions for degreening

Parameter Lean Rich 

Time (sec) 300 300 

Space velocity (1/h) 30,000 30,000 

Temperature (°C) 575 575 

NO (ppm) 300 0 

SO2 (ppm) 0 0 

CO (%) 0 2.5 

H2O (%) 8 8 

CO2 (%) 8 8 

O2 (%) 8 0 

N2 Balance Balance 

Table 2. Experimental conditions for catalyst characterization before sulfation/desulfation.

Parameter Lean Rich 

Time (sec) 56 4 

Space velocity (1/h) 30,000 30,000 

Temperature (°C) Varied Varied 

NO (ppm) 300 0 

SO2 (ppm) 0 0 

CO (%) 0 2.5 

H2O (%) 8 8 

CO2 (%) 8 8 

O2 (%) 8 0 

N2 Balance Balance 

For evaluation of temperature effects, a fresh 

core was used for each desulfation temperature. 

Conditions during sulfation and desulfation are 

listed in Table 3. It is important to note that no oxygen was included in the desulfation gas stream. This 

both aided in controlling the temperature of the 

cores during desulfation and eliminated the need to 

inject extra reductant to remove the oxygen. 

The total sulfur release is shown in Figure 1. As 

shown in Figure 2, even a small portion of residual 

sulfur will adversely impact subsequent NOx reduction performance. High-efficiency NOx removal at 

lower temperatures in Figure 2 was due to the mildness of the NOx performance cycle rather than to a 

lack of impact due to sulfur poisoning. 

Table 3. Experimental conditions for sulfation/desulfation.

Parameter Sulfation Desulfation 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

55 

Time (min) 37 15 

Space velocity (1/h) 30,000 30,000 

Temperature (°C) 550 Varied 

NO (ppm) 0 0 

SO2 (ppm) 40 0 

CO (%) 0 4.0 

H2O (%) 8 8 

CO2 (%) 8 8 

O2 (%) 8 0 

N2 Balance Balance 

Figure 1. Total sulfur released during desulfation as a 

function of desulfation temperature. 

Figure 2. NOx performance after desulfation at various 

temperatures. 

Similar degreening cycles and test methods were 

used in preparing fresh cores for evaluation of the 

impact of reductant concentration. The impact of 

reductant concentration on total sulfur release is 

shown in Figure 3. All cores were fresh except for 

the 0.25% CO core, which had been previously used 

for the 4% CO data. Temperature was 550°C for 

each of these experiments. The decrease in sulfur 

release with increasing reductant at low concentration was likely due to CO poisoning of the oxidation 

Figure 3. Total sulfur released during desulfation as a 

function of CO concentration at 550°C. 

sites on the platinum catalyst. The high-temperature 

NOx reduction performance of cores desulfated with 

0.5% and 2% CO was similar, and both were better 

than the core desulfated at 1% CO. 

The impact on desulfation of the sulfur quantity 

loaded on the catalyst was investigated at 550°C and 

at 2% CO. As total sulfur on the catalyst increased, 

sulfur released during desulfation increased nearly 

proportionally. Desulfation time effects were investigated at 550°C with 2% CO. As desulfation time 

increased, sulfur release increased and the release 

rate decreased (see Figure 4). 

Figure 4. Total sulfur released during desulfation as a 

function of desulfation time at 550°C and 2% CO. 

Microwave-heated Catalytic Soot 

Oxidation

As discussed in the previous annual report, a 

microwave TGA was constructed and calibrated so 

that microwave effects on catalyzed oxidation of 

soot could be investigated. Calibration to conventional TGA results was done with a microwave susceptor supporting the soot specimen so that heating 

of the soot was primarily though conduction and 

convection (see Figure 5). The presence of the susceptor allowed heating of the specimen without 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

56

Figure 5. Comparison of microwave and conventional 

TGA results for uncatalyzed soot oxidation. 

significant thermal gradients; but in the absence of a 

susceptor, thermal gradients were present. 

Figure 6 shows the result of heating a mixture of 

CuO and carbon with microwaves on a quartz (nonsuscepting) specimen support. Each specimen in the 

photograph was heated for a longer time than the 

specimen immediately to its left. The specimen on 

the left was not heated. The specimen in the middle 

has a mottled color, indicating that carbon has been 

oxidized. Note that the lighter colors are present toward the center of the pile of powder in the basket. 

(CuO alone is a red-brown color.) The specimen on 

the right is entirely red-brown in the middle of the 

pile, while the edges remain black. This indicates 

that the temperature gradients measured through infrared thermography (Figure 7) are indicative of 

where the soot is burning. 

Figure 6. Mixtures of CuO and soot heated to various 

temperatures in microwave TGA. 

It is not surprising that the edges of the specimens are cooler because of the higher surface area 

there relative to the mass available to absorb microwaves. Because of this effect, the fraction of soot 

burned (when no susceptor other than the soot is 

Figure 7. Thermography of specimen during heating in 

microwave TGA. 

present) is a function of the soot present at the beginning of the experiment. (See Figure 8.) 

Figure 8. Soot oxidation vs. time for various soot loads in 

microwave TGA.

Despite this drawback of the microwave TGA 

system, it could be used for evaluating the usefulness of potential microwave catalysts because large 

effects were being looked for. Although a decrease 

in soot oxidation temperature of as little as 25 or 

50°C would be useful in a real system, the added 

cost and complexity of microwave heating would be 

justified only by a much larger decrease in soot ignition temperature. 

Figure 9 shows TGA results for mixtures of soot 

with CuO and LiFe5O8. Both oxides slightly 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

57 

Figure 9. Catalyzed soot oxidation with microwave and 

conventional heating. 

catalyze soot oxidation in conventional and microwave heating. It is important to note that in Figure 8, 

surface temperatures from microwave TGA are being compared with bulk temperatures from conventionally heated TGA. (Bulk temperatures in 

microwave heating can be substantially higher than 

surface temperatures.) The best point to compare the 

microwave and conventional results is at the beginning of oxidation. This best represents the temperature required for ignition. 

While both CuO and LiFe5O8 lower soot oxidation temperatures relative to uncatalyzed oxidation, 

the difference between microwave and conventional 

temperatures required for ignition is less than 50°C. 

Taking into consideration the surface temperature 

measurement and temperature gradients likely present in the microwave TGA, these data provide no 

reasonable basis to believe that microwave heating 

significantly improves catalytic oxidation of soot. 

Summary and Conclusions

An examination of the kinetics of desulfation on 

a commercial NOx adsorber catalyst was begun. The 

information will be useful for designing desulfation 

control strategies. The impact of microwave heating 

on soot oxidation was investigated. Based on measurements made with microwave and conventionally 

heated TGA, it is not likely that microwave heating 

significantly lowers the temperature of catalyzed 

soot oxidation with the catalysts examined in this 

work. 

Publications

Hongmei An, Caitlin Kilroy, and Paul J. 

McGinn, “An Examination of Microwave Heating to 

Enhance Diesel Soot Combustion,” submitted to 

Applied Catalysis B, Environmental. 

J. Kong, M. Henrichsen, and A. J. Shih, “Infrared Thermometry Measurement of Temperature Distribution in The Microwave Regeneration of Diesel 

Particulate Filters,” International Journal of Engine 

Research, in press. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

58
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

59 

G. Durability of Diesel Particulate Filters 

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 

Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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 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 to 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 

• Completed the stress analysis of DPF beams subjected to flexural loading. The importance of accounting for 

the cellular structure of the component has been identified. 

• Obtained the elastic properties of DPFs using a dynamic mechanical analyzer. 

Future Direction 

• Organize a round-robin test program to assess the precision of four-point bending test methods to determine the 

flexural strength of DPF ceramic substrates. 

• Assess the susceptibility of DPF ceramic substrates to slow-crack growth using dynamic fatigue and doubletorsion test methods. 

• Implement probabilistic design tools to predict the reliability and durability of DPF ceramic substrates subjected to arbitrary thermomechanical histories. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

60

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 2004 and 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 structure and 

allows the particles inside to break down (or oxidize) into less harmful components. DPFs reduce 

emissions of PM by 60 to 90% and emissions of hydrocarbons and carbon monoxide by 60 to 90%. 

A typical DPF consists 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 

exhaust gases carrying carbon particles flow through 

the fine pores of the walls, the carbon particles are 

filtered out. Porosity values in the range of 60% 

heighten filtration efficiency to more than 90% 

while reducing gas-flow resistance so as not to affect 

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 carbon particles collect on 

the ceramic walls of the DPF, 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. The burning produces carbon dioxide and water vapor that 

pass through the filter. This process, called regeneration, results in a cleaner filter. The regeneration 

process depends upon exhaust temperature, oxygen, 

NOx content, time, and PM levels.

The key to the successful use of DPFs is to reliably regenerate the filter (e.g., burn the PM that the 

filter traps or collects). 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. As a result, to burn soot in air, an ac 

Figure 1. Corning’s cordierite-based DPFs. 

tive system is required, i.e., one that increases the 

temperature of the exhaust using some external heat 

source. But if an active system is not carefully controlled, or if the level of PM collected on the filter 

walls is too high, a filter may 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 

corresponds to the useful life of the component and 

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, in which the failure rate increases 

with time as a result of aging phenomena. Aging 

phenomena include thermal and mechanical fatigue, 

corrosion, and creep deformation, among others. 

Figure 2. Reliability bathtub 

curve. 

Because of accelerated wear-out, the time at the 

onset of wear-out is often regarded as the useful life 

of the component The reliability bathtub curve can 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

61 

be used as a good descriptor of how the failure rate 

of DPFs evolves over time. Infancy failures of DPFs 

could be related to manufacturing or process defects 

(e.g., 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, the chemistry of the environment, 

and/or chemical and microstructural changes in the 

material due to long-term exposure to elevated temperatures in the exhaust environment and the presence of compounds used for aftertreatment 

processes. 

The objective of this project is to develop life 

prediction methods for DPFs and, by using those 

methods, to design more durable and 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 experience demanding thermomechanical 

conditions during service, including thermal shocks 

resulting from rapid heating/cooling and stresses that 

arise from temperature gradients. The approach to be 

followed in this project includes identifying and implementing test techniques for non-destructive 

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. The properties to be analyzed 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 experimentallydetermined strength data, stress analyses of the 

component using finite-element analysis, and the 

selection of appropriate failure criteria. The durability (service life) of the component can also be predicted using this framework by considering the 

mechanisms responsible for the degradation of material strength, such as slow-crack growth or creep. 

The implementation of these methodologies has 

been successful in considering dense structural ceramic components. However, there is limited work 

on their use to analyze non-structural, porous ceramic components. 

Results

To use CARES to predict the reliability of porous ceramic structures, it is necessary, through thorough fractography, to identify the flaws that control 

the strength of the material. This often amounts to 

identifying strength-limiting flaws that reside in the 

volume of the material and on the surface of the 

component. Often, more than one population of 

flaws can be identified in the volume or on the surface. Although fractographic analyses are essential 

for identifying strength-limiting flaws as part of the 

probabilistic design methodologies, the porous nature of the material and the cellular structure of the 

component have made fractographic analysis effectively impossible. This is an important issue that will 

be researched further in the future. 

One of the first tasks carried out was focused on 

analyzing the flexural behavior of beams of various 

sizes obtained from ceramic substrate DPFs. Special 

attention was given to determining flexural strength. 

The maximum tensile stress in a beam when it is 

subjected to bending is given by 

σmax = M c

I (1) 

where I is the second moment of inertia of the crosssectional area of the beam, M is the moment, and c is 

the distance from the neutral axis to the bottom of 

the beam, which is subjected to the maximum tensile 

stress. Figure 3 depicts a schematic of such a beam 

subjected to 4-point bending (dimensions are in 

inches), and Figure 4 illustrates the calculation of 

the maximum tensile stress when the cellular 

structure of the beam is either taken into account or 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

62

Figure 3. Schematic of DPF ceramic substrate 1×1 

beam subjected to 4-point bending. 

neglected. It was found that when the cross-sectional 

structure of the beam is neglected, the flexural 

strength is underestimated by 25%. These results 

were verified through a finite-element stress analysis 

(Figure 5). 

The discrepancy in determining the flexural 

strength of beams when the cellular structure of the 

beam is neglected is further magnified when the 

cross-sectional area of the beam increases. Figures 6–8 depict the case for a 2×2 beam. In this case, 

the flexural strength is underestimated by 45% when 

the cellular structure of the beam is neglected. These 

Figure 4. Stress analysis for DPF ceramic substrate 1×1 

beam subjected to 4-point bending. 

Figure 5. Results of stress analysis using finite elements 

for a 1×1 DPF ceramic substrate beam subjected to 4-

point bending. 

Figure 6. Schematic of DPF ceramic substrate 2×2 beam 

subjected to 4-point bending. 

Figure 7. Stress analysis for DPF ceramic substrate 2×2 

beam subjected to 4-point bending. 

Figure 8. Results of stress analysis using finite elements 

for a 2×2 DPF ceramic substrate beam subjected to 4-

point bending. 

results were also verified using a finite-element 

stress analysis. 

As indicated, one of the objectives of this project is to implement probabilistic design methodologies to analyze porous ceramic components with 

complex cellular structures. One of the questions 

that needs to be answered is to what extent the 

meso-structure of the component (cellular structure) 

can be accounted for, neglected, or simplified, in 

order to achieve accurate results with numerical expediency. These experimental results indicate that it 

is necessary to account for the cellular structure of 

the material, but in order to achieve numerical expediency, local-global techniques will be investigated. 

Another important consideration in implementing probabilistic design methodologies requires accurate knowledge of the properties of the material, 

including their elastic properties. The elastic properties of porous DPF test specimens have been determined with a high degree of accuracy through the 

use of a dynamic mechanical analyzer. Figure 9 depicts the experimental set-up, which consists of a 

3-point bending fixture in which the middle roller is 

attached to an actuator. A small, cyclic sinusoidal 

signal is used to stress the test specimen, and the 

response is recorded to determine the dynamic 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

63 

Figure 9. Three-point bending fixture for determining 

elastic properties of DPF ceramic substrates by using a 

dynamic mechanical analyzer. 

modulus of the material. Figure 10 depicts the crosssectional area of the beams evaluated, while Table 1 

lists the values obtained for the storage and loss 

moduli of the material. The sensitivity of this test 

technique makes it ideal for monitoring damage evolution, as determined by changes in elastic properties. 

Figure 10. Cross-sectional area of 4×2 DPF ceramic substrate beam. 

Table 1. Dynamic modulus properties of ceramic substrate DPF. 

Sample No. of 

tests 

Storage 

modulus 

(GPa) 

Loss 

modulus 

(GPa) 

Tan δ

1 4 14.7 ±0.7 0.21±0.03 0.01±0.00 

2 6 13.3 ±0.89 0.25±0.07 0.02±0.00 

3 5 13.4 ±0.9 0.23±0.04 0.02±0.00 

4 5 15.2 ±0.7 0.22±0.04 0.02±0.01 

5 5 14.9 ±0.6 0.23±0.03 0.02±0.00 

Average 14.2 ±1.1 0.23±0.05 0.02±0.00 

Summary

Test techniques have been identified and implemented to determine the physical and mechanical 

properties of DPF ceramic substrates. Specifically, 

values for the elastic modulus and flexural strength 

have been determined by dynamic mechanical 

analysis and 4-point bending, respectively. A roundrobin testing program involving Corning, Cummins, 

and Oak Ridge National Laboratory has been 

planned and will be carried out to assess the precision of test methods for determining strength and 

elastic properties. Work is under way to implement 

probabilistic design methodologies to predict the 

reliability and durability of DPFs when they are subjected to an arbitrary thermomechanical history. 

References

1. United States Environmental Protection 

Agency, Office of Transportation and Air Quality, 

“Questions and Answers on Using a Diesel Particulate Matter Filter in Heavy-Duty Trucks and Buses,” 

EPA420-F-03-017, June 2003. 

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, National 

Aeronautics and Space Administration—Glen Research Center, Life Prediction Branch, 1993. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

64
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

65

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 

Mark Andrews, John Grassi, David Longanbach, and Matt Fletcher 

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: Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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: Caterpillar Inc., Peoria, IL

Objective 

• Design and fabricate prototype engine valves from silicon nitride and titanium aluminide (TiAl) materials that 

are 30% lighter than steel valves and provide a 200% increase in service lifetime and a 10% increase in fuel efficiency. 

Approach 

• Select γ-TiAl and silicon nitride materials for this project because of their high corrosion resistance and ability 

to maintain superior strength at elevated temperatures. 

• Use an optimization routine in conjunction with a probabilistic approach to design valves (although the traditional deterministic method is practical for traditional metals, a probabilistic approach is necessary for highhardness materials). 

• Develop welding techniques to join TiAl to a titanium alloy in the fabrication of TiAl valves. 

• Perform bench tests and engine tests to verify the robustness of valve design. 

Accomplishments 

• Successfully optimized friction welding between TiAl valve heads and Ti-6V-4Al valve stems. 

• Conducted bench tests on silicon nitride valves with varying quality of surface finish; valves with good surface 

finish performed well and valves with poor surface finish failed in an anticipated mode. 

• Completed flexure testing of silicon nitride material specimens with different finish machining parameters; results yielded Weibull moduli between 10 and 25. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

66

Future Direction 

• Perform bench tests on finish-machined silicon nitride engine valves; then perform engine testing in collaboration with the National Transportation Research Center. 

• Procure additional valve blanks and flexure specimens from a second supplier; finish-machine, bench test, and 

engine test to evaluate strength and durability. 

• Continue developing ceramic valve life prediction model using results from bench tests as input. 

Introduction

Valve train components in heavy-duty engines 

operate under high stresses and temperatures and in 

severely corrosive environments. Advanced ceramics and emerging intermetallic materials are highly 

resistant to corrosion and oxidation 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 (about 1/3 that of production alloys) will 

lead to lower reciprocating valve train mass that 

could improve fuel efficiency. This research and 

development program is an in-depth investigation of 

the potential for use of these materials in heavy-duty 

engine environments. 

The overall valve train 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 computer-based lifetime prediction 

models and validated in rig bench tests and shortand long-term engine tests. After proof of concept is 

established with valves, this design approach will be 

applied to other components made from hightemperature materials. 

Approach

In an effort to develop the tools necessary to design effectively with high-strength, lightweight, brittle materials, a three-pronged approach was adopted. 

First, computational analysis was performed to optimize thermal and mechanical stresses within the 

valve during operating conditions. This initial design 

was cast into silicon nitride and TiAl valve blanks 

by several suppliers. The second phase of this approach is a characterization effort intended to evaluate valve performance on rig tests and engine tests. 

Third, a parallel path of study investigating the effects of damage caused by various machining techniques is being pursued. In order to minimize the 

probability of a brittle failure caused by surface defects, optimal surface finish must be achieved. 

Work reported in the FY 2003 report outlines 

the design process used to adapt the current baseline 

metallic valve design to a design more suited for a 

ceramic material. For example, the head dimensions 

were changed to reduce the stress concentration experienced in the fillet radius. Several silicon nitride 

valve blanks were procured and finish-machined to 

these specifications. Because of the high cost of 

TiAl, only the heads of the valves were cast and finish-machined. After rough machining, the TiAl 

heads were successfully friction-welded to cylindrical stock of a titanium alloy (Ti-6V-4Al). Subsequently, they will be finish-machined before rig and 

engine testing. 

Results

The mechanics of a structural brittle material are 

inherently different from those of a traditional metal. 

Whereas a metal (current-production valve material) 

may be adequately modeled as a continuum material, a brittle material (proposed valve material) is 

susceptible to microscopic flaws or cracks. A primary failure mode of a structural ceramic is the 

propagation of a surface crack initiated during machining. Therefore, it is imperative to understand the 

influence of various machining parameters on material strength and reliability. 

Two types of valves were bench tested with a 

motor-driven valve impact rig: valves with a good 

surface finish (NT551 material) and a valve with a 

poor surface finish (SN235P material). Several 

NT551 valves survived 1000-hour rig tests incurring 

minimal damage (Figure 1). However, there appear 

to be engine deposits on the fillet radius and on the 

contact regions (composition and nature of the deposits has not yet been determined). 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

67

Figure 1. NT551 engine valves subject to 1000 hours of impact / wear tests. Note the deposits in the fillet radius and the 

contact regions (boxed regions) 

In contrast to the good performance of the 

valves with acceptable surface finish, the performance of a poorly finished valve (SN235P material) 

was less than acceptable. This does not imply a difference in the quality of the materials but rather the 

importance of proper machining and surface finish. 

The keeper groove was not finished to specification, 

resulting in a rough surface with significant microcracking. Within one hour of rig testing, the valve 

fractured at the keeper groove. Microscopic analysis 

indicated that the fracture origin coincided with the 

region of severe machining damage (Figure 2). 

Nondestructive analysis performed by J. G. Sun at 

Argonne National Laboratory confirmed that the 

failure was a result of a pre-existing surface defect 

(Figure 3). Note that the white line corresponds to 

the location in the keeper groove where the flaw 

originated. 

It was therefore decided to discontinue the 

bench rig testing until the appropriate surface finish 

was provided for all portions of the valve. The result 

of a subsequent investigation suggested that there 

may be an unacceptably low static friction coefficient between the valve and the keeper locks, resulting in a misapplication of the load from the valve 

spring. A study was launched to determine if the 

frictional force is inadequate for this application. 

To further investigate the importance of surface 

finish to material strength, an optimization study 

was performed on cylindrical material (SN235P) 

flexure specimens with varying machining parameters. The test specimens (8 mm diam × 75 mm 

length) were subjected to three different finishing 

machining processes before they were strengthtested in four-point flexure. The tests were conducted at the Oak Ridge National Laboratory 

(ORNL) through a facility user agreement. A specialized four-point test fixture was designed and 

built for the cylindrical test specimens by A. 

Wereszczak of ORNL. All tests were conducted at 

ambient temperature conditions at the displacement 

rate of 0.5 mm/min, as specified in the ASTM C1161 standard. 

Figure 4 shows the characteristic strength and 

uncensored Weibull moduli of 12 different surface 

finish conditions. The characteristic strength shown 

in Figure 4 suggests that the inherent strength of the 

SN235P material is met regardless of machining 

condition. However, it is apparent that some of the 

finishing processes improve the data variance, as 

indicated by the high Weibull modulus values in 

Figure 4. A plausible explanation for the increase in 

the Weibull modulus values is that some finishing 

processes induce an appreciable compressive residual stress while others do not. The fractured samples 

will be examined using X-ray diffraction at ORNL 

to investigate the role of residual stress in the materials’ strength behavior. 

Conclusions

An investigation into the use of advanced materials for valve train applications is under way. The 

three-pronged approach combines probabilistic design techniques, rig and engine testing, and a systematic study of the effects of machining damage on 

reliability material. Rig testing exposed the importance of generating components with low-defect sur- 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

68

Figure 2. Fracture surface from fractured SN235P valve keeper groove observed from axial view: (left) optical micrograph, (right) scanning electron microscope image. 

Figure 3. Laser scanning nondestructive evaluation image of the keeper notch of a poorly finished SN235P valve, showing a line of possible machining damage at the top of the image. 

faces, and a study examining the effect of machining 

conditions on identical material specimens reached 

the same conclusion. 

In addition to this fundamental study, several 

other steps will be taken to further examine the feasibility of fabricating a commercially viable ceramic 

or intermetallic valve. Arrangements are being made 

to engine test the silicon nitride valve prototypes 

near ORNL. Also, a new test rig will be rebuilt that 

will allow examination of both thermal and mechanical effects on experimental valves; and a more 

detailed life prediction model will be developed to 

examine the effect of accumulated damage on ceramic and intermetallic valves. Present and future 

efforts will help establish an understanding of how 

to design and manufacture nontraditional lightweight materials for valve train and other heavyduty diesel engine applications. 

2.5mm 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

69

Figure 4. Results from four-point flexure test of SN235P 

cylindrical test specimens for various finish machining 

conditions: (top) uncensored Weibull characteristic 

strength; (bottom) uncensored Weibull moduli. Note that 

regions at the top of each column represent the 95% confidence interval of the data. 

0

200

400

600

800

1000

1 2 3 4 5 6 7 8 9 10 11 12

Characteristic Strength (MPa)

Cylindrical Specimen Data Set

0

5

10

15

20

25

30

35

40

1 2 3 4 5 6 7 8 9 10 11 12

Weibull Modulus

Cylindrical Specimen Data Set
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

70
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

71 

B. Engineered Surfaces for Diesel Engine Components

M. Brad Beardsley and Jesus Chapa-Cabrera 

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: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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: Caterpillar Inc., Peoria, Illinois 

Objective 

• Evaluate new laser technologies such as surface dimpling, cleaning, and laser-assisted spraying to enhance adherence and increase coating strength. 

• Develop phosphate-bonded composites for internal thermal management coatings. 

• Evaluate new quasicrystalline materials as potential thermal barrier and wear coatings. 

Approach 

• Evaluate laser treatment techniques by treatment of small samples. 

• Examine the resulting sample microstructures to determine the potential of the treatment method. 

• Measure the mechanical and physical properties of phosphate-bonded composites and use them to determine 

their potential for use as thermal barrier coatings (TBCs). 

• Evaluate the thermal stability of quasicrystalline materials using diffusion couples to assess the stability of the 

aluminum-based quasicrystalline materials at engine operating temperatures. 

Accomplishments

• Completed cost analysis of laser pre- and post-treatments and selected post-laser “tacking” as the method for 

treating coatings to increase adherence. 

• Developed phosphate-bonded composite coatings that survived initial thermal cycling to 650oC with good residual adherence to the substrate. 

• Sprayed quasicrystalline coatings using the high-velocity oxygen fueled technique and began evaluating their 

thermal conductivities. 

Future Direction 

• Install a 2000-W Q-switched laser in Caterpillar’s thermal spray facility to investigate the effects of laser surface ablation on the adherence and density of thermal-sprayed coatings. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

72

• Prepared phosphate-bonded coatings with high surface adherence that have shown promise as high-temperature 

adhesives as well as TBCs. Additional formulations are under development. 

• Conduct additional oxidation studies as well as limited in-cylinder engine testing for the aluminum-based quasicrystalline coating, which has shown limited stability at 900°C in diffusion couples. At temperatures below 

500°C, the material shows promise to retain the low thermal conductivity of the quasicrystalline state.. 

Introduction

Engine testing of thermal sprayed coatings has 

demonstrated their potential benefit as thermal barriers and as wear coatings to reduce fuel consumption, 

reduce wear, and reduce component temperatures.1,2,3 The durability of thermal sprayed coatings, 

particularly TBCs, remains the major technical challenge to their implementation in new engine designs. 

New approaches to coating design and fabrication 

will be developed to aid in overcoming this technical 

hurdle. 

Three approaches to new TBCs have been pursued: (1) laser treatments applied to plasma-sprayed 

TBCs to modify their structures and adherence to 

enhance durability, (2) new quasicrystalline materials capable of being sprayed using a high-velocity 

oxygen-fuel (HVOF) process that produces denser 

coatings with higher bond strengths, and (3) new 

phosphate-bonded composites that allow for inner 

diameters to be coated and that have low thermal 

conductivity and high adherence. 

Laser Treatments. Two types of laser treatments have been investigated. The laser-assisted 

plasma spray technique uses the power applied by 

the laser to completely melt the sprayed materials on 

the substrate to allow for metallurgical bonding and 

higher adherence of the coating. A second laser 

treatment, laser ablation, uses a high-power, shortpulsed laser to ablate the surface oxide on the substrate and interlayer of the sprayed material to increase the bonding of the coatings. Initial samples 

were produced using laser techniques that melted the 

sprayed materials to demonstrate the effect of the 

coating structure. Consultation with the Fraunhofer 

Institute in Dresden, Germany, provided sufficient 

confidence in the process to proceed with installation of a 4-kW YAG laser in Caterpillar’s thermal 

spray laboratory for further process development. A 

review of the laser ablation technique with the University of Technology at Belfort-Montbeliard 

(UTBM) in France provided sufficient data to show 

that the technique has promise for improving the 

adherence of sprayed coatings. Caterpillar is acquiring a Q-switched laser capable of producing the energy needed for the ablation process to proceed with 

this evaluation. 

Quasicrystalline Materials. This new class of 

metallic materials exhibit low thermal conductivity 

with high thermal expansion as a result of their 

unique, aperiodic crystal structure. The material 

chosen for evaluation is an aluminum-based quasicrystal, Al71Co13Cr8Fe8, shown previously to be capable of 1100°C operation in a turbine engine.4

Initial diffusion couples have been run to aid understanding of the thermal stability of the alloy at operating temperatures. 

Phosphate-bonded Composites. Phosphatebonded composites consist of filler material bonded 

together by a phosphate. The compositions of both 

the filler and the phosphate binder can be widely 

varied and, for this reason, so can the properties of 

the composite. Phosphate-bonded composites are 

well known and have been in use for many years as 

refractory mortars and cements; high-temperature, 

corrosion-resistant coatings; temporary bone replacement; fast-cure paving cement; and dental cement. 

Of the well-known phosphates, most are not 

suitable for high-temperature use. These include all 

of the hydrated crystal forms, such as those used for 

bio-replacement; low-temperature cements; and paving compounds, which simply decompose at elevated temperatures. Certain other non-hydrated 

alkali and alkaline phosphates are equally unsuitable 

because of their tendency to form low-meltingtemperature eutectics in some environments. Metal 

phosphates have the highest potential for long-term 

chemical durability in high-temperature applications 

such as internal combustion and turbine engines. 

Approach 

Laser Treatments. Initial sprayed samples were 

treated using a key-hole laser beam to demonstrate 

the resulting coating structure. The samples were 10 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

73 

mm in diameter with a substrate thickness of 3 mm. 

Coating thickness was approximately 0.5 mm with 

two types of coatings treated. One type of sample 

was the bond coating only with no ceramic top coat. 

The second coating type was a bi-layer bond 

coat/ceramic. The bond coating was a standard Ni20%Cr-6%Al-0.5%Y, and the ceramic was 8% yttria-zirconia. The resulting laser-treated coating microstructures were analyzed for depth of penetration 

and cracking. 

The laser ablation treatment was evaluated by 

reviewing the prior work preformed by UTBM in 

France. This work demonstrated good adherence for 

coatings applied to aluminum and titanium substrates. As these materials are highly reactive, the 

laser ablation done just prior to the coating application would remove the inherent surface oxide, allowing for better substrate wetting. Limited work has 

been done on steel and cast iron materials, but sufficient technical promise was shown for Caterpillar to 

invest in the Q-switch laser required for further 

process evaluation. The lasers will be delivered to 

Caterpillar in December 2004. 

Quasicrystalline Materials. Owing to the metallic nature of the Al71Co13Cr8Fe8 quasicrystal, 

HVOF spraying techniques can be used to produce 

coatings with this material. The HVOF process produces higher density and better bonding of the coating than the plasma spray process by using high 

particle velocities at relatively low particle temperatures to “peen” or “forge” the particles onto the substrate surface. The low temperature of the process 

prevents ceramics from being sprayed using this 

technique. 

HVOF coatings of Al71Co13Cr8Fe8 were produced for diffusion couple testing. Four diffusion 

couples were produced: (1) the quasicrystal and a 

steel substrate; (2) the quasicrystal, a Ni-17Cr-6Al0.5Y bond coat, and a steel substrate; (3) the quasicrystal, a Ni-31Cr-11Al-0.6Y bond coat, and a steel 

substrate; and (4) the quasicrystal, a Fe-26Cr-8Al0.4Y bond coat, and a steel substrate. The diffusion 

couples were produced by spraying a 12.5-mm-diam 

by 19-mm steel substrate with 0.5-mm-thick layers 

of the bond coatings and quasicrystal. The bond 

coatings were applied first, followed by the diffusion 

couple. 

The diffusion experiments were conducted by 

sealing each type of sample in evacuated quartz 

tubes. Temperatures of 500, 700 and 900°C with 

times of 25, 100 and 500 hours were used. The samples were water-quenched after annealing. Samples 

were then mounted and polished for examination by 

optical microscopy, scanning electron microscopy, 

and microprobe analysis. 

Phosphate-bonded Composites. Research efforts on metal-phosphate composites at Caterpillar 

have been aimed at developing seal coatings, corrosion resistant coatings, and, more recently, thermal 

insulating coatings. Thin, corrosion-resistant coatings have been applied on various surfaces and have 

been successfully tested in various applications. 

Metal-phosphate adhesives have been used at Caterpillar for assembling prototype aftertreatment components for testing at temperatures around 800°C. 

Qualitative testing between phosphate-based and 

commercially available adhesives has shown the 

superior adhesive strength of the phosphate-based 

adhesives; however, quantitative characterization of 

the strength of phosphate-based adhesives was 

deemed necessary to assess their potential for hightemperature adhesive applications and as an intermediate step for the development of robust coatings. 

Binder preparation and characterization. The 

basic stoichiometry of the phosphate material under 

consideration is 1.4 Al2O3: 0.25 Cr203: 3 P2O5. The 

binder preparation consists of continuously stirring 

orthophosphoric acid at 110°C and dissolving chromium trioxide, which is later reduced by adding sucrose; finally, alumina is added to the desired 

stoichiometry until it is fully dissolved. Additional 

chemistries, including the use of different chemical 

species and different proportions, are being explored. 

Adhesive preparation and characterization. 

Adhesive composites consisting of 68% packed 

alumina with average particles of 0.7 μm were produced, filling the interstitial space with different 

phosphate binders. For lap shear testing, 1 × 1 × 0.5 

blocks of 90% pure alumina were used to make 

bonded, stepped specimens. The mixtures had a 

thick slurry consistency and were applied by brush 

on a section of the surface of the blocks. The 

brushed halves of two blocks were then aligned to 

form a step, obtaining a consistent alignment with an 

aluminum fixture that was built to produce five 

specimens at a time. Next, the fixture containing the 

samples was introduced to a curing oven at 300°C. 

A small amount of pressure perpendicular to the 

joint was applied at this stage. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

74

Results 

Laser Treatments. Laser tacking of the zirconia 

ceramic induces high amounts of cracking and void 

formation (Figure 1). Laser tacking of the bond coating material was more successful, in that cracking 

was suppressed during solidification of the material 

because of the higher ductility of the metal (Figure 2). This indicates that the laser-assisted plasma 

processing will need to be applied to thinner layers 

of the ceramic coating as it is built up during spraying to reduce the residual stress induced. Layer 

buildup during spraying usually proceeds at 

0.025 mm or less for the ceramic layers. If the laser 

beam interaction time can be adjusted to provide 

sufficient heat input during coating application, it 

should be possible to build the ceramic layer without 

detrimental cracking. 

Figure 1. Laser tacked ceramic using medium laser 

power showing melting of the zirconia layer with crack 

and void formation. 

Quasicrystalline Materials. HVOF spraying of 

the quasicrystalline material produced a dense, wellbonded structure (Figure 3). The stability of the quasicrystalline material at 700°C when in contact with 

the steel substrate material was fair, and a minor interaction zone was created after 500 hours (Figure 

4). Cracking in the diffusion couple samples is due 

to the high stress state caused during water 

Figure 2. Laser tacked bond coating at medium laser 

power showing melting of the bond coat material with 

fusing into the coating and oxide formation under melt 

pool. 

Figure 3. HVOF quasicrystalline material showing dense, 

well bonded coating structure. 

Figure 4. Quasicrystal/substrate diffusion couple after 

500 hours at 700oC showing small reaction zone at interface. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

75 

quenching of the samples after exposure. At the 

higher 900°C temperature, the reaction zone with the 

substrate increased substantially, as expected (Figure 

5). When coupled with the nickel-based bond coatings, large reaction zones were developed at the 

700°C temperature (Figure 6), probably as a result 

of the higher reactivity of the aluminum with the 

nickel. As with the steel substrate, a smaller reaction 

zone was found with the iron-based bond coat (Figure 7). Microprobe analysis of the diffusion couple 

samples is under way to better understand the interaction of the materials at temperature. Oxidation 

samples will be made and exposed at similar temperatures to determine the effect of oxidation on the 

material stability. 

Figure 5. Quasicrystal/sbustrate diffusion couple after 25 

hours at 900oC showing small reaction zone at interface. 

Figure 6. Quasicrystal Ni-31Cr-11A1-0.6Y diffusion 

couple after 500 hours at 700oC showing large reaction 

zone at interface. 

Figure 7 Quasicrystal/Fe-26Cr-8A1-0.4Y diffusion couple after 500 hours at 700oC showing a large reaction 

zone at interface. 

Phosphate-bonded Composites. Evaluation of 

the shear strength of the different metal-phosphatebased adhesives and two commercial adhesives was 

conducted following the ASTM D401 standard. The 

fixture used for this test is shown in Figure 8. 

Figure 8. Fixture used for measuring the shear strength of 

the adhesives as per ASTM D401. 

Data obtained from the shear lap testing are 

shown in Figure 9. Results are presented for the 

phosphate adhesive Ad7, which is the phosphate 

adhesive that yielded the highest strengths on the 

first round of experiments. Strength data of two 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

76

Figure 9. Measured strengths of adhesive produced inhouse, Ad7, compared with commercially available adhesives Com-1, Com-2. 

commercial adhesives, designated as Com1 andCom2, are shown for comparison. The error bars 

represent the standard deviation. These preliminary 

results suggest that phosphate adhesives can be produced with strengths of at least the same level as 

commercial adhesives. Testing with other phosphate 

formulations is continuing. Selected formulations 

showing the highest shear strengths will be used for 

continuing coating development. 

Conclusions 

Laser Treatments. Sufficient promise has been 

demonstrated for laser surface treatments to proceed 

with moving a 4-kW YAG laser into Caterpillar’s 

thermal spray laboratory for further evaluation of the 

laser-assisted plasma spray process and to procure a 

Q-switched laser system for evaluation of the laser 

ablation method. Moving the YAG laser and procuring the Q-switched laser are being done independently of the current program funding. 

Quasicrystalline Materials. The reaction zones 

of the substrate/quasicrystalline and Fe-26Cr-8Al0.4Y/quasicrystalline diffusion couple samples indicate some promise for limiting the interdiffusion of 

the materials in a TBC structure by designing the 

coatings to have interface temperatures of below 

500°C. The ability to use a graded design may be 

limited by the reactivity of the quasicrystalline material with the bond coating materials used for grading. Further evaluation will be necessary to 

understand what types of graded designs can be 

made. 

Phosphate-bonded Composites. The current 

phosphate-bonded composite shows much promise 

as a high-temperature adhesive, and commercialization of the material for this use will be pursued. 

References

1. R. A. Miller, Assessment of Fundamental Materials Needs for Thick Thermal Barrier Coatings 

(TTBC’s) for Truck Diesel Engines, Department of 

Energy/National Aeronautics and Space Administration/21749-1, NASA TM-103130, 1990. 

2. M. B. Beardsley, and H. J. Larson, Thick 

Thermal Barrier Coatings for Diesel Components, 

DOE/NASA/0332-1, National Aeronautics and 

Space Administration CR-190759, 1992. 

3. M. B. Beardsley, P. G. Happoldt, K. C. Kelley, E. F. Rejda, and D. F. Socie, “Thermal Barrier 

Coatings for Low Emission, High Efficiency Diesel 

Engine Applications,” SAE Tech Paper Series, 

1999-01-2255, Government/Industry Meeting, 

Washington, D.C., 1999. 

4. 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). 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

77 

C. Processing and Characterization of Structural and Functional Materials for 

Heavy Vehicle Applications 

J. Sankar, Z. Xu, and S. Yarmolenko 

Department of Mechanical Engineering 

North Carolina A & T State University 

1601 East Market Street 

Greensboro, NC 27411 

(336) 256-1151, ext. 2282; fax: (336) 256-1153; e-mail: sankar@ncat.edu 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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: North Carolina A&T State University, Greensboro, North Carolina 

Objectives 

• Produce yttria-stabilized zirconia (YSZ) thin films using combustion chemical vapor deposition technique for 

solid oxide fuel cells. 

• Investigate the nucleation of YSZ and its controlling parameters. 

• Study the evolution of the thin film microstructure. 

• Enhance the film deposition rate. 

Approach 

• Deposit YSZ thin films using combustion chemical vapor deposition (CCVD). 

• Characterize the films using scanning electron microscopy (SEM) and X-ray diffraction. 

• Study the effects of substrate temperature and metal concentration on the nucleation density. 

• Examine the evolution of the thin film microstructure by experiments and stochastic simulation. 

• Optimize the operation parameters of the CCVD system. 

Accomplishments 

• Enhanced the nucleation density 

• Developed a stochastic model to simulate the evolution of the thin film microstructure. 

• Established a set of conditions for high film growth rates. 

Future Direction 

• Study YSZ film deposition on porous electrodes. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

78

Introduction

YSZ is an oxygen-ion-conductive material that 

has conventionally been used as an electrolyte in 

solid oxide fuel cells (SOFCs) and in oxygen sensors 

for pollution and safety monitoring, control, and 

automation of industrial processes, and energy 

conservation.1

 Because of the relatively low 

electrical conductivity of YSZ (about 0.1 S·cm−1 at 

1000°C), most of the ohmic loss is due to the 

resistance of the electrolyte. To reduce the ohmic 

loss and increase the power generating efficiency of 

the fuel cell, it is desirable to use a thin film 

electrolyte. On the other hand, a temperature of as 

high as 1000°C is usually employed to obtain 

reasonable ion conductivity for YSZ material. This 

high temperature makes material selection difficult 

and increases the product cost greatly. Reducing the 

thickness of the electrolyte would increase the 

conductivity at lower temperatures. 

YSZ thin films usually have been synthesized by 

CCVD or electrochemical vapor deposition 

(EVD).2,5 [Should this ref callout be 2–5? No 

callouts found for refs. 3 and 4.] The vacuum 

systems used by these techniques limit the sizes of 

the parts and increase the cost of the products. 

Liquid fuel CCVD is a promising technique to 

produce YSZ thin films for industrial applications 

because CCVD works in an open atmosphere and 

offers the potential for conformal deposition of films 

on non-flat surfaces. Moreover, because no vacuum 

chamber is needed for the process, there is no size 

restriction on parts on which the film can be 

deposited. Because of the relatively high precursor 

concentration in liquid fuel compared with the fuel 

usually used for low-pressure CVD processes, 

CCVD is expected to have a higher deposition rate. 

Therefore, CCVD is a promising technique for 

producing YSZ thin films for industrial SOFC 

applications at a lower cost than any other 

CVD/EVD technique. 

Electrolytes for SOFCs must be highly dense to 

avoid any cross flow of oxygen or fuel. The main 

purpose of this research is to produce dense YSZ 

thin films through the study of nucleation and 

microstructure evolution. Reducing the cost of fuel 

cell manufacturing by using thin film technology 

and a high film growth rate is of great interest. 

Methods to enhance the growth rate of YSZ have 

been studied. 

Approach

Thin films of YSZ to be used as electrolytes for 

SOFCs have been deposited using CCVD. The 

deposition system was developed at North Carolina 

A&T State University. Metal-organic reagents were 

dissolved in an organic solvent (toluene). The 

solution was atomized into small aerosols with a 

nebulizer, mixed with oxygen, and then ignited by a 

pilot flame. Thin films then were deposited on 

substrates placed downstream of the aerosol flame. 

Nucleation study of YSZ was performed by 

varying the substrate temperature and metal 

concentration in the solvent. At design conditions, 

depositions were conducted for different lengths of 

time. The microstructures of these depositions were 

observed and analyzed, and the grain growth was 

studied with stochastic modeling. Efforts to enhance 

the film growth rate were carried out through study 

of the effects of thermophoresis and metal 

concentration. All the present depositions were 

performed on mirror-polished single-crystal silicon 

substrates. The phases of the deposited films were 

examined with X-ray diffraction, and their 

morphology was characterized by SEM. 

Results

Nucleation 

Effect of substrate temperature. The 

temperature of the substrate is an important factor 

that controls the nucleation density, defined as the 

number of aggregates created per unit area. To be 

able to use the image analysis software Image Pro to 

compute the nucleation density, the nucleation time 

of the YSZ particles was set for 130 s. The 

depositions were conducted at substrate 

temperatures of 800, 900, 1000, 1100, and 1200oC 

and a metal concentration of 1.25E-03 M. The asgrown samples were characterized with SEM as 

shown in Figure 1. 

Counting the particle number in the specific area 

with the application software showed that the 

nucleation rates ranged over an order of magnitude, 

from 1010 to 1011 cm−2

, in the tested temperature 

range. The nucleation rate variation versus the 

substrate temperature is illustrated in Figure 2. The 

variation of the nucleation rate with the substrate 

temperature is characterized by two tendencies: 

When the substrate temperature is between 1000 and 

1200°C, the nucleation rate decreases with the 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

79

Figure 1. Micrographs of YSZ nuclei nucleated at various 

substrate temperatures: (a) 800°C, (b) 900°C, (c) 1000°C, 

(d) 1100°C, and (e) 1200°C under conditions of 

nucleation time of 130 sec and total metal concentration 

of 1.25× 10−3

 M. 

Figure 2. YSZ nucleation rate as a function of the 

substrate temperature according to the information in 

Figure 1. 

substrate temperature; in the substrate temperature 

range of 800 to 1000°C, the nucleation rate increases 

with the substrate temperature. The former 

nucleation rate tendency was well accepted. As far 

as the author knows, the latter tendency has not been 

reported. It is proposed that at low substrate 

temperatures, the material species in the gas phase 

condense into a solid phase and coagulate into 

particles. They rebound from the surface instead of 

becoming adsorbed when they hit the substrate 

surface. This process also leads to depletion of 

species in the gas phase that is needed for CVD 

and then results in a decrease in the nucleation rate. 

The images in Figure 1 also reveal that on samples 

nucleated at low temperatures, there are many large 

nodular particles that support the above statement. 

Effect of metal concentration. Another method 

used to enhance the nucleation rate was to increase 

the deposition flux, i.e., the metal concentration, for 

our experimental system. At low concentrations of 

5×10−4 and 1.25×10−3

, only isolated particles 

presented on the substrate surface after depositions 

of 130 s, whereas almost continuous films were 

obtained when the concentration was increased to 

3×10−3

 M and up. At high total metal concentrations, 

the particle size reached 30–50 nm. From this set of 

experiments, it can be concluded that the lowest 

metal concentration, 5.5×10−4 M, is not suitable if a 

high nucleation rate is demanded. 

Structural Evolution

Experimental study. To grasp the physics of 

the evolution process of the YSZ thin film, the grain 

sizes at various stages of the processing time were 

measured for statistical post-processing. Figure 3 

shows the micrographs of YSZ particles/crystallites 

at different processing times. In Figure 3(a), the 

nuclei can hardly be seen after only 70 s of 

processing, except some large particles that are 

assumed to be contamination. From Figures 3(b) to 

(d), with the increase of processing time, the sizes of 

the particles increase; however, the number of 

particles is reduced. This phenomenon can be 

interpreted by the mechanisms of coarsening and 

coalescence of the growing particles. With increased 

processing time, the sizes of the particles increase. 

Some secondary nucleation and growth on the large 

particles can be noticed. It is also reasonable to 

assume that the secondary nucleation should take 

place on the substrate before all the surface of the 

substrate is covered with particles. The film consists 

of both (111) and (100) oriented crystals. The 

difference in orientation originated at the stage when 

the particles were very small. 

According to the obtained mean particle radii, 

particle growth rate can be estimated by plotting the 

mean diameter versus the growing time, as shown in 

Figure 4. The particle growth rate is approximately 

linear during the time period studied (130 to 430 s). 

The intercept of the line on the time axis is about 

33.3 s, which shows the incubation time for 

nucleation. With prolonged deposition time, 

continuous thin films can be obtained. 

Stochastic simulation. The number and size of 

the particles were obtained with the image 

processing software. After statistical post- 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

80

Figure 3. Microstructure of the YSZ particles/crystallites 

at different processing times: (a) 70 s, (b) 130 s, (c) 190 s, 

(d) 250 s, (e) 310 s, (f) 370 s, (g) 430 s, (h) 490 s, and 

(i) 550 s on Si(100) substrates, at a substrate temperature 

of about 1200°C, metal concentration of 1.25×10−3

 M.

Figure 4. Particle growth rate vs. processing time from 

the data obtained from the micrographs as shown in 

Figure 3. 

processing, a comparison of cumulative distribution 

functions (CDFs) of grain radius normalized by the 

mean radii is given in Figure 5. As can be seen from 

Figure 5, there is a distinct dispersion among these 

CDFs plotted in the normalized space. Thus the use 

of a normal grain growth model cannot accurately 

capture the evolution process of grains in YSZ thin 

films. The deviation from the normal grain growth 

model is mainly attributed to the orientationdependent grain boundary energies and mobility. 

Figure 5. Comparison of CDFs of grain radius in 

normalized space. 

The lognormal distribution has been used 

extensively as the initial distribution of grain size.6,7

To validate the use of the lognormal model, a 

comparison of the initial CDF of grain radius at 

t=130s with its equivalent lognormal distribution is 

shown in Figure 6. As shown in Figure 6, the 

equivalent lognormal distribution can accurately 

capture the true initial grain size distribution. A 

stochastic model was developed by taking the radii 

of the neighboring grains and the orientationdependent constants into consideration. The 

computational efforts were greatly reduced by 

dividing the entire population of grains into several 

sub-groups based on the level of their percentiles. It 

was assumed that all grains within each sub-group 

have identical grain growth behavior. The 

orientation-dependent grain growth constants were 

determined from curve fitting of the experimental 

data. An example of comparing model prediction 

with the experimental observations is shown in 

Figure 7. 

Growth Rate

Thermophoresis is a thermal-gradient–directed 

flow of material from high-temperature to lowtemperature regions. Thermophoresis has been 

found to be a strong factor in some CVD 

environments.8

 A temperature gradient across a 

diffusion boundary layer causes a thermophoresis 

effect. In this case, if the flame were hotter than the 

substrate, on average, the gas molecules from the 

substrate would have a smaller velocity than the gas 

molecules traveling toward the substrate. This 

creates a driving force for YSZ clusters formed in 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

81

Figure 6. Illustration of accuracy of using an equivalent 

lognormal distribution for statistical characterization of 

initial grain size distribution. 

Figure 7. Comparison of model prediction with measured 

CDF data of grain radius at t=430s. 

the direction toward the substrate. A larger 

temperature gradient adjacent to the substrate 

increases the thermophoresis effect. The study of the 

effect of thermophoresis was carried out by film 

depositions at different substrate-to-nuzzle distances 

from 51 mm (far into the flame) to 83 mm (out of 

the visible end of the flame). The film thickness 

measured on SEM images is plotted versus the 

substrate-to-nozzle distance in Figure 8 on a solid 

line. The data can be fit by an exponential equation. 

With an increase in the substrate-to-nozzle distance, 

precursor concentration in the flame will be 

attenuated because of the expansion of the flame. 

Normalization was done to eliminate the effect of 

concentration dilution on the film growth rate. The 

normalized thickness data were again plotted in the 

Figure 8. Film thickness as a function of the substrate-tonozzle distance, line 1 for the original measured data, line 

2 for the normalized data. 

same graph of the originally measured data, as is 

shown in Figure 1 by curve 2 in the dashed line. It is 

apparently noticeable that after normalization, the 

growth rate of the film still follows the exponentially 

decaying mode with the substrate-to-nozzle distance. 

In other words, the parameter substrate-to-nozzle 

distance does play a role in film growth. 

The effect of the metal concentration on the film 

growth rate is shown in Figure 9. Within the range 

of the concentration employed in our experiments, a 

linear relationship was obtained between the film 

growth rate and the concentration. By observing the 

morphologies of the samples, the films had wellcrystallized and faceted particles as long as the metal 

concentration was less than 4.25×10−3

 M. Beyond 

this limitation, the film was in a cauliflower-like 

structure. 

Conclusions

According to the investigations we have 

performed, the CCVD technique has been 

demonstrated to be a promising method for film 

processing of YSZ electrolyte. Nucleation density 

can be improved by controlling the substrate 

temperature and the metal concentration 

appropriately. The analysis of the grain growth 

demonstrated that after the initial stage, the grain 

growth does not follow the lognormal mode. A 

stochastic model has been developed to predict the 

grain size in the film at any deposition time. The 

experiments confirmed the enhancement of the film 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

82

Figure 9. Film growth rate versus metal concentration in 

the liquid solution. 

growth rates by employing the effect of 

thermophoresis and increasing the metal 

concentration. 

References

1. S. P. S. Badwal, “Ceramic Superionic 

Conductors,” p. 570 in Materials Sciences and 

Technology, A Comprehensive Treatment, R. W. 

Cahn, P. Haasen, and E. J. Kramer, eds., Vol. 11, 

1994. 

2. G. Garcia, J. Casado, J. Llibre, and A. 

Figueras, J. Crystal Growth 156, 426 (1995). 

3. S-C. Hwang and H-S. Shin, J. Am. Ceram. 

Soc. 82(10), 2913 (1999). 

4. U. B. Pal and S. C. Singhal, J. Electrochem. 

Soc. 137, p. 937 (1990). 

5. L. G. J. De Haart, Y. S. Lin, K. J. de Vries, 

and A. J.Burggraaf, J. Eur. Ceram. Soc. 8(1), 59 

(1991). 

6. V. Y. Novikov, Acta Mater. 47(18), 4507 

(1999). 

7. C. V. Thompson, Annu. Rev. Mater. Sci. 20, 

245 (1990). 

8. W. Bai, K. L. Choy, N. H. J. Stelzer, and J. 

Schoonman, Solid State Ionics 116, 225. 

Publications

1. Z. Xu, J. Sankar, and S. Yarmorlenko, 

“Yttria-Stabilized Zirconia Coatings Produced Using 

Combustion Chemical Deposition,” Surface and 

Coating Tech., in print. 

2. Z. Xu, C. Hilton, Bobby Watkins, Sergey 

Yarmolenko, and J. Sankar, “Thin YSZ Electrolyte 

Film Depositions on Dense and Porous Substrates”, 

paper IMECE2003-43330, 2003 ASME 

International Mechanical Engineering Congress and 

RD&D Expo, 2003. 

3. Z. Xu, J. Sankar, S. Yarmolenko, and Q. Wei, 

“Nucleation And Growth of Yttria-Stabilized 

Zirconia Thin Films Using Combustion Chemical 

Vapor Deposition,” pp. 509–514 in Materials 

Research Society Symposium Proceedings, Vol. 756, 

2003. 

4. Z. Xu, S. Yarmolenko, and J. Sankar, 

“Enhancement Of YSZ Thin Film Deposition Rate 

In CCVD,” pp. 861–862 in Proceedings of 10th 

International Conference on Composites 

Engineering, New Orleans, July 20–26, 2003. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

83

D. NDE Development for Ceramic Valves 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: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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 a laser scattering nondestructive evaluation (NDE) method for detection and characterization of subsurface defect/damage in silicon-nitride ceramic valves so NDE data can be used to determine costeffectiveness and reliability of ceramics valves for diesel engines. 

• Identify strength-limiting flaws and failure mechanisms in cylindrical silicon-nitride specimens due to inherent 

material defects or machining damage. 

Approach 

• Continue optimization of an automated laser scattering NDE system for fast scanning of entire valves. 

• Inspect as-processed valves to determine the level of machining damage and evaluate bench-tested valves for 

accumulated damages. 

• Correlate NDE data with fracture test results for machined cylindrical silicon nitride specimens to identify fracture initiation defects and defect type, depth, and severity. 

Accomplishments 

• Successfully tested a photomultiply tube (PMT) optical detector in a laser scattering system. The PMT detector 

is more sensitive and >6 times faster than the current semiconductor optical detector. 

• Evaluated ten NT551 silicon-nitride valves that were bench tested for 100, 500, and 1000 hours. Analyzed 

processing and accumulated damages in the valve head sections. 

• Inspected 48 as-processed SN235P silicon nitride valves. Determined levels of machining damage in various 

sections of the valves and identified severe damage that caused valve failure. 

• Demonstrated laser scattering detection sensitivity for fracture initiation defects in machined cylindrical 

SN235P silicon nitride specimens. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

84

Future Direction 

• Implement the PMT detector in the automated NDE valve-scan system and use optical fibers for collecting and 

delivering scatter light to the detector. Further investigate other detectors that may have higher speed for this 

application. 

• Inspect re-machined SN235P valves for initial damage level due to machining. Characterize accumulated damage by periodic inspection of these valves during planned rig/engine tests of various operating conditions and 

durations. 

• Develop new theoretical, experimental, and quantitative image-processing methods for optimized characterization of strength-limiting subsurface defects/damage (size and depth) due to corrosion, fatigue, impact, and wear 

in silicon nitride ceramics. 

Introduction

Advanced ceramics are leading candidates for 

high-temperature engine applications that offer improved fuel efficiency and engine performance. 

Among them, silicon nitrides are being evaluated as 

valve train materials for diesel and natural gas engines. However, material defects and process/operation-induced damage in the subsurface may 

significantly degrade the reliability and performance 

of these ceramics. These defects are typically less 

than 100 microns in size and within 200 microns 

under the surface. To detect and characterize these 

defects, Argonne National Laboratory (ANL) developed a laser scattering NDE method to measure detailed subsurface microstructures for ceramics. The 

objective of this research is to demonstrate that this 

method can be used to assess/evaluate the costeffectiveness and reliability of ceramic valves for 

diesel engines. The primary effort in FY 2004 was to 

evaluate accumulated impact/wear damages in 10 

NT551 silicon nitride valves tested in a bench rig for 

100, 500, and 1000 hours and to inspect 48 asprocessed (machined) SN235P silicon nitride valves 

for initial machining quality. Another effort was focused on NDE identification of fracture initiation 

defects in machined cylindrical SN235P silicon nitride specimens. This research is collaborated with 

Caterpillar, Inc. 

Approach

The critical region of ceramic components for 

structural applications is near-surface (usually to a 

depth of <200 μm). The common types of defects in 

this region are mechanical ones, such as cracks, 

spalls, inclusions, and voids. Because ceramics are 

partially translucent to light, a laser scattering 

method based on cross-polarization detection of optical scattering originated from the subsurface can be 

used for noncontact, nondestructive measurement of 

subsurface microstructure in such materials. By 

scanning the entire surface (flat or curved) of a ceramic component and constructing a twodimensional (2D) scatter image, 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 ceramic valves, an automated laser scattering system was developed by ANL for scanning 

entire valve surfaces. 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 2D scattering image data are used to 

identify the location, size, and relative severity of 

subsurface defects/damage. This system can therefore be used for inspection of the quality of asprocessed valves and for evaluation of accumulated 

damage in ceramic valves from rig/engine tests. 

Results 

Evaluation of Bench-Tested NT551 Valves 

Ten NT551 silicon nitride valves (Figure 1) 

were tested for 1000 h in a valve bench rig at Caterpillar. These valves were seated with metallic seat 

inserts at matched or mismatched angles, or with 

eccentric seat inserts. The bench test matrix is listed 

in Table 1.1

 These valves were examined by the 

automated laser scattering system at accumulated 

durations of 100, 500, and/or 1000 h as shown in 

Table 1. NDE data for each valve at different test 

durations were compared and analyzed to determine 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

85

Table 1. Bench test matrix and NDE evaluation intervals for NT551 valves 

Valve number 100-h 

valve-seat 

400-h 

valve-seat 

500-h 

valve-seat 

Total 

hours 

NDE at accumulated 

bench-test hours 

CV04, CV12 30-30 30-30 30-30 1000 500 

CV19, CV20 45-45 45-45 45-45 1000 500, 1000 

CV18, CV25 45-30 45-30 45-30 1000 500, 1000 

CV08, CV15 30-45 30-ECC 30-ECC 1000 100, 500, 1000 

CV14, CV16 30-ECC 30-ECC 30-ECC 1000 100 

Notes: numbers for valve-seat are their respective angles; ECC = eccentric seat insert. 

Figure 1. Ten NT551 silicon nitride valves. 

any new damage or growth of existing defects in the 

contact and fillet radius surfaces. 

Figure 2 shows the laser scattering image of the 

valve-head section of NT551 valve CV19 after a 

1000-h bench test. The scanned axial length was 

30 mm (the top 4-mm region is the contact area) and 

is aligned in the vertical direction. The scan resolution was 10 microns in the axial direction and 

11 microns in the circumferential direction. Typical 

of all NT551 valves, NDE data for CV19 indicated 

no significant accumulated damage from the bench 

tests, except surface wear within the contact surface. 

Figure 3 shows the detailed image of the contact 

surface. The reduction of scattering intensity within 

the wear scars (two darker horizontal stripes) was 

due to surface damage of increased roughness and 

contamination (embedded metal particles from the 

metal seat insert). Several segmented subsurface 

flaws of high optical-scattering intensity near the top 

Figure 2. Laser scattering scan image for NT551 valve 

CV19 after 1000-h bench test. 

Figure 3. Laser scattering image of the contact surface of 

NT551 valve CV19. 

edge are indicated in Figure 3. One of the detected 

edge flaws is shown in Figure 4 and compared with 

the surface photomicrograph. This flaw is pre-spall 

damage that was probably generated during valve 

processing. Other flaws within the subsurface are 

mostly material porosity of <80 μm in size. Valve 

CV19 also has several machining-damage marks in 

Figure 4. Detailed laser scattering image and photomicrograph of a subsurface edge flaw. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

86

the fillet radius surface. In addition, subsurface 

damage, near a scratch, that was not present after a 

500-h test was seen after a 1000-h test (Figure 5). 

Figure 5. Detailed laser scattering image and photomicrograph of a subsurface scratch damage. 

Inspection of As-Processed SN235P 

Valves

Forty-eight as-processed (machined) SN235P 

silicon nitride valves were inspected by the valvescan NDE system. These were diesel and natural-gas 

intake and outlet valves, each with a slightly different axial profile. Figure 6 shows a laser scattering 

image of a typical SN235P diesel intake valve. Spatial resolution in both scan directions is ~10 micron. 

Significant machining damage was found in the fillet radius and nearby stem regions, as shown in detailed images in Figure 7. The damage is seen as 

lines of high optical-scattering intensities along the 

machining direction. This result is consistent with 

the machining conditions for these valves. As indicated in ref. 1, because of an error in the machining 

process, the valve fillet and keeper notch regions 

were only rough-machined, while the contact surface and the majority of the stem regions were finemachined. The high damage level in the fillet region 

as detected by the NDE would likely reduce valve 

strength. 

Figure 6. Laser scattering image of a typical SN235P 

diesel intake valve. 

Figure 7. Detailed laser scattering images in various regions, as indicated in Figure 6. 

Machining damage in the valve keeper notch region was not anticipated to be significant. However, 

when a diesel valve was tested in the bench rig at 

Caterpillar, it was soon broken near the upper edge 

of the keeper notch. In comparison, ten NT551 

valves were tested in the same rig for 1000 h without 

a single failure. From scanning electron microscopy 

analysis, Caterpillar determined that the fracture origin was machining damage. NDE image data for a 

typical SN235P valve notch are shown in Figure 8. 

Very severe machining damage was found around 

the entire circumference of the upper edge, which is 

at the same location of the notch failure. The damage appears to be segmented; each segment is similar to the edge flaw of NT551 valve, as shown in 

Figure 4. The intensity of this damage is much 

higher than that within the notch which, however, is 

already considered to be more significant than typical rough machining damage. Because all machining 

damage is likely to be median cracks, these results 

indicate that laser scattering intensity may be correlated with crack depth and may predict failure. 

Figure 8. Laser scattering image in keeper notch region 

of a typical SN235P valve. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

87

Detection of Machining Damage and 

Fracture Origins 

Three sets of cylindrical SN235P silicon nitride 

specimens were examined by the laser scattering 

method to determine machining damage. Similar to 

the results for SN235P valves (Figures 6–7), NDE 

data showed damage level and distribution at different machining conditions. In addition, a color variation was observed on some rod surfaces, as seen in 

Figure 9. This coloration change is probably due to 

non-uniform mixing of powder materials during material processing, and it may have some effect on the 

mechanical property of the rods. 

Figure 9. Color variation in a machined rod. 

NDE data were correlated with destructive test 

results to determine the fracture initiation defect for 

these cylindrical SN235P specimens. The fracture 

surfaces were first examined to locate the fracture 

flaw site. This site was then mapped in the tensile 

(or cylinder) surface. By comparing the flaw location with a previously acquired laser scattering image for the surface, the fracture initiation defect can 

be identified in the NDE image. Figure 10 shows the 

fracture surfaces of a SN235P rod in which the fracture initiation defect can be easily observed. The 

correlation of the surface photomicrograph with the 

laser scattering image is seen in Figure 11. As indicated, the fracture initiation defect is not the largest 

in the region. This is because the strength test used a 

four-point bending fixture. The maximum stress is 

therefore localized within an axial line in a particular 

azimuth angle. Only those defects positioned within 

the maximum stress line will grow and cause fracture. The fracture initiation defect in this specimen 

appears to be a high-porosity material defect. 

Conclusions 

An automated laser scattering system developed 

by ANL was used to evaluate ten NT551 valves at 

accumulated bench-test durations of 100, 500, 

Figure 10. Fracture surfaces of a fractured SN235P 

rod. 

Figure 11. Photomicrograph and laser scattering image on tensile surface of a fractured SN235P rod. 

and/or 1000 h. NDE data for each valve at different 

test durations were analyzed to determine any new 

damage or growth of existing defects. In general, no 

significant subsurface damage from the bench tests 

was detected. Existing flaws, such as machining 

damage and inherent material defects, did not grow 

under the various test conditions and durations. 

Damage from the bench tests was mainly in the 

valve contact surface, including wear scars and an 

unknown deposit for some valves.1

 The NDE data 

will be correlated with the retained strength from 

destructive tests of these valves. 

Forty-eight as-processed (machined) SN235P 

valves were inspected by the valve-scan system. Because of an error in the machining process, the valve 

fillet and keeper notch regions were only roughmachined, while the contact and stem surfaces were 

fine-machined. The NDE data showed significant 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

88

machining damage in the rough-machined fillet radius and nearby stem regions. More severe damage 

was found at the upper edge of the keeper notch. 

The severe damage at the keeper notch was responsible for a valve failure, and the high damage level 

in the fillet region would likely reduce valve 

strength. Based on these NDE results, it was recommended that all SN235P valves be re-processed 

to remove all rough machining damage before being 

tested in an engine or a bench rig. 

The laser scattering method was also used to 

characterize defects/damage in cylindrical SN235P 

silicon nitride specimens that were machined by different vendors at various machining conditions. 

NDE data identified damage level and distribution at 

different machining conditions, and a color variation 

in base material that was probably due to material 

processing. In addition, NDE data were demonstrated for identification of fracture initiation defects 

for these cylindrical SN235P specimens. 

References

1. M. J. Andrews, “Lightweight Valve Train 

Materials,” quarterly progress report for January 1 

through March 31, 2004, Heavy Vehicle Propulsion 

Materials Program, U.S. Department of Energy. 

Publications/Presentations

J. G. Sun, J. Zhang, and M. J. Andrews, “Laser 

Scattering Characterization of Subsurface Defect/Damage in Silicon Nitride Ceramic Valves,” to 

be published in Proceedings of the American Ceramic Society's Annual Cocoa Beach Conference, 

Cocoa Beach, FL, Janury 25-30, 2004. 

Patents

J. G. Sun, “Device and Nondestructive Method 

to Determine Subsurface Microstructure in Dense 

Materials,” invention report ANL-IN-02-086, DOE 

case no. S-101,670, U.S. patent filed in December 

2003. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

89 

E. Durability of Diesel Engine Component Materials 

Peter J. Blau with Jun Qu and John J. Truhan, Jr*. 

Oak Ridge National Laboratory 

P. O. Box 2008, Mail Stop 6063 

Oak Ridge, TN 37831-6063 

(865) 574-5377; fax: (865) 574-6918; e-mail: blaupj@ornl.gov 

*University of Tennessee, Knoxville 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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 Number DE-AC05-00OR22725; 

Subcontractor: University of Tennessee, Knoxville, Tennessee 

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 coating technologies. 

Approach 

• Identify diesel engine components that need advanced materials or surface treatments in order to ensure durability while providing low-friction behavior. 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 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 (~ 600–700°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. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

90

• 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. 

• Prepared a model for scuffing tendency that considers lubricant characteristics and solid material characteristics. 

Future Direction 

• Use scuffing damage maps and test data to support the development of the new scuffing model and to obtain a 

better understanding of the evolution of localized surface damage. Compare results of the model with laboratory test data on metals, ceramics, and composite materials. 

Introduction 

The diesel engine industry faces important challenges to improve fuel efficiency in the face of increasingly strict emissions regulations. These 

challenges can be addressed by modifying the engine design and electronic control systems for engines, and by developing exhaust gas aftertreatments. Such modifications affect the mechanical, thermal, and chemical environments to which 

the engine materials are subjected; and the current 

structural materials may not perform as well as they 

did in previous designs. 

The objective of this effort is to enable the selection and development 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 test method development, microstructural analysis of candidate materials, mapping 

the effects of applied parameters on surface damage, 

and modeling. The focus on EGR components and 

fuel injector plungers was based on discussions with 

diesel engine manufacturers. Before developing tests 

to evaluate materials for improved durability, however, it was necessary to conduct a tribosystem 

analysis to understand the conditions under which 

the surfaces of these components materials must perform in an operating diesel engine. The nature of 

contact damage to engine components was reviewed 

to ensure that laboratory test methods would adequately reproduce that kind of damage. Then the test 

development, data analysis, and modeling tasks were 

begun. 

Approach 

In FY 2001, based on the definition of several 

key durability problems, a test method was developed to study the high-temperature friction and wear 

characteristics of candidate EGR system materials. 

The testing system we built continues to be used to 

evaluate metal alloys, ceramics, coatings, and other 

experimental materials for that application. In 

FY 2002, this 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 diesel engine 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 3-dimensional (3D) scuffing 

map was developed to graphically depict conditions 

for scuffing initiation and propagation in time and 

space domains. It was applied to explore the effects 

of surface finish and sliding velocity on materials 

and coatings for fuel injector systems. Such scuffing 

maps are intended to aid designers to select more 

durable materials and to determine optimum surface 

finishes. Near the end of FY 2004, a new scuffing 

model was developed that integrates boundary film 

characteristics with material properties. Chemical 

analyses of scuffed surfaces were conducted to investigate the wear modes and tribochemical film 

compositions in a fuel-lubricated environment, underpinning the basis for the model. Plans for 

FY 2005 involve refining and improving the model 

using experimental data concerning the friction and 

surface damage of advanced materials for durable 

diesel engine components. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

91 

Results

As reported previously, a pin-on-twin pins test 

was developed to evaluate the scuffing tendencies of 

diesel fuel injector plunger materials in commercial 

diesel fuel and in a low-sulfur fuel. That test was 

used to rank a variety of promising new materials 

such as cermets and ceramics. Data were in the form 

of exposure time prior to the onset of scuffing. In 

considering the data, it was quite clear that many 

different variables could affect scuffing, and there 

was a need for a better method to display the data 

for visualizing combinations of variables. Therefore, 

work in FY 2004 involved studying the current data 

and developing a format to display it. 

After the relevant issues were considered, a 3D 

friction-scuffing mapping technique was developed 

to present surface damage progression in the space 

and time domains. The concept is based on the assumption that an increase of the friction force at a 

certain location corresponds to lubricant film breakdown and surface morphology changes at that location. Scuffing is associated with asperity-level or 

grosser surface changes, such as plastic deformation 

and material displacement. Therefore, we might expect the rise in the plowing contribution of the friction force to reflect an increased incidence of surface 

deformation and hence become an indicator of scuffing. This correlation was verified by optical surface 

morphology examinations and electron microscopy. 

An example of a 3D friction-based scuffing map 

is provided in Figure 1. The change of friction coefficient relative to that of the first stroke in the test is 

portrayed as the vertical axis in Figure 1. One horizontal axis represents the location along the length 

of the stroke, and the orthogonal axis represents the 

cumulative time of sliding contact. Shaded horizontal bands, defined in the Figure 1 legend, represent 

ranges of friction coefficient. The test used to construct Figure 1 used a reciprocating pin-on-twin pins 

configuration and was conducted on self-mated AISI 

52100 steel in ultra-low-sulfur fuel (Jet A aviation 

fuel). By defining a threshold change in friction coefficient to signal the initiation of scuffing (e.g., an 

increase of 10% over the initial friction coefficient), 

it is possible to use the friction-scuffing map to indicate the location and progress of scuffing damage. 

Figure 1. A 3D friction-scuffing map for self-mated bearing steel is plotted in space and time domains. 

Three-dimensional friction-scuffing maps like 

that in Figure 1 were subsequently used to construct 

scuffing transition diagrams to portray the effects of 

surface roughness and sliding velocity on scuffing 

resistance for diesel fuel injector materials, such as 

AISI 52100 steel and transformation-toughened zirconia (TTZ). 

Figure 2 depicts the time until the first instance 

of scuffing for specimens with different initial surface roughness in ultra-low-sulfur fuel. Scuffing 

usually initiates locally, generally at the ends of the 

stroke; and surface damage may or may not spread 

along the whole sliding stroke during the course of 

the test. In Figure 2, the ‘local scuffing’ curve refers 

to the moment of the onset of local scuffing, and the 

‘global scuffing’ curve implies when the whole 

stroke was scuffed. The region between these two 

curves represents the period of scuffing propagation 

from local spots to the whole stroke. This outsideinward spreading is a consequence of the sliding 

motion. The velocity of our pin-on-twin reciprocating test follows a sinusoidal wave form. The sliding 

speed starts at 0 at the turnaround points and reaches 

a maximum at the stroke’s midpoint. Scuffing, if it 

occurred, generally appeared at the stroke ends first, 

because there is less effective lubrication there and 

the lack of lubrication leads to more solid-to-solid 

contact. 

Compared with self-mated steel, zirconia against 

steel performed better at each surface finish 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

92

Figure 2. Scuffing transition diagrams for 52100 steel 

and TTZ in ultra-low-sulfur fuel. 

level. Smoother contact surfaces tend to produce 

thinner lubricant films and more chance for incidental contact, but rougher surfaces have higher local contact pressure under boundary lubrication conditions. Either case may cause the lubricant to fail. 

Therefore, there is likely to be an optimum surface 

finish to provide the best scuffing resistance for a 

given lubricated sliding system. The steel specimens 

with intermediate roughness outperformed other two 

surface finish levels (Figure 2a), while rougher surfaces showed higher scuffing resistance for TTZ 

against steel (Figure 2b). The composite roughness*

of the current-production zirconia plunger and 

52100 steel bore in heavy-duty diesel engines is 

about 0.27 μm. To meet stricter diesel engine emis-

 

*

 Composite roughness is defined as the square root of the 

sum of the squares of the arithmetic average surface 

roughnesses of both mating surfaces. 

sions standards, superior surface finish and tighter 

tolerances are required to seal the higher injection 

pressures. Better surface finish would reduce the 

scuffing resistance of the injector system, according 

to the scuffing transition diagrams in Figure 2b. This 

necessitates the investigation of new materials or 

surface engineering processes to resist scuffing under the more stringent operating conditions. 

Under boundary lubrication, tribochemical films 

formed on contact surfaces play an important role in 

scuffing behavior. Depending on the specifics of the 

tribosystem, these in-situ formed films can be either 

protective or detrimental.1

 In the current investigation, the surface films formed in our fuel-lubricated 

scuffing tests have been examined and analyzed using optical microscopy, scanning electron microscopy (SEM), energy dispersive spectroscopy (EDS), 

and scanning auger microprobe (SAM). 

For zirconia/steel pairs with relatively smooth 

surfaces, a tribo-film could easily be observed under 

a light optical microscope on both sides of the contact surfaces after scuffing. The brown or darkreddish color suggested the presence of iron oxides 

such as FeO and/or Fe2O3. EDS analysis confirmed 

that oxygen was present in this layer. The scuffing 

process of zirconia against steel seemed to be a 

competition between the formation and removal of 

this tribochemical film. Although the oxide layer 

may protect the original surfaces from abrasive wear 

to some extent, its non-uniformity and the resulting 

adhesion and spallation problems make it undesirable for fuel-injection systems with very tight geometric clearances. 

To better understand the scuffing transition in 

fuel-lubricated systems, the tribochemical surface 

films at an early scuffing stage were analyzed using 

a PHI 680 SAM. Figure 3 shows the sputtering 

depth profiles of elemental compositions for the surface films on and off the wear scar on a bottom steel 

pin (against steel) in Jet A fuel. There was a relatively thick (about 80 nm) iron-oxide-rich layer on 

the worn surface, compared with that (< 4 nm) on 

the unworn surface. This demonstrated that tribochemical excitation, such as frictional heating and 

mechanical deformation, significantly accelerated 

the oxidation process.2

 The surface film formed in 

#2 diesel fuel showed a thinner oxide-rich layer with 

lower oxygen concentration than the film formed in 

Jet A fuel. This implied that the diesel fuel “cooled” 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

93 

Figure 3. SAM sputtering profiles of surface films that 

were formed during tests in ultra-low-sulfur fuel. 

the contact area more effectively, which probably 

helped to resist scuffing. The diesel-fuel-lubricated 

surface also had a higher non-carbide carbon concentration, which suggested that possibly more 

polymeric materials or organometallic compounds, 

thought to be generally protective, were deposited 

on the contact surfaces. These observations helped 

to explain the higher scuffing resistance in #2 diesel 

fuel than in Jet A fuel. 

In late FY 2004, a new model for scuffing was 

developed. It was based on the recognition that a 

sequence of interfacial events needs to occur to set 

the stage for scuffing. First, the lubricant film must 

cease to effectively protect and separate the surfaces; 

next, the solids in contact must begin to accumulate 

damage to the point that significant deformation, 

roughening, and elevated friction occur. These phenomena have been embodied in a model for scuffing 

that was described in a project milestone report (see 

publications list) and will be further evaluated and 

refined in the light of additional experiments to be 

done in FY 2005. 

Conclusions

• 3D friction-scuffing maps have been developed 

to depict scuffing initiation in both time and 

space domains. These maps were used to construct scuffing transition diagrams that correlate 

scuffing resistance to surface finish and sliding 

velocity. 

• In the range of surface roughness tested in this 

study, zirconia against steel showed the scuffing 

resistance to be proportional to the composite 

surface roughness, while self-mated steel exhibited optimum performance with an intermediate 

surface finish. 

• A thinner oxide layer and a higher non-carbide 

associated carbon concentration were detected 

on the #2 diesel-fuel-lubricated surfaces. This 

implied lower temperature and more protective 

organometallic compounds formed on the contact area. These observations could partially explain the higher scuffing resistance in normal 

diesel fuel than in ultra-low-sulfur fuel 

• A combination of experimental data, surface 

observations, and derived scuffing transition 

diagrams led to the development of a new scuffing model. This model will be tested and validated in the year to come. 

References

1. S.M. Hsu and R. S. Gates, “Boundary Lubrication and Boundary Lubricating Films,” in Modern Tribology Handbook, ed. B. Bhushan, Vol. I, 

CRC Press, Boca Raton, pp. 455–492, 2001. 

2. G. Heinicke, Tribochemistry, Carl Hanser 

Verlag, Munchen-Wien-Berlin, 1984, p. 97. 

Publications/Presentations

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, 

FL, October 28, 2003. 

P. J. Blau, J. J. Truhan, and J. Qu, Scuffing 

Transition Diagrams for Fuel Injection System Materials, Project Milestone Report, Oak Ridge National Laboratory, March, 19, 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,” International Journal of Engine Research (in press). 

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, (in press). 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

94

P. J. Blau with 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, September, 17, 2004. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

95

F. Life Prediction of Diesel Engine Components 

H. T. Lin, T. P. Kirkland, A. A. Wereszczak, and M. K. Ferber 

Oak Ridge National Laboratory 

Oak Ridge, TN 37831-6068 

(865) 576-8857; fax: (865) 475-6098; e-mail: linh@ornl.gov 

Jeremy Trethewey and M. J. Andrews 

Caterpillar, Inc. 

Peoria, IL 61656-1875 

(309) 578-0056; fax: (309) 578-2953; e-mail: trethewey_jeremy_s@cat.com 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-2476; e-mail: sid.diamond@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 Number: DE-AC05-00OR22725 

Subcontractor: Caterpillar, Inc., Peoria, Illinois 

Objective 

• Generate database and characterize damage mechanisms of candidate advanced ceramics and intermetallic alloys. 

• Apply and verify probabilistic life prediction and components design and verification for advanced diesel engine components. 

Approach 

• Evaluate the dynamic fatigue and rotary bending fatigue performance of candidate silicon nitride ceramics and 

titanium aluminide (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 linked to the longterm mechanical performance and reliability of ceramics and intermetallic alloys. 

• Predict the failure probability and reliability of complex-shaped components subjected to application conditions 

via the use of life prediction codes. 

Accomplishments 

• Completed mechanical properties of prototype NT551 silicon nitride exhaust valves after 500-h bench rig test 

• Completed dynamic fatigue study for Kyocera SN235P silicon nitride, fabricated from the same powder batch 

for exhaust valve blanks, after 100-h oil immersion test. 

• Completed development of dynamic fatigue database for a commercial-grade silicon nitride, SN147-31E, acquired from Ceradyne, Inc 

• Completed mechanical property verification for flexure test bend bars machined from Kyocera SN235 exhaust 

valve blanks for mechanical performance verification. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

96

Future Direction 

• Characterize microstructure and mechanical properties of Kyocera SN235P silicon nitride exhaust valves after 

engine field demonstration at the National Transportation Research Center (NTRC) and verify the probabilistic 

component design and life prediction. 

• Develop a dynamic fatigue database for Ceradyne SN147-31N silicon nitride fabricated from the same power 

batch for exhaust valve blanks. 

• Verify mechanical properties for Ceradyne SN147-31N silicon nitride machined from the exhaust valve blanks. 

• Develop the mechanical database for the specimens extracted from diesel particulate filter (DPF) substrates for 

long-term mechanical reliability 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, long-term durability, and reduction in nitrogen oxides and CO exhaust emission as 

required in the 21st Century Truck Program. This 

interest has focused primarily on research aimed at 

characterization and design methodology development (life prediction) for advanced silicon nitride 

ceramics and TiAl alloys in order to manufacture 

consistent and reliable complex-shaped components 

for diesel engine applications. The valid prediction 

of mechanical reliability and service life is a prerequisite for the successful implementation of these advanced materials as internal combustion engine 

components. 

This research project has three primary goals, 

which contribute toward successful implementation: 

• 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 

components fabricated from them 

• 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 the material suppliers and component end-users 

to refine and optimize the processing parameters to 

achieve consistent mechanical reliability, and to 

validate the probabilistic design and life prediction 

of engine components made from these advanced 

materials. 

Approach

All silicon nitride test bend bars were longitudinally or transversely machined according to ASTM 

standard C1161 from production billets purchased 

from material suppliers.1

 Flexure testing was conducted in ambient air in four-point bending using 

20/40- mm, α-SiC, semi-articulating fixtures at temperatures ranging from 20 to 1000°C and at a stressing rate 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; 

the 0.003 MPa/s test condition was chosen to measure the change in slow crack growth (SCG) susceptibility at elevated temperatures. Pneumatic actuators 

were programmed to produce the desired loading 

rate (and corresponding stressing rate) via a personal 

computer. Load was continuously measured as a 

function of time, and flexure strength was calculated 

using ASTM standard C1161. The accumulated 

strength data were then further analyzed. The 

strengths for each test set were fit to a twoparameter Weibull distribution using the program 

CERAMIC,2

 which uses the maximum likelihood 

estimation advocated in ASTM C1239.3

 Reported 

results are uncensored because fractography analysis 

was not conducted in detail to identify strengthlimiting flaws for all of the bend bars tested. Following the dynamic fatigue test, both optical and scanning electron microscopy analysis were carried out 

on fracture surfaces and polished cross-sections of 

selected bend bars to characterize the fracture as 

well as degradation mechanisms. X-ray analysis was 

also carried out to evaluate the possible phase 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

97

changes resulting from oxidation during dynamic 

fatigue testing or after long-term exposure to simulated engine environments, which could possibly 

cause degradation in mechanical performance and 

reliability. 

Results

Dynamic Fatigue Response of Kyocera 

SN235P Silicon Nitride 

Kyocera SN235P silicon nitride was downselected by Caterpillar for engine demonstration because of its excellent resistance to slow crack growth 

and to a simulated diesel engine environment at elevated temperatures. To ensure a consistent microstructure and chemistry in the secondary phase, and 

thus mechanical reliability, of the valve blank materials, bend bars were machined from co-processed 

billets for long-term oil immersion as well as exhaust gas study. Bend bars were also machined from 

the valve blank itself for mechanical comparison. 

The results obtained would also allow Caterpillar to 

verify its probabilistic component design and life 

prediction tasks. The test silicon nitride samples 

were placed in platinum crucibles and covered with 

commercial-grade 10W30 engine oil. The crucibles 

with oil-covered silicon nitride bend bars were 

heated in a furnace at 600°C for approximately 

30 minutes to ash the oil. After completion of oil-ash 

conversion, the samples covered with 1–2 mm oil 

ash powder were then heat-treated at 850°C for 

1000 h in ambient air. In addition, SN235P bend 

bars were exposed to exhaust gas for 1000 h. The 

results previously generated for the as-machined 

SN235P samples, co-processed with the exhaust 

valve blanks, showed a very high fatigue exponent 

(N ~ 93) at 850°C, indicative of no susceptibility to 

the SCG process.4,5 Therefore, the dynamic fatigue 

tests for exposed SN235P samples were carried out 

only at 850°C and 0.003 MPa/s in air according to 

ASTM C1465 because of the limited number of 

specimens. 

Results of tests carried out at 20 and 850°C and 

at 30 MPa/s showed that the SN235P valve blank 

material exhibited very consistent mechanical characteristic strength and Weibull modulus with respect 

to those SN235P materials evaluated previously in 

the program. Dynamic fatigue tests at 850°C and 

0.003 MPa/s show that both of the oil-immersed as 

well as exhaust-gas exposed SN235P silicon nitride 

specimens exhibit an inert characteristic strength 

comparable to those obtained from the as-machined 

samples tested under the same condition (Table 1). 

Also, results show that there is no difference in 

characteristic strength values between oil-ash–

exposed and exhaust-gas–exposed samples. In addition, both of the exposed SN235P samples showed a 

higher Weibull modulus than the as-machined samples. The higher Weibull modulus of exposed samples could arise from the surface sealing effect due 

to the formation of oxide scale after exposure. The 

dynamic fatigue results suggest that the SN235P-CP 

exhibits excellent corrosion resistance to diesel engine environments and no susceptibility to the SCG 

process at test temperatures. Thus the SN235P material would meet the application criteria for exhaust 

valve components of advanced diesel engines. 

Dynamic Fatigue Response of Ceradyne 

SN147-E Silicon Nitride

Studies of dynamic fatigue properties of a commercial-grade silicon nitride, i.e., SN147-31E, 

manufactured by Ceradyne Advanced Ceramic, Inc., 

were completed this year. The objective of this study 

is to extend the database generation efforts to other 

potential candidate silicon nitride ceramics, especially those manufactured by domestic material suppliers. The SN147-31E was processed with sintering 

additives (i.e., Al2O3 and Y2O3) similar to those employed for SN147-31N that were evaluated and reported previously. However, the SN147-31E 

contains a crystalline secondary phase achieved by 

post–heat treatment. Thus it is anticipated that 

SN147-31E would exhibit a higher temperature mechanical reliability than SN147-31N because of the 

presence of the crystalline secondary phase. 

Dynamic fatigue results showed that the longitudinally machined SN147-13E exhibited a characteristic strength that is 20% lower than that obtained 

from the SN147-31N machined and tested under the 

same conditions (Table 2). The lower strength of 

SN147-31E could be due to the change in internal 

residual stress resulting from the crystallization 

process of the secondary phase(s) by post–heat 

treatment. In addition, the transversely machined 

SN147-31E showed a minor decrease (~8%) in 

characteristic strength as compared with the SN147-

31N. The minor difference in strength between 

SN147-31E and –31N, which were transversely 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

98

Table 1. Summary of uncensored Weibull and strength distributions for Kyocera SN235P silicon nitride 

specimens longitudinally machined per ASTM C1161from co-processed billets and exhaust valve blanks 

before and after exposure to simulated diesel engine environments. Data generated for SN235P evaluated 

previously in the program are used for reference. 

Material 

# of 

spmns. 

tested 

Stressing 

rate 

(MPa/s) 

Temp. 

(°C) 

Uncens. 

weibull 

modulus 

± 95% 

Uncens. 

weibull 

modulus 

Uncens. 

chrctstic 

strength 

(MPa) 

± 95% 

Uncens 

chrctstic 

strength 

(MPa) 

SN235P 10 30 20 32.6 

18.8 

49.6 792 

773, 

810 

SN235P 10 30 850 30.3 

16.5, 

48.1 673 

643, 

691 

SN235P 10 0.003 850 23.6 

12.7, 

38.8 641 

662, 

641 

SN235PCPB 30 30 20 11.94 8.58. 

15.94 820 

794, 

847 

SN235PCPB 30 0.003 20 14.56 10.93, 

18.64 741 

721, 

761 

SN235PCPB 30 30 850 20.56 15.16, 

26.87 684 

671, 

697 

SN235PCPB 30 0.003 850 19.44 14.49, 

25.01 621 

608, 

633 

SN235PEVB 15 30 20 17.58 11.39, 

25.07 895 

866, 

924 

SN235PEVB 15 30 850 10.42 6.64, 

15.23 732 

691, 

772 

SN235PCPB-Oil 15 0.003 850 30.40 19.47, 

44.19 660 

647, 

672 

SN235PCPBExhGas 12 0.003 850 25.07 

15.55, 

36.81 677 

659, 

694 

CPB: specimens were machined from co-processed billets of exhaust valve blanks. 

EVB: specimens were machined from exhaust valve blanks. 

machined, could be due to the similarly dominant 

strength-limiting flaws, i.e., surface machining 

flaws. However, SN147-31E exhibited a lower 

Weibull modulus than that obtained for SN147-31N. 

On the other hand, results at 20°C and 0.003 MPa/s 

showed that both longitudinally and transversely 

machined SN147-31E exhibited little or no decrease 

in characteristic strength as compared with results 

obtained at 30 MPa/s, as shown in Table 2. Results 

also show that the Weibull moduli were not sensitive 

to the stressing rate. In addition, both materials exhibited high fatigue exponents at room temperature 

(Figures 1 and 2), indicative of no susceptibility to 

the SCG process. However, results at 850°C showed 

that the SN147-31E exhibited a relative low fatigue 

exponent of 56 with respect to SN147-31N 

(N ~ 131), indicative of a susceptibility to the SCG 

process at test temperatures (Figure 3). The low fatigue exponent obtained for SN147-31E suggests an 

incomplete crystallization of the glassy phase during 

post–heat treatment. The presence of these residual 

glassy phase(s) could play an important role in the 

high-temperature mechanical performance and susceptibility to SCG. 

The study of dynamic fatigue behavior in fourpoint bending for transversely machined SN147-

31N silicon nitride after 1000 h of exposure to an oil 

ash environment was also completed during this reporting period. The current study was carried out to 

understand the effect of long-term exposure to a 

simulated diesel engine environment on the mechanical performance of SN147-31N. As a result of 

the limited number of bend bars, the dynamic fatigue tests for exposed SN147-31N samples were 

carried out only at 850°C and 0.003 MPa/s in air 

according to ASTM C1465. Results of dynamic 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

99

Table 2. Summary of uncensored Weibull and strength distributions for Ceradyne SN147-31E silicon 

nitride specimens transversely machined per ASTM C1161. Data of SN147-31N machined longitudinally as well as transversely are used for reference. 

Material 

# of 

spmns. 

tested 

Stressing 

rate 

(MPa/s) 

Temp. 

(°C) 

Uncens. 

Weibull 

modulus 

± 95% 

Uncens. 

Weibull 

modulus 

Uncens. 

chrctstic 

strength 

(MPa) 

± 95% 

Uncens 

chrctstic 

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-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-

31N1000hOil* 15 850 0.003 15.54 10.31, 

21.65 600 

578, 

622 

*1000h Oil: specimens were transversely machined and exposed to an oil ash environment for 1000h.

fatigue tests at 850°C and 0.003 MPa/s showed that 

the exposed SN147-31N silicon nitride exhibited an 

inert characteristic strength comparable to results 

obtained from the as-machined samples tested under 

the same condition (Table 2). Also, there is no difference in Weibull modulus between oil-ash exposed 

and as-machined samples. The dynamic fatigue results suggest that the SN147-31N exhibits excellent 

corrosion resistance in diesel engine environments, 

similar to those observed for SN235 and GS44 silicon nitride ceramics evaluated previously. Therefore, one could conclude that silicon nitride 

ceramics, in general, exhibit superior corrosion resistance in diesel engine environments compared 

with those metallic alloys. 

Mechanical Properties of Diesel Particulate Filters

This study, part of a cooperative research and 

development agreement between Cummins and Oak 

Ridge National Laboratory, is carried out to characterize the mechanical properties of a DPF substrate 

and develop analytical tools for predicting long-term 

reliability and durability. Miniature test samples 

have been successfully machined from the substrates 

of DPFs acquired from Corning (Figure 4a) during 

this reporting period. The samples have nominal 

dimensions of 0.20 × 0.90 × 15 mm. A miniature test 

fixture has also been designed and machined (Figure 4b). The alumina test fixture has inner and outer 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

100

Figure 1. Failure strength vs stressing rate curve of 

SN147-31E longitudinally machined and tested at 20°C in 

air. 

Figure 2. Failure strength vs stressing rate curve of 

SN147-31E transversely machined and tested at 20°C in 

air. 

spans of 1 and 2 mm, respectively. Figure 5 shows 

the uncensored Weibull strength distribution of DPF 

samples. Results indicate that the DPF material exhibits low characteristic strength as well as Weibull 

modulus because of its porous microstructure. The 

strength data will be also generated from the fullsize DPF samples for comparison. The data generFigure 3. Failure strength vs stressing rate curve of 

SN147-31E longitudinally machined and tested at 850°C 

in air. 

 

Figure 4. (a) Photos show a section of diesel particulate 

filter, and sample machined from it. (b) Photos show the 

micro test system and miniature test fixture. 

ated for the miniature test samples from DPF substrates would be used for life prediction. 

References

1. American Society for Testing and Materials, 

“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, PA, 1999. 

2. Life Prediction Methodology for Ceramic 

Components of Advanced Heat Engines, Phase 1, 

ORNL/Sub/89-SC674/1-2, prepared by AlliedSignal 

102

103

0.001 0.01 0.1 1 10 100

Failure Strength (MPa)

Stressing Rate (MPa/s)

Dynamic Fatigue of SN147-31E Si3N4

Uncensored Flexure Strength at 850°C

N ~ 56

102

103

0.001 0.01 0.1 1 10 100

Stressing Rate (MPa/s)

Dynamic Fatigue of SN147-31E Si3N4

Longitudinally Machined

Uncensored Flexure Strength at 20°C

Failure Strength (MPa)

N ~ 130

102

103

0.001 0.01 0.1 1 10 100

Stressing Rate (MPa/s)

Dynamic Fatigue of SN147-31E Si3N4

Transversely Machined

Uncensored Flexure Strength at 20°C

Failure Strength (MPa)

N ~ 338

1 mm
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

101

Figure 5. Failure strength vs stressing rate curve of 

SN147-31E longitudinally machined and tested at 20°C in 

air. 

Engines, Department of Energy Office of Transportation Technologies, Washington, D.C., 1995. 

3. American Society for Testing and Materials, 

“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, PA, 

1999. 

4. 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). 

5. H. -T. Lin, T. P. Kirkland, A. A. Wereszczak, 

and M. J. Andrews “Effect of Long-Term Oil Immersion Test on Mechanical Reliability of Candidate 

Silicon Nitride Ceramics for Diesel Engine Applications,” pp. 261–273 in Ceramic Transactions, Vol. 

142, “Silicon-Based Structural Ceramics for the 

New Millennium” (2003). 

Publications

1. H. -T. Lin, T. P. Kirkland, A. A. Wereszczak, 

and M. J. Andrews, “Effect of Long-Term Oil Immersion Test on Mechanical Reliability of Candidate 

Silicon Nitride Ceramics for Diesel Engine Applications,” pp. 261–273 in Ceramic Transactions, Vol. 

142, “Silicon-Based Structural Ceramics for the 

New Millennium” (2003). 

2. H. T. Lin, T. P. Kirkland, A. A. Wereszczak, 

and M. J. Andrews, “Strength Retention of Silicon 

Nitride After Long-term Oil Immersion Exposure,” 

submitted to J. Mater. Sci.

-6

-5

-4

-3

-2

-1

0

1

2

1 10 100

Max. Lik. Fit

Strength Data

ln ln ( 1 / ( 1 - P f ) )

Failure Stress (MPa)

Corning Diesel Particulate Filter

Uncensored Strength Distribution

20°C - 0.01 mm/s - As-processed surface

Probability of Failure, P f , (%)

99.9

99.0

90.0

50.0

20.0

10.0

2.0

5.0

1.0

0.5

m = 4.98 (3.02, 7.28)

σθ

 = 21.70 MPa (18.70, 25.00)

n = 10 specimens

95% Confidence

Bands Shown

ASTMC1161-B

3-point-bend

1.5 mm span
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

102
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

103 

G. Low-Cost Manufacturing of Precision Diesel Engine Components 

Principal Investigator: P. J. Blau 

Oak Ridge National Laboratory 

P.O. Box 2008, Oak Ridge, TN 37831-6063 

(865) 574-5377; fax: (865) 574-6918; e-mail:blaupj@ornl.gov 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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 

• Develop and demonstrate optimized, cost-effective fabrication processes for producing precision components 

for use in diesel engines. 

• Develop and demonstrate optimized, cost-effective methods to detect, characterize, and minimize machininginduced damage in advanced materials for engine components. 

Approach 

• Collaborate with manufacturers and suppliers of heavy vehicle engine components to identify enabling technologies for the cost-effective machining of precision components, such as valves, valve guides, and fuel injection system components, from advanced materials. 

• Utilize Oak Ridge National Laboratory’s (ORNL’s) instrumented machine tools and dimensional metrology 

and surface characterization instruments to correlate the microstructures of advanced materials with their machining characteristics. 

• Work with a consortium of software developers and component manufacturers to develop improved models for 

high-speed machining of titanium alloys. 

Accomplishments 

• Developed a method to measure the wear of super-abrasive grinding wheels that use “witness blocks” of easyto-machine material to profile worn and unworn areas of the wheel. 

• Machined silicon nitride engine valves from dimensionally-characterized starting blanks provided by Caterpillar Corporation. 

• Initiated a new project to correlate the microstructure and hard particle concentration of cermets with their machining behavior. 

• Joined and participated in a consortium led by Third-Wave Systems (TWS) to model high-speed machining of 

titanium alloys. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

104

Future Direction 

• Continue to work with the TWS machining consortium to develop computer models and to provide supporting 

data for those models to enable low-cost machining of advanced, lightweight engine materials. 

• Survey the state of the art and conduct systematic studies of the effects of hard particle concentration on the 

machining characteristics and surface quality of advanced composites based on metal-bond ceramic particles. 

Introduction 

Manufacturing involves the conversion of raw 

materials into finished, useful components and 

products. Historically, it represents a vital industrial 

component of the economic security of the United 

States. Manufacturing begins with processed materials, such as metals, ceramics, polymers and composites, and then shapes, joins, and finishes them to 

meet the functional requirements of machines, machine parts, and structures. Manufacturing comprises 

a broad range of engineering disciplines that are applied to an even wider range of unit processes. It is 

necessary, therefore, to focus on an area of manufacturing that, in the present case, enables both the 

near-term and the longer-term introduction of advanced materials and machining methods for diesel 

engine components into the commercial marketplace. 

For more than 150 years, since the Industrial 

Revolution, U.S. industry has invested hundreds of 

billions of dollars in the development of machining 

technology and the purchase of equipment. However, with the egress of manufacturing to foreign 

sources in recent years, a significant amount of capital equipment and the knowledge base to develop 

new machining technology to enable the introduction of advanced materials has eroded. Experienced 

machinists have retired and, with fewer manufacturing jobs in the United States, they are not being replaced by motivated, young trainees from the next 

generation. The machine tool industry in the United 

States also has suffered significant losses and many, 

if not most, new machine tools are made by foreign 

sources. It is vital, therefore, to avoid full dependence on foreign sources and to maintain a base of 

machining science in the United States. 

Approach 

The approach taken in this project involves three 

areas of machining technology: 

1. Machining engine parts from advanced materials, such as ceramic engine valves. 

2. Modeling machining processes with state-of-theart software. 

3. Characterizing the machining response of 

newly-developed materials, such as composites. 

Projects in each of these areas are separately 

discussed in the Results section. 

Results

Machining of silicon nitride diesel engine 

valves. Silicon nitride valve blanks were received by 

ORNL from Caterpillar, and their straightness and 

diameters were measured using a coordinate measuring machine. The appropriate dimensions to use, 

considering several valve designs and the coordinate 

measuring machine data for the blanks, were then 

selected. A set of programming instructions was developed for use with the Weldon grinder. This program was then tested and verified using a bar of 

machinable alumina silicate ceramic. Profiles on the 

“dummy part” were checked with the coordinate 

measuring machine. Particular attention was paid to 

the dimensions of both the bevel and the blend section between the valve stem and head. The finished 

valve is shown in Figure 1. 

Participation in the machining consortium. In 

FY 2004, a new collaboration began with TWS of 

Minneapolis. TWS has been leading a U.S. Air 

Force Phase II Small Business Innovative Research 

project, “Improved Titanium Machining Processes,” 

to develop an advanced machining modeling package based on finite element modeling (FEM). TWS 

has established a consortium of aerospace and 

ground transportation component manufacturers that 

are interested in improving machining processes for 

titanium alloys. Benefits from membership in the 

consortium are 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

105 

Figure 1. Silicon nitride diesel engine valve machined at 

ORNL. 

1 Members share the latest technologies for titanium machining from other consortium members, such as TWS, Kennametal, Caterpillar, 

Stork Fokker (Netherlands), and EADS (Germany). 

2. Members have the use of the advanced FEMbased machining modeling package called AdvantEdgeTM. 

3. Participants help define the direction of the machining technology being developed, specifically, working with the University of Michigan 

to develop and verify models for drilling, boring, and face milling of titanium, in addition to 

having access to other models in the current 

modeling package. 

4. ORNL helps the consortium by conducting material characterization to help select new cutting 

tool materials and tool coatings for improved 

machining of titanium. 

5. New modeling capabilities will not be limited to 

titanium but will represent a new in-house resource for materials development programs and 

the High Temperature Materials Laboratory Machining, Inspection, and Tribology User Center. 

The basis for our interest in the TWS consortium stems from recent technology that promises to 

reduce the cost of titanium raw material1–3 and therefore expand titanium’s potential use as a lightweight 

material in engines and brakes.4

 However, the difficulty of high-speed machining of titanium alloys is a 

major technical barrier. Because of the metal’s poor 

thermal conductivity (6.8 W/ 

m-K), the cutting heat cannot be dispersed efficiently, and the tool tip temperature is much higher 

than that generated in cutting most other metals. The 

high tool temperature promotes the diffusion tool 

wear and limits the cutting speed. Serrated chips 

with adiabatic shear band are usually produced in 

titanium machining because of the thermalmechanical instability. The chip shear band formation creates fluctuation in cutting forces, leading to 

tool chipping. Therefore, research is needed in titanium machining, including advanced process modeling and development of new tool materials/coatings. 

Jun Qu of ORNL and Paul Becker of the University of Tennessee attended the 2004 TWS AdvantEdge international users’ conference in March 

in Gaithersburg, MD, and Jun also took the short 

course in modeling metal cutting offered by TWS in 

April. 

Output from the model. The TWS modeling 

package AdvantEdge can be used to improve cutting 

speeds and evaluate tooling performance by simulating and predicting cutting forces and tool/workpiece 

/chip temperatures. AdvantEdge may also be used to 

address other important issues, such as chip breakage, tool wear, and residual stresses. 

Simulations were first run on several steel and 

aluminum alloys using different tool materials, tool 

geometries, and cutting parameters to become familiar with the modeling technique and test the feasibility of AdvantEdge. Then the focus was shifted to 

titanium alloys. For example, AdvantEdge was used 

to simulate performance in cutting (turning) Ti-6Al4V alloy with different tool inserts (bare WC-Co and 

TiAlN coatings with thicknesses of 3 and 6 μm). 

The cutting speed was set to 195 m/min with a 

0.381-mm feed and a 1.02-mm depth of cut. The 

steady-state temperature maps of the cutting areas 

are shown in Figure 2. Figure 3 displays the ex- 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

106

Figure 2. Temperature maps of coated and uncoated tools 

cutting Ti-6Al-4V. 

Figure 3. Rake face temperature profiles. 

tracted rake face temperature profiles. Apparently, 

the temperature at the uncoated tool tip is higher 

than that of the coated tips. It is interesting to note 

that the tip with the 6-μm coating is hotter than the 

tip with the 3-μm coating. This implies that a thicker 

coating may not improve tool performance. Thicker 

coatings also increase the production cost and are 

more vulnerable to spalling. In fact, 3 μm is the current industry standard thickness for TiN and TiAlN 

tool coatings. Besides having lower cutting temperatures, coated tools generate smoother cutting surfaces, as illustrated in Figure 3. 

Tooling studies. In addition to process modeling, new tool materials for titanium are being explored. One potential tool material is a TiC 

(particles)–Ni3Al (matrix) composite that was initially developed for heavy-duty diesel fuel injector 

plungers. It contained 40–50 vol % TiC and 50–60 

vol % Ni3Al. Experimental work at the University of 

Michigan has proved that this composition is inadequate for cutting tools. Rather, TiC–Ni3Al materials 

with high concentrations of TiC (90–95 vol %) are 

being prepared and will be tested in FY 2005. 

The physical properties of five production tool 

materials and coatings are also being characterized. 

The coating thicknesses and microstructures of each 

material were examined by both scanning electron 

microscopy and scanning acoustic microscopy, a 

novel nondestructive method for observing micrometer-sized surface flaws. The coating hardness 

and modulus of elasticity have been measured by 

nanoindentation. The thermal conductivity and 

thermal diffusivity of the tool coatings significantly 

affect the cutting performance and tool life. The 

thermophysical properties of these thin coatings will 

be measured using a newly developed photoacoustic 

technique at the Microscale Thermophysical Property Laboratory at Purdue University. 

In titanium machining, the high tool wear, particularly diffusion wear, significantly limits the cutting speed and tool life. Measurement and 

characterization of the tool wear are upcoming to 

investigate the wear mechanisms and explore an effective method to quantify the wear at the tool tips. 

Machining of hard composite materials.

Composite materials are used because they can draw 

from the advantageous properties or two or more 

materials to create a new material with properties 

better than those of either constituent alone. They 

find a wide variety of uses in structural and nonstructural applications. Composite materials in 

which the various phases in the microstructure differ 

greatly in hardness represent one of the most challenging problems for machining in cases when it is 

necessary to achieve high-precision surface finishes 

and minimal surface damage. Fine particles of grit 

embedded in the grinding wheel can produce differential grinding of the hard and soft phases. Pull-out 

of the hard particles may leave holes and other artifacts that can affect surface quality. In the coming 

year, a series of cermets with different concentrations of hard phases will be prepared by ORNL and 

used to study the behavior of cermets and metal matrix composites. 

In preparation, ORNL is conducting a survey of 

machining practices for drilling, turning, and grind-
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

107 

ing cermets and metal matrix composites. Then the 

cermets will be machined under a range of conditions to study how multi-phase composites respond 

to varying levels of aggressiveness in material removal parameters. The focus will be on understanding the relationship between microstructure, hard 

particle concentration, and surface quality in machined composite materials. 

 

Conclusions 

• Two diesel engine valves were machined from 

silicon nitride materials supplied by Caterpillar. 

In preparing these valves, it was necessary to 

compensate for a lack of straightness in the stem 

portion of the blank. A precision coordinate 

measuring machine first was used to characterize the starting blank, and special mounting 

methods then were used to ensure that the final 

part could be prepared from imperfectly shaped 

stock. 

• A new effort in computer-assisted modeling of 

high-speed machining processes for titanium 

was initiated through participation in a consortium with TWS. Initial results are encouraging, 

and the software provides a new capability to 

ORNL’s machining program. 

• In FY 2005, there will be a new activity directed 

toward understanding the machining characteristics and resulting surface quality of composite 

materials. This work will support the development of precision machining processes for difficult-to-machine materials with relatively hard 

and soft phases in the microstructure. 

References

1. E. H. Kraft, Summary of Emerging Titanium 

Cost Reduction Technologies, U.S. Department of 

Energy/Oak Ridge National Laboratory, 2004. 

2. B. E. Hurles and F. H. Froes, “Lowering the 

Cost of Titanium,” AMPTIAC Quarterly, 6(2) 3–9 

(2002). 

3. P. C. Turner and J. S. Hansen, “Progress 

Toward Low-Cost Titanium,” Advanced Materials 

and Processes, 143(1) 42–43 (1993). 

4. E. H. Kraft, Opportunities for Low Cost Titanium in Reduced Fuel Consumption, Improved 

Emission, and Enhanced Durability Heavy-Duty 

Vehicles, U.S. Department of Energy/Oak Ridge 

National Laboratory, 2002. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

108
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

109

H. Cost-Effective Machining of Titanium Materials 

Albert J. Shih 

Mechanical Engineering 

University of Michigan 

Ann Arbor, Michigan 48109-9379. 

(734) 647-1766; fax: (734) 936-0363; e-mail: shiha@umich.edu 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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: University of Michigan, Ann Arbor, Michigan 

Objective 

• Investigate new technologies for cost-effective machining of titanium (Ti) alloys. 

• Gain in-depth knowledge of the effects of process parameters using three-dimensional (3D) finite element modeling (FEM). 

• Study machining using a production tool with complicated geometry. 

• Investigate tool wear mechanisms and investigate new tool materials and coatings. 

Approach 

• Experimentally validate force and temperature in 3D FEM. 

• Conduct parametric study of the machining parameters using the validated model. 

• Investigate the tool wear mechanisms and develop new tools and coatings. 

Accomplishments 

• Conducted turning and drilling experiments on Ti. 

• Conducted 3D FEM of Ti turning and validated the cutting force and chip thickness with experimental measurements. 

• Investigated the effects of cutting speed, a limiting factor for productivity in Ti machining. 

• Conducted tool wear tests. Used the laser profilometry to study tool wear mechanism in Ti machining. 

Future Direction 

• Expand the drilling research to Ti drilling, an important and bottleneck process for engine manufacturing. 

• Explore advanced tool materials and coatings for Ti machining. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials

110

Introduction 

Titanium and its alloys are lightweight, 

corrosion-resistant, high-temperature materials. New emerging technologies have been 

developed to reduce the cost of producing raw 

Ti material and have brought the potential for 

low-cost Ti and for automotive applications.1,2

While reducing the material cost can go a long 

way toward increasing the usage of Ti, high 

machining cost is another major technical barrier that needs to be overcome. To further reduce the cost of effective machining of Ti, a 

better understanding of the underlying mechanisms of the Ti machining process is necessary through better process and performance 

models, thermal modeling, and finite element 

simulation.3

 There is a lack of research in 3D 

FEM, which can provide detailed information 

about the stress, strain, and temperature distributions in the tool and workpiece; surface integrity of the part; and mechanics of chip curl. 

Advancements in machining process modeling 

have demonstrated the capability to provide 

reliable predictions of the performance of the 

cutting process and the influence of the process parameters on the product quality.4

 

Approach 

Experimental Setup 

The machining experiments were conducted on a Lodge and Shipley 30-hp computer numerically controlled lathe. A gradetwo commercially pure (CP) Ti rod, 31.8 mm 

in diameter, was used. A left-handed tool 

holder, Kennametal CTAPL-163D, was used 

to hold the triangular-shaped insert. The tool 

holder controls the orientation of the tool and 

workpiece, which is defined by the lead, back 

rake, and side rake angles in turning.5

 The lead 

angle and back rake angle were 0°, the side 

rake angle was 5°, and the relief angle was 6°. 

A Kistler 9257A 3-axis piezoelectric dynamometer was used to measure cutting forces 

on the tool insert in three directions, which are 

designated as tangential force (Ft), radial force 

(Fr), and axial force (Fa), respectively. The 

force signal was processed using the charge 

amplifiers and recorded by a PC-based dataacquisition system. 

FEM of 3D Turning of Titanium 

The Third Wave Systems AdvantEdgeTM 3D machining simulation software was used to model the Ti 

machining process. The updated-Lagrangian finite 

element method with continuous remeshing and adaptive meshing techniques was applied.6

 The coupling of 

the thermal and mechanical modeling of the tool and 

workpiece deformation was applied.7 Figure 1 shows 

the initial 3D finite element mesh using a tool with a 

corner radius. An XYZ coordinate system is defined to 

show the relative orientation of the tool and cutting 

forces. The X-axis is the cutting direction. The Y-axis 

is the radial direction, which is parallel to the major 

cutting edge. The Z-axis is parallel to the axial direction of the rod workpiece. The 4-node, 12-degree-offreedom tetrahedral finite element was used to model 

the workpiece and tool. The whole model includes 

approximately 100,000 nodes. The top and back surfaces of the tool are fixed in all directions. The workpiece is constrained in vertical (Z) and lateral (Y) 

directions on the bottom surface and moves at the cutting speed in the horizontal direction (–X) toward the 

stationary tool. 

Figure 1. Initial finite element mesh for 3D turning 

model (unit: mm). 

To validate the FEM, a baseline cutting test with 

four cutting speeds (24.4, 48.8, 97.5, and 195 m/min) 

and two feeds (0.254 and 0.381 mm/rev) was conducted at 1.02 mm depth of cut. Cutting forces were 

measured and compared with 3D finite element simulation results. One additional experiment was performed at a small depth of cut (0.254 mm) with a 

48.8 m/min cutting speed and 0.254 mm/rev feed. All 

cutting tests were conducted dry without using coolX

Y

Z Workpiece

Tool 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

111

ant. Nine simulations were conducted with the 

process parameters matched to those of the 

cutting experiments. The initial temperature 

was 20oC. 

Tool Wear Test 

In the tool wear test, a tool insert was used 

to cut the Ti for an extended period of time to 

investigate the wear mechanism. Table 1 lists 

the corner radius and coating used in tool inserts for the wear test. 

Table 1. Inserts and feed used for tool wear 

test (195 m/min cutting speed and 1.02 mm 

depth of cut) 

Typea

 TPG322 TPG320 

Corner radius 

(mm) 0.8 0.1 

No coating TiAlN Coating TiB2 TiB2 

Feed (mm/rev) 0.254 0.0508 

a

Kennametal tool number 

The tool with a small, 0.1-mm corner radius was tested with a 0.254 mm/rev feed. The 

tip fractured prematurely. A smaller feed 

(0.0508 mm/rev) was therefore used for the 

tools with a 0.1-mm corner radius. 

The flank wear was measured using an optical microscope. The width of flank wear land 

was recorded. The crater wear was measured 

using a laser profilometry (Rodenstock RM 

600). The 3D surface topography of the tool 

rake face was recorded, and the volume loss 

on the tool rake surface was calculated. 

Results 

Experimental Validation of Cutting 

Forces 

Figure 2 shows the comparison between 

the FEM (open symbol) and experimental 

measured (solid symbol) tangential, radial, 

and axial forces. For the tangential forces Ft, 

the finite element method generally underestimated by less than 15% with less discrepancy at high cutting speeds. For the radial and 

axial forces (Fr and Fa), the discrepancy was 

0

100

200

300

400

500

600

0 50 100 150 200

Cutting speed (m/min)

Cutting forces (N)

tangential force tangential force

radial force radial force

axial force axial force

Finite element 

modeling

Experimental 

measurement

0

100

200

300

400

500

600

0 50 100 150 200

Cutting speed (m/min)

Cutting forces (N)

tangential force tangential force

radial force radial force

axial force axial force

Finite element 

modeling

Experimental 

measurement

Figure 2. Comparison of cutting forces: (a) 0.254 

and (b) 0.381 mm/rev feed. 

less than 10% between the finite element simulation 

and experimental measurement results. 

Although some discrepancies exist, the cutting 

force measurements validate the FEM results and set 

the direction for further refinement to improve the accuracy in modeling. The feed has a significant effect 

on the cutting force. At a high cutting speed, cutting 

forces were reduced slightly. 

Cutting Speed Effects 

Tool temperature is a key factor that accelerates 

tool wear and limits the cutting speed and productivity 

in Ti machining. Experimental measurement of tool 

temperature is difficult. However, a finite element 

simulation can provide a quick and accurate prediction 

of the tool temperature under various cutting conditions. The effect of cutting speed on tool temperature 

was investigated. 

Y
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

112

An example of the tool and workpiece 

temperature distributions of the nine 3D finite 

element cutting simulations is shown in Figure 3. The process parameters for the selected 

example are 195 m/min cutting speed, 

0.254 mm/rev feed, 1.02 mm depth of cut, and 

6.35 μm tool edge radius. The chip curl and 

the engagement of the tool and workpiece are 

shown in Figure 3(a). Figure 3(b) shows the 

tool temperature distribution. The high temperature is concentrated around the straight 

major cutting edge and the round corner of the 

tool. The peak temperature is about 480°C. 

(a) 

(b) 

Figure 3. Temperature distributions: 

(a) workpiece and chip and (b) tool. 

Figure 4 shows the peak temperature on 

the tool rake face of the eight baseline cutting 

experiments at two feeds and four cutting 

speeds. 

200

250

300

350

400

450

500

0 50 100 150 200

Cutting speed (m/min)

Peak tool temperature (oC)

0.254 0.381 

 Feed 

(mm/rev)

Figure 4. Peak tool temperature as a function of 

cutting speed. 

The peak tool temperature is independent of the 

feed. This is comprehensible, since the very sharp tool 

cutting edge radius (6.35 μm) was used. The peak tool 

temperature increased significantly from about 260°C 

at a 24.4 m/min cutting speed to 480°C at 195 m/min. 

This demonstrates the effect of high-speed machining 

on tool temperature. 

Chip Curl 

This study explores the use of 3D FEM to predict 

the chip curl in turning. Figure 5(a) shows an experimental picture of the chip curl in machining of CP Ti 

with a 0.254-mm depth of cut, 48.8 m/min cutting 

speed, 0.254 mm/rev feed, and 0.8-mm tool corner 

radius. The change in cutting speed along the chip 

thickness is only 1.6%. The small change in cutting 

velocity across the chip thickness results in the up-curl 

dominated pattern. The effect of side-curl can also be 

seen to create the pitch of the curled chip. The tubular 

type helix chip is formed by rotational and translational motion along an axis parallel to the tool rake 

face. 

The finite element method is applied to simulate 

cutting using the right-handed tool. Figure 5(b) shows 

a close-up view of a curled chip 15.2 mm in length in 

a 3D finite element simulation. Qualitatively, the morphology of the chip curl obtained from experimentation (Figure 5(a)) and FEM (Figure 5(b)) shows a 

similar pattern. The chip flow directions also match 

each other in FEM and experimental observation. 

Figure 5. Chip curl with 0.254-mm depth of cut: 

(a) tubular chip in the experiment and (b) close-up 

view of the curled chip in simulation. 

However, the finite element method underpredicts the radius and pitch of the chip curl by about 

50%. One possible cause for this discrepancy is the 

contact of the chip with the tool, tool holder, and/or 

workpiece during cutting. Another possible cause is 

the external force acting on the chip from the weight 

of the chip after cutting. Since chip formation is con-
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

113

tinuous, the chip is pulled by its own weight to 

make a larger radius and pitch on the curled 

chip. 

Tool Wear Test 

Figure 6 shows the tangential forces (Ft) 

and the maximum width of flank wear land of 

the tool inserts. 

0

100

200

300

400

500

0 50 100 150 200

Time (s)

Tangential forces (N)

TiAlN coating Uncoated 

TiB2 coating TiB2 coated

0

0.2

0.4

0.6

0 50 100 150 200

Time (s)

Maximum width of flank 

wear land (mm)

Tool corner radius 0.1 0.8

Feed (mm/rev) 0.0508 0.254 

Figure 6. Tangential forces and the maximum 

width of flank wear land in Ti turning. 

The tool life was determined at the time 

when the sharp increase in cutting forces occurred. For tools with a 0.8-mm corner radius, 

the life of uncoated and TiB2-oated tools was 

69 and 199 s, respectively. For tools with a 

0.1-mm corner radius, the life of the TiAlN 

and TiB2-coated tools was 15 and 115 s, respectively. TiAlN is a new type of coating, 

developed by Kennametal for interrupted cutting of Ti. It was not successful in the continuous turning of CP Ti. The flank wear of a 

TiB2-coated tool with a 0.8-mm corner radius 

was measured by an optical microscope. A 

gradual increase in the maximum width of 

flank wear land can be seen in Figure 6. For 

the uncoated tool and tools with a 0.1-mm 

corner radius, because the tools were severely 

blurred with build-up material, no obvious 

flank wear could be observed. 

Laser profilometry was applied to measure the crater wear. Figure 7 shows the SEM micrograph and the 

corresponding 3D laser profilometry measurement of 

the crater wear on the rake face of the TiB2-coated tool 

with a 0.8-mm corner radius. The shape of crater wear 

can be observed. From the 3D contour shape of wear 

land, the volume of crater wear is calculated. 

Figure 7. Crater wear of worn TiB2-coated tool insert 

with 0.8-mm corner radius: (a) SEM micrograph and 

(b) 3D laser profilometry image. 

Laser profilometry can also be used to quantitatively measure the shape of the build-up edge at the 

tool tip. Examples of the TiAlN and TiB2-coated tool 

with a 0.1-mm corner radius are shown in Figure 8. 

The TiAlN obviously had a larger build-up edge, 

which resulted in the premature failure of the tool and 

the increase in the cutting force in Figure 6. 

Figure 8. SEM micrograph and 3D laser profilometry 

image of the build-up edge at the tip of tool inserts 

with 0.1-mm corner radius: (a) TiAlN coated and 

(b) TiB2 coated. 

The friction coefficient of two types of tool coating was investigated. By dividing the cutting forces 

into the force components normal and parallel to the 

rake face, the coefficient of friction, μ, can be calculated by the following equations with 0o lead angle. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

114

2 2 ( sin cos ) P = Ft α + Fa α + Fr (1) 

N = Ft cosα − Fa sinα (2) 

N

P

μ = (3) 

where 

P is the force parallel to the rake face,

Nis the force normal to the rake face, 

 α is the tool rake angle. 

The coefficient of friction of a TiAlNcoated tool with a 0.1-mm corner radius was 

about 0.84, which was larger than the coefficient fractions of 0.52 and 0.34 of the TiB2-

coated tool with 0.1- and 0.8-mm corner radii, 

respectively. 

Drilling 

Drilling modeling is more complex than 

turning modeling. 3D drilling modeling is ongoing in close cooperation with Third Wave 

Systems. Figure 9 shows the preliminary result provided by Third Wave Systems. The 

chip was generated by drilling Al-6061 using 

a 9.35-mm-diameter WC-Co drill at 

1.5 mm/rev feed and 40,000 rpm. The drill has 

a 30° helix angle, a 118° point angle, a 130° 

chisel edge angle, and a 2-mm web thickness. 

The modeling of Ti drilling is currently under 

development. 

Figure 9. FEM of 3D drilling with chip generation. 

Conclusions 

The 3D finite element simulation of turning CP Ti was validated by comparing cutting 

forces with reasonable agreement. The effect of cutting speed on peak tool temperature was investigated. 

The application of finite element simulation to study 

the chip curl was explored. Qualitative agreement on 

the type of chip curl and chip flow direction were 

achieved for turning with a small depth of cut. The 

feasibility of using the finite element method to model 

the complicated 3D machining processes was demonstrated. The tool wear test demonstrated that a large

tool corner radius and TiB2 tool coating are more favorable to tool life than a small corner radius and 

TiAlN coating. 

References

1. B. E. Hurless and F.H. Froes, “Lowering the 

Cost of Titanium,” The AMPTIAC Quarterly 6(2), 

3–9 (2002). 

2. E. H. Kraft, Summary of Emerging Titanium 

Cost Reduction Technologies, ORNL/Sub/ 

4000023694/1, Oak Ridge National Laboratory, December 2003. 

3. X. Yang, C. R. Liu, “Machining Titanium and 

Its Alloys,” Machining Science and Technology 3(1), 

107–139 (1999). 

4. K. F. Ehmann, S. G. Kapoor, R. E. Devor, I. 

Lazoglu, “Machining Process Modeling: A Review,” 

Journal of Manufacturing Science and Engineering

119, 655–663 (1997). 

5. D. A. Stephenson and J. S. Agapiou, Metal Cutting Theory and Practice, Marcel Dekker, 1996. 

6. T. D. Marusich, M. Ortiz, “Modeling and Simulation of High-Speed Machining,” International Journal for Numerical Methods in Engineering 38(21), 

3675–3694 (1995). 

7. A. J. Shih, “Finite Element Simulation of Orthogonal Metal Cutting,” ASME Journal of Engineering for Industry 117, 84–93 (1995). 

Publications/Presentations

R. Li and A.J. Shih, “Finite Element Modeling of 

3D Turning of Titanium,” presented at the 2004 

ASME IMECE, Anaheim, California, November 

2004. 

R. Li and A. J. Shih, “Finite Element Modeling of 

3D Turning of Titanium,” accepted for publication in 

International Journal of Advanced Manufacturing 

Technology. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

115

I. Advanced Cast Austenitic Stainless Steels for High-Temperature Components 

P. J. Maziasz, J. P. Shingledecker, 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 Manager: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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: Caterpillar Inc., Mossville, Illinois 

Objectives 

• Continue and expand commercial scale-up of new CF8C-Plus cast austenitic stainless steel. 

• Establish a more complete database for tensile creep and fatigue properties. 

• Determine resistance to aging, fatigue, and thermal fatigue to enable redesign or prototype casting trials of 

stainless diesel exhaust components. 

Approach 

• Produce 500-lb air-melt heats of static and centrifugally cast CF8C-Plus at MetalTek International or other 

stainless steel foundries. 

• Obtain a more complete tensile and fatigue properties database on CF8C-Plus at Caterpillar. 

• Determine improvements in resistance to thermal fatigue at Caterpillar and to creep at ORNL of CF8C-Plus 

relative to standard CF8C steel and other typical exhaust component alloys. 

• Determine the most effective new alloying additions to CF8C-Plus to maximize strength above 800°C and then 

down-select an alloy for commercial scale-up. 

Accomplishments 

• Received a 2003 R&D 100 Award for CF8C-Plus cast stainless steel. 

• Selected CF8C-Plus for further commercial scale-up based on superior aging, fatigue and thermal resistance at 

up to 850o

C compared with CN-12 or CN-12-Plus. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

116

• Produced new commercial static and centrifugally cast smaller heats of CF8C-Plus and a larger heat for a prototype component trial for a different application. 

Future Direction 

• Characterize microstructural changes after selected high temperature testing or aging; complete initial aging and 

expand aging to new alloys and impact properties evaluation. 

• Evaluate the effects of different commercial-scale casting processes on properties of CF8C-Plus steel. 

• Complete alloy modification studies of CF8C-Plus to boost strength above 800o

C, and commercially scale-up 

and test the most promising new heat. 

Introduction

Advanced heavy truck diesel engines must continue to have higher fuel efficiency, as well as reduced exhaust emissions, without sacrificing 

durability and reliability. More demanding normal 

duty cycles require exhaust manifolds and turbocharger housing materials to withstand temperatures 

ranging from 70°C to above 750°C. Such materials 

must withstand both prolonged, steady hightemperature exposure and rapid and severe thermal 

cycling. New emissions reduction technology and 

transient power excursions can push temperatures in 

these critical components even higher. 

Current diesel exhaust components are generally 

made from SiMo ductile cast iron, and higher engine 

temperatures push such materials well beyond their 

current strength and corrosion limits. The first Oak 

Ridge National Laboratory (ORNL)/Caterpillar 

(CAT) cooperative research and development 

agreement (CRADA) produced systematic data 

comparing a heat-resistant cast CN12 stainless steel 

with SiMo cast iron for more demanding diesel exhaust component applications. It clearly identified 

the tensile, creep, and fatigue strength advantages of 

standard CN12 stainless steel at 550–700°C and 

above. That first CRADA project also developed 

new lab-scale heats of modified CN12 and modified 

CF8C steels with better creep strength and significantly better aging resistance and thermal fatigue 

resistance than standard CN12. A second 

ORNL/CAT CRADA project began several years 

ago for commercial scale-up of these new modified 

cast stainless steel heats, and for development of a 

more systematic and thorough database required by 

designers to redesign stainless components and qualify them for trial component production. In 2003, the 

initial commercial scale-up data demonstrated clear 

advantages of the new CF8C-Plus cast steel relative 

to the other alloys for exhaust components and other 

applications, and this work won a 2003 R&D 100 

Award. In 2004, commercial scale-up focused on the 

new CF8C-Plus cast stainless steel and expanded to 

include more commercial heats, several different 

casting processes (static sand-casting, centrifugal 

casting), and alloying modifications to increase 

maximum strength at the highest temperatures. 

Approach

Two commercial stainless steel foundries initially produced 500-lb heats of CN12-Plus and 

CF8C-Plus steels for testing and evaluation compared with similar standard steels. One of those 

commercial foundries produced additional new 

static-cast and centrifugally-cast heats of CF8CPlus. Caterpillar Technical Center obtained tensile, 

fatigue, and thermal-fatigue data on these new heats 

of CF8C-Plus steel, while ORNL obtained initial 

creep-rupture data and aged specimens for various 

times at 700–850°C. ORNL also cast a new set of 

lab-scale heats to explore the effects of adding Al, 

B, Cu, and W to CF8C-Plus to maximize strength 

and creep resistance at 800–850°C. Results of the 

testing will be used to expand commercial scale-up 

of CF8C-Plus and then to down-select a heat of the 

new modified CF8C-Plus for commercial scale-up. 

Technical Progress

Tensile, fatigue, and thermal fatigue testing on 

the initial commercial heats of CF8C-Plus and 

CN-12 Plus were completed earlier at Caterpillar. 

Tensile data clearly showed CF8C-Plus to be as 

strong as standard CN12 and CN-12-Plus cast steels 

and much more ductile, with no additional heat 

treatment required after casting. CF8C Plus has an 

“engineered microstructure,” which includes a very 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

117

stable austenite parent matrix phase that is free of 

the δ-ferrite typically found in standard cast CF8C. 

The creep strength of CF8C-Plus at high temperatures comes from nano-scale dispersions of NbC 

precipitates that form and remain stable and are 

much finer than those found in standard CF8C steel. 

Because CF8C-Plus has no δ-ferrite, it is also free of 

σ-phase relative to standard CF8C during aging or 

creep at 650–850°C. 

Creep rupture stress versus rupture life, expressed as the Larson Miller Parameter (LMP), for 

standard CF8C and CF8C-Plus tested at 750–850o

C 

is plotted in Figure 1, together with data from highSiMo and Ni-resist cast irons tested at lower temperatures. Clearly the CF8C-Plus cast stainless steel 

has about twice the creep strength of standard CF8C 

steel, and both are much stronger than either highSiMo or Ni-resist cast irons. Figure 2 shows the significant advantage the CF8C-Plus alloy has in lowcycle fatigue resistance compared with standard CN12 steel at 800°C, which was a factor in choosing 

CF8C-Plus for further commercial scale-up for diesel exhaust component applications. The fatigue behavior of cast CF8C-Plus with a very coarse as-cast 

grain structure also compares well with wroughttype 348 stainless steel (similar alloy composition) 

with a much finer grain size. CF8C-Plus was deliberately designed to have a combination of good 

strength and ductility at both higher and lower temperatures, in order to achieve its best fatigue and 

thermal fatigue resistance. Thermal-mechanical fatigue (TMF) testing of these initial commercial heats 

was completed this year, together with aging of 

various specimens. There was a significant advantage for CF8C-Plus relative to CN-12 or high-SiMo 

cast iron in TMF testing to 760°C or higher, and 

data will be shown next quarter. 

Additional lab heats of CF8C-Plus were cast at 

ORNL to define the effects of other minor alloying 

additions, including B, Al, Cu and W additions, and 

establish their limits for commercial heats. Creeprupture test screening at 750–850°C showed that 

copper and tungsten were most effective at improving creep resistance, as shown for creep testing at 

750°C in Figure 3. Based on these data, a modified 

CF8C-Plus with copper and tungsten additions was 

selected for the next step of commercial scale-up 

and testing. 

Figure 1. Larson Miller parameter plot of creep rupture 

stress for high-SiMo and Ni-resist cast irons, standard 

CF8C, and the new CF8C-Plus cast stainless steels tested 

in air at various temperatures. 

Figure 2. Isothermal strain-controlled fatigue data versus 

cycles to failure in air at 800°C, for commercial scale-up 

heats of standard CN-12, CN-12-Plus, and CF8C-Plus 

cast stainless steels. 

Finally, this year, MetalTek International produced additional static sand-cast and centrifugallycast heats of CF8C-Plus for testing at Caterpillar and 

ORNL. That testing began this year and will be 

completed next year. MetalTek International has 

also produced several heats of a new CF8C-Plus 

with copper and tungsten additions, selected from 

the ORNL lab-scale heat data on alloy modifications, and testing will begin next year. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

118

Figure 3. Creep-rupture testing in air at 750°C and 140 

MPa of new ORNL lab heats of baseline CF8C-Plus steel 

and the same steel with additions of copper or tungsten.

Conclusions

ORNL and Caterpillar have chosen CF8C-Plus 

cast stainless steel as the focus for further scale-up 

and development, based on the initial commercial 

scale-up results of CN-12-Plus and CF8C-Plus. The 

development of CF8C-Plus was recognized with a 

2003 R&D100 Award. Several static-cast and centrifugally-cast heats of CF8C-Plus were produced by 

MetalTek International for further testing for diesel 

engine exhaust component applications and other 

high-temperature structural component applications. 

Another set of ORNL lab-scale heats of CF8C-Plus 

were made to test the effects of minor alloying additions on improved creep resistance above 800°C, 

and copper and tungsten additions were found to be 

most effective. New commercial heats of CF8C-Plus 

and CF8C-Plus with copper and tungsten were produced by MetalTek this year and will be tested next 

year. 

Publications/Presentations

P. J. Maziasz and M. J. Pollard, “HighTemperature Cast Stainless Steel,” Advanced Materials & Processes (ASM-International), 161, 57-59 

(October 2003). 

P. J. Maziasz, “Development of CF8C-Plus: A 

New Cast Austenitic Stainless Steel with Improved 

Heat-Resistance and Performance above 650°C,” 

invited presentation at ASM Peoria Chapter Meeting, April 5, 2004. 

M. J. Pollard, P. J. Maziasz, and J. P. Shingledecker, “Development of Low-Cost Cast Austenitic 

Stainless Steel for Diesel Engine and Gas Turbine 

Components,” Project Review of DOE Programs at 

Caterpillar, April 2004. 

P. J. Maziasz, J. P. Shingledecker, N. D. Evans 

and M. J. Pollard, “Advanced Cast Austenitic 

Stainless Steels for Diesel Engine and Gas Turbine 

Components,” presentation at DOE FreedomCAR 

and Vehicle Technologies Program Review at 

ORNL, June 15, 2004. 

Special Recognitions/Awards/Patents

ORNL and Caterpillar received a 2003 R&D 

100 Award for “CF8C-Plus Cast Stainless Steel for 

High-Temperature Performance,” presented at the 

R&D Awards Banquet in Chicago on October 16, 

2003. 

P. J. Maziasz, T. McGreevy, M. J. Pollard, C. 

W. Siebenaler, and R.W. Swindeman, “Heat and 

Corrosion Resistant Cast CN-12 Type Stainless 

Steels with Improved High Temperature Strength 

and Ductility,” U.S. patent application 

2003/0084967, filed July 15, 2002.. 

P. J. Maziasz, T. McGreevy, M. J. Pollard, C. 

W. Siebenaler, and R.W. Swindeman, “Heat and 

Corrosion Resistant Cast CF8C Stainless Steels with 

Improved High Temperature Strength and Ductility,” U.S. patent application 2003/0056860, filed 

July 15, 2002. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

119

J. Interfacial Control of Deformation Twinning in Creep-Deformed TiAl/Ti3Al 

Nanolaminate 

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; e-mail: hsiung1@llnl.gov 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

(805) 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: Lawrence Livermore National Laboratory, Livermore, California 

Objectives 

• Exploit thermomechanical-processing techniques to fabricate TiAl/Ti3Al nanolaminate composites at sizes from 

lamella width down to nanometer-length-scales. 

• Characterize the microstructure and elevated-temperature creep resistance of the nanolaminate composites. 

• Investigate the fundamental relationships among microstructures, alloying additions, and mechanical properties 

of the nanolaminate composites to achieve the desired properties of the composites for high-temperature 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. %). 

• 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). 

Accomplishments 

• Collaborated with Oak Ridge National Laboratory (ORNL) (C.T. Liu) to fabricate TiAl/Ti3Al laminate composites with high tungsten content using hot-extrusion processing techniques. 

• Characterized and measured the effect of tungsten addition on creep resistance of the TiAl/Ti3Al laminate composites. 

Future Direction 

• Continue to collaborate with ORNL to fabricate an oxidation- and heat-resistant class of TiAl/Ti3Al laminate 

composites with high niobium content (>10 at %) using the hot-extrusion processes. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

120

• Continue to investigate the effects of alloying additions and deformation (mechanical) twinning on the microstructural stability and creep resistance of the nanolaminate composites at elevated temperatures of up to 850°C. 

Introduction

One of the unique deformation substructures of 

TiAl/Ti3Al laminate composite is the formation of 

deformation twins (DTs) within γ lamellae. The 

twinning phenomena have been found to be significantly promoted within ultrafine lamellar TiAl as a 

result of refined lamellar spacing presumably because the increment of lamellar interfaces provides 

even more nucleation sites for twinning. Although it 

has been well known that deformation twinning can 

be activated by the homogeneous glide of 1/6<11 2 > 

twinning dislocations on the {111} planes, the underlying twinning mechanism still remains unclear. 

To better design lamellar alloys for high-temperature 

applications, it is important to understand and gain 

insights regarding the role of lamellar interfaces in 

the twinning process, as well as the mechanical behavior of the alloys. Accordingly, this investigation 

was conducted to elucidate the deformation twinning 

mechanisms in TiAl/Ti3Al nanolaminate. 

Approach

A TiAl/Ti3Al nanolaminate composite was fabricated at ORNL by a hot extrusion process, which 

involves hot-extrusion of a cast TiAl alloy at 

1350°C. After extrusion, the alloy was stressrelieved at 900°C in a vacuum (∼ 10−4

 Pa) for 2 h. 

Creep tests were conducted in a dead-load creep machine with a lever arm ratio of 16:1. For the current 

study, the deformation substructures of the specimen 

were investigated [tested at 760°C, 138 MPa (creep 

strain: 0.25%); 760°C, 518 MPa (creep strain: 

3.6%); and 815°C, 420 MPa (creep strain: 1.7%)]. 

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 the tested samples 

were examined using a JEOL-200CX TEM 

equipped with a double-tilt goniometer stage. Images of dislocations were mostly recorded using 

weak–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:1

 (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

Microstructure 

Figure 1 is a bright-field TEM image showing a 

typical edge-on microstructure within a TiAl (γ)-

Ti3Al (α2) nanolaminate. In general, the material 

contains two types of interfaces:2

1. The γ/α2 interphase interface, which has a 

usual orientation relationship (0001)α2 || (111)γ and 

<11 2 0>α2 || <11 0>γ. 

2. The γ /γ twin-related interface, which includes 

true-twin (180° rotational) and pseudo–twin (60 

and/or 120°rotational) interfaces. 

Here, the width of the α2 layers ranges from 10 

to 50 nm, and the width of the γ layers ranges from 

150 to 300 nm. Figure 2 is a WBDF TEM image 

showing a typical dislocation substructure within the 

nanolaminate. Both lattice dislocations (LDs) within 

the γ layer and a high density of interfacial dislocations (IDs) on inclined interfaces can be clearly 

seen. The density of the IDs is much greater than 

that of the LDs, and the LDs are primarily threading 

dislocations, which terminate their two ends at the 

interfaces. While the IDs on semi-coherent γ /α2 and 

γ /γ pseudo-twin interfaces are 1/6<112> or 

1/3<112> type misfit dislocations,3

 those on the γ /γ

true-twin interface are mainly 1/6[11 2 ] type twinning dislocations or geometry necessary dislocations 

for accommodating the departure of the true-twin 

interface from the exact (111) twin plane. 

Deformation Twinning and Proposed 

Mechanisms

When the nanolaminate was creep-deformed at 

760°C and 518 MPa, a deformation substructure 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

121

Figure 1. Bright-field TEM image showing a lamellar 

grain from an edge-on orientation. 

Figure 2. Weak-beam dark-field TEM image 

showing a typical dislocation structure of TiAl 

nanolaminate. 

associated with DTs within γ layers was developed. 

Typical examples of the formation of (111) [211]-

type DTs within the nanolaminate are shown in Figures 3 (a) and (b). Note that one of the twin lamellae 

was still growing between two lamellar interfaces, 

and its growth would be eventually blocked by the 

lower interface. This observation suggests that the 

interfaces are preferred nucleation sites for DTs, 

presumably resulting from the high local stresses 

caused by the pileup of interface dislocations. Accordingly, it is proposed that deformation twinning 

in the TiAl-Ti3Al nanolaminate can be viewed as a 

stress relaxation process to relief the local stress 

concentration caused by the pile-up of interfacial 

dislocations during deformation. The effective stress 

(τe) at the tip of the pile-up of n dislocations can be 

evaluated by τe = nτi [4], 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 

Figure 3. Dark-field (a) and bright-field 

(b) TEM images showing several (111) type 

deformation twins formed a growing process 

toward another interface. Schematic illustration 

(c) of the nucleation of a (111) type DT from a 

γ/α2 interface, where b1, b2, and b3 denote the 

interfacial, stair-rod, and twinning dislocations, 

respectively. 

based upon a stair-rod cross-slip mechanism.5,6 As 

for an example of the (111)-type DTs formed in the 

nanolaminate, the corresponding dislocation reaction 

(dissociation) is proposed to be 1/6[121] (111) →

1/6[011] (100) + 1/6[11 2 ] ( 1 11). The (111)-type DT 

is accordingly formed by a successive cross-slip of 

the twinning dislocations 1/6[112] on the (111)

plane. leaving the stair-rod dislocations 1/6[011] on 

the (100) plane. Twin (stacking) faults are subsequently formed on the interfaces between the γ layer 

and the DTs. This is schematically illustrated in Figure 3 (c). 

The formation of stair-rod dislocations at the intersections between the DTs and the α2 layer is evi-
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

122

denced in Figure 4, where the array of 1/6[011] 

stair-rod dislocations become 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 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 (F) 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 them the repulsive force 

(F) between the interfacial (Shockley) and stair-rod 

dislocations is considered to be rate-controlling. 

That is, a critical (minimum) stress (τc) is required 

to activate the dissociation reaction for twinning. 

Figure 4. Paired WBDF images demonstrating the existence of the array of 1/6[011] stair-rod dislocations at the 

intersections (indicated by arrows) between the (111)-

type DT and α2

 layer: (a) Invisible at g = 200 (g•b = 0), 

(b) visible at g = 021, Z (zone axis) ≈ [01 2]. 

Conclusion

The role of interfaces in deformation twinning 

of TiAl/Ti3Al nanolaminate has been investigated. 

Since the multiplication of lattice dislocations within 

both γ and α2 lamellae becomes very limited at a 

low stress level, the motion of interfacial dislocations (i.e., interface sliding) becomes an important 

deformation mode. Impinged lattice dislocations 

were observed to impede the movement of interfacial dislocations, which move in a cooperative fashion along the lamellar interfaces. The impediment of 

dislocation motion subsequently causes a dislocation 

pile-up in front of the obstacle as creep strain accumulates. When the laminate deforms at a high stress 

level, deformation twinning becomes a predominant 

deformation mode. The deformation twinning is 

suggested to be a stress relaxation process for the 

concentration of stress at the head of each dislocation pile-up. An interface-controlled twinning 

mechanism based upon a stair-rod cross-slip dislocation reaction is proposed and verified. 

References

1. J. W. Edington, Practical Electron Microscopy in Materials Science, Van Nostrand Reinhold, 

New York, 1976. 

2. M. Yamaguchi and Y. Umakoshi, Progress in 

Materials Science 34, 1 (1990). 

3. L. M. Hsiung and T.G. Nieh, Mater. Sci. & 

Eng. A239–240, 438 (1997). 

4. J. D. Eshelby, F. C. Frank and F. R. N. Nabarro, Phil. Mag. 42, 351 (1951). 

5. Fujita and T. Mori, Scripta Metall. 9, 631 

(1975). 

6. T. Mori and H. Fujita, Acta Metall. 28, 771 

(1980). 

Publications/Presentations

1. L. Hsiung, “Interface Sliding and Migration 

in an Ultrafine Lamellar Structure,” invited keynote 

presentation to the International Conference on 

Computational and Experimental Engineering & 

Sciences, Madeira, Portugal, July 28, 2004. 

2. Andrea M. Hodge, L. Hsiung, T. G. Nieh, 

“Creep of Nearly Lamellar TiAl Alloys Containing 

1.0 and 2.0 at % W,” presented at Materials Science 

and Technology 2004, September 27, 2004. 

3. L. M. Hsiung, T.G. Nieh, “Microstructures 

and Properties of Powder Metallurgy TiAl Alloys,” 

Mater. Sci. & Eng. A364, 1 (2004). 

4. A. Hodge, L. M. Hsiung, T. G. Nieh, “Creep 

of Nearly Lamellar TiAl Alloy 

Containing W,” Scripta Mater. 51, 411 (2004) . 

5. Luke L. Hsiung and Clyde L. Briant, Segregation of Tungsten to Interfaces in Lamellar TiAl , 

Lawrence Livermore National Laboratory, 

April 6, 2004. 

6. Luke L. Hsiung and Andrea M. Hodge, Effects of W on Creep Deformation of Nearly Lamellar 

TiAl, Lawrence Livermore National Laboratory, 

July 16, 2004 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

123 

K. Laser Surface Texturing of Lubricated Ceramic Parts 

Peter J. Blau and Jun Qu 

Oak Ridge National Laboratory 

P. O. Box 2008, Mail Stop 6063 

Oak Ridge, TN 37831-6063 

(865) 574-5377; fax: (865) 574-6918; e-mail: blaupj@ornl.gov 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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, TN 

Prime DOE Contract Number: DE-AC05-00OR22725 

Objectives 

• Determine the effectiveness of using regular patterns of micro-scale dimples produced by a laser to reduce the 

friction coefficient of lubricated surfaces of ceramics. 

• Determine the extent to which the laser surface texturing (LST) process affects the structure of the ceramic 

materials to which the process is applied. 

Approach 

• Select the proper tests and analyses needed to reveal potential friction reduction benefits of dimpled ceramic 

surfaces for use in lubricated sliding contacts. 

• Obtain transformation-toughened zirconia (TTZ) and silicon carbide ceramic materials, design and fabricate 

test specimens, and provide them to Surface Technologies, Ltd., for laser dimpling. 

• Characterize the ceramic materials after LST, using a variety of techniques, including electron microscopy, 

nanoindentation, and scanning acoustic microscopy (SAcM). 

• Conduct friction tests in lubricants with different viscosities to determine the extent to which the laser surface 

treatment can reduce friction. Compare results for reciprocating motion, like that in engine components, with 

results for unidirectional sliding, like that found in rotary bearings or seals. 

Accomplishments 

• Conducted reciprocating friction tests of LST ceramics in various lubricants, including diesel oil, mineral oil, 

and water to determine which frictional conditions benefit more from LST. 

• Investigated the use of SAcM to identify near-surface micro-fractures due to the laser-induced thermal pulses 

used to form the dimples. 

• Conducted transmission electron microscopy (TEM)of zirconia specimens and determined that a difference in 

crystal structure occurs in shallow zones affected by LST. 

• Obtained nanohardness data on the thermally affected zones in the material below dimples. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

124

• Established the sensitivity of LST effects to precise surface alignment, and the need for sufficiently high sliding 

velocity to detect frictional effects of the dimpled surfaces. In our experience, achieving friction reduction by 

LST in reciprocating sliding contacts requires taking extraordinary measures in alignment and apparatus 

design. 

• Explored the effects of solid lubrication with graphite powder on the friction of LST zirconia surfaces under 

reciprocating motion. There was no apparent improvement due to dimpling. 

Future Direction 

• Summarize the results of this study in a final report. 

Introduction 

Improving the fuel efficiency of diesel engines, 

while still enabling them to meet requirements for 

emissions, presents both design and material 

challenges. Reducing parasitic frictional losses in 

engines can be achieved by a combination of 

strategies such as 

• redesigning the engine components 

• reformulating the lubricants 

• improving methods of lubricant filtration and 

delivery 

• reducing churning losses in fluids 

• changing the operating conditions of the engine 

• substituting more durable, low-friction materials 

• altering the finish or microscale geometry of the 

bearing surfaces 

This project specifically addresses the last 

approach by investigating the potential of a 

technique called LST (developed by I. Etsion of 

Surface Technologies Ltd., in Israel). 

LST was successfully demonstrated by reducing 

the friction and wear of water pump seals in which 

the mating surfaces were wide and flat. Work by 

Japanese investigators on water-lubricated silicon 

carbide, dimpled by different methods, showed the 

importance of both dimple spacing and running-in 

on the extent to which dimpling was able to reduce 

the friction coefficient.1

 

Initial studies at Oak Ridge National Laboratory 

(ORNL) in FY 2002–2003 involved conducting 

reciprocating sliding friction tests with flat-sided 

(both dimpled and non-dimpled) cylindrical rods 

oscillating on a flat ceramic specimen. Water, 

mineral oil, and diesel oil were used as lubricants. 

By carefully analyzing the data, friction lowering 

effects of LST were seen. In FY 2003, 

microstructural characterization of laser-surfaceprocessed ceramics continued, as did the friction 

experiments. It was eventually concluded that the 

benefits of LST for reciprocating, narrow contacts 

could be detected only for very precisely aligned 

surfaces, and even that was difficult. In fact, the 

difficulty of detecting LST effects without special 

experimental procedures was confirmed by a paper 

by Kato2

 in which more than a day of waterlubricated sliding was required to wear-in the 

surfaces of the ceramics sufficiently to reveal the 

frictional benefits of LST. Unspectacular test results 

on liquid-lubricated LST ceramics during FY 2003 

prompted a change in approach for FY 2004. 

During the past year, the focus was shifted from 

using liquid lubricants to filling the dimples with 

solid lubricants, in this case, graphite. The 

hypothesis to be tested was that LST-produced 

dimples in surfaces might help catch abrasive debris 

and at the same time re-supply solid lubricating 

material to both lower friction and extend the life of 

ceramic sliding surfaces. This annual report 

summarizes the experiments designed to explore the 

latter approach. 

Approach 

This project is being conducted in parallel with a 

complementary effort at Argonne National 

Laboratory (ANL) on LST metals and LST used in 

conjunction with coatings. ORNL efforts have 

focused on ceramic materials. The approach is twopronged: (1) to conduct controlled reciprocating 

friction experiments on LST ceramics to evaluate the 

effects of both velocity and lubricant type on friction 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

125 

and (2) to characterize the effects of LST on both the 

microstructure and the near-surface micromechanical properties of ceramics. To accomplish 

the latter, various techniques are being applied—for 

example, nano-indentation hardness and elastic 

modulus determination, SAcM, and TEM and 

scanning TEM (TEM). Results of acoustic 

microscopy and nanoindentation tests were 

described in the previous annual report. TEM 

images and diffraction patterns of the laser-affected 

zone at the dimple bottoms will be provided in the 

final report to be prepared during FY 2005. They 

show indications that the crystal structure of TTZ is 

affected at a nanometer scale in a thin region near 

the surface of the dimple.

Results

In FY 2004, experiments were conducted to test 

the hypotheses that dimples in LST surfaces could 

serve both as reservoirs to replenish worn-away 

lubricant and as traps for any harmful, abrasive 

debris that happened to enter the interface. A flaton-flat reciprocating sliding configuration was used 

(see Figure 1). TTZ surfaces in three conditions, asground (non-dimpled), laser-dimpled, and dimpled 

and polished, slid against ground Gall-ToughTM 

stainless steel (upper specimen) in an ambient air 

environment. Pressure-indicating plastic film was 

used to help pre-align the two contact surfaces 

before each test. The nominal contact area of the 

sliding contact was 10 × 4 mm. A constant 10-N 

normal force, 2-Hz reciprocating frequency, and 10-

mm stroke length were used. The lubricant was 

Dixon Ticonderoga 1651 graphite powder. A layer 

of graphite powder was distributed on the zirconia 

surface before each test, and no additional lubricant 

was supplied thereafter. The level of friction force 

was monitored using a high-speed data acquisition 

system. Worn surfaces were examined to assess their 

morphology and to study the distribution of graphite 

and debris created by prolonged sliding contact. 

Results of the first graphite-lubricated test 

conducted on LST zirconia against Gall-ToughTM 

were promising. Without re-lubricating, the initially 

formed graphite lubricating film lasted nearly 90 

hours before failure (see Figure 2a). At the 

beginning of the test, the graphite powder quickly 

formed a lubricating film under pressure from the 

Figure 1. Flat-on-flat reciprocating sliding test setup 

showing the dead-weight load. 

mating surfaces. Since the two sliding surfaces were 

fully separated by a compressed graphite film, the 

frictional response reflects the resistance to sliding 

from adhering graphite particles. Initially, the 

friction coefficient remained relatively low at ~0.15. 

However, since this was an open system that 

allowed particles to escape, the graphite particles 

were gradually wiped off and the film became 

progressively thinner. When the lubricating film 

began to fail, asperity contact occurred between the 

mating surfaces. In this second stage, the tangential 

force resulted from a combination of metal/ceramic 

asperity contacts and the shear of the incomplete 

graphite film. The friction coefficient in this stage 

therefore fluctuated more noticeably. Eventually the 

lubricating film collapsed, leading to large areas of 

steel-on-zirconia in nominally unlubricated contact. 

The friction coefficient quickly transitioned to a 

much higher level, 0.5–0.6 (i.e., stage 3 in 

Figure 2a). This final frictional transition was used 

as the criterion for lubrication failure. 

The worn LST TTZ surface is shown in 

Figure 2b. Patches of wear debris, evidently 

containing both iron oxides and metallic particles, 

were observed. Many dimples were filled with wear 

debris, while others still retained some graphite 

powder. As hypothesized, the dimples acted both as 

reservoirs for the lubricant and as traps for wear 

debris. 

Although the first test results seemed promising 

for solid lubricant-filled dimples, subsequent tests 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

126 

0

0.1

0.2

0.3

0.4

0.5

0.6

0 20 40 60 80 100

Time (hour)

Friction Coefficient

Dimpled ziconia against flat GT steel

(a) Friction trace 

(b) Worn surface 

Figure 2. Results for LST TTZ reciprocating against 

Gall-ToughTM, a commercial stainless steel lubricated by 

graphite powder: (a) friction trace, (b) TTZ surface 

appearance (optical microscope image). 

designed to verify the repeatability of frictional 

behavior gave inconsistent results. Frictional 

increases that suggested film failure occurred much 

earlier in the latter tests. A repetition of the first test 

failed after only 8 hours, compared with more than 

85 hours for the first run. The two causes for this are 

alleged to be (1) slight variations in the initial 

alignment of the specimens and (2) an alignment 

change due to the effects of fixture vibration and 

loosening during sliding. These minute, but not 

insignificant, misalignments turned the relatively 

large, rectangular area of initial contact into a more 

concentrated line or point contact, which in turn 

caused the lubricating film to break through much 

more quickly. Considering the three stages of 

frictional behavior shown in Figure 2, initial 

alignment controls the length of the first sliding 

stage in which the friction is fairly low, the sliding is 

relatively smooth, and the vibration is minimal. 

When the lubricating film starts to fail and frictional 

behavior enters the second stage, the higher and 

more erratic level of friction induces more vibration. 

Higher vibration loosens the fixtures, changing the 

alignment and accelerating lubricant failure. After 

this behavior was observed, modifications were 

made to the equipment to eliminate vibrationinduced misalignment. Tests were then rerun on 

both the dimpled and non-dimpled surfaces. 

As shown in Figure 3, the lubricating films 

survived fairly long in both tests. Although the lowfriction stage did not last as long as on the first run, 

the second stage, with its partially lubricated 

surface, persisted for more than 100 hours. After 115 

hours of sliding, the friction behavior of dimpled 

and non-dimpled tests diverged. In the test of the 

dimpled surface, the friction coefficient became 

unstable and quickly transitioned to 0.5, indicating 

lubrication failure. However, the friction coefficient 

for the non-dimpled surface fluctuated between 0.2 

and 0.3 for another 50 hours or more, implying that 

it could survive under partial coverage by lubricant 

better than the dimpled specimen. Thus for solid 

lubrication by thin films, smoother surfaces seemed 

to perform better than dimpled surfaces. Perhaps if 

the surface wore down, it would expose more solid 

lubricant from the dimples to replenish the film, but 

in the case of ceramics such as zirconia and silicon 

carbide, their high wear-resistance would not allow 

this to happen. 

0

0.1

0.2

0.3

0.4

0.5

0.6

0 20 40 60 80 100 120 140 160 180

Sliding Time (hour)

Friction Coefficient

Dimpled zirconia

Flat zirconia

Figure 3. Friction traces for laser dimpled and nondimpled zirconia surfaces after stabilizing the test 

fixtures. 

A final report that summarizes the current 

research findings on LST ceramics will be prepared 

in FY 2005. It will include both tribology results and 

microstructural characterization. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

127 

Conclusions

Although dimples can serve as reservoirs for 

graphite powder and traps for wear debris, they did 

not appear to effectively re-supply the lubricating 

film. In fact, the smoother non-dimpled surfaces 

sometimes remained at a low-friction condition 

longer than the LST dimpled ones. 

The present results on solid lubricated surfaces 

have reinforced earlier conclusions from 

experiments with liquid lubricants: contact 

alignment is critical for optimizing the effects of 

dimpled surfaces. With utmost care in establishing 

good initial alignment and fixture stability, the 

lubricating film that formed from an initial 

application of graphite powder can be made to 

survive more than 100 hours of sliding (10-N load, 

2-Hz frequency, 10-mm stroke length). But even 

with added care in testing, LST offered no apparent 

advantages over non-dimpled, smooth surfaces 

tested under graphite lubrication. 

References

1. X. Wang, K. Kato, K. Adachi, and K. 

Aizawa, “The Effect of Laser Texturing of SiC on 

the Critical Load for the Transition of Water 

Lubrication Mode from Hydrodynamic to Mixed,” 

Tribology International, 34, 703–711 (2001). 

2. K. Kato, “Water Lubrication of Textured 

Silicon Carbide,” presented at the conference on 

Boundary Lubrication for Transportation, Copper 

Mountain, CO, August 3–8, 2003. 

Publications/Presentations

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, 

FL, October 28, 2003. 

P. J. Blau with J. Qu and L. Riester, Laser 

Surface Dimpling Effects on the Frictional Behavior 

and Microstructure of Transformation-Toughened 

Zirconia and Silicon Carbide Ceramics, Milestone 

Progress Report, submitted October 2003. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

128 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

129 

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

P. Schwartz, S. M. Abkowitz, and S. Abkowitz 

Dynamet Technology, Inc. 

Eight A Street, Burlington, MA 01803 

(781) 272-5967; fax: (781) 229-4879; e-mail: sabkowitz@dynamettechnology.com 

DOE Technology Development Manager: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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: 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 highdensity barstock. 

• Demonstrated that billets of 100% Ti-6Al-4V machine turnings can be successfully cast from billets in the 

sintered condition, without the additional step of HIP of the billet. The resulting castings demonstrated tensile 

strength and ductility that met or exceeded the typical properties of Ti-6Al-4V castings. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

130

Future Direction 

• Produce an additional extrusion billet, using specially selected and processed high-quality Ti-6Al-4V turnings 

(according to the process now developed), that will maintain oxygen content at 0.25% maximum. 

• Increase the parameters for consolidation. A higher HIP cycle (temperature, hot pressing temperature, and time) 

will be employed to produce a higher-density feedstock (full density will be the target). 

• Increase the extrusion reduction ratio so that a 5-in.-diameter billet will be extruded to a 1-in.-diameter bar 

(25:1 extrusion ratio). 

• Evaluate the resulting test bars for chemistry and tensile properties. 

Introduction

Dynamet Technology is developing 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 as extrusion billet, forging stock, or casting 

electrodes. This process promises to produce billets 

that are a low-cost alternative to billets produced 

from ingot. 

Figure 1: Dynamet’s CHIP process for producing lowcost feedstock. 

The use of low-cost billet made by Dynamet’s 

process in manufacturing Ti alloy parts for heavyduty 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 benefit from the metal’s advantages as a 

structural material. In addition to the energy, 

environmental, and economic benefits that result 

from vehicle weight reduction, the recycling of Ti 

turnings through Dynamet’s process would result in 

significant energy reduction because it would supply 

low-cost Ti for new applications that would not 

require new Kroll process sponge as a raw material. 

Future Direction

To complete the promising aspect of this highly 

desirable concept, it is necessary to successfully 

extrude Ti-6Al-4V turnings to produce barstock 

meeting the conventional Ti-6Al-4V mechanical 

properties. Based upon what we have learned thus 

far, it is believed that such successful extruded 

product can be demonstrated. A supplementary 

program to accomplish this goal has been designed. 

At the completion of this supplementary study, a 

detailed effort would be proposed to lead to 

commercialization of a low-cost specific grade of 

Ti-6Al-4V manufactured using processed turnings 

emanating from Ti-6Al-4V machining operations. 

This effort will not only generate a low-cost product 

but also use an energy-efficient and environmentally 

beneficial manufacturing process. 

Publications

Dynamet Technology, Inc., Ti-6Al-4V Billet 

Feedstock Manufacture and Evaluation—Final 

Report, Burlington, MA, July 15, 2004.
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

131

M. High-Density Infrared Surface Treatment of Materials for 

Heavy-Duty Vehicles 

T. N. Tiegs, F. C. Montgomery, D. C. Harper, 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: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-2476; e-mail: sid.diamond@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-resistant coatings on metal substrates. 

• Use lighter or more cost-effective bulk materials with coatings applied to surfaces where improved properties 

are required. 

Approach 

• Examine approaches to surface modification that would be of interest for materials for heavy-duty vehicles. 

• First examine the application and formation of adherent, wear-/corrosion-resistant coatings. 

• Base the initial tests on hardmetal compositions applied onto iron-based parts that are currently used in diesel 

engines. 

Accomplishments 

• Demonstrated the ability to produce adherent hardmetal-based coatings with low trapped porosity on three different metal substrates. 

• Showed improved wear resistance for coatings produced using HDI. 

Future Direction 

• Examine the relationships between processing parameters, microstructural development, and wear resistance. 

Introduction 

HDI technology is relatively new in the materials processing area and is gradually being exploited. 

The HDI processing facility at Oak Ridge National 

Laboratory (ORNL) uses a unique technology to 

produce extremely high power densities of up to 

3.5 kW/cm2

 with a single lamp. Instead of using an 

electrically heated resistive element to produce radiant energy, a controlled and contained plasma is 

used. 

Since the technology is relatively new, its utility 

in the surface treatment of materials for applications 

in heavy-duty vehicles is being explored. In most 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

132

cases, wear resistance, corrosion resistance, or high 

strength is necessary only in selected areas of the 

part that are 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 have the appropriate surface 

properties only where required. In addition, the HDI 

approach would be more cost-effective than other 

competitive processes such as physical vapor deposition. 

Results

Earlier work showed that adherent coatings of 

hardmetal compositions could be applied to a variety 

of metal substrates. Cross-sections of the coatings 

and base alloys revealed some large porosity within 

the coatings and at the interface with the underlying 

alloy. These bubbles are believed to be porosity that 

is trapped during the melting of the coating. To 

minimize the presence of the bubbles, the infrared 

(IR) treatment procedure was modified to expose the 

coating twice to cause the coating to remelt. The 

idea was that this procedure would allow the bubbles 

to escape during the second melting. A similar technique has been used before with silica refractories to 

reduce trapped porosity. The effects of double exposure on the properties of the coating and underlying 

base metal were mixed. 

Samples of WC-(Ni-Cr-Fe) on 4140 steel are 

shown in Figures 1 and 2 for one-step and two-step 

HDI processing. No reduction in the presence of the 

trapped porosity was observed. However, the level 

of trapped porosity was decreased by allowing additional time during a one-step IR exposure, as shown 

in Figure 3. Evidently, the additional time during 

liquid formation allowed the trapped gas to escape. 

Some cracks in the coating were observed with the 

additional exposure times. In another sample, the 

double exposure resulted in a reduction of the 

trapped porosity, as shown in Figures 4 and 5. The 

additional exposure also appeared to increase the 

mixing between the coating and the underlying base 

metal. 

As indicated, previous results had shown that 

adherent coatings of hardmetal compositions could 

be applied to a variety of metal substrates (D-2 tool 

steel, 4140 alloy steel, and cast iron). Selected samples with visually adherent and uniform coatings 

were wear-tested under dry sliding conditions using 

a block-on-disc tribometer. The tests were conFigure 1. Cross-section of WC-(Ni-Cr-Fe) coating on 

4140 steel after one-step HDI processing. IR exposure 

consisted of a lamp power of 900 amps for 2 sec. 

Figure 2. Cross-section of WC-(Ni-Cr-Fe) coating on 

4140 steel after two-step HDI processing. IR exposure 

consisted of a lamp power of 900 amps for 2 sec, 15 sec 

non-exposure, followed by 900 amps for 1 sec. 

ducted with 4.5 m/s linear speed and a 60-N contact 

load. 

The results are summarized in Figures 6 through 

9. The wear test results on the D-2 tool steel samples 

showed that the HDI-coated samples exhibited significantly less wear than the as-received material. 

The expanded view shown in Figure 7 reveals that 

the WC and TiC bonded with Ni3Al exhibited less 

wear than the other coated samples. A similar trend 

is also present in the 4140 alloy steel samples shown 

in Figure 8, where the TiC-Ni3Al coating showed the 

best wear behavior. The cast iron samples actually 

showed that the WC-Ni coating performed worse 

than the as-received material in the wear test. How 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

133

Figure 3. Cross-section of WC-(Ni-Cr-Fe) coating on 

4140 steel after one-step HDI processing. IR exposure 

consisted of a lamp power of 900 amps for 3 sec. 

Figure 4. Cross-section of WC-Ni3Al coating on D2 tool 

steel after one-step HDI processing. IR exposure consisted of a lamp power of 900 amps for 2 sec. 

ever, as observed before, the Ni3Al bonded sample 

exhibited the best wear behavior. 

Conclusions

It was determined that adherent hardmetal-based 

coatings could be produced using cost-effective 

slurry deposition followed by boning with the HDI 

lamp. Hardmetal-based coatings with reduced porosity were fabricated on three different metal substrates. The wear resistance of the HDI-coated 

samples was significantly better than that of the asreceived materials. In addition, coatings using Ni3Al 

as the binder phase had superior behavior in all of 

the wear tests. 

Figure 5. Cross-section of a WC-Ni3Al coating on D2 

tool steel after a two-step HDI processing exposure. IR 

exposure consisted of a lamp power of 900 amps for 2 

sec, 15 sec non-exposure, followed by 900 amps for 1 sec. 

Figure 6. Wear test results on as-received and HDIcoated D-2 tool steel samples. As shown, the coated samples exhibited significantly less wear than the as-received 

material. 

Publications/Presentations

T. N. Tiegs, “Ceramic-Metal Coatings Applied 

Using High Intensity Infrared Heating,” presented at 

the American Ceramic Society Pacific Coast Conference, Seattle, WA, September 13–16, 2004. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

134

Figure 7. Wear test results on as-received and HDI-coated D-2 

tool steel samples. 

Figure 8. Wear test results on as-received and HDIcoated 4140 alloy steel samples. 

Figure 9. Wear test results on as-received and HDIcoated cast iron samples. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

135

N. High-Temperature Aluminum Alloys 

Yong-Ching Chen and Randy Stafford 

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: Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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 Number: DE-AC05-00OR22725 

Subcontractor: Cummins Inc., Columbus, Indiana 

Objective 

• Develop high-temperature aluminum alloys with adequate properties and shape capability for turbocharger 

compressor wheels and housing applications. 

Approach 

• Develop ternary-phase compositions by modeling the equilibrium phase diagram, making 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 attainable by this method. 

• Evaluate a nanophase particulate-reinforced aluminum alloy patented by Chesapeake Composites, Inc. 

Accomplishments 

• Accurately predicted and verified the binary-phase diagrams of Al-Y and Al-Yb using ThermoCalc. 

• Studied the age-hardening behavior in NASA 388-T5 alloy. The improvement in high-temperature strength of 

the NASA 388 alloy was found to be not significant enough to pursue further. 

• Demonstrated impressive high-temperature tensile strength, fatigue strength, and thermal stability with the dispersion-strengthened composite (DSC) material. 

Future Direction 

• Use ThermoCalc to generate the Al-Y-Yb ternary system and identify optimal compositions for casting trials 

and material property characterization. 

• Conduct fatigue tests, creep tests, notch sensitivity tests, and machinability studies of DSC material. 

• Fully characterize the DSC material 

 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

136

Introduction

The new emission requirements for diesel engines mandate that the turbocharger compressor be 

operated at a significantly higher temperature and 

pressure. This condition makes standard castable 

aluminum alloys unusable as new-generation turbocharger materials 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 high-temperature 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 the investigation of three different paths for providing 

strength retention in aluminum alloys at high temperatures. Path 1 uses a ternary-phase aluminum 

alloy using 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 was performed to determine if 

there are other rare earth (or other metal) elements 

that produce beneficial properties at a reasonable 

cost. The main goal of this path is to generate the 

phase diagram of Al-Y-Yb (aluminum-yttriumyterrbium), using ThermoCalc software. Since there 

are very 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 elevatedtemperature properties reported by NASA are attractive, but Cummins prefers to use an alloy with a 

lower silicon content. Eck Industries purchased the 

license for the NASA-developed technology for 

high-silicon casting alloys 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. 

The nanophase particulates at 50 vol % provided 

adequate high-temperature strength in early experiments; however, the particulate-loaded alloy could 

only be forged or squeeze cast, so its complex-shape 

capability was limited. The mechanical behavior of 

this DSC material is being fully characterized. Casting modifications will be investigated to determine 

the shape capability for the alloy. 

Results

Path 1: Ternary-Phase Aluminum Alloy Development 

In the previous report, the binary-phase diagrams of Al-Fe and Al-Mn and a ternary-phase diagram of Al-Fe-Mn were generated using 

ThermoCalc. The versatility of the software has 

been clearly demonstrated in the plotting of the 

phase diagrams. A major aspect of this project is to 

be able to examine the minute details of any region 

of interest, which happens to be the aluminum-rich 

region of the phase diagram. This type of plot simplifies the task of accurately determining the limits 

of the phase fields in the aluminum-rich corner (or 

any region of interest) of the phase diagram 

Although the software features a wide spectrum 

of thermodynamic models, databases, and modules, 

it requires accurate assessment of the elements in 

order to perform phase equilibrium/diagram calculations. The Al-Y-Yb alloy poses a peculiar challenge, 

because there has been limited assessment of rare 

earth elements, especially alloys containing Yb. Although it is not presented, an extensive literature 

search has been conducted on the thermodynamic 

assessment of Y and Yb elements. In spite of this 

fundamental problem, the binary-phase diagrams of 

Al-Y and Al-Yb have been generated (using ThermoCalc) and validated. Figures 1 and 2 present the 

phase diagrams of Al-Y and Al-Yb, respectively. 

With the prediction of Al-Y and Al-Yb systems, 

work will begin on predicting the Al-Y-Yb ternary 

system. Upon deciding on the actual composition of 

interest, a material supplier will be asked to cast a 

prototype of the material. A detailed characterization 

of the material will be carried out at the Florida 

A&M University–Florida State University College 

of Engineering. Microcharacterization of the material will rely on scanning electron microscopy 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

137

Phases 

1 Liquid 

2 Fcc_A1 

3 Bcc_A2 

4 Hcp_A3 

5 Al3Y 

6 Al2Y 

7 AlY 

8 Al2Y3 

9 Al3Y5 

10 AlY2

Figure 1. The Al-Y binary phase diagram generated using 

ThermoCalc. 

Figure 2. The Al-Yb binary phase diagram generated 

using ThermoCalc. 

(SEM) and transmission electron microscopy 

(TEM), while the mechanical properties will be 

evaluated by tensile tests. 

Path 2: NASA 388-T5 Aluminum Alloy 

The microstructure of 388-T5 Al alloy with 

NASA modification was characterized. The results 

indicate that compositional modification leads to 

some refinement in the size of the silicon particles. 

A variety of coarse precipitates—Cu2Mg8Si6Al5, 

(Cu,Fe,Mn,Ni,V)3Si2Al15, (Cu,Ni)2Al and CuAl2—

were observed. The aluminum matrix was dominated by the presence of plate-shaped precipitates of 

the metastable θ’ phase. These precipitates appear to 

be mainly responsible for the strengthening. In addition, very fine spherical precipitates having the L12

structure and presumably the Al3(Ti,Zr) chemistry 

were noted. Figure 3 shows the SEM microstructure 

of NASA 388-T5 alloy. 

Figure 3. SEM micrograph showing microstructure of the 

NASA-modified 388 T5 alloy. Several indicated coarse 

intermetallic precipitates can be noted. 

The tensile strength of the thermally soaked and 

unsoaked NASA 388-T5 material increased slightly 

from room temperature and reached a peak at 300°F, 

after which it dropped almost linearly with the increase in temperature. The fracture morphology correspondingly changed from brittle to ductile/dimpled 

rupture, and microvoids containing second-phase 

particles in their interior were present. Thermal 

soaking at high temperatures for times up to 500 

hours led to some degradation of the tensile properties. 

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 with 

those of the current production materials, C355-T61 

and 354.0-T61 alloys. A slight improvement in the 

tensile strength was realized only at test temperatures above 300oC for the NASA 388-T5 alloy, 

compared with the strength of the current materials. 

The degree of improvement was deemed to be insignificant compared with 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. 

Path 3: Particulate-Reinforced Alloy 

This task involves a nanophase particulateloaded aluminum alloy patented by Chesapeake 

Composites. A patented, low-cost, liquid metal infiltration process is used to produce a billet form ready 

for secondary operations. This composite material 

combines the enhanced elevated-temperature 

Cu S 2Mg8Si6Al5

Si 

(Cu,Fe,Mn,Ni,V)3Si2Al15

(Cu,Ni)2Al 

CuAl2
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

138

strength, toughness, and ductility of dispersionstrengthened alloys with the stiffness and low coefficient of thermal expansion of metal matrix composites. It is claimed that this composite can be 

readily turned using tungsten carbide tooling and 

drilled and tapped using high-speed steel tools. Potential applications include pistons, compressor 

wheels, and engine components. 

As-infiltrated billets of the DSC with a 40 volume fraction of nanoscale Al2O3 particles in either 

1090Al matrix or 2024Al matrix were supplied by 

Chesapeake Composites for testing. Tensile specimens were prepared from the billets and thermally 

soaked at 204, 260, 316, and 371°C for 500 hours 

prior to testing. Figure 4 shows the tensile test results for the DSC material at elevated temperatures. 

The tensile strength of the DSC material decreased 

nearly linearly from ~58 ksi at room temperature to 

~20 ksi at 371°C, the latter value being impressively 

high for an aluminum alloy. It was noted that essentially no reduction in the tensile strength of the DSC 

material was observed. This indicated that the DSC 

material exhibited a very good thermal stability even 

after long-tem usage at high temperatures. The properties of the 2024Al-DSC material are comparable to 

those of the 1090Al-DSC material over the entire 

temperature range, suggesting that the second-phase 

precipitates expected in the former have little or no 

influence and that the properties are dominated by 

the presence of a high volume fraction of submicron-scale Al2O3 particles. 

Tensile Strength of Dispersion Strengthened Composites Material 

0

10

20

30

40

50

60

70

0 50 100 150 200 250 300 350 400 450

Test Temperature (oC)

Tensile Strength (ksi)

1090Al-DSC- No Thermal Soaking

1090Al-DSC after 500 hr. Soaking

2024Al-DSC- No Thermal Soaking

2024Al-DSC- after 500 hr. Soaking

Figure 4. Tensile strength of DSC material at elevated 

temperatures. 

Rotating beam fatigue tests of 2024Al-DSC and 

A354-T6 alloy were conducted at 260°C (500°F). 

Figure 5 shows the fatigue test results. The fatigue 

Fatigue Curves of 2024 Al-DSC at 260oC (500oF)

0

5

10

15

20

25

30

1E+05 1E+06 1E+07 1E+08

Number of Cycles to Failure, Nf

Stress Amplitude, ksi

2024 DSC

A354-T6

Figure 5. Rotating beam fatigue strength of DSC material 

and 354-T6 alloy at 260°C. 

strength of the 2024Al-DSC material was determined to be far superior to that of A354-T6 alloy, 

which is a premium aluminum alloy. A very impressive 19 ksi fatigue strength at 107

 cycles was obtained in the 2024Al-DSC material. The fatigue 

strength of A354-T6 alloy was determined to be 

about 10 ksi at 107

 cycles. 

The microstructure of the DSC material (asreceived and following thermal treatment) was characterized by TEM. The aluminum matrix grains 

were found to contain a moderate density of dislocations. The nanoparticles were found to have the 

structure and composition of α-Al2O3. The particles 

are nano-sized (30–100 nm) in the as-synthesized 

material, but they coarsen to 100–1000 nm following high-temperature thermal treatment, independent 

of time or temperature. After the initial holding 

times, little or no change in the sizes of the particles 

occurred at longer times and independent of temperature. This points to great thermal stability of 

these materials and a high, nearly constant hardness 

as a function of time and temperature. The TEM results are consistent with the SEM observations reported previously. 

Conclusions

Three different paths were adopted to develop 

and evaluate high-temperature aluminum alloys for 

turbocharger compressor wheels 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 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

139

tensile strength was realized only at test temperatures above 300oC for the NASA 388-T5 alloy compared with the current materials. The degree of 

improvement was deemed to be insignificant compared with 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 applications. The DSC material exhibited very impressive 

high-temperature tensile and fatigue strength. Excellent thermal stability was confirmed with SEM and 

TEM studies. The DSC material has high potential 

to be used to replace the current C355 alloy to 

achieve improved material strength and thermal stability. Other important material properties such as 

creep resistance, notch sensitivity, and machinability 

will be further evaluated. 

Publication/Presentations

P. Prasad, Y-C. Chen, and V. K. Vasudevan, 

“Microstructure and Mechanical Properties of New, 

Cast Aluminum Alloys for High Temperature Diesel 

Engine Applications,” in Symposium of Challenges 

for High temperature Alloys in Aerospace, Landbased Gas Turbines, Power and Transportation Systems, ASMI Fall Meeting, Columbus, OH, October 

18–20, 2004. 

P. Prasad, Y-C. Chen, and V. K. Vasudevan, 

“Characterization and Aging Studies of New, Cast 

Aluminum Alloys for High Temperature Diesel Engine Applications,” in General Abstracts Session on 

Mechanical Behavior and Characterization, 

TMS/MST 2004 Fall Meeting, New Orleans, September 26–29, 2004. 

P. Prasad, P. C. Becker, Y-C. Chen, and V. K. 

Vasudevan, “Microstructure of New, Cast Aluminum Alloys for High Temperature Diesel Engine 

Applications,” in General Abstracts Session on High 

Temperature Alloys, TMS/MST 2003 Fall Meeting, 

Chicago, November 7–11, 2003. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

140
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

O. Applications of Titanium Alloys for Heavy-Duty Vehicles 

P. C. Becker 

Oak Ridge National Laboratory 

P.O. Box 2008, MS-6068 

Oak Ridge, Tennessee 37831-6068 

(865) 576-4844; fax:(865) 574-6098; e-mail:beckerpc@ornl.gov 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax:(202) 586-1600; e-mail:sid.diamond@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 

• Determine if it is feasible to substitute a high-strength titanium (Ti) casting alloy for the heavier gray cast iron 

in a heavy-duty diesel engine head and block.

• Determine if it is feasible to substitute a high-strength aluminum (Al) casting alloy for the heavier gray cast 

iron in the same engine.

• Determine if it is feasible to substitute a high-strength magnesium (Mg) casting alloy for the heavier gray cast 

iron in the same engine.

Approach 

• Use mechanical and physical property data for finite element analysis (FEA) to compute the temperatures, 

stresses and fatigue cover factors (safety factors) at various locations on the head and block of a 5.9L Cummins 

ISB engine, run at maximum designed power (305 hp), for three lightweight casting alloy systems. 

• Compare these data to those computed for a production engine (all cast iron engine) to determine the feasibility 

of substituting any or all of these alloy systems for cast iron. (Ricardo, Inc. was contracted to conduct the FEA.) 

Note that fatigue analyses assumed that the engine was always stressed at the maximum for the life of the engine. In reality, this stress level is much higher than for typical duty cycles for this engine. 

Accomplishments 

• Determined the following through the FEA. 

⎯ Cast Ti-6Al-4V can be used as a direct substitution for gray cast iron in the engine block without any 

design changes as long as a cast iron liner is used. Ti was far superior to cast iron from a fatigue perspective. 

⎯ Cast Ti-6Al-4V can be used to replace cast iron in the engine head, but the valve bridge temperatures are 

too high (700°C); thus, without some design or materials changes, the head will not function well when 

made of Ti. 

⎯ To preserve the weight reduction and superior strength benefits of Ti for the block, without any design 

changes, the analyses demonstrated that an engine having either a cast iron or an Al head (15% overall 

engine weight saving) in combination with a Ti block with cast iron liners is feasible. 

141
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

⎯ The use of a Be-Cu alloy insert in the valve bridge regions of the Ti head, in combination with up-rating 

the engine to 450 hp, was shown to be a feasible way of reducing the temperatures in this region. The superior mechanical properties of Ti would enable the engine to operate at a higher power rating while 

achieving lighter weight (11% overall engine weight saving). The Ti block had no problem withstanding 

the higher stresses created by the up-rating. 

⎯ A high-strength Al casting alloy, A354-T6, could be used as a substitute for gray cast iron in both the 

head and block of the engine running at 305 hp if a gray cast iron cylinder liner were used. This would 

result in a 20% overall engine weight saving. Since Al does not have a fatigue endurance limit, the use of 

Al would be duty-cycle limited. 

⎯ A high-strength Mg casting alloy, WE 43, could be used as a substitute for gray cast iron in both the head 

and block of the engine running at 305 hp if a gray cast iron cylinder liner were used. This would result 

in an overall engine weight saving of 33%. 

⎯ Although the head temperature at the valve bridges is about 50°C higher than the recommended maximum for WE 43, it is considered a feasible substitute for gray cast iron if design changes are made to 

lower the temperature. 

Future Direction

• Seek funding from either a major diesel engine manufacturer or a major customer to construct and test engines. 

A potential customer is the Department of Defense, since there is a strong desire to reduce the weight and/or increase the power of engines in military vehicles.

Introduction 

Many programs are focused on reducing the cost 

of Ti powder and ingot by 50%. If any of these programs reaches this target, the cost of Ti might be 

become attractive enough for use in heavy duty diesel engines for certain applications where weight 

reduction is critical. In addition, there is a desire on 

the part of the military both to reduce the weight of 

transportation and to reduce the magnetic signature. 

This project explores the feasibility of replacing the 

cast iron in diesel engine heads and blocks with Ti, 

which is about 40% lighter than cast iron and is 

nonmagnetic. Using a Ti-6Al-4V casting alloy in 

place of gray cast iron in just these two components 

could save hundreds of pounds of weight, depending 

on the size of the engine. In addition, because of its 

superior strength and fatigue properties compared 

with cast iron, Ti affords the opportunity to upgrade 

the power of the engine without extensive design 

changes. It is recognized that Ti has some few disadvantages compared with cast iron, including lower 

thermal conductivity, poorer machineability, and 

higher cost. 

The approach used with Ti was also used to 

model the same engine with either an aluminum or a 

magnesium casting alloy. Thus this project studied 

the application of the three major lightweight metals 

systems to determine if it were feasible to replace 

gray cast iron in the engine head and block with alloys of one or all of these metals. 

Table 1 shows the weight savings that can be 

achieved when Ti, Al, or Mg is used to replace gray 

cast iron in a Cummins ISB 5.9-L diesel engine. 

Table 1. Weight savings produced by replacing cast 

iron with cast lightweight alloys in a Cummins ISB 

diesel engine (total engine wt. of about 1000 lbs)

Material Wt. saving (lb) 

Ti-6-4 head and block 

Ti-6-4 head and block 

with Cu inserts and cast 

iron (CI) liners 

Al(A354) head and Ti 

block with CI liners 

Al head and Al block 

with CI liners 

Mg (WE43) head and 

block with CI liners 

147 

107 

149 

200 

335 

Note: The liners weigh 7 lb each; the Cu inserts weigh 

1.5 lb each. 

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Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

Results

Ti Alloy Feasibility Study 

This project employed FEA to model a Cummins 5.9-L, B-series diesel engine running at 

305 hp. The temperatures and stresses were computed at various locations on the head and block 

castings for both the cast iron currently used and the 

Ti-6-4 alloy. Since the production engine is a parent 

bore engine (no cylinder liners), cast iron liners were 

incorporated into the model when evaluating the 

engine made with lightweight alloys. Based on the 

modeling and a comparison of the two materials, the 

feasibility of using Ti-6-4 in place of gray cast iron 

was determined. The modeling work was contracted 

to Ricardo, Inc. which has in-house the model for 

this engine and the experience to perform the necessary computer operations and interpret the results. 

As a result of the initial modeling, additional 

modeling iterations were needed to obtain the optimum combination of materials for the head and 

block. For example, a cast iron head or an Al head 

coupled with a Ti block was modeled. Even a copper 

alloy insert to the valve bridges of the Ti head was 

modeled in conjunction with up-rating of the engine. 

Table 2 compares Ti with gray cast iron in terms of 

major advantages and disadvantages. 

Figures 1 and 2 show the FEA model summary 

and block mesh; Table 3 shows the engine operating 

conditions. 

Tables 4 and 5 show the materials properties of 

gray cast iron and Ti-6Al-4V, respectively. Table 6 

shows the thermal conductivity data for gray cast 

iron, Ti, and Al. It is obvious that Ti has a much 

lower conductivity than the other two materials, and 

Al has very high thermal conductivity. 

The temperatures in the flame face of the head 

were computed for the all cast iron engine (production engine), labeled Fe/Fe/Fe, and for three alternatives. The first was an iron head on a Ti block with a 

cast iron liner (Ti/Fe/Fe); the second was a Ti head 

on a Ti block having a cast iron liner (Ti/Fe/Ti); the 

third was an Al head on a Ti block having a cast iron 

liner (Ti/Fe/Al). As can be seen from Figure 3, the 

temperatures in the valve bridge areas of the Ti head 

are high (709°C). The temperatures in the same location on the three other combinations are low 

enough not to cause problems. The temperatures in 

the other parts of the head were considered low 

enough in all the combinations that they posed no 

problem. Figure 4 shows the temperature determinaTable 2. Advantages and disadvantages of using Ti in 

place of gray cast iron 

Advantages of Ti Disadvantages of Ti 

1. Weight savings 

—Cummins B-Series 

• 100 lb saving 

from the machined block 

• 47 lb saving 

from the whole 

engine. (total 

dry wt. of 1036 

lb) 

• 14% weight 

saving from the 

whole engine 

(total dry wt. of 

1036 lb) 

2. Performance (room 

temp) 

• UTS of TI-6-

4:125 ksi 

• UTS of grey 

cast iron (CI): 28 

ksi 

• Fat. Str. of CI: 

10 ksi 

• E for Ti-6-4: 17 

msi 

• E for CI: 15 msi 

3. Non-magnetic 

1. Cost 

—Ti is about $10/lb 

and must be vacuum 

investment cast 

- Gray CI <$ 

0.20/lb. and is 

green sand cast in 

air 

2. Machinability 

—Ti has much lower 

throughput and much 

higher tool wear 

3. NVH attenuation 

4. Thermal 

 conductivity (300EC) 

—KCI: 49 W/m-k 

—KTi: 12 W/m-k 

Figure 1. FEA model summary. 

143
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

Figure 2. Block mesh. 

Table 5. Mechanical properties of Ti-6Al-4V used in 

this study 

Comparison of mechanical properties of Ti-6AI-4V 

from the literature vs. values used in this study 

Literature 

(Crucible steel) 

Temp (°C) UTS (Mpa) Y.S. (Mpa) 

21 1063 973 

315 814 725 

538 635 531 

This study 

(Mil Handbook 5, annealed, HIPed castings) 

Temp.(°C) UTS ,(Mpa) Y.S. (Mpa) 

21 862 820 

315 608 502 

538 431 362 

Table 3. Engine operating conditions for the purpose 

of modeling 

Number of cylinders 6 

Engine displacement 5.59L 

Rated speed 2900 rpm 

Rated power 305 hp 

Cylinder pressure 160 bar 

Table 6. Thermal conductivity of gray cast iron, Al 

and Ti at various temperatures 

Temperature-dependent conductivity 

Aluminum data supplied by Cummins 

K (W/m-K) 

Temperature 

(EC) Cast iron 

Titanium 

(Ti6AI4V) 

Aluminum 

(A354T6) 

−18 6.7 

21 168.8 

27 7.3 

38 7.4 

50 170.9 

93 7.6 

100 48.8 175.5 

149 8.1 

150 177.6 

200 47.8 179.3 

204 8.7 

260 9.6 180.8 

300 46.8 

316 10.4 

371 11.2 

400 45.8 

427 11.9 

482 12.7 

50 44.8 

538 13.5 

Table 4. Materials properties of gray cast iron used in 

the Cummins engine 

Block – gray cast iron —data supplied by Cummins 

—Young’s modulus = 103.4 GPa 

—Poisson’s ratio = 0.25 

—Ultimate tensile strength = 189.6 MPa 

—Tensile yield strength = 136.5 MPa 

—Compressive yield strength = 284.4 MPa 

—Fatigue strength = 72.39 MPa 

Head – gray cast iron– data supplied by Cummins 

—Young’s modulus = 103.4 GPa 

—Poisson’s ratio = 0.25 

—Ultimate tensile strength = 206.8 MPa 

—Tensile yield strength = 148.9 MPa 

—Compressive yield strength = 317.2 MPa 

—Fatigue strength = 110.3 MPa 

tions for the tops of the liners in the engine block, 

determined for the same material combinations. 

Table 7 summarizes the temperatures in the key 

locations on the head and block for the four head 

and block material combinations. 

from (0,0) through (δm, δa) to the point where it intersects the boundary of the safe working area, to the 

length of the line from (0,) to (δm, δa). The factor 

will be greater than 1 for points within the safe 

working area, and less than 1 everywhere else. Ricardo’s acceptable cover factor for this type of application is 1.5 or greater.

Figure 5 shows how the fatigue calculations 

were predicted based on modified Goodman diagrams and where on the diagram the cover factors 

(safety margins) are determined. Fatigue cover factor is defined as the ratio of the length of the line 

144
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

Table 7. Temperatures at key locations in the head and block for four material combinations 

Summary of temperatures @ areas of interest 

All temperature in degrees centigrade 

From previous iterations 1 2 3 

Location All iron All titanium 

Ti block 

Fe liner 

Fe head 

Ti block 

Fe liner 

Fe head 

Ti block 

Fe liner 

Fe head 

Top of block 227 410 228 287 194 

Top of liner - - 248 289 214 

Cylinder head exhaust valve bridge 370 710 381 709 220 

Intake valve seat 357 550 360 551 249 

Exhaust valve seat 444 741 455 741 326 

Figure 3. The computed flame face temperatures of the 

engine head for four head and block material combinations. 

Figure 4. The computed temperatures for the tops of the 

cast iron liners for four head and block material combinations. 

Figure 5. Goodman high cycle fatigue calculations. 

Table 8 shows a summary of the computed 

stresses and fatigue cover factors for the cast iron 

head from the production engine and for the Ti head 

on a Ti block with a cast iron liner. The table shows 

many key locations on the heads. 

It can be seen from Table 8 that the Ti head has 

more fatigue resistance than the cast iron head of the 

production engine in all of the key locations of the 

head, and that there are many locations on the production engine head that have a cover factor of less 

than 1. Even with the low cover factors, the production engine survives in service with a high degree of 

reliability. This is because, in this study, the analyses assumed that the components see only the 

maximum operating conditions. In real life, the engine does not operate at this maximum all the time. 

Thus the stresses are lower for a large portion of the 

duty cycle of the engine. This study can be considered a worst-case scenario. 

145
Table 8. Computed stresses and fatigue cover factors in the head for the combinations Fe/Fe/Fe and Ti/Fe/Ti 

Cast iron 

Location 

(lwj = lower water jacket) 

(uwj – upper water jacket) 

Description Node # 

Minimum stress 

 (N/m2) 

Maximum stress 

 (N/m2) 

Mean stress 

 (N/m2 

Alternating 

stress 

(N/m2) 

A uwj fillet between bolt box 

and valve guide 

3759825 Cyl 5 156.3E+6 Cyl 2 279.3E+6 217.8E+6 61.5E+6 

B lwj floor fillet between exh. 

Port and inj. bore 

3785105 Cyl 5 221.6E+6 Cyl 2 242.4E+6 232.0E+6 10.4E+6 

C lwj floor fillet to exhaust port 3783535 Cyl 5 202.5E+6 Cyl 1 228.7E+6 215.6E+6 13.1E+6 D lwj roof fillet to head bolt boss 3784266 Cyl 1 107.2E+6 Cyl 2 219.1E+6 163.2E+6 56.0E+6 E uwj rof fillet to head bolt boss 3759206 Cyl 6 121.4E+6 Cyl 2 217.0E+6 169.2E+6 47.8E+6 F lwj floor fillet to exhaust port 3780615 Cyl 3 136.0E+6 Cyl 2 207.4E+6 171.7E+6 35.7E+6 G intake port to valve seat area 3879902 Cyl 2 190.5E+6 Cyl 1 206.8E+6 198.7E+6 8.2E+6 H lwj roof fillet to head bolt boss 3776708 Cyl 3 186.1E+6 Cyl 2 198.8E+6 192.5E+6 6.4E+6 I exterior on exhaust side near 

flame deck 

3793591 Cyl 2 160.2E+6 Cyl 1 184.2E+6 172.2E+6 12.0E+6 

J uwj roof fillet to head bolt 

boss 

3777863 Cyl 1 109.2E+6 Cyl 2 181.1E+6 145.2E+6 36.0E+6 

K uwj roof fillet to head bolt 

boss 

3779640 Cyl 6 123.3E+6 Cyl 2 174.0E+6 148.7E+6 25.4E+6 

146

L lwj roof fillet to head bolt boss 3775260 Cyl 1 108.7E+6 Cyl 2 170.1E+6 139.4E+6 30.7E+6 M uwj fillet between head bolt 

boss exhaust port 

3780476 Cyl 1 132.2E+6 Cyl 3 169.8E+6 151.0E+6 18.8E+6 

N exterior pocket @ top deck 

near bolt boss 

3791833 Cyl3 89.8E+6 Cyl 2 161.7E+6 125.7E+6 36.0E+6 

O uwj roof fillet to top deck 3779648 Cyl 6 141.8E+6 Cyl 2 160.3E+6 151.1E+6 9.3E+6 P lwj fillet – valve guide to fuel 

drilling 

3787973 Cyl 6 113.8E+6 Cyl 3 158.4E+6 136.1E+6 22.3E+6 

Q top deck rib between guides 3783137 Cyl 6 85.9E+6 Cyl 2 158.4E+6 122.2E+6 36.2E+6 R lwj floor between exh. Ports 

and inj, bore 

3785319 Cyl 2 107.9E+6 Cyl 3 156.9E+6 132.4E+6 24.5E+6 

S top deck exterior 3783291 Cyl 6 115.6E+6 Cyl 2 156.7E+6 136.2E+6 20.6E+6 T exterior pocket @ top deck 

near head bolt boss 

3792836 Cyl 1 112.2E+6 Cyl 2 155.3E+6 133.8E+6 21.6E+6 

U uwj roof to coolant hole 3778568 Cyl 5 110.5E+6 Cyl 2 155.2E+6 132.9E+6 22.4E+6 V lwj fillet – exh. guide to exh. 

port 

3780566 Cyl 5 85.0E+6 Cyl 2 151.4E+6 118.2E+6 33.2E+6 

W lwj floor fillet to exhaust port 3776526 Cyl 3 129.3E+6 Cyl 2 150.2E+6 139.8E+6 10.5E+6 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

Figure 6 shows the results of the fatigue analyses of the block in cast iron and Ti, and Table 9 

summarizes the results of the fatigue prediction 

study for various locations on the block for both cast 

iron and Ti. As can be seen, there is only one small 

location on the block with a low cover factor for 

both materials, although the Ti value is more than 

four times that of the cast iron. This location is so 

small that it should be easy to increase the cover 

factor by adding some material in that location in the 

casting. 

Figure 6. Description of fatigue cover factors (safety factors). 

One approach for dealing with the unacceptably 

high temperatures in the valve bridge regions is to 

replace the whole head with either cast iron or Al. In 

the former case, the only weight saving would be 

from using Ti for the block. In the latter case, a further weight saving would be achieved because Al is 

57% lighter than Ti but less durable at high temperatures and fatigue stresses. Another approach, which 

would preserve the advantages of Ti, would be to 

make the head of Ti but replace the valve bridge regions with an insert of a material that has a high 

thermal conductivity and high elevated temperature 

strength. 

To this end, a literature study was conducted of 

materials of known high thermal conductivity and 

high strength. A copper alloy was considered to be 

the best candidate insert material. The insert design 

in the valve bridge region of the head was therefore 

modeled by FEA; but at the same time, the engine 

was up-rated to 450 hp to see if the superior mechanical properties of Ti-6-4 would enable the engine to be rated at a much higher horsepower than 

current production engines. Figure 7 illustrates the 

concept. It shows an Al head with cast iron inserts 

showing where the copper alloy insert would be located. 

Figure 7. The FEA model of one cylinder of the block 

showing the only location that has a low cover factor for 

both materials studied 

To determine if the insert strategy described was 

feasible, the physical and mechanical properties of a 

2% Be-Cu alloy were used. Figure 8 shows a graph 

of ultimate tensile strength vs. temperature for Ti6Al-4V, A354-T6, 2% Be-Cu, and gray cast iron. 

The 2% Be-Cu alloy has excellent elevated temperature strength and exceeds or equals Ti up to about 

350°C. The thermal conductivities of the four key 

alloys shown in Figure 8 are shown in Figure 9. The 

design of the BE-Cu insert is shown in Figure 10. 

Figure 8. Photo of an Al engine showing cast iron inserts 

in the fire deck as a way of showing how a Cu alloy insert 

would be placed in a Ti head. The actual engine has four 

valves instead of the two shown here. 

147
Table 9. Summary of the computed stresses and fatigue cover factors for the cast iron and the Ti blocks Cast iron 

Location Description Node # Minimum stress (N/m2) Maximum stress (N/m2) Mean stress (N/m2 Alternating stress (N/m2) 

AH interior fillet to top deck @ end face 6187310 Cyl 1 758.0E+6 Cyl 4 1.1E+9 944.0E+6 186.0E+6 AI oil hole edge to brg. shell 6468000 Cyl 1 130.0E+6 Cyl 6 206.0E+6 168.0E+6 38.0E+6 AJ corner of side inlet to water jacket 6215232 Cyl 6 17.1E+6 Cyl 2 129.0E+6 73.1E+6 56.0#+6 

AK internal fillet to head bolt boss 6173813 Cyl 3 70.7E+6 Cyl 1 149.0E+6 109.9E+6 39.2E+6 

148

AL corner of side inlet to water 

jacket 

6214185 Cyl 5 70.8E+6 Cyl 1 139.0E+6 104.9E+6 34.1E+6 

AM ‘liner’ to top deck 6166821 Cyl 3 107.0E+6 Cyl 1 139.0E+6 123.0E+6 16.0E+6 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

Figure 9. Ultimate tensile strength vs. temperature for 

four key diesel engine head and block alloys. 

Figure 10. Thermal conductivity vs. temperature for four 

alloys. 

Figure 11 shows the location of the insert in the 

flame face. 

The FEA model was modified so that part of the 

back face of the insert makes direct contact with the 

coolant (Figure 12). 

The engine with the Be-Cu insert was modeled 

with the engine up-rated to 450 hp from the standard 

305 hp. This means that the peak cylinder pressure 

was increased from 2320 to 3191 psi. The block 

model showed that the Ti-6Al-4V alloy allowed the 

block to function acceptably at the new rating. The 

head temperatures were reduced in valve bridge regions from 709°C at 305 hp to 545°C at 450 hp 

(Figure 13). The temperature in the region of the 

valve seats is still about 50°C too high, but considering that this model assumes that the engine is runFigure 11. Solid model of the Be-Cu flame face insert. 

Figure 12. Location of the insert in the flame face. 

Figure 13. Model of the back face of the insert showing 

where the insert makes direct contact with the coolant 

(darker gray). 

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FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

ning at 450 hp all the time, this temperature is considered acceptable from a feasibility standpoint. 

Up-rated engine parameters: 

• Maximum cylinder pressure increased from 160 

to 220 bar (3191 psi). This was the only parameter changed. 

― Power increased by 48% to 450 hp (current 

ISB—305 hp) 

― All other engine operating conditions unchanged 

• Speed: 2900 rpm 

• AFR: 24.4:1 

• Compression ratio: 17:1 

― Flame face thermal boundary conditions 

have been modified to account for material 

changes. However, the baseline coolant side 

boundary conditions (obtained from computational fluid dynamics analysis at 305 hp) 

are unchanged because the coolant flow remains the same as a result of the engine 

speed. 

Figure 14 shows a graph of the thermal expansion coefficients vs. temperature of the Ti alloy and 

the Be-Cu alloy. Note that the level Be-Cu is more 

than an order of magnitude greater than the Ti alloy. 

This leads to a high stresses at the interface between 

the Ti and the insert, as shown in Figures 15 and 16. 

Figure 14. Description of up-rated engine parameters. 

Figure 15. FEA model of the temperature distribution in 

the valve bridge area of the flame face for the Ti head 

rated at 305 hp and the Ti head with a Be-Cu insert. 

Figure 16. Coefficient of thermal expansion vs. temperature for Ti and Be-Cu. 

The thermal results for the top of the block cylinder are shown in Figure 17. This region is the only 

location where the temperatures get high enough to 

be of concern. 

Thermal Results Summary

• Using Ti-6Al-4V to replace cast iron in the engine block is feasible. 

• Using Ti-6Al-4V to replace cast iron in the head 

poses some problems associated with the poor 

thermal conductivity of Ti compared with gray 

cast iron. 

150
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

Figure 17. FEA model of the Ti alloy lower water jacket 

showing the principal stresses generated at the interface 

between the back face of the insert and the Ti as a result 

of the differences in thermal expansion coefficients of the 

two materials. 

― The valve bridge regions of the fire deck get 

to about 700°C, which is too high for acceptable durability. 

― The valve seats see too high a temperature 

for durable use. 

• To take advantage of the Ti alloy’s favorable 

properties for the block (lower weight and 

higher strength), the engine could be constructed 

with a Ti block and either a gray cast iron head 

or an aluminum head. This leads to the following limitations: 

― The weight advantages of using Ti for the 

engine are limited if a cast iron head is used. 

― A larger weight saving can be achieved by 

using Al for the head, but durability may be 

limited. 

A novel approach to solving the high temperatures in the head of the “all Ti engine” was proposed: using a high-thermal-conductivity alloy as an 

insert in the valve bridge regions. 

― A 2% Be-Cu alloy was selected as the insert 

material for the purpose of FEA modeling. 

― Rather than model the insert at the normal 

305 hp rating, the insert was modeled in an 

engine that was up-rated to 450 hp. This was 

done to determine if the use of Ti would allow engines normally designed for cast iron 

to be used at higher power ratings. 

• The Be-Cu alloy insert resulted in a valve bridge 

temperature of 545°C when the engine was run 

at 450 hp (3190 psi peak cylinder pressure). 

Since this modeling run assumed that the engine 

was run at these values for its whole life, which 

is not the normal way an engine is used, it is 

concluded that the insert strategy is feasible for 

an up-rated Ti engine. Note that it is possible 

that increased coolant flow and some design 

changes may be needed to guarantee the durability of the engine. 

• Owing to the large difference in the coefficients 

of thermal expansion of the Be-Cu and Ti alloys, 

large stresses are generated at the interface between the insert and the Ti lower water jacket. 

Conclusions

• FEA modeling of stress, fatigue, and temperature shows that cast Ti–6Al-4V is a feasible replacement for gray cast iron in heavy-duty 

diesel engine blocks with cast iron liners, 

whether the engine is rated at the normal (305 

hp) or a 50% higher level (450 hp).

• FEA modeling of stress, fatigue, and temperature shows that cast Ti–6Al-4V is a feasible replacement for gray cast iron in heavy-duty 

diesel engine heads rated at 305 hp if the high 

valve bridge temperatures can be overcome. 

• FEA modeling shows that using a cast iron head 

or an Al head in conjunction with a Ti block is a 

feasible way of taking advantage of the high 

strength of Ti for the blocks without concern for 

high temperatures in the head for an engine rated 

at 305 hp. However, if a gray cast iron head is 

used, there will not be much of weight saving 

for the engine. If the head is made from Al, a 

great weight saving can be achieved, since Al is 

54% lighter than Ti. 

• FEA modeling shows that if a 2% Be-Cu insert 

is used to replace Ti in the valve bridge regions 

in the flame face for an engine that was up-rated 

by 50% (450 hp), the high temperatures in the 

flame face are reduced substantially, and it is 

feasible to use it at high power ratings. 

• The high stresses generated at the interface between the insert and the Ti lower water jacket, 

caused by the differences in coefficients of 

151
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

152

thermal expansion, can be overcome by some 

design changes. 

Aluminum and Magnesium Feasibility 

Study

Aluminum 

The FEA modeling for these studies was for the 

engine rated at 305 hp, which is the maximum for 

the production engine. 

Thermal analysis 

Table 10 provides the mechanical and physical 

properties of gray cast iron and the aluminum alloy, 

A354-T6 used in this modeling study. 

Figure 18 shows the calculated temperatures in 

the valve bridge region of the fire deck for the production cast iron engine and for the engine having 

an Al head and block. Owing to the very high thermal conductivity of Al, the maximum temperature in 

the valve bridge region is only slightly above 200°C, 

well below Ricardo’s recommendation for maximum temperature in this region. Figure 19 shows the 

calculated temperatures for the top of the engine 

block for both the cast iron production engine (parent bore) and the Al engine. Again, the temperatures 

at the top of the Al block are well below any level of 

concern. 

Figure 18. FEA model of the Be-Cu insert showing the 

stresses generated in the Be-Cu insert caused by the differences in thermal expansion coefficients between the Ti 

alloy head and the insert when the engine is run at 450 hp. 

Table 11 shows a summary of temperatures at 

key locations in the head and block for cast iron and 

Figure 19. Computed temperatures at the top of the cylinder liner for the Ti engine rated at 305 hp and at 450 hp. 

Table 11. Computed temperatures at key locations in the 

head and block for cast iron and aluminum 

Location 

All-cast-iron 

engine 

All-aluminum 

engine 

Top of block 227 165 

Top of liner 230 192 

Cylinder head 

exhaust valve 

bridge 

370 210 

Intake valve seat 357 249 

Exhaust valve seat 444 325 

Al. FEA modeling of high cycle fatigue was performed on the Al head and block and compared with 

modeling for cast iron. Figure 20 shows an example 

of the model for cover factors (safety margin) 

greater than 1 (no failure). It can be seen that the 

aluminum compares very favorably with the cast 

iron. 

FEA modeling of fatigue was performed for the 

block. The cover factors at key locations on the aluminum block were compared with those locations on 

the cast iron production block (see Table 12). With 

the exception of one location, the cover factors for 

the aluminum block were equal to or greater than 

those for the cast iron block. The one exception is 

the region from the bearing saddle to the block fillet 

(location AJ), shown in Figure 21. The cover factor 

in this region can easily be increased by making the 

fillet region thicker. The higher fatigue cover factors 

of Al compared with cast iron are caused by lower 

temperatures, lower thermal stresses, and the higher 

strength of aluminum. The low cover factors near 

the head bolt bosses are caused by high compressive 
Table 10. Mechanical and physical properties of gray cast iron and A354-T6 Gray cast iron (liner insert) Temperature (Centi- grade) Young’s modulus (GPa) Poisson’s ratio Thermal conduc- tivity (Wm-K0 Thermal expan- sion (:m/m-C) Ultimate tensile strength (MPa) Compressive yield strength (MPa) Tensi

yield str

(MPa

20.0 117.2 0.25 48.80 1.10E-05 189.6 284.4 136.

200.0 47.80 1.10E-05 300.0 46.80 400.0 45.80 1.25E-05 500.0 44.80 Aluminum (A3540-T6) – Data supplied by Cummins, Inc. Temperature (Centi- grade) Young’s modulus (GPa) Poisson’s ratio Thermal conduc- tivity (Wm-K0 Thermal expan- sion (:m/m-C) Ultimate tensile strength (MPa) Compressive yield strength (MPa) Tensi

yield stre

(MPa

20.0 76.1 0.33 168.80 2.03E-05 265.2 233.8 241.50.0 170.90 100.0 175.50 2.09E-05 150.0 62.1 0.33 177.60 213.3 216.7 205.

175.0 65.6 0.33 204.1 212.6 201.200.0 66.7 0.33 179.80 2.16E-05 202.5 215.6 199.260.0 54.5 0.33 180.80 142.7 142.2 142.300.0 2.24E-05 

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FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

Figure 20. FEA model of temperatures in the valve 

bridge region of the fire deck for the production cast iron 

engine and the Al engine. 

Table 12. Fatigue cover factors at key locations 

of the cast iron and aluminum blocks 

Cast iron Aluminum 

Location Description 

Fatigue 

cover factor 

Fatigue 

cover factor 

AE* interior fillet to 

top deck @ end 

face 

0.12 0.13 

AF oil hole edge to 

brg. shell 

0.66 1.59 

AG corner of side 

inlet to water 

jacket 

0.87 1.04 

AH internal fillet to 

head bolt boss 

0.89 0.89 

AI corner of side 

inlet to water 

jacket 

0.98 1.26 

AJ Bearing saddle 

to block fillet 

1.87 0.94 

AK Bearing saddle 

to block fillet 

1.88 1.06 

*Model geometry found not to be representative of actual 

casting 

stresses induced by bolt loads. Washers under the 

bolt heads to reduce embedment are a possible solution to this issue. 

Conclusions

• Aluminum can be a viable substitute for cast 

iron in this and many other diesel engines, reFigure 21. FEA calculations of temperature at the top of 

the engine block for the cast iron production engine and 

for the Al block. 

sulting in at least an overall engine mass reduction of about 20%. 

• Some design changes would be needed for the 

Al engine to 

― support the cast iron liner 

― increase the stiffness of the cylinder head 

top deck 

― modify the interference fits between valve 

seat inserts and the head 

• Since Al has no fatigue endurance limit, it 

would be prudent to determine acceptable life 

targets so as not to have cracking after many 

miles or hours of use. 

Magnesium 

An FEA study of Mg, identical to the one for Al, 

was conducted to determine if Mg is a feasible replacement for cast iron in heavy duty diesel engine 

heads and blocks. The FEA analysis was performed 

on the same Cummins ISB engine at 305 hp, but 

with cast iron cylinder liners. 

From previous analyses, it was considered important that a Mg alloy be selected that had maximum elevated temperature properties. A highperformance Mg casting alloy, WE-43, was selected 

for the study. Table 13 contains the composition of 

the alloy along with T-6 heat treatment details. 

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Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

155

Table 13. Composition and heat treatment 

of WE-43 Mg alloy 

• Composition 

• Yttrium 

 3.7 – 4.3 

- Rare Earths 

(Ce,La,Nd) 

- 2.4 – 4.4 

- Zirconium 

 0.4 min. 

• Optimum heat treatment 

- Solution treat for 8 h at 

525°C. 

- Air cool, hot water or 

polymer quench. 

- Age for 16 h at 250°C, 

air cool 

Figure 22. FEA model of the cast iron and aluminum 

heads showing fatigue cover factors in key locations. 

Table 14 shows a comparison between the mechanical and physical properties of cast iron and 

WE-43, while Figure 22 is a graph of the thermal 

conductivities of cast iron and WE-43 vs. temperature. 

The FEA of the valve bridge region of the head 

is shown in Figure 23 for production cast iron and 

for WE-43. It can be seen that although the temperature of the valve bridge region in the Mg head is 

lower than for the gray cast iron, it is 60°C above the 

maximum recommended use temperature for this 

alloy. Figure 24 shows the FEA temperature predictions for the top of the block. It shows that there are 

no temperature issues with using Mg for the block. 

Fatigue cover factors were calculated for the Mg 

head and compared with those for the production 

cast iron head. The data are shown in Table 15. With 

one exception, location N, the cover factors for Mg 

(WE-43) are the same or higher than for the cast iron 

and are nearly the same at location N. FEA analysis 

of fatigue cover factors for the Mg block at key locaFigure 23. FEA model of the cast iron and aluminum 

blocks showing cover factors at key locations. 

Figure 24. Thermal conductivity vs. temperature for WE43 and grey cast iron. 

tions is summarized in Table 15. The cover factor 

for the Mg block is higher or the same in all key locations, compared with the cast iron block. Location 

AE is low no matter which material is used: cast 

iron, Ti, Al, or Mg. It can be raised by adding more 

material to that location. 

Conclusions

FEA has shown that it is feasible to substitute a 

high-performance Mg casting alloy, WE-43, for 

gray cast iron in heavy-duty diesel engine heads and 

blocks, for a potential overall engine weight savings 

of over 30%. However, to do this, certain issues 

must be addressed: 

• Compressive yielding is likely to occur in the 

valve bridge areas of the flame face of the head 
Table 14. Mechanical and physical properties of gray cast iron and Mg (WE-43) Gray cast iron (liner insert) Temperature (Centi- grade) Young’s modulus (GPa) Poisson’s ratio Thermal conduc- tivity (Wm-K0 Thermal expan- sion (:m/m-C) Ultimate tensile strength (MPa) Compressive yield strength (MPa) Tensi

yield stre

(MPa

20.0 117.2 0.25 48.80 1.10E-05 189.6 284.4 136.

200.0 47.80 1.10E-05 300.0 46.80 400.0 45.80 1.25E-05 500.0 44.80 Magnesium (WE43-T6). Temperature (Centi- grade) Young’s modulus (GPa) Poisson’s ratio Thermal conduc- tivity (Wm-K0 Thermal expan- sion (:m/m-C) Ultimate tensile strength (MPa) Compressive yield strength (MPa) Tensi

yield stre

(MPa

20.0 44.1 0.27 51.30 2.67E-05 265.0 185.0 185.0 

150.0 42.0 0.27 64.80 2.92E-05 250.0 175.0 175.0 200.0 39.0 0.27 69.50 245.0 170.0 170.0 250.0 36.0 0.27 73.80 3.00E-05 220.0 160.0 160.0 300.0 36.0 0.27 77.90 160.0 120.0 120.0 350.0 81.70 400.0 85.20 

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Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

Table 15. Fatigue cover factors at key locations 

in the head for cast iron and Mg (WE-43) 

(lwj = lower 

water jacket) 

(uwj = upper 

water jacket) Cast iron Aluminum 

Location Description 

Fatigue 

cover factor 

Fatigue 

cover factor 

A uwj fillet between 

bolt boss and 

valve guide 

0.53 0.58 

B lwj floor fillet 

between exh. port 

and inj. bore 

0.61 1.01 

C lwj floor fillet to 

exhaust port 

0.65 1.07 

D lwj roof fillet to 

head bolt boss 

0.68 0.75 

E uwj roof fillet to 

head bot boss 

0.69 0.73 

F lwj floor fillet to 

exhaust port 

0.72 1.05 

G intake port to 

valve seat area 

0.72 1.50 

H lwj roof fillet to 

head bolt boss 

0.75 0.98 

I exterior o exhaust 

side near flame 

deck 

0.81 1.14 

J uwj roof fillet to 

head bolt boss 

0.82 0.94 

K uwj roof fillet to 

head bolt boss 

0.86 0.93 

L lwj roof fillet to 

head bolt boss 

0.88 1.18 

M uwj fillet between 

head bolt boss 

exhaust port 

0.88 1.04 

N exterior pocket @ 

top deck near 

head bolt boss 

0.92 0.88 

O uwj roof fillet to 

top deck 

0.93 0.98 

P lwj fillet – valve 

guide to fuel drilling 

0.94 1.57 

Q top deck rib between guides 0.94 0.98 

R lwj floor between 

exh. Ports and inj. 

bore 

0.95 1.34 

T Exterior pocket @ 

top deck near 

head bolt boss 

0.96 1.08 

U uwj roof to coolant hole 0.96 1.01 

V lwj fillet – exh. 

guide to exh. port 

0.98 1.09 

W lwj floor fillet to 

exhaust port 

0.99 1.47 

as a result of high temperatures coupled with 

high stresses. Low cycle fatigue or creep could 

be the result. 

• Ways to reduce the temperature in the valve 

bridges to below 250°C must be explored. Some 

ideas are 

― design changes to reduce the stresses in the 

area 

― thinning of the fire deck to improve cooling 

― higher coolant flow to accommodate the 

heat flow 

― slight de-rating of the engine so less heat 

and lower stresses would be generated in the 

area 

• Other design changes should be explored in order to modify the interference fit between the 

valve seat inserts and the head and to bolster the 

support for the cast iron liner. 

• An alternative to making the design changes is 

to use an Al head on the Mg block. 

• Mg is well known as a metal that promotes galvanic corrosion. Should the engine contain substantial amounts of Mg, then the issues of 

galvanic corrosion must be explored. 

Conclusions of Overall Study 

• FEA modeling of a Cummins ISB 5.9-L diesel 

engine has shown that all three lightweight alloys, Ti-6Al-4V, A354-T6 (Al), and We-43 

(Mg), are feasible substitutes for gray cast iron 

in a heavy duty diesel engine. The modeling incorporated a gray cast iron liner in the block for 

the lightweight alloys because the production 

cast iron engine has no liner (parent bore). The 

liner acts in two ways: it provides a viable tribological surface for the piston rings, and it 

makes the study more generic and less enginespecific. 

• Each lightweight material posed unique challenges to making the material work in this engine:

― The Ti alloy could be used as a direct substitute for cast iron in the block when the engine was rated at either 305 hp, as designed, 

or up-rated to 450 hp.

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FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

― The Ti alloy posed a problem in the valve 

bridge region of the fire deck of the head 

because of its low thermal conductivity. It 

was shown that if a Be-Cu insert were employed in that region, it would be feasible to 

use Ti when the engine was up-rated to 450 

hp. This strategy would allow for a 10% 

weight saving and an up-rating of 50%.

― Using the Al alloy for the head and block 

would allow for a 20% weight saving. FEA 

of the engine rated at 305 hp showed that Al 

is a feasible substitute for gray cast iron. 

Note that since Al has no fatigue endurance 

limit, life cycle analysis needs to be performed to determine when to use this material.

― It is feasible to substitute a highperformance Mg alloy for gray cast iron in 

this engine rated at 305 hp. However, the 

temperatures and stresses in the valve bridge 

areas of the fire deck in the head are slightly 

too high and need to be reduced. This can be 

done by increased coolant flow, some design 

changes, or a slight de-rating of the engine.

Presentations

Presentation to TACOM, March 2004. 

Presentation to Cummins Engine Co., April 

2004 

Presentation to Caterpillar, April 2004. 

Presentation at a low-cost Ti users conference 

sponsored by the Defense Advanced Research Projects Agency, April 2004. 

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Heavy Vehicle Propulsion Materials FY 2004 Progress Report

159

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

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: Dr. Sidney Diamond

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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-ACo5-00OR22725

Objective

• Characterize and model contact-induced damage mechanisms in ceramics and link them to microstructure and

ultimately to wear performance, the optimization of machining, rolling contact fatigue, etc.

• Characterize the influence that independent parameters such as grain size, loading rate, confinement or residual

stresses, temperature, and field effects have on the evolution of contact-induced damage.

Approach

• Develop and use instrumented static and dynamic indentation and instrumented scratch test methods to interrogate quasi-plasticity and fracture processes in ceramics.

• Model contact-induced quasi-plasticity and fracture processes.

• Link damage mechanism activity to wear performance, machining, contact fatigue, etc.

• Develop strategies for engineering control (e.g., confinement, temperature) that promote the dominance of one

mechanism (e.g., one that causes strengthening or hardening) over others.

Accomplishments

• Established dynamic indentation and instrumented scratch testing facilities.

• Nearly completed indentation rate effect studies on several silicon nitride ceramics.

• Developed an ANSYS model that enables the study of indenter depth-of-penetration in a target material while

considering elastic and yield properties of the indenter and target material, indenter diameter, etc.).

Future Direction

• Characterize the influences that grain size and residual/confinement stress have on damage evolution.

• Link indentation-induced damage evolution to that produced by scratch testing.

• Add tension-compression strength/yield anisotropy effects to indentation model.
FY 2004 Progress Report Heavy Vehicle Propulsion Materials

160

Introduction

More ceramic components could be used in engines and transportation systems if they could be

confidently used in more wear-related applications

and if they could be manufactured and machined

faster (i.e., more cost-effectively). This would occur

if contact-induced damage mechanisms (which 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 testing. Understanding the

competition and interaction among those processes

is a critical part of this project as well. Though more

traditional structural ceramics (e.g., Si3N4 and ZrO2)

under consideration for diesel engine components

are the primary focus, other classes of ceramics and

brittle materials that could provide greater understanding of damage evolution fundamentals are also

studied. These include cubic oxides and non-oxides,

nanoceramics and nanocermets, piezoelectric ceramics, and micaceous ceramics.

Through quantifying the loads or stresses that

initiate quasi-plasticity and fracture, contact-induced

damage mechanisms in ceramics are characterized;

modeled; and linked to microstructure and ultimately to wear performance, optimization of machining, rolling contact fatigue, etc. Additionally,

the influence of independent parameters such as

grain size, loading rate, confinement or residual

stresses, temperature, and field effects on the evolution of contact-induced damage are examined. A

thorough understanding of the competition of fracture and quasi-plasticity in ceramics will enable improved and faster means of manufacturing ceramic

components and surface engineering (e.g., ductile

regime machining) and will maximize mechanical

performance in cases when surface conditions

(e.g., bending) or surface-located events (e.g., wear,

impact) are service-life limiters in engine and transportation system components.

Approach

FY 2004 was the first full year of this project,

and numerous parallel efforts were initiated. The

project has five primary elements; three represent

ongoing and near-term efforts, and the other two

represent longer-term efforts. The first element is to

develop and use instrumented static and dynamic

indentation and instrumented scratch test methods

(see Figure 1) to interrogate quasi-plasticity and

fracture processes in ceramics. Second, contactinduced quasi-plasticity and fracture processes are

modeled. Third, ceramic materials are characterized

that either are of strategic relevance to heavy-duty

diesel engine components or serve as model materials for the examination of damage evolution. Fourth,

damage mechanism activity is linked to wear performance, machining, contact fatigue, etc. Finally,

strategies are developed for engineering control

(e.g., confinement, temperature) that promote the

dominance of one mechanism (e.g., one that causes

strengthening or hardening) over others.

Results

In order to quantify quasi-plasticity and fracture

processes, instrumented static and dynamic indentation test systems were established in FY 2004, as

well as instrumented scratch testing capabilities.

Each of the instruments is shown in Figure 1a–c, and

are all in use to experimentally characterize quasiplasticity and fracture processes. A dedicated acoustic emission (AE) system is interfaced with all these

instruments and is used to detect damage processes

that result during indentation or scratch testing.

An example of a Hertzian indent is shown in

Figure 2. The depth of the crater-like indent is an

indication that quasi-plasticity was initiated. A circular ring crack and a median crack (2 and 8 o’clock

positions) are indications that at least two different

crack systems were initiated. The loads and indenter

depth-of-penetration associated with the initiation of

the damage mechanisms are measurable using the

static instrumented indentation test system (IITS)

and the dynamic indentation test systems (DITS).

Their values are compared with those for other ceramics and interpreted with respect to traditionally

measured mechanical properties, such as hardness,

fracture toughness, and tensile strength.

AE testing and analysis was an integral part of

the indentation testing because it can be efficiently

used to identify the loads associated with the initiations of cracks (i.e., fracture processes). Measured

load-unload histories and the loading predicted by

finite element analysis (FEA) provide a reference to

the AE interpretation. For example, for the spherical

indent-generated ring cracks formed in the concentric pattern shown in Figure 3a, the AE spectrum
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

161

Figure 1. Facilities established included instrumented

static (a) and dynamic (b) indentation test systems and a

scratch tester (c).

(Figure 3b) shows that numerous acoustic events

had occurred, with some events having greater

amounts of energy than others (e.g., the event that

occurred at ~ 175 N). Additionally, acoustic events

Figure 2. Hertzian indent illustrating quasi-plasticity

(dimple) and fracture processes (ring and median

crack). Indent was generated with a 2.25-mm-diam

diamond indenter and 300 N maximum load.

did not occur until a load of ~ 48 N, indicating that

the material behaved elastically below that load. The

load at which divergence of the black and red curves

(i.e., where permanent damage is introduced) occurred in the indenter depth-of-penetration load

curve matches fairly well with the load where

acoustic signals begin to be produced, and the load

where the highest-energy acoustic peak occurred

matches the load on the experimental loading curve

(black curve) where the slope arguably begins to

increase (i.e., a relatively large increase in specimen

compliance).

The DITS and instrumented (pendulum) scratch

tester were interfaced with a high-speed data acquisition system, and experimental familiarity was

gained with both testers. Gas gun settings, shape

pulsing, and overall test approaches were optimized

with the dynamic indenter to produce controllable

low-load dynamic indentations (i.e., a few to tens of

Newtons). Also, an ability to test specimens having

a nominal geometry of 3×4×4 mm (a specimen cut

from conventional ASTM C1161B bend bars—a

geometry that is readily available for many materials

in the Ceramic Science and Technology Group) was

established with the scratch tester. The specimen

holder that came with the instrument (which also

houses two load cells for normal and thrust force

measurements) was designed to test specimens having the geometry of a 1⁄4-in. (6.35-m) cube—an inconvenient size.

An indentation model was developed using

ANSYS FEA software that predicts the stress state
FY 2004 Progress Report Heavy Vehicle Propulsion Materials

162

Figure 3. Hertzian indent of a soda lime glass (a), and (b)

its measured load-unload waveform compared with the

predicted loading waveform and the acoustic emission

history associated with the generated crack pattern. The

indent was generated with a 2.0-mm-diam diamond indenter and 300 N maximum load.

in a material subjected to Hertzian (spherical) indentation. An illustration of the model is shown in

Figure 4 with a predicted response shown in Figure 5. Independent parameters that may be considered in the model are elastic properties and hardness

(i.e., yielding) of the target or indenter material, indenter diameter, coefficient of friction, confinement

pressure, and, of course, applied indenter load. InFigure 4. ANSYS Hertzian indentation model.

Figure 5. Example of predicted indentation response. The

total indent-depth-of-penetration (top curve) is the sum

of elastic (middle curve) and permanent deformation

(bottom curve).

terpretation of the von Mises and First Principal tensile stresses is enabling the study of the competition

between plasticity and fracture processes in the test

material (linkable to machining and wear phenomena), and the results are being correlated to experimental observation.

μ-FEA, a microstructure-based FEA computer

program, was developed and its copyrighting is under way through ORNL’s Technology Transfer and

Economic Development organization. μ-FEA is a

graphics-user-interfaced LabVIEW executable program that serves as a pre- and post-processor to the

ANSYS FEA solver, and it enables stress analyses

of both real (e.g., imaged with scanning electron

microscopy or optical microscopy) and simulated

material microstructures. An example of an analyzed

microstructure and its subsequent residual stress calculation is shown in Figures 6a and b. For one of its

applications, μ-FEA is being used to study stress

concentrators, deformation, and microcracking
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

163

Figure 6. Digital microstructure of (a) coating and (b) μFEA residual stress evaluation induced from thermal expansion mismatches between material constituents.

around pores and large grains and in secondary

phases in materials subjected to a Hertzian contact

stress gradient.

Silicon nitride specimens of Ceralloy 147-31N,

NBD200, SN101C, NC132, and TSN-03NH were

sectioned and either underwent or are undergoing

instrumented static and dynamic indentation and

instrumented scratch testing. An apparent rate dependence (i.e., quasi-plasticity rate dependence) on

Vickers Hardness for Ceralloy 147-31N is shown in

Figure 7. There is a rationale behind the choice of

testing each silicon nitride composition. The 147-

31N grade is presently used as a cam roller follower

material, so interest exists in quantifying its indentation and scratch responses. NBD200, SN101C, and

TSN-03NH are ball bearing grades, so their testing

is relevant to the examination of rolling contact fatigue behavior. NBD200 is a NIST standard reference material for Knoop hardness, and NC132 is a

Figure 7. Static and dynamic Vickers hardness of Ceralloy 147-31N as a function of indentation rate.

NIST standard reference material for fracture toughness, so they serve as references for performance

comparison among all the silicon nitrides.

To examine the effect of grain size on damage

due to cracking and quasi-plastic yielding generated

with instrumented indentation and scratch testing,

three different mean-grain-sized (2–3, 15, and

25 μm) aluminas were acquired (Ceradyne, Costa

Mesa, CA). Microstructures of the three materials

are shown in Figure 8. The alumina has 99.9% purity, has magnesium as a sintering aid, and has relatively narrow size distributions. Alumina was

chosen as the model material in which to examine

grain size effects because of the ease associated with

measuring residual stresses on it using optical fluorescence, and because of the immediate availability

of such an analytical tool. Residual stresses will be

measured in, around, and under generated indents

and scratches, and the results will likely prove to be

valuable for interpreting grain size effects on cracking and quasi-plasticity.

Figure 8. Microstructures of 99.9% alumina with average

grain sizes of (left) 2–3 microns, (middle) 15 microns,

and (right) 25 microns.

The effects of confinement (i.e., pressure) on instrumented static and dynamic indentation, instrumented scratch testing, and the competition of
FY 2004 Progress Report Heavy Vehicle Propulsion Materials

164

fracture and quasi-plastic damage mechanism behaviors are under examination. The 99.9% alumina

is being machined, and metal sleeves/tubes will be

shrink-fitted about it to produce a radial confinement. The residual stress field on the surface of the

radially confined alumina will be quantified using

optical fluorescence; and the static, dynamic, and

instrumented scratch testing will commence on it.

Because the aluminas described will also serve as

the model materials for this examination, grain size

will also be an independent parameter in this confinement effect study. If confinement introduces

beneficial effects related to damage associated with

indentation or scratch testing, then the characterization of that effect will lead to its implementation into

grinding or wear applications.

Experimental strategies are being pursued to enable the study of the effect of temperature on instrumented static and dynamic indentation and

instrumented scratch testing. The characterization of

this effect will have a direct link to optimizing heatassisted machining of ceramic materials or promoting minimized surface damage that will result in

improved wear resistance and strength.

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 are under investigation, and

rate effects on hardness (i.e., quasi-plasticity) are

evident. An ANSYS model was developed that enables the study of indenter depth-of-penetration in a

target material while considering elastic and yield

properties of the indenter and target material, indenter diameter, etc. A microstructure-based FEA

software (μ-FEA) was developed to enable stress

analysis characterization at the microstructural scale

and to help with the interpretation of quasi-plastic

and fracture processes.

FY 2004 progress provides a nice springboard

for planned FY 2005 work. The influences that grain

size and residual/confinement stress have on damage

evolution will be characterized. Indentation-induced

damage evolution produced by scratch testing will

be linked to that generated from static and dynamic

indentation. Tension-compression strength/yield

anisotropy effects will be added to the indentation

model and will enable the study of how (or if) hardness is linked to quasi-plasticity and how (or if) tensile strength and fracture toughness are linked to

fracture processes.

Publications/Presentations

“Evaluation of Ceramic Deformation Processes

Through Hertzian Indentation,” presented at the

Tank and Automotive Research Development and

Engineering Center-ORNL-Army Research Laboratory Advanced Materials Meeting, Network Computing Services, Minneapolis, February 5, 2004.

“ORNL Characterization of Ceramics for Armor, Transportation, and Energy Applications,” presented at Cercom, Inc., Vista, CA, February 25,

2004.

“ORNL Characterization of Ceramics for Armor, Transportation, and Energy Applications,” presented at Ceradyne Inc, Costa Mesa, CA, February

26, 2004.

“Evaluation of Ceramic Deformation Processes

Through Instrumented Hertzian Indentation,” presented at 2004 Interagency Coordinating Committee

on Structural Ceramics, National Science Foundation, Arlington, VA, April 14, 2004.

“Evaluation of Ceramic Deformation Processes

Through Hertzian Indentation,” presented at SaintGobain, Worcester, MA, August 3, 2004..

Copyrights

μ-FEA, a microstructure-based finite element

analysis computer program, was developed and is

undergoing copyrighting through ORNL’s Technology Transfer and Economic Development organization. μ-FEA is a graphics-user-interfaced

LabVIEWTM 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 scanning electron microscopy

or optical microscopy) and simulated material microstructures.
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

165 

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: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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: 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 mechanical properties of the nanostructured chips. 

• Develop methods for converting the nanostructured chips into particulates suitable for processing into bulk 

forms. 

• Develop powder processing methods, including shear-based densification (e.g., extrusion, rolling, forging) and 

infiltration processing, for creating monolithic and composite materials in bulk form with nanocrystalline microstructure. 

• 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. 

• Achieved typical grain sizes in the chips of 50–700 nm, depending on the strain and temperature induced by the 

machining. 

• Consolidated cylindrical bulk samples of nanostructured Al-6061T6/pure aluminum composites by powder extrusion. Densification in excess of 95% has been obtained while retaining the hardness of the nanocrystalline 

structures. 

• Fabricated bulk samples of nanostructured Al-6061T6, with small amount of polymer as binder, by infiltration 

and pressing. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

166 

Future Direction 

• Continue development of shear-based densification and infiltration methods to produce bulk monolithic and/or 

composite alloy samples with nanocrystalline microstructure, using the chips as precursor materials. 

• Characterize microstructure and properties of the bulk forms to demonstrate retention of the nanostructures and 

enhanced mechanical properties. 

• Determine tensile strength, ductility and density of the bulk forms to demonstrate integrity of the consolidated 

alloy samples. 

• Identify 1–2 applications in the transportation sector for prototype development.

Introduction

It has been widely reported that nanostructured 

materials, composed of nanometer-scale grains, have 

properties significantly different from those of conventional materials.1,2 These properties include 

higher hardness, strength and ductility, and superplasticity at relatively low temperatures and sintering temperatures that are several hundred degrees 

below those of microcrystalline powders.3,4 Widespread use of these materials, including bulk forms 

(components and solid bodies) with nanocrystalline 

microstructure, has been sought; however, the cost 

of creating these materials, frequently quoted as being more than $100 per pound, and the inability to 

produce them in tonnage quantities with reproducible properties have restricted their broad application.5

This project builds on the recent discovery of a 

low-cost means of producing nanostructured materials in essentially any alloy.6,7 The discovery that 

chips produced during normal machining operations 

are composed entirely of nanocrystalline structures 

of high hardness (strength) provides a low-cost, direct way of making these materials in large volumes. 

Comminution of the nanostructured chips will enable large-scale production of nanocrystalline particulate, which can be converted into nanostructured 

bulk forms using powder processing methods. Alternatively, the chips and particulates may be used 

as continuous or discontinuous reinforcements in 

metal or polymer matrices to create advanced composite bulk forms. 

Approach

Chips were produced by machining, ranging 

from soft alloys [e.g., oxygen-free high-conductivity 

(OFHC) copper, aluminum 6061-T6[ to highstrength alloys (e.g., Inconel, 52100 steel); and the 

microstructure and hardness of the chip and bulk 

(workpiece) material was characterized.7,8 The bulk 

materials were annealed in an inert atmosphere before machining so that he initial grain size in all 

cases was greater than 20 μm. Typical dimensions of 

chip samples examined were 100–3000 μm width, 

100–1000 μm thickness, and at least 5 mm length, 

with the smaller chips coming from the higherstrength steels. The machining conditions were selected so that the temperature increase in the deformation zone was minimal for all materials. 

Transmission electron microscopy (TEM) was 

used to characterize the nanoscale defect structures 

in the metal chips and to measure grain size, grain 

size distribution, and grain misorientation. Specimens for TEM examination were prepared from 

chips produced under well-defined deformation 

conditions (strain, strain rate, temperature) by standard mechanical, ion beam, and electrochemical 

thinning techniques. Bright-field imaging coupled 

with electron diffraction analysis was conducted in a 

scanning TEM at Purdue and at the High Temperature Materials Laboratory (HTML) at Oak Ridge 

National Laboratory (ORNL). The hardness of the 

chip and bulk samples was measured by Vickers and 

nano-indentation. Uni-axial tests were carried out on 

miniature copper tension samples prepared from the 

chips. 

Preparation of bulk forms from nanostructured 

particulate derived from the machining chips was 

initiated using powder extrusion and infiltration 

processing. For powder extrusion, nanostructured 

Al6061-chips were converted into particulate using 

attrition milling. The resulting particulate was mixed 

with varying weight percentages of commercially 

available, gas-atomized aluminum powder; this mixture was powder-extruded through steel dies at extrusion ratios of 10–60. The extruded samples were 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

167 

prepared in the form of cylinders of up to 10 mm in 

diameter and 5–25 mm in length. 

Results 

Figure 1 shows a TEM image of a typical Al 

6061-T6 chip. Shown in the inset is the corresponding selected area diffraction pattern, typical of a 

polycrystalline structure. The grain size is seen to be 

in the range of 50–80 nm. Misorientation measurements carried out using convergent beam electron 

diffraction in the TEM have established that more 

than 50% of the grain boundaries in the Al 6061-T6 

chips are large-angle; the fraction of the large-angle 

boundaries is greater in chips produced at higher 

levels of strain. Evidence for mechanical twinning 

has been seen in some of the highly strained chips, 

akin to that reported elsewhere.9

 Similar nanocrystalline structures have been observed in a variety of 

alloys, including high-strength 52100 and M2 steels 

(Figure 2). Ultrafine sub-grain structures, nanoscale 

grains, and dynamic recrystallization have been 

demonstrated in a variety of materials by varying the 

conditions of strain and temperature during chip 

formation. An example of such nanostructural 

changes can be seen in the sequence of TEM pictures in Figure 3 of OFHC copper chip samples produced with different levels of shear strain (~ 3–20). 

A switchover from sub-grain to nanoscale, equiaxed 

grain structures at the higher levels of strain (Figure 3) appears to be a consequence of the onset of 

dynamic recrystallization at a strain of ~ 13. This 

switchover, which appears controllable in different 

materials by varying the deformation conditions, has 

not been realized in a single stage of deformation in 

any of the frequently studied severe plastic deformation processes. These observations suggest that machining also offers a unique framework in which to 

systematically explore the formation of nanocrystalline structures by large strain deformation. 

Table 1 gives a summary of the hardness values 

measured on the chips and the bulk samples. The 

chips are seen to be up to three times as hard as the 

corresponding bulk samples prior to machining. The 

hardness of the 52100 steel chip is 1310 kg/mm2

, 

which is comparable to that of patented steel wire, 

one of the hardest and strongest steel structures 

known.10 As has been reported in many studies,1–4

nanostructured alloys have significantly greater 

strength than their microcrystalline counterparts. 

Uni-axial tensile tests carried out on miniature 

Figure 1. Nanostructured Al 6061-T6 

Figure 2. 52100 steel showing nanoscale ferrite grains 

and carbide particles. 

OFHC copper chip samples have shown the increases in strength to parallel those in hardness. 

These increases in chip strength and hardness are 

much higher than those that result from extrusion, 

forging, and rolling.11

Figure 4 shows an optical micrograph of a partially formed OFHC copper chip, as yet not separated from the bulk.6

 Superimposed on the 

micrograph are Vickers hardness (strength) values 

recorded at different locations. The hardness shows 

a steep increase over a very narrow deformation 

zone between the chip and the bulk. This zone, 

where chip formation occurs, is characterized by 

very large deformations, as highlighted by the flow 

lines in the chip in Figure 4. While grains are clearly 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

168 

Figure 3. Microstructure of OFHC copper chips at different levels of shear strain (γ). Features typical of onset of 

dynamic recrystallization are visible in the chip produced 

with a shear strain of ~13. 

Figure 4. Partially detached OFHC copper chip. Note 

large hardness increase across shear plane. 

visible in the bulk OFHC copper sample, no grains 

can be resolved in the chip or near the deformation 

zone, consistent with TEM observations that the microstructure is nanoscale. 

Figure 5a shows a cylindrical sample of Al6061-

T6/Al composite produced by powder extrusion. 

This composite is composed of 50% by weight of 

nanostructured 6061 chip particulate blended with 

an equal weight of aluminum powder. Density 

measurements showed the sample to be consolidated 

to 97% of the theoretical density. Figure 5b shows a 

transverse cross-section of this sample; the porosity 

is seen to be negligible, consistent with the measured density. Figure 5c shows the distribution of 

hardness over the cross-section. The higher hardness 

values of ~160 VHN are similar to those measured 

in the nanostructured chip particulate, suggesting 

that the nanostructure is retained. The remainder 

most likely correspond to that of the aluminum 

powder strain-hardened by the extrusion. The early 

observations pertaining to hardness, density, and 

microstructure, made on the powder-extruded samples, are promising vis-à-vis densification and 

nanostructure retention. 

We have also been successful recently at producing bulk Al6061-T6 samples from nanostructured chip particulate via an infiltration/pressing 

approach using a small amount (<10%) of epoxy as 

binder. 

Conclusions 

Nanostructured materials with grain sizes as 

small as 70 nm have been produced by controlled 2-

dimensional machining in a wide variety of metals 

and alloys. The hardness of these alloys is up to 

three times that of their micro-crystalline counterparts. Preliminary results on consolidation of these 

materials using shear-based densification and infiltration processes are promising. A variety of lowtemperature consolidation protocols have been identified for producing bulk forms using these nanostructured chips as precursor materials. These include 

powder extrusion, rolling, and forging; infiltration 

processing; dynamic compaction; rapid prototyping; 

and spraying. In the next year, intensive efforts are 

planned on using these processes to address the 

grand challenge of consolidating these nanostructured materials into high-performance bulk forms. 

Property and performance studies will be initiated 

on these consolidated samples to determine indices 

of relevance to transportation applications. Close 

collaboration with the HTML is envisioned on various aspects of this work. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

169 

Figure 5. Characteristics of powder-extruded Al-6061T6/Al composite: (a) typical extruded samples, (b) microstructure 

of a transverse cross section, and (c) hardness distribution determined by Vickers indentation. 

References

1. H. Gleiter, Progress in Materials Science 33, 

1 (1989). 

2. R. W. Siegel, Scientific American 275 74 

(1996). 

3. L. Lu, M. L. Sui, and K. Lu, Science 287, 

1463 (2000). 

4. X. McFadden, R. S. Mishra, R. Z. Valiev, A. 

P. Zhilyaev, and A. K. Mukerjee, Nature 398, 684 

(1999). 

5. H. Breisen, A. Fuhrmann, and S. E. Pratsinis, 

“Nano-Powder and Industrial Applications,” MRS 

Proc. 520, 3 (1998). 

6. T. L. Brown, S. Swaminathan, S. 

Chandrasekar, W. D. Compton, A. H. King, and K. 

P. Trumble, J. Matls. Res. 17 (2002). 

7. S. Chandrasekar, W. D. Compton, T. N. Farris, and K. P. Trumble, “Method of Forming Nanocrystalline Structures and Product Formed Thereof,” 

U.S. Patent, 6,706,324, March 16, 2004. 

8. S. Swaminathan, S., T. L. Brown, M. Ravi 

Shankar, B. C. Rao, K. P. Trumble, W. D. Compton, 

and S. Chandrasekar, “Large Scale Manufacturing of 

Ultra-fine grained Materials,” Ultrafine Grained 

Materials III, Ed. Y. T. Zhu, T. G. Langdon, R. Z. 

Valiev, S. L. Semiatin, D. H Shin, and T. C. Lowe, 

TMS (The Minerals, Metals and Materials Society), 

2004. 

10. Y. M. Wang, M. W. Chen, F. H. Zhou, and 

E. Ma, Nature 419, 912–915 (2002). 

11. A. Kelly and N. H. MacMillan, Strong Solids, Clarendon, Oxford, UK, 1986. 

12. W. A. Backofen, Deformation Processing, 

Addison Wesley, Reading, Massachussets, 1972. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

170 

Publications/Presentations

S. Swaminathan, S. Chandrasekar, W. D. Compton, K. P. Trumble, and A. H. King, “Microstructural Refinement in Single-Phase Copper Solid 

Solutions by Machining,” p. 8.27 in Nanoscale Materials and Modeling-Relations Among Processing, 

Microstructure and Mechanical Properties, Materials Research Society (MRS) Proceedings, 821, 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,” p. 9.5 in Nanoscale Materials and 

Modeling-Relations Among Processing, Microstructure and Mechanical Properties, Materials Research 

Society (MRS) Proceedings, 821, 2004. 

S. Swaminathan, T. L. Brown, M. Ravi Shankar, 

B. C. Rao, K. P. Trumble, W. D. Compton and S. 

Chandrasekar, “Large Scale Manufacturing of Ultrafine Grained Materials,” pp. 161–166 in Ultrafine 

Grained Materials III, The Minerals, Metals and 

Materials Society (TMS) Proceedings, 2004. 

T. L. Brown, S. Swaminathan, B. C. Rao, R. F. 

Kezar, S. Chandrasekar, W. D. Compton, K. P. 

Trumble, and A. H. King, “Machining as a Method 

for Studying Effects of Very Large Strain Deformation,” pp. 167–172 in Ultrafine Grained Materials 

III, The Minerals, Metals and Materials Society 

(TMS) Proceedings, 2004. 

S. Chandrasekar, W. D. Compton, and K. P. 

Trumble, “Low-cost Nanocrystalline Materials for 

Structural Components,” presented at the U.S. Air 

Force Propulsion Environmental Working Group 

Meeting, Las Vegas, April 2004. 

S. Chandrasekar, W. D. Compton, “Low-cost 

Nanocrystalline Materials for Structural Components,” presented at the SAE Inner Value Symposium—Nanotechnology for the Automotive Sector, 

invited talk, Carnegie Mellon University, Pittsburgh, 

May 2004. 

S. Chandrasekar, “Emerging Manufacturing 

Technologies,” presented at the State of Indiana Advanced Manufacturing Summit, Purdue University, 

West Lafayette, IN, May 2004. 

S. Chandrasekar, W. D. Compton, K. P. Trumble, and A. H. King, “Large-scale Manufacturing of 

Nanocrystalline Materials,” Indian Institute of Technology, Madras, India, July 2004. 

S. Chandrasekar, W. D. Compton, K. P. Trumble, and A. H. King, “Large-scale Manufacturing of 

Nanocrystalline Materials,” presented at the Timken 

Global Engineering Center, Bangalore, India, July 

2004. 

Patents

S. Chandrasekar, W. D. Compton, T. N. Farris, 

and K. P. Trumble, “Method of Forming Nanocrystalline Structures and Product Formed Thereof,” 

U.S. Patent, 6,706,324, March 16, 2004. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

171

R. Deformation Processes for Next-Generation Ceramics 

P. F. Becher 

Oak Ridge National Laboratory 

P.O. Box 2008, MS-6068 

Oak Ridge, Tennessee 37831-6068 

(865) 574-5157; fax: (865) 574-6098; e-mail: becherpf@ornl.gov 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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 

• Assess the plastic deformation behavior in nanocrystalline oxide ceramics and its applicability to fabrication of 

complex shapes. 

• Characterize the ability to deform nanocrystalline ceramics at temperatures well below 1000°C through control 

of microstructure and external parameters (e.g., electric fields, environment). 

• Characterize thermal and mechanical properties to assess the 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 

to develop advanced colloidal technology-based processing to fabricate monolithic bodies and sintering technology to fabricate dense monolithic nanocrystalline bodies (i.e., grain sizes <100 nm). 

• Conduct both compressive creep and plastic deformation studies as a function of temperature, environment, and 

applied electric fields on nanocrystalline oxides with grain sizes of <100 nm. 

Accomplishments 

• Established a novel pressureless sintering method to fabricate dense monolithic zirconia bodies with selected 

grain sizes in the range of 20 to 90 nm. 

• Developed data on the deformation of dense zirconia ceramics with a grain size of 350 nm to serve as a baseline 

for comparison of future observations. Initial observations have been obtained on the enhancement of deformation of this dense zirconia ceramic with sub-micron-sized grains from applied electric fields through collaboration with North Carolina State University researchers. 

Future Direction 

• Exploit the ability to sinter monolithic nanocrystalline oxides to full density to characterize the enhanced plastic 

deformation below 1000°C in zirconia ceramics having grain sizes below 100 nm. The goal is to explore the 

minimum temperature at which permanent deformation can still be achieved. 

• Evaluate conditions required for optimizing deformation-shape forming. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

172

• Examine room temperature fracture and wear properties to seek approaches for developing ceramics with 

greater wear resistance. 

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 shape 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. Complexshape-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 hotworked at temperatures below 1000°C not only to 

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 compared with ceramics. 

In ceramic systems, processing to tailor microstructures needs to be combined with novel concepts 

that offer the potential to enhance deformation. One 

such concept that has been of considerable interest is 

superplastic deformation of ceramics with submicron-size grains at temperatures above 1000°C.1

Recent studies suggest that reducing the microstructure to achieve grain sizes of <100 nanometers could 

lead to permanent/plastic deformation in ceramics at 

quite modest temperatures.2,3 Indeed, classical deformation behavior suggests significant decreases in 

yield stresses at a fixed temperature or decreases in 

temperature at which deformation can be induced 

via the reduction in grain size.4

 Typically, temperature, strain rate or stress, and grain size are the parameters to alter the deformation or creep rates in 

single-phase oxides. Recent studies have shown that 

the application of electric fields can reduce the yield 

stress or increase the creep rates of ceramic systems.5

 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 ceramic monoliths. 

Approach

This project involves two major thrusts. The 

first is a collaborative effort in the processing and 

sintering of nanocrystalline ceramics with researchers in the Synthesis of Nanocrystalline Ceramics 

task at Pennsylvania State University. The goal is to 

develop the technology to form “green” (unfired) 

bodies of nanocrystalline particles with densities 

well in excess of 50% of the theoretical density, 

which will allow us to devise pressureless sintering 

profiles to produce dense monolithic bodies with 

grain sizes of <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 

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-size materials include 

the dispersion of the nanosize particles and control 

of gas evolution from the very-high-surface-area 

compacts during sintering. Typically, nanometersize particulates suffer from a high degree of agglomeration, which results in poor particle packing 

(green densities < 40% of theoretical); this 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, it 

should be possible to produce green bodies with 

densities of more than 60% of the theoretical by 

control of interparticle forces. This control is 

achieved through appropriate chemical additives to 

suspensions of nanoparticulates, combined with filter pressing or pressure casting of the high-solids-
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

173

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 samples with sub–100-nm grain size.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 allow undesired 

grain growth to proceed.10 Heating to a temperature 

sufficient to achieve densities of ≥75% of theoretical, and then rapidly lowering the temperature to 

obtain sufficient boundary diffusion to remove pores 

trapped on boundaries while minimizing grain 

growth, has been only partially successful. Long 

hold times at the lower temperature were still required.11 Here we seek to optimize the processing of 

the green bodies and use modified sintering cycles to 

achieve dense oxides with sub–100-nm grain sizes. 

Deformation Behavior in Nanostructured 

Materials

To evaluate the influence of nanocrystalline microstructures, baseline deformation data will be established through studies of high-purity, dense 

zirconia ceramics with sub-micron-size grains. 

These studies will then be extended to ceramics with 

<100 nm grain size. This will include 

• nanostructured ceramics, to explore for the softening effects observed in nanocrystalline metals 

by evaluating both (1) the creep response over a 

range of temperatures and applied stresses and 

(2) the yield stress as a function of strain rate 

and temperature 

• 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 phenomena 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 

Advanced colloidal processing in the task on 

Synthesis of Nanocrystalline Ceramics has provided 

monolithic green bodies of both yttria-doped and 

undoped zirconias with densities from 55 to 68 % of 

theoretical and an average particulate size of 15 nm. 

Advanced sintering cycles now allow the production 

of dense monolithic zirconias with grain sizes well 

below 100 nm (Figure 1). As noted, sintering in a 

vacuum promotes densification at lower temperatures than sintering in air. At this point, we assume 

this is due to lower capillary pressure within pores 

and faster removal of any residual organics or moisture used in the colloidal processing. 

Figure 1. Pressureless sintering of colloidally processed 

zirconias yields dense ceramics with grain sizes < 100 nm 

(density—closed symbols; open symbols—grain size). 

Based on the data shown in Figure 1, it is obvious 

that the addition of yttria inhibits the densification of 

nanocrystalline zirconia. Recent work indicates that 

space charge effects due to yttrium segregation may 

inhibit Zr+4 lattice diffusion,12,13 which would explain the lower-than-expected creep deformation 

observed in yttria-doped zirconia with grains of sub–

100-nm size (Figure 2).14 While grain boundary diffusion is a critical factor in the removal of pores 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

174

Figure 2. Microstructure of ZrO2—1.7 mol % Y2O3 ceramic sintered in air at 1150°C. 

during the densification of nanoparticulate bodies, 

lattice diffusion might be a limiting parameter in 

sintering. 

Deformation Studies 

In the first year of this project, a fully dense, 

hot-isostatically-pressed zirconia—2 mol % yttria 

ceramic with a grain size of ~ 350 nanometers (Figure 3)—was selected for studies of compressive deformation. These include compressive creep and 

constant strain rate deformation studies conducted in 

air at temperatures above 1000°C to obtain measurable deformation. These results will serve as a baseline for assessing the effectiveness of increasing the 

plasticity by producing microstructures of sub–100-

nm grain size. In addition, data on the tetragonal 

zirconia ceramic with 350-nm grain size can be used 

to correlate with data in the literature on zirconias 

with sub-micron grain size. 

The creep behavior illustrates the creep rate dependence on applied stress above 1100°C with stress 

exponents of 1.7 to 1.9, while at 1100°C the stress 

exponent is ~ 0.6 (Figure 4a). The activation energy 

ranges from 625 to 645 kJ/mol with a discontinuity 

at 1100°C with a stress of 30 MPa (Figure 4b). The 

stress exponent and the activation energy values are 

similar to those reported in the literature for highpurity zirconias with grains of sub-micron size;14

however, the test temperatures were typically greater 

than 1200°C. The extent of permanent deformation 

was limited by the displacement range of the contact 

extensometers, which require repositioning for displacements of 150 to 200 microns, or by the length 

Figure 3. Fracture surface microstructure of 350 nm grain 

sized tetragonal zirconia—2 mol. % yttria ceramic used to 

establish baseline data on deformation behavior. 

Figure 4. Applied stress (a) and temperature (b) dependence of the compressive creep rates of submicron grain 

sized ZrO2–2 mol% Y2O3 in air. 

of the test in the case of those conducted at 1100°C. 

However, permanent strains of ≥2 % were obtained 

based on measurements prior to and after testing.

Initial assessments of the compressive yield 

stresses of this sub-micron-grain-size zirconia were 

obtained in air. Data at 1300°C reveal a typical decrease in the yield stress as the stress rate is reduced 

(Figure 5a). At the same time, the yield stress increased significantly when the temperature was reduced to 1200°C for comparable stress rates 

(Figure 5b). Thus permanent strains under dynamic 

loading in air were limited to <2 %. On the other 

hand, observations reveal that extensive plastic 

strain (~ 45%) could be achieved at 1200°C with the 

application of a dc field of 1 kV/cm in this submicron-grain-size zirconia. In view of the recent observations of plasticity in nanocrystalline zirconia 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

175

Figure 5. Compressive yielding of 350-nm-grain sized 

ZrO2; 2 mol % Y2O3 exhibits typical decrease in yield 

stress due to reduction in stressing rate at 1300°C in air 

(a) and an increase with decrease in test temperature (b). 

ceramics at 1000°C,3

 there would appear to be significant potential for enhanced plasticity in 

nanocrystalline zirconias. The demonstrated ability 

in the current task to fabricate dense monolithic 

nanocrystalline zirconia ceramics will now allow us 

to evaluate the plastic deformation behavior of 

doped and undoped zirconia ceramics over a range 

of grain sizes (i.e., 30 to 100 nm). 

Conclusions

Advances in the processing of monolithic 

nanoparticulate zirconia bodies with green densities 

in excess of 55% of theoretical, through our collaboration with the task on Synthesis of Nanocrystalline 

Ceramics, now allow us to fabricate dense monolithic bodies of both yttrium-doped and undoped 

zirconias with grain sizes well below 100 nm by 

pressureless sintering. This is a substantial innovation and provides the critical path forward to assessing the plastic deformation behavior of nanocrystalline oxides, as well as other properties (e.g., wear, 

strength) that would prove advantageous for improvements in heavy vehicle component fabrication 

and performance. 

Baseline data on the deformation response of 

zirconia ceramics are being established using a fully 

dense 350-μm grain-size tetragonal zirconia containing 2 mol % yttria. Results show that this ceramic 

can be plastically deformed in compression to permanent strains in excess of 2% at stresses below 100 

MPa and temperatures of 1200 to 1300°C in air. 

Current work is focusing on the preparation of compression samples with sub–100 nm grain size to begin evaluating the deformation behavior of 

nanocrystalline (<100 nm) ceramics. 

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). 

4. M. F. Ashby and R. A. Verrall, “DiffusionalAccommodated Flow and Superplasticity,” Acta 

Metall. 21, 149–163 (1973). 

5. H. Conrad, “Space Charge and the Dependence of the Flow Stress of Ceramics on an Applied 

Electric Field,” Scripta Mater. 44, 211–16 (2001). 

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) 2649–2658 

(1992). 

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 (20002). 

11. I. W. Chen and X. –H. Wang, “Sintering 

Dense Nanocrystalline Ceramics Without FinalStage Grain Growth,” Nature 404, 168–71 (2000). 

12. D. Gomez-Garcia, C. Lorenzo-Martin, A. 

Munoz-Bernabe, and A. Dominguez-Rodriguez, 

“Model of High-Temperature Plastic Deformation of 

Nanocrystalline Materials: Application to Yttria 

Tetragonal Zirconia,” Phys. Rev. B67(14), 144101-1 

to –8 (2003). 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

176

13. A. Dominguez-Rodriguez, F. GutierrezMora, M. Jimenez-Melendo, J. L. Routbort, and R. 

Chaim, “Current Understanding of Superplastic Deformation of Y-TZP and Its Application to Joining,” 

Mater. Sci. and Eng. 302, 154–61 (2001). 

14. M. Jimenez-Melendo and A. DominguezRodriguez, “Like-Metal Plasticity of Fine-Grained 

Y2O3-Stabilized Zirconia Ceramics,” Phil. Mag. A 

79(7) 1591–1608 (1999). 

Publications

C. H. Hsueh and P. F. Becher, “Effects of 

Dopant Segregation on Lattice-Diffusional Creep of 

Nanocrystalline Ceramics,” Phys. Rev. B, accepted 

for publication. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

177 

S. Synthesis of Nanocrystalline Ceramics 

Dr. James H. Adair, Principal Investigator 

Christopher J. Szepesi, Graduate Assistant 

The Particulate Materials Center, The Pennsylvania State University 

 249 Materials Research Laboratory 

University Park, PA 16802 

(814) 863-6047; fax (814) 863-9704; e-mail: jha3@psu.edu 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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: The Pennsylvania State University, University Park, Pennsylvania

Objectives 

• Synthesize well-dispersed, concentrated aqueous suspensions of nominally 8–10 nm tetragonal ZrO2. 

• Form nano-grained green bodies from concentrated ZrO2 suspensions. 

• Dry nano-grained green bodies while avoiding crack formation. 

• Sinter nano-grained green bodies to high theoretical density while maintaining nano-scale grain structure. 

Approach 

• Complex zirconyl nitrate with bicine in aqueous solution, which is then hydrothermally treated to precipitate 8–

10 nm tetragonal ZrO2. 

• Increase ZrO2 yield by increasing reagent concentrations. 

• Form green bodies using constant-rate filter pressing. 

• Dry green bodies under controlled humidity conditions or through a solvent exchange procedure. 

• Fire green bodies at Oak Ridge National Laboratory using transient sintering cycles.1

 

Accomplishments 

• Hydrothermally synthesized 10–20 nm tetragonal ZrO2 using a procedure developed by R.A. Kimel and J.H. 

Adair. 

• Increased ZrO2 yield from about 25 g/L to 100 g/L by increasing reagent concentrations and modifying the 

dispersant concentration to accommodate the higher surface areas in solution. 

• Routinely made nano-grained green bodies 3⁄4 in. in diameter and 1–4 mm thick using constant-rate filter 

pressing. 

• Successfully dried nano-grained green bodies using controlled humidity drying. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

178

• Successfully dried nano-grained green bodies by exchanging entrained water with acetone containing an 

acryloid binder. 

• Sintered nano-grained green bodies to 100% theoretical density while maintaining grain sizes below 100 nm. 

Future Direction 

• Continue to produce nano-grained green bodies with constant-rate filter pressing. 

• Compare constant-rate filter pressing to constant-pressure filter pressing. 

• Optimize solvent exchange drying procedure to optimize green densities. 

• Produce larger-diameter green bodies for sintering and mechanical testing. 

Introduction 

The main objective of this research was to 

manufacture bulk, dense ceramics with grain sizes of 

less than 100 nm. The material chosen for this 

purpose was ZrO2 and Y2O3-stabilized ZrO2 because 

of its mechanical properties and its ease of 

production in the appropriate size scale. 

Previously, nanoscale ZrO2 and Y2O3-stabilized 

ZrO2 have been synthesized using a protocol 

developed by Kimel and Adair 2–4 in which zirconyl 

nitrate (and yttrium nitrate) is complexed with bicine 

in aqueous solution through alkaline pH adjustment. 

Hydrothermal treatment at 200°C for 8 hours 

allowed particles of the tetragonal phase with sizes 

of between 8 and 15 nm to precipitate. The product 

suspension was then washed with an aqueous 

solution of oxalic acid to remove any remaining 

bicine and disperse the suspended particles. 

However, solid yield was only about 25 g/L using 

this method. 

By increasing the concentration of reagents, 

yield was increased to about 50 g/L. The limiting 

factor in this protocol became the solubility of both 

bicine and the organic base used to adjust pH 

(tetraethylammonium hydroxide, TEAOH). To 

address this issue, dry zirconyl nitrate and bicine 

were mixed with an appropriate amount of solid 

organic base, tetramethylammonium hydroxide 

pentahydrate (TMAOH). Because of the water 

content of both the zirconyl nitrate and TMAOH, the 

mixture became an aqueous solution, which was 

hydrothermally treated to produce 10–30 nm ZrO2. 

Using this method, ZrO2 yield was increased to 

100 g/L. 

Approach

The current task in this research has been the 

production of green bodies for sintering studies. This 

was accomplished with constant rate filter pressing, 

in which plungers moving at a constant rate apply a 

load to a concentrated suspension. Solution is then 

removed through filters at the bottom of the solution 

chamber (see Figure 1). 

Once the wet pellets are removed from the filter 

press die, they must be carefully dried to avoid 

cracking within the body. This can be done either 

with controlled humidity or by exchanging water 

within the pellet with another liquid of lower surface 

tension (solvent exchange). 

Results

With the current filter pressing protocol, pellet 

production has been increased to about four pellets 

per 4 hours (from four pellets per 12–18 hours). 

Humidity drying, the previously preferred method, 

has been limited as a result of difficulties with 

equipment. Therefore, solvent exchange drying has 

been investigated. 

Initially, four solvents were chosen: 

methylethylketone (MEK), 40% MEK–60% ethanol 

(by weight), acetone, and isopropanol. Each of these 

solvents was pushed through a filter-pressed pellet 

immediately after pressing concluded. Upon 

removal from the die, the solvent-exchanged pellets 

were all cracked but dry. Density measurements by 

the Archimedes method revealed a marked increase 

in density (see Table 1). However, the dry pellets 

maintained little green strength. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

179 

Table 1. Archimedes density measurement results of 

solvent-exchange–dried ZrO2 pellets 

Solvent 

Density 

(g/cm3) 

Theoretical density

(%) 

Methylethylketone 

(MEK) 

3.99 66 

40% 

Methylethylketone – 

60% Ethanol (by 

weight) 

4.01 68 

Isopropanol 3.07 51 

Acetone 3.61 60 

It was then proposed that a binder be added to 

the solvent to add green strength to the dry body. An 

acetone-soluble acryloid binder (Rohm and Haas), 

which is typically used in dry pressing, was chosen. 

When the acetone-binder solution was pressed 

through wet filter-pressed pellets, cracks were still 

present, but to a smaller degree. Cracking was 

reduced further by introducing the acetone-binder 

solution to pellets that were not fully consolidated. 

Archimedes density determination on fragments of 

these pellets reveals, again, that pellets are achieving 

a high density after drying with the solvent 

exchange process (Table 2). Also, the high densities 

achieved with acryloid binder present require 

recalculation once the actual concentration of binder 

is determined by controlled thermolysis. It is likely 

that the very large values (in some cases greater than 

70% theoretical density) will be in the 60% 

theoretical density range when the concentration of 

binder is known and accounted for in the density 

calculation. Whether cracking can be eliminated 

completely by beginning the solvent exchange 

drying process earlier in the filter pressing procedure 

is currently under investigation. 

Conclusions

The research focus has centered recently on 

increasing yield during synthesis and producing 

green bodies for sintering studies. By modifying the 

previous synthesis procedure, ZrO2 yield was 

increased fourfold. Pellet pressing times have also 

been decreased significantly. Solvent exchange 

drying was demonstrated to be an effective method 

of drying filter-pressed green bodies to a high 

density. Efforts are continuing to eliminate cracks 

after drying with the solvent exchange procedure. 

Table 2. Archimedes density measurement results of solventexchange–dried ZrO2 pellets (dried with a solution of 

acetone and an acryloid binder) 

Sample 

number Binder concentration Density 

(g/cm3)* 

Theoretical 

density* (%)

a3 0% 4.32 72 

a1 2% (relative to 

weight of solids) 4.48 74 

a2 3% (relative to 

weight of solids) 4.17 69 

b1 0% 

4.71 (very 

small 

sample) 

78 

b2 1% (relative to 

weight of solution) 4.65 77 

b3 2% (relative to 

weight of solution) 4.45 74 

b4 5% (relative to 

weight of solution) 4.13 69 

c1 1% (relative to 

weight of solution) 4.30 71 

c2 2% (relative to 

weight of solution) 4.48 75 

c3 5% (relative to 

weight of solution) 3.95 66 

References

1. I. W. Chen and X. H. Wang, “Sintering Dense 

Nanocrystalline Ceramics without Final-Stage Grain 

Growth,” Nature 404, 168–171 (2000). 

2. R. K. Kimel, Aqueous Synthesis and 

Processing of Nanosized Yttria Tetragonally 

Stabilized Zirconia, Ph.D. Thesis, The Pennsylvania 

State University, August 2002. 

3. R. K. Kimel and J. H. Adair, “Aqueous 

Synthesis of Well-Dispersed Less than 10 nm Yttira 

Tetragonally Stabilized Zirconia at 200°C by 

Precipitation from Homogeneous Solution Using 

Complexation Chemistry,” submitted to J. Am. 

Ceram. Soc. 

4. R. K. Kimel and J. H. Adair, “Recovery and 

Dispersion of 8 nm Y-TZP Utilizing a Nanoscale 

Analog to the Protective Colloid Approach,” 

submitted to J. Am. Ceram. Soc. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

180
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

181

T. Development of Titanium Component Applications in Heavy Duty 

Diesel Engines 

Jesus Chapa-Cabrera, Principal Investigator 

John Grassi, Co-Investigator 

Caterpillar Inc. 

Tech Center 854-80 TC-E 

14009 Old Galena Rd. 

Mossville IL, 61552 

(309) 578-8347; fax: (309) 578-2953; e-mail:chapa-cabrera_jesus_g@cat.com

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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: Caterpillar Inc., Peoria, Illinois 

Objectives 

• Design and fabricate a cost-competitive diesel engine turbocharger using lightweight titanium materials that reduces both fuel consumption and transient emissions. 

Approach 

• Design a series turbocharger for use on multiple engine platforms. The turbocharger consists of one turbo wheel 

and two compressor wheels that are attached to a single rotating drive shaft. 

• Use this compact design to simplify complex multi-turbocharger systems presently installed in different engine 

platforms. 

• Use titanium aluminide (TiAl) for the turbine wheel and a titanium alloy for one of the compressor wheels. 

Accomplishments 

• Procured cast TiAl turbine wheels and test materials from two suppliers. 

• Finalized subcontracts with two vendors. 

• Completed three iterations of simulation and friction welding experiments to produce turbine-to-shaft joints. 

Produced welds with higher joint strength than the base TiAl turbine wheel. 

• Established a nondestructive inspection methodology to assess robustness of the joints. 

• Conducted preliminary turbocharger bench tests using base nickel alloy and TiAl turbines. 

Future Direction 

• Conduct a complete redesign of the wheel to increase the overall strength and create a robust joint. 

• Complete detailed material characterization of two TiAl alloys processed by counter-gravity and gravity casting. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials

182

• Perform turbocharger bench tests and engine tests. 

• Assess the commercial viability of using titanium alloys for turbocharger applications for heavy-duty diesel engines. 

Introduction

Turbochargers on diesel engines play an integral 

role in meeting emission regulations and controlling 

fuel economy. A recently designed series turbocharger by Caterpillar will use TiAl and titanium 

alloys for the turbo wheel and compressor wheel, 

respectively. The dual turbocharger has two complete systems that are connected in series, as shown 

in an ACERT® engine in Figure 1, while the proposed turbocharger design is one system with two 

compressor wheels attached to a single drive shaft. 

Lightweight titanium materials are important technologies for this turbocharger design and are anticipated to provide faster response time (reduced 

transient emissions) and better fuel economy. 

Figure 1. ACERT® engine. The two turbochargers in series are circled.

Approach

Caterpillar has selected candidate TiAl alloys for 

a series of characterization tests and development of 

a joining process. These tests will examine mechanical and physical properties of selected TiAl alloys. 

Friction weld trials will determine the optimal parameters for joining the TiAl wheel to the drive 

shaft. Based on the turbocharger validation tests, 

modifications to the turbo wheel will be modeled 

and the performance will be assessed. Bench tests 

will be completed before engine tests are scheduled. 

Results

Two of the three parties that were initially identified as potential subcontractors for the project have 

completed contract agreements with Caterpillar. After further consideration, it was determined that the 

third vendor would participate in the program solely 

as a materials supplier. 

Finite element modeling of the existing HEAT®

turbocharger design was completed. The objective 

of this analysis was to evaluate the critical stresses 

that develop during an extreme operating cycle and 

compare the results for TiAl with the nickel-based 

alloy Inconel 713, which is typically used in similar 

applications. The modeled cycle consisted of going 

from a cold start to peak torque, followed by low 

idle conditions. Identical operating conditions and 

wheel design were used in both analyses; thus the 

different results can be attributed only to differences 

between the two materials. The code ABAQUS®

was used for the analysis, using loads and boundary 

conditions obtained from aerodynamic simulations. 

Only an angular portion of the wheel and shaft was 

modeled, owing to symmetry, as shown by the mesh 

in Figure 2. This wedge consists of a blade, a section 

of the wheel hub, and a section of the shaft. Temperatures and stresses were obtained from the analyses, and the results showed that transient peak 

stresses on the wheel are reduced by approximately 

50% by replacing the Inconel 713 with TiAl. Critical 

stresses result from superimposing centrifugal and 

thermal effects, and the significant reduction in operating stresses corresponds to the lower density and 

higher thermal diffusivity of TiAl relative to Inconel 713. 

Thirty TiAl wheels produced by the countergravity casting method were procured, in addition to 

more than 30 wheels produced by a second supplier 

using a gravity investment casting process. Analysis 

of the castings produced with these two techniques 

is ongoing to determine the optimum and most costefficient process for the application. Test bars for 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

183

Figure 2. Mesh of an angular section of the wheel 

modeled in ABAQUS®. Using TiAl reduces the 

operating stresses by approximately 50% relative to 

Inconel 713.

welding experiments and material characterization 

were procured in addition to the wheel castings. 

Test bars were machined from the wheel and 

shaft materials. The surfaces of the TiAl wheels 

were conditioned for friction welding trials. The 

cross section of the TiAl bars used for welding experiments was identical to that of the wheels. Joints 

were produced between the TiAl and the shaft material using bar-to-bar and shaft-to-wheel configurations. The friction welding process parameters were 

varied to produce the joints, and the resulting bar-tobar jointed specimens were tested to failure in bending. A combination of parameters was established 

that produced bar-to-bar joints with a strength higher 

than that of the TiAl base material. Friction welding 

experiments showed a large dependency of the quality of the weld on the different process parameters. 

In addition, mass changes were shown to have a significant impact on the strength of the welds. This 

effect was evident after TiAl bars were welded to 

bars of the shaft material, producing joints with 

strength higher than that of the base TiAl. The conditions that produced strong bar-to-bar welds were 

used to make two prototype shafted wheels. Weak, 

flawed joints with a strength of approximately 50 

MPa resulted; this was attributed to the higher heat 

dissipation from the weld area due to the increase in 

thermal inertia and resulting sharper temperature 

gradients. 

Simulation was conducted to increase understanding of the sensitivity of the joint strength to the 

different process parameters. Modeling of the friction welding process was conducted using commercial simulation tools and customized user 

subroutines. Analyses were conducted using systematic combinations of rotational speed, frictional and 

forge loads, and frictional time and evaluating the 

effect of these factors on the resulting temperatures 

and thermal stresses. A proportional relation between the severity of the temperature gradients and 

the thermal inertia was observed. A higher thermal 

inertia increases the severity of the temperature gradients, which leads to higher thermal stresses and 

cracking. This correlates with cracking observed on 

the heat-affected zone of the shafted wheels. 

Simulation results were used to define improved 

processing conditions for an additional round of 

welding experiments; and 16 additional joints, including bar-to-bar and shaft-to-wheel joints, were 

produced. The jointed specimens were heat-treated 

to relieve stresses, machined to remove the flash, 

and subjected to non-destructive evaluation (NDE). 

After NDE, the specimens were tested to failure in 

tension, and the tensile strengths were recorded. 

Subsequently, fractography was conducted to determine the nature of the failures and to correlate with 

features observed with NDE. The strength of the 

joints was increased in this round of tests, producing 

shafted wheels that, when tested in tension, failed 

away from the weld on the base TiAl material. 

However, the strength recorded was lower than expected for TiAl because of stress concentration. A 

parametric study is being conducted to modify the 

back face geometry of the turbine, with the purpose 

of increasing the overall strength of the component 

by reducing stress concentrations and thermal 

stresses. The results of the parametric study will be 

used for another iteration of welding experiments 

with the objective of completing the development of 

the joint. 

Two different NDE techniques were used to 

identify subsurface defects on the weld and the heataffected zone. A specimen with a poor weld was 

selected for initial evaluation to establish a baseline 

by identifying and characterizing features that are 

not desirable in sound welds, such as incomplete 

bonding and radial cracking. The first technique employed was eddy current, using a setup for automatic 

scanning by attaching the specimen to a turntable 

and controlling the vertical position of the probe. A 

prototype TiAl wheel shaft undergoing eddy current 

inspection is shown in Figure 3. The second 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials

184

Figure 3. Shafted TiAl turbine wheel during NDE inspection.

technique was immersed ultrasound inspection, setting the specimen up in a similar automated configuration. Ultrasound inspection, with its finer 

resolution, showed not only the features observed 

with the eddy current but also smaller cracks. Features observed with both techniques were correlated 

with features observed on fractographic analysis, 

positively identifying defects on the weld plane as 

well as radial cracks. 

A prototype HEAT® turbocharger was built using an Inconel 713 turbo wheel and tested on an internally designed Caterpillar gas stand as well as on 

an engine. The purpose of these prototype tests was 

to validate the compression capacity of the HEAT®

turbocharger design and to obtain baseline data for 

comparison with a HEAT® turbocharger to be built 

using a TiAl turbo wheel. Steady and transient test 

results indicate that the HEAT® turbocharger likely 

has the potential to meet the performance and emission requirements of the engine after redesign of the 

wheel. 

The decision was made to use one of the prototype TiAl wheel shafts produced during the second 

round of experiments to assemble a HEAT® turbocharger unit and to obtain all performance information possible in gas stand testing, knowing of the 

limited strength of its weld. This test was intended 

as a proof of concept and in no way intended to be a 

reliability or durability checkpoint. This HEAT®

turbocharger unit was tested in the gas bench, and 

data were collected at two speed lines above 

60,000 RPM. The weld failed at 330°C after 9 hours 

of operation and immediately following an engine 

surge. After failure, the loose wheel continued rotating inside the housing, impacting the leading edges 

with the housing, and four blade tips fractured. The 

wheel did not burst, even though it failed below its 

brittle-to-ductile transition temperature. Post-test 

inspection of the wheel revealed that the damage on 

the TiAl wheel blades was similar to that observed 

on Inconel wheels after a weld failure, even with the 

severity of impacts it received. These results 

strengthened the position of the entire team that TiAl 

has great potential for this application. Joints with 

strengths close to three times the strength of the 

specimen tested in the gas stand have been produced 

after iterations of experiments and simulation. Further gas stand testing will be conducted after fully 

optimizing these joints. 

Joined by a representative from a turbocharger 

manufacturer, the Advanced Materials Technology 

and Air Systems team conducted on-site initial assessments of the TiAl casting processes of two suppliers. Test materials, including wheels and bars, 

have been procured from both suppliers and are under evaluation. 

Conclusions

Titanium aluminide is an important material for 

less complex and more efficient air systems for Caterpillar engines. The lightweight nature of the material reduces the inertia of rotating components, 

enabling more compact and reliable turbocharger 

designs that will potentially lead to increased fuel 

efficiency and reduced transient emissions. An additional benefit of TiAl, demonstrated with finite element analyses, is the reduction of stresses during 

operation by 50% relative to Inconel 713, as a result 

of lower centrifugal loads and higher thermal diffusivity. 

Initial assessment of two foundries was completed, and TiAl wheel castings and test bars were 

procured from both suppliers. Materials obtained 

from these suppliers have been used to develop a 

process to join the TiAl wheel to shaft materials. An 

iterative approach for this development has been 

applied, conducting welding experiments and simulation. Optimization of the geometry of the wheel 

and shaft and of friction welding parameters is ongoing to increase the overall strength of the component. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

185

A methodology for conducting NDE of the 

wheel-shaft joints has been developed using eddy 

current and automated immersed ultrasound. Features observed on fractographic examination of 

specimens tested in tension have been positively 

correlated to features observed with the NDE techniques used. 

A shafted TiAl turbine wheel was assembled 

into a turbocharger and tested on a gas stand rig. 

This test was intended as a proof of concept and not 

as a reliability or durability checkpoint. Damage on 

the wheel was similar to that observed on failures of 

Inconel 713 wheels. These results strengthened the 

position of the team that TiAl has great potential for 

this application. Joints with strength close to three 

times the strength of the specimen tested in the gas 

stand have already been produced after iterations of 

experiments and simulation. Further gas stand testing will be conducted after fully optimizing these 

joints. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials

186
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

187

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: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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 

• Characterize the rolling contact fatigue (RCF) performance of ceramics and tribological coatings that are under 

consideration for use in rolling element components (e.g., hybrid bearings) 

• Determine the effects of subsurface damage, properties such as (static and dynamic) hardness and fracture 

toughness, and the stress state from Hertzian contact on RCF performance. 

• Model hypothetical microstructures that will promote 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. Correlate test results. 

• Vary machining-induced subsurface damage in a silicon nitride (Si3N4) and correlate that to measured RCF performance. 

• Provide new understanding of RCF performance/microstructure relationship to ceramic manufacturers that will 

serve to promote microstructure engineering and improved RCF performance. 

Accomplishments 

• Established 3BOR RCF facility. 

• Initiated formal collaboration with Bournemouth University (BU) in the United Kingdom to relate performances measured with 3BOR and three-ball-on-ball (3BOB) RCF test techniques. 

• Defined machining conditions and designed and procured test coupons. Initiated testing. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials

188

Future Direction 

• Conduct domestic 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. 

• Conduct international interlaboratory study to compare performances measured with different RCF test 

methods. 

Introduction

The use of ceramics in rolling elements shows 

many practical advantages over traditional hard and 

bearing-grade steels. Current demands on loadbearing contacts in all kinds of machinery are leading to developments aimed at running them at high 

speeds with minimum vibration, hostile environments, increased unit loads, and restricted lubrication. The design and manufacture of such contacts 

are at the limit of established technology; hence the 

need for ceramic contacts. 

Si3N4 is one class of ceramic that has combined 

properties that are most suited for rolling contact 

conditions. Experimental programs with Si3N4 contacts using full-scale and bench tests have produced 

some understanding of RCF modes failure and durability; however, the effects of grinding or finishing 

rate on the subsurface integrity and hence RCF durability and failure modes are poorly understood. 

To address these issues, this project is characterizing 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, cam roller followers) (see Figure 1). 

There are several different test methods to measure 

RCF performance, so this project also works toward 

reconciling their measured differences so that their 

results can be validly pooled. The study of the effects of subsurface damage, properties such as (static 

and dynamic) hardness and fracture toughness, and 

the stress state from Hertzian contact on RCF performance is a natural outcome. Based on those 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—

measure RCF performance and reconcile results that 

Figure 1. Schematic of a hybrid bearing system that 

includes silicon nitride balls. Source: 

www.cerbec.com.

 

were generated using different RCF test techniques; 

examine the effects that subsurface damage and microstructure have on RCF performance; and iteratively work with ceramic manufacturers and communicate RCF results and interpretations those 

manufacturers can then use to develop lower-cost 

Si3N4 compositions, improve Si3N4 machinability, 

improve RCF performance, or combinations thereof. 

A multi-task, 3-year project involving BU and 

the Oak Ridge National Laboratory (ORNL) was 

initiated in the last quarter of FY 2004; this collaboration will serve as a key component of the project. 

A BU Ph.D. student will be a primary performer of 

these tasks and will alternate on-site research at BU 

and ORNL over the life of the project, spending 

about half of the time at each institution. This student will work under the supervision of Prof. Mark 

Hadfield at BU and the project principal investigator 

at ORNL. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

189

RCF performance of Si3N4 will be interrogated 

using 3BOB and 3BOR (NTN–Bower) RCF test 

methods conducted at BU and ORNL, respectively; 

and RCF performance measured with each technique 

will be linked. Additionally, the generated subsurface damage in cylindrical rods and spheres machined under the “same” machining conditions will 

be compared to assess/verify the sought equivalence. 

Ceradyne’s Ceralloy 147-31N will serve as the 

model Si3N4. A critical element of this project will 

be the identification, characterization, and interpretation of subsurface damage; these will be accomplished through a variety of inspection means and 

test techniques. Rotating flexure testing, optical and 

scanning electron microscope fractographies, dye 

penetration, optical coherent tomography, scanning 

acoustic microscopy, and residual stress analysis 

will be used. Machining damage in Ceralloy 147-

31N, which was systematically characterized and 

exploited by uniaxial flexure testing in a recent 

study,1,2 will provide a useful guide to RCF sample 

preparation. Indentation and scratch testing and finite element analysis (FEA) will also be used to 

supplement the characterization and interpretation of 

results. Finally, the student, project principal investigator, and Hadfield will integrate the results and 

interpret the effect of subsurface damage on RCF in 

Si3N4. 

Results

RCF performance is being interrogated as a 

function of pre-existing subsurface damage in Ceralloy 147-31N Si3N4. Specimens were ground in one 

of three ways: 100-grit roughing and 600-grit finishing; 180-grit roughing and 600-grit finishing; and 

the conventional method for machining RCF specimens (1200-grit finish). Table 1 provides additional 

details of the machining. 

The effects of subsurface damage in monolithic 

ceramics are critical to RCF performance, but that 

relationship has largely been taken for granted and is 

little understood. To help resolve the issue, numerous test coupons were prepared or designed to facilitate the quantification of a critical independent parameter in this project: depth of machining damage. 

Ceralloy 147-31N RCF (Figure 2), half-RCF (Figure 

3), rotary bend strength (RBS) (Figure 4), half-RBS 

(Figure 5), and ASTM C1161B test specimens have 

been prepared and are undergoing testing. All five of 

those specimen geometries were machined according to the conditions outlined in Table 1. The RBS, 

half-RCF, half-RBS, and ASTM C1161B (uniaxial 

4-point-bend flexure) specimens are being tested in 

parallel with the RCF specimens because flexure 

tests are efficient at quantifying the depth of subsurface damage (via companion fractography) and at 

exploiting the effects of sub-surface damage (on 

strength in this case). The RBS specimen rotates 

during its monotonic loading to fracture, so its entire 

gage section is subjected to outer-fiber tensile 

stresses. This will enable the study of strength-size 

scaling effects. The half-RCF and half-RBS specimens were strength-tested in four-point-flexure 

(Figure 6). The strength calculation for the half-RCF 

(i.e., half-cylinder) geometry is straightforward; 

however, the calculation of fracture stress for the 

half-RBS geometry from the fracture load is not. 

FEA of the half-RBS specimen (Figure 7) was used, 

and failure load was linked to maximum outer fiber 

tensile stress in the specimen’s gage section. 

Weibull strength statistics for the half-RBS specimen are listed in Table 2. Strengths among the three 

sets were not appreciably different. 

Table 1. Ceralloy specimens were prepared by one of three finishing conditions 

Finish Step Wheel Removal 

(in.) 

Removal per pass 

(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 

Fine 

1. Roughing 

2. Induce damage 

3. Finishing 

Accepted practice 

180 grit 

600 grit 

 0.004 

 0.0005 

 0.001 

 0.001 

 0.0001 

RCF-conventional 

Use the “accepted” practice for RCF test bar finishing (Ref PO# SB 1341-02-N1669). Fine grinding using 1200-grit diamond, per Siddiqui e-mail of February 18, 

2004 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials

190

Figure 2. Schematic of RCF specimen geometry. 

This specimen is used in 3BOR RCF tests.

Figure 3. Schematic of half-RCF specimen geometry. The four-point-flexure flexure strength distribution and depth of machining-induced subsurface 

damage were measured.

Figure 4. Schematic of RBS specimen geometry. 

This specimen is being subjected to rotary bending 

and monotonically loaded to fracture—a new test 

method developed in this program.

Figure 5. Schematic of half-RBS specimen geometry. The four-point-flexure flexure strength distribution and depth of machining-induced subsurface 

damage were measured.

Figure 6. Four-point-bend fixture used to flexurestrength-test cylindrical or half-RCF or half-RBS 

specimens.

(a)

(b)

Figure 7. Oblique view (a) of FEA stress analysis of 

half-RBS specimen. Symmetry was used to simplify the 

model, so only one-half of the half-RBS specimen was 

analyzed. Four-point-flexure failure load was used to 

calculate the outer fiber tensile stress (b) that existed at 

fracture.

Fractography of these specimens (and all the other 

flexure specimens) is presently under way to examine the depth of machining-induced subsurface damage. 

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 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

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Table 2. Half-RBS flexure strength results 

Finish Number of test 

specimens 

Average strength 

(MPa) 

Standard deviation (MPa) Characteristic 

strength (MPa) 

Weibull 

modulus 

Coarse 10 715.6 99.1 753.6 8.1 

Fine 10 756.4 51.9 775.2 15.1 

RCF-Conv. 10 732.0 79.6 765.1 11.3 

147-31N are used to explore how contact loading 

damage (as a function of loading rate) is affected by 

sub-surface damage. The competition of quasiplastic damage and cracking processes in Ceralloy 

147-31N is being interrogated as a function of the 

depth of the sub-surface damage and compared with 

polished material as well. NBD200 Si3N4 (NIST 

standard reference material for Knoop hardness) and 

NC132 Si3N4 (NIST standard reference material for 

fracture toughness) are also being tested; specimens 

of those materials are presently undergoing indentation and scratch testing, the results of which will 

serve as a performance reference for comparing the 

performance of Ceralloy 147-31N and whatever additional ceramics are ultimately tested in this program. SN101C and TSN-03NH (both bearing grades 

of Si3N4) are undergoing instrumented indentation 

and scratch testing as well, and their performance 

will be compared with that of the other listed silicon 

nitride compositions. Raman spectroscopy, which 

can 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. 

Figure 8. PiezoRaman spectroscopy profile of Raman peak position as a function of distance from the 

edge in a SN101C ball. Higher peak position is proportional to residual compressive stress, so this 

measurement shows that a gradient in compressive 

stress exists as a function of depth. 

An interlaboratory study of research institutions 

and universities within the United States is being 

organized whereby RCF testing will be performed 

on samples of Si3N4 and the results combined and 

compared. Colorado State University and the Air 

Force Research Laboratory have so far given verbal 

commitments of participation. 

In addition to the collaboration with BU, interactions exist with Germany’s Bundesanstalt für Materilforschung und–prüfung, Japan’s National Institute of Advanced Industrial Science and Technology, and the United Kingdom’s National Physics 

Laboratory to compare RCF performance. These 

interactions exist under the auspices of the International Energy Agency Annex III. Working group 

participants met in January in Cocoa Beach, FL; the 

next meeting is tentatively set for June 2005 in 

Kobe, Japan, coincident with the International Tribology Conference. 

Dialogue with ceramic bearing manufacturers is 

critical to maintaining relevance 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 at 

ORNL, and the evaluation of ceramic RCF performance is now under way. A formal collaboration was 

initiated with BU for 3BOB RCF testing of ceramics; it will enable the eventual reconciliation of RCF 

performance measured with those two techniques. 

The effects of subsurface damage on RCF performance are being scrutinized. Several supplemental 

characterization routes (e.g., flexure strength testing 

and subsurface damage interrogation, damage 

evaluation from Hertzian indentation, piezoRaman 

spectroscopy) are used to assist in the interpretation 

of RCF performance. Domestic and international 

interlaboratory RCF studies are planned to compare 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials

192

and verify ORNL measurements with those generated at other laboratories and institutes. Finally, new 

under-standing of the RCF performance/microstructure relationship will be provided to 

ceramic manufacturers that will serve to promote 

microstructure engineering and improved RCF performance. 

References

1. J. Kang, R. T. Cundill, and M. Hadfield, “The 

Consequences of Aggressive Lapping Processes on 

the Surface Integrity of HIPed Silicon Nitride Bearing Balls,” Tribology in Environmental Design 

2000, Professional Engineering Publishing, London. 

2. G. D. Quinn, L. K. Ives, and S. Jahanmar, 

“On the Fractographic Analysis of Machining 

Cracks in Ground Ceramics: A Case Study on Silicon Nitride,” NIST Special Publication 996, National Institute of Standards and Technology, U.S. 

Department of Commerce, May 2003. 

Presentations

A. A. Wereszczak, “IEA Annex III Working 

Group Meeting on Rolling Contact Fatigue,” presented at the 28th International Cocoa Beach Conference on Advanced Ceramics and Composites, 

Cocoa Beach Hilton, Cocoa Beach, FL, January 27, 

2004. 

A. A. Wereszczak, “ORNL Characterization of 

Ceramics for Armor, Transportation, and Energy 

Applications,” presented at Cercom, Inc, Vista, CA, 

February 25, 2004. 

A. A. Wereszczak, “ORNL Characterization of 

Ceramics for Armor, Transportation, and Energy 

Applications,” presented at Ceradyne Inc, Costa 

Mesa, CA, February 26, 2004. 

A. A. Wereszczak, “Evaluation of Ceramic Deformation Processes Through Hertzian Indentation,” 

presented at Saint-Gobain, Worcester, MA, August 

3, 2004. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

193

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: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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, TN 

Prime DOE Contract Number: DE-AC05-00OR22725 

Objectives 

• Facilitate the integration of new technologies into the transportation sector by implementing research that validates the applicability of these technologies 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

Accomplishments 

• Created a web site that provides the international technical community with information about (1) the mission 

and strategy of the IA, (2) details of the various annexes, and (3) a number of technical reports generated by 

collaborative research tasks conducted under Annex II. The site is at http://ia-amt.ornl.gov/index.html. 

• Initiated Annex IV, A Cooperative Program on Integrated Engineered Surface Technology. Technical efforts 

associated with this annex are under way. 

• Initiated Annex V, Light-Weighting of Materials. Planning efforts are under way. 

• Canada and the United Kingdom joined the IA. Canada will participate in the lightweighting materials activity, 

while the United Kingdom will participate in Annex III. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

194

Future Direction 

• Develop plan for lightweighting materials annex and present to the Executive Committee for approval. 

• Initiate test plan for Annex V. 

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 materials 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) the 

identification of promising new technologies for improving materials performance and (2) the development of specialized characterization techniques for 

validating the applicability of this technology to improve material properties while maintaining acceptable life cycle costs. 

At present the active contracting parties for the 

IA-AMT are 

• Germany—Bundesanstalt für 

Materialforschung und –prüfung (BAM) 

• Canada—Materials Technology Laboratory, 

CANMET 

• United Kingdom—Department of Trade and 

Industry 

• United States—U.S. Department of Energy 

Approach

In the area of 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 

cost; (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 the 

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 measuring 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 

on 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 highpressure 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 to the interface, are unproven. Item 3 addresses this limitation. In a similar 

fashion, as surface modification technologies mature, proven characterization techniques will be required to validate their performance (Item 4). 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

195

Results

The activities related to performance improvement and performance evaluation are covered in 

4 annexes. Annex II focuses on pre-standardization 

(i.e., performance evaluation) issues related to the 

deployment of structural ceramics in transportation 

activities. Specific applications include ceramic diesel exhaust values, cam roller followers, timing 

plungers, etc. Results from these subtasks have led 

to the optimization of techniques for (1) characterization of powder properties, (2) quantification of the 

green-state characteristics, and (3) evaluation of mechanical performance. This extensive set of results 

generated in Annex II has been used in the establishment of standards [via the American Society for 

Testing and Materials (ASTM), Japan Industrial 

Standards (JIS), the Committee for European Normalization (CEN), the International Organization for 

Standards (ISO)] and National Institute of Standards 

and Technology (NIST) guidelines (see the publications list), which in turn have benefited the entire 

ceramics community. 

Table 1 summarizes some of the existing standards that have benefited from the work conducted 

in this IA. In the case of room-temperature flexural 

strength, the test 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 IA–AMT round robin (Subtask 5–Annex II) 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, and 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 lessons learned. 

ISO 15490 evolved from two standards (ASTM 

C1273 and JIS 1606); and the IA–AMT 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 IA–

AMT subtask on thermal shock was completed. Although this standard follows a more conventional 

approach to thermal shock by using water quenching 

and standard MOR bars, insights garnered from the 

IA–AMT work are used in providing guidance to 

users in notes and discussions. 

Current activities are focused on the evaluation 

and consolidation of nanocrystalline ceramics. Participants include the United States and Germany. 

Table 1. Standards that have benefited from the IA subtasks 

Property JIS ASTM CEN ISO 

Flexural strength: RT R1601-95 C1161-02 EN 843-1:95 14704:2000 

Flexural strength: HT R1604-95 C1211-98 prEN 820-1 DIS 17565 

Statistical analysis R1625-96 C1239-00 ENV 843-5:97 CD 20501 

Fractography C1322-02 prENV843-6 

Tensile strength R1606-95 C1273-00 15490:2000 

Sample preparation for the determination 

of particle size distribution of ceramic 

powders 

R1619:95 C1282-00 EN 725-5:96 14703:2000 

Surface area R1626-96 C1274-00 EN 725-6:96 DIS 18757 

Particle size distribution of powder by laser diffraction method R1629:97 TC206 NP02 

Thermal shock R1615-93 C1100-98 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 C1368-00 TC206 PWI 07 

Bending fatigue: HT R16xx:01 

RT = room temperature; HT = high-temperature
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

196

Annex III also focuses on performance evaluation with emphasis on the evaluation of contact 

damage. International activities include a formal collaboration between Mark Hadfield of Bournemouth 

University (BU) in the United Kingdom and Oak 

Ridge National Laboratory (ORNL) in the United 

States. The study of the effects of subsurface damage (e.g., that results from component machining) on 

the rolling contact fatigue (RCF) performance of 

silicon nitride (Si3N4) is the primary theme. A BU 

Ph.D. candidate will serve as the cornerstone of the 

collaboration and alternate on-site research at BU 

and ORNL over the life of the project. The RCF performance of Si3N4 will be interrogated using ball-onthree-ball (Plint TE92) and three-ball-on-rod (NTN) 

RCF tests conducted at BU and ORNL, respectively. 

The Ph.D. candidate will attempt to link RCF performances measured with each technique. Additionally, the generated subsurface damage in cylindrical 

rods and spheres machined under the “same” mching 

conditions will be compared to assess/verify the 

sought equivalence. 

A second international activity involves the 

preparation of a written report detailing techniques 

on RCF testing in each of the participating countries. 

Participants include Japan (AIST, KOYO-Seiko, and 

NSK), Germany (BAM), United Kingdom (NPL), 

and the United States (ORNL). The report will be 

completed at the end of 2004. 

In 2005 and 2006, RCF studies involving both 

international and domestic participants will be implemented. Details of this international activity will 

be developed at the next IEA Annex III RCF Working Group meeting, tentatively set for June 2005 in 

Kobe, Japan—coinciding with the International Tribology Conference. Attendance of participants from 

Japan and the United Kingdom has been confirmed. 

The primary goal is to reconcile differences in the 

various test techniques and hardware. Results, including recommendations, will be forwarded to the 

appropriate standards setting organizations. 

The effort conducted within the United States 

consists of an interlaboratory study of RCF testing 

involving research institutions and universities. 

Colorado State University and the U.S. Air Force 

Research Laboratory have so far given verbal commitments of participation. 

Annex IV, A Cooperative Program on Integrated 

Engineered Surface Technology, focuses on the implementation of a vertically integrated engineered 

surface concept: combining surface texturing, thin 

films, and lubrication to create an engineered surface 

appropriate for different applications, with an overall 

objective of reducing friction and increasing durability. Although engineered surfaces can play a dominant role in energy conservation and utilization, 

several technical challenges must be addressed before such technology can be fully utilized. One 

broad area of crucial importance and mutual benefit is the development of international standards 

and practices for textured surfaces under different 

application conditions. Characterization and mechanical properties of an integrated engineered 

surface in terms of friction reduction, load-bearing 

capability, fatigue mechanisms, and life prediction 

are important issues that need to be addressed. Activities under Annex IV shall include various forms 

of characterization, evaluation methods, modeling, 

bench-scale testing, and demonstration field trials. 

The implementation of lightweighting strategies 

in the transportation sector represents a major approach to increasing fuel efficiency and reducing 

emissions of greenhouse gases. Accordingly, a number of countries (see Table 2) are involved in extensive research activities to develop and validate advanced materials and manufacturing technologies for 

significantly reducing vehicle body and chassis 

weight. Materials under consideration include highstrength steels, magnesium, aluminum, titanium, 

metal-matrix composites, plastics, and ceramics, 

where applicable. Annex V, Light-Weighting of Materials, focuses on methods required for both performance improvement and evaluation, which are of 

interest to the participants. 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 cost; 

(6) data on energy impacts of lightweight materials; 

and (7) shared information on research programs on 

lightweight materials. In terms of recent achievements a draft plan outlining the scope of work, 

which covers areas of mutual interest, was prepared 

for discussion at the executive committee meeting in 

October 2004. Participants included the United 

States, Canada, and Australia (membership pending). 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

197

Table 2. Summary of lightweighting materials programs 

Country Program 

United States FreedomCAR and Vehicle Technologies Program 

http://www.eere.energy.gov/vehiclesandfuels/technologies/materials/index.shtml 

Canada AUTO21 Network of Centres of Excellence 

http://www.auto21.ca/materials_e.html 

Canadian Lightweight Materials Research Initiative (CLiMRI) 

http://climri.nrcan.gc.ca/default_e.htm 

Australia Light Metals Action Agenda 

http://www.industry.gov.au/content/itrinternet/cmsindexpage.cfm 

Cooperative Research Centre for Cast Metals Manufacturing (CAST) 

http://www.cast.crc.org.au/ 

Conclusions

The IA–AMT has made significant progress in 

expanding its scope, as evidenced by the addition of 

two new annex (IV and V). Both annexes are expected to generate new members. Given this progress, a formal request has been made to the IEA to 

extend the IA–AMT for a period of 3–5 years. 

Publications/Presentations

Implementing Agreement for a Programme of 

Research and Development on Advanced Materials 

for Transportation Applications (IA-AMT), Strategic 

Plan, March 2004 (available at http://iaamt.ornl.gov/. 

International Energy Agency Implementing 

Agreement for a Programme of Research and Development on Advanced Materials for Transportation Applications, Annual Report 2003, February 

2004 (available at http://ia-amt.ornl.gov). 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

198
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

199

C. Testing Standards 

George Quinn 

National Institute of 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: Dr. Sidney Diamond 

(202) 586-8032; fax: (202) 586-1600; e-mail: sid.diamond@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: National Institute of 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 refinement, or develop new test 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 

• Completed pre-standardization work on the flexural strength testing of split cylinders. A report is in preparation. Completed and fully documented all work on the fractographic characterization of grinding damage 

cracks. 

• Completed a series of papers and made presentations. 

• Completed about half of a NIST Guide to Practice on Fractographic Analysis of Brittle Materials. 

• Revised, refined, and improved a variety of current ASTM and ISO standards. 

• Helped found a new ASTM subcommittee on applications standards. 

• Worked to correct and refine an ASTM standard specification for silicon nitride bearing balls. 

Future Direction 

• Write up the split cylinder test method and present the results to ASTM Committee C-28, Advanced Ceramics, 

for review. The error analyses for flexural strength testing of cylindrical rod specimens will be completed and a 

draft standard prepared for ASTM. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

200

• Complete the Guide to Practice on Fractographic Analysis. 

• Resume pre-standardization testing on the diametral compression strength test with the goal of standardization. 

Introduction 

This project creates new test methods that will 

facilitate the use of advanced materials in heavyduty propulsion systems. Much of the work is for 

brittle materials such as ceramics, for which classical 

mechanical test methods developed for metals 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 test methods or invent 

new ones that will allow engineers and researchers 

to measure mechanical properties with good accuracy and precision. Formal test method standards are 

our primary objective. Sound test methods and highquality databases will enhance the credibility of new 

materials and encourage engineers to use them in 

advanced heat engines. 

Figure 1. Silicon nitride bearings. 

Approach

Over the course of this program, we have formulated or contributed to the development of 17 ASTM 

and ISO standards. These have covered a range of 

topics including hardness, flexural strength, fracture 

toughness, fractography, elastic modulus, and 

Weibull strength distribution parameter determination. We have even contributed to a ceramic material 

specification for silicon nitride ball bearing materials. A spin-off benefit of this work has been that the 

generic test method standards have been used to create several biological ceramic materials specifications for dental and hip joint implant ceramics. 

Work continues on new alternative test methods. 

A major element of this program has been the 

refinement of fractographic techniques to facilitate 

the detection and characterization of strengthlimiting flaws in both laboratory-scale test specimens and components. The strength of ceramic 

components or parts is controlled by microscopic 

flaws, which may either be intrinsic to the manufacture or be introduced during grinding and finishing. 

We also continued our work on several new mechanical test methods that are designed to measure 

mechanical properties of cylindrically shaped components or test coupons. Formal standard test methods are on the books for rectangular bend bars or 

tension-strength specimens, but these are not suitable for cylindrical parts (Figure 2). Simple methods 

for cylindrical parts are needed. Three methods we 

are working on are split cylinder flexural strength 

tests, rod flexural strength tests, and diametral compression strength tests. 

Figure 2. Round specimens. 

Important refinements and corrections were 

made to several existing standards during the past 

year. Thus “standards maintenance” activity is necessary to incorporate improvements and lessons 

learned from this and other projects. 

Results 

A major task was completed that involved fractographic characterization of grinding-induced 

cracks on the surface of finished ceramics. Figure 3 

shows an example of a grinding crack in a rod 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

201

Figure 3. Fracture surface of a broken ceramic rod test specimen. The 

fracture origin is in the smooth 

semicircular region on the bottom of 

the figures. The origin is a “Vmachining” crack (arrows in the 

lower figure) created when the rod 

was ground to final dimensions. 

Similar cracks can limit strength in 

components 

strength test specimen. Several major reports and 

technical papers were prepared and presented to distribute this important information about grindinginduced damage. Lessons learned were incorporated 

into several ASTM standards. 

An ASTM standard for fractographic analysis, 

C 1322, has already been adopted; but a more userfriendly Guide to Practice for Fractographic Analysis is also in preparation. The formal standard outlines the optimum procedures that engineers and 

fractographers should follow in performing fractography, but it assumes the analyst has moderate experience. The Guide to Practice will be a tutorial 

that will help the lay engineer learn the art of fractographic analysis of brittle materials and help transform what may seem to be a mysterious art into a 

routine engineering practice. It will have hundreds 

of example fractographs. Figure 4 is an example illustration showing how grazing incident angle lighting can help bring out fracture markings on fracture 

surfaces. The new Guide to Practice will include a 

set of recommendations for measuring 

Figure 4. An example illustration from the Guide to 

Practice for Fractographic Analysis of Brittle Materials. This schematic shows vicinal, or low angle 

grazing illumination, which can create shadows that 

highlight features on the rough fracture surface of a 

broken test piece such as shown in the previous figure. Fractography can be effectively applied to understand the causes of fracture in the laboratory and 

in service components. 

fracture mirror sizes. Fracture mirrors are relatively 

smooth zones surrounding fracture origins in ceramics, and their size can be directly correlated to the 

stress in the specimen or component at the instant of 

fracture. Fracture mirror analysis is powerful tool 

for forensic analysis since it can calculate the fracture stress even if the mode of loading is unclear or 

unknown. 

Substantial progress was made this year on prestandardization for the split cylinder strength test 

method. Short, stubby cylindrical parts are difficult 

to test for strength. A simple solution is to split the 

cylinder lengthwise and test the halves in a modified 

common flexural strength test fixture, as shown in 

Figures 5 through 7. Figure 8 shows a fracture surface of one of the rods. We had difficulty identifying 

the precise nature of the flaws in this material, but 

we eventually deciphered several fracture origins. A 

paper is in preparation on our findings. The Weibull 

strength distribution for one particular batch of material is shown in Figure 9. The findings of this work 

will be presented to ASTM Committee C-28, Advanced Ceramics, for review for suitability as a 

standard test method. Results will also be coordinated with the appropriate ceramic part manufacturers and engine companies. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

202

Figure 5. A ceramic split cylinder specimen 

(white) in a common three-point flexure strength 

fixture. This is a simple way to measure the 

strength of the ceramic part, which would be difficult to do for the intact part due to the high forces 

that would be necessary to break it. The finished 

round outer part surface is tested, and the cut surface which is on the top (compression side) of the 

bending loading, does not affect the results. 

Figure 6. A cylindrical zirconia ceramic engine 

part. The part is a fuel injector pin for a heavy 

duty diesel engine. The left shows the intact part 

and the right shows a part cut in half. The latter is 

easier to test for strength. 

Figure 7. Expanded view of the bend fixtures for 

a split cylinder (white specimen). Cradles hold the 

bottom of the specimen and distribute the load 

evenly. Load is applied to the specimen top (the 

cut surface) by the middle roller which is shown 

held in place by rubber bands on the subassembly 

on the top right of the photo. The middle roller is 

free to articulate or rock to match the specimen 

surface. Extra parts are shown in the fore and 

background. 

Figure 8. Optical fractograph showing the fracture 

surface of a broken split cylinder zirconia specimen. 

The arrow marks the fracture origin. Fractographic 

analyses was difficult in this material which had very 

rough fracture surfaces. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

203

Figure 9. Weibull plot for split zirconia fuel injector 

pin parts. 

We have made substantial progress in refining 

testing methods for measuring the flexural strength 

of full cylindrical rod specimens. Figure 10 shows a 

full-size rod specimen which, unlike the parts shown 

in earlier figures, is longer and therefore can be 

tested more easily. Cradles must be used to apply the 

forces evenly to the specimen. In FY 2004, we made 

progress on the analytical error analysis for this configuration. We expect to finish this work in the upcoming year, paving the way for standardization. 

Figure 10. Flexural strength test fixture for a rod 

specimen. 

Pre-standardization work on the diametral compression strength test specimen is planned for the 

upcoming year (Figure 11). Small pill-shaped 

specimens are loaded on their rims and split by tensile stresses. This is a convenient testing configuration, which in principle is ideally suited for small 

cylindrical parts; but there are some unsolved problems. We have completed some preliminary work 

and a review of the literature, which has revealed 

some of the problems with the method. 

Figure 11. Diametral compression strength test. 

Finally, we refined and updated a number of the 

standards created earlier in this program, including 

ASTM hardness, C 1326 and C 1327; ASTM fracture toughness, C 1421; ASTM fractographic analysis, C 1322; ASTM silicon nitride bearing specification, F 2094; and ISO 14704 and ISO 17565 on 

flexural strength at room temperature and high temperatures, respectively. As new information becomes available or shortcomings are revealed, it is 

prudent to update and correct these formal standards. 

 

Conclusions 

We have made progress in expanding the suite 

of test methods available for evaluating the properties of brittle ceramic materials. Our goal is to improve the test method and data base infrastructure 

that facilitates the incorporation of new advanced 

materials in heavy-duty diesel engines. Substantial 

progress was made in improving fractographic 

methods for detecting grinding-induced cracking 

and in developing test methods for measuring the 

strength of cylindrical parts. ASTM and ISO standards have been adopted, refined, and updated. 

 

Publications/Presentations

G. D. Quinn, L. K. Ives, and S. Jahanmir, “Fractography Reveals Machining Cracks,” Bul. Amer. 

Ceram. Soc. 82(7) 11 (2003). 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

204

G. D. Quinn, L. K. Ives, and S. Jahanmir, “Machining Cracks in Finished Ceramics,” to be published in Fractography of Advanced Ceramics, II, 

ed. J. Dusza, TransTech Publishing, Zurich, 2005. 

G. D. Quinn, L. K, Ives, and S. Jahanmir, “Machining Damage Cracks: How to Find and Characterize Them by Fractography,” pp. 383–394 in Ceram. Eng. Sci Proc. 24(4), 2003. 

G. D. Quinn, R. J. Gettings, and L. K. Ives, “A 

Standard Reference Materials for Vickers Hardness 

of Ceramics and Hardmetals,” to be presented at the 

IMEKO Hardness Conference, National Institute of 

Standards and Technology, November 11–12, 2004. 

G. D. Quinn, J. Eichler, U. Eisele, and J. Rödel, 

“Fracture Mirrors in a Nanoscale 3Y-TZP,” J. Amer. 

Ceram. Soc. 87(3) 513–516 (2004). 

A. B. Kounga Njiwa, T. Fett, J. Rödel, and G. 

D. Quinn, “Crack Tip Toughness Measurements on 

Sintered Reaction-Bonded Si3N4,” J. Amer. Ceram. 

Soc. 87(8) 1502–1508 (2004). 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

205

D. 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 

DOE Technology Development Area Specialist: Dr. Sidney Diamond 

(202) 586-8032; Fax: (202) 586-1600; e-mail: sid.diamond@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: National Institute of Standards and Technology, Gaithersburg, Maryland 

Objectives 

• Organize an international cooperative research program on 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 the understanding and appropriate models to explain the texturing effects on frictional characteristics. 

Develop appropriate thin films and coatings to achieve a synergistic and complementary relationship with texturing to enhance performance. 

• Discover and develop a surface chemistry for protective films and coatings that works in synergy with the coatings. 

Approach 

• Determine the effect of size, shape, pitch, and patterns of surface textural features on friction under (1) a highspeed, 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 an individual approach alone. 

• Concurrently, organize an international cooperative research program under the auspices of the IEA to pool resources and share this energy conservation technology worldwide. 

Accomplishments 

• Developed a new design principle for surface texture design to enable friction reduction under boundary lubrication conditions. 

• Developed in situ instrumentation to allow direct observation of interface flow patterns with and without surface textural features. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

206

• Visited several countries to establish agreement to join the IEA cooperative research activity. 

• Visited various domestic companies and universities to organize a U.S. working group on integrated surface 

technology. 

• Obtained final approval from the IEA Executive Committee on a new Annex on integrated surface technology 

in March 2004. 

Future Direction 

• Organize a U.S. national working group under the IEA banner to conduct joint research and provide information for international exchange. 

• Organize an international working group and arrange a meeting at a location and time convenient to all participants. 

• Add a modeling component to determine why some textural features function and some do not. The results will 

lead to design guidelines. 

• Initiate the synergistic effect of thin films and coatings to protect textures. 

• Explore the tribochemistry of various thin films and coatings. 

Introduction

Frictional losses are inherent in all mechanical 

components in relative motion. The ability to control 

friction enables many technologies. Over the years, 

materials, lubricants, and surface modifications have 

provided the means to control and reduce friction to 

enhance energy efficiency in automotive and truck 

applications. A new avenue is needed to continue to 

increase the efficiency of energy utilization. Recently, dimples have been demonstrated to achieve 

friction reduction and durability enhancement in 

conformal contacts such as seals. The effect of dimples depends strongly on pattern design, contacting 

materials, and lubricant properties. A full understanding of these dependencies has yet to be developed. This research opens up a new avenue for increasing energy efficiency by combining surface 

textures, thin films and coatings, and the associated 

chemistry to protect surface textures. The objective 

of this project is to identify critical pattern features 

that control friction under a broad range of contact 

conditions and to develop models to explain these 

relationships. 

Concomitant to this research, international cooperative research under the auspices of the IEA is 

planned to pool resources worldwide to accelerate 

this technology development. The results of this cooperative research will impact energy conservation 

efforts worldwide. Toward this end, the United 

Kingdom, Germany, Finland, Sweden, Israel, and 

Japan have agreed to participate in this activity under IEA Annex IV. 

Approach

During FY 2004, experiments were conducted 

to examine various surface textures on steel surfaces 

using photolithography and chemical etching. Using 

the same area coverage (% 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 those experienced by surface seals. Results indicated that (1) surface feature shape had great influence on friction reduction, and (2) friction reduction 

also depended strongly on the orientation of the surface features with respect to the direction of sliding. 

The results suggested that the conventional theory of 

hydrodynamic lift is not adequate to explain these 

observations, and a new theory is needed. Therefore, 

an in situ observation of the lubricant flow pattern 

will be developed to observe how a specific feature 

would affect the friction. A modeling effort will also 

be initiated in FY 2005. 

To achieve energy savings, friction reduction 

needs to be achieved over a wide range of contact 

conditions existing in gears, transmissions, and engines. Therefore, the approach was to move into 

higher and higher loads and slower and slower 

speeds. New instrumentation and methods needed to 

measure the effect of textures will be developed. 
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

207

New design principles for textural features will need 

to be developed to meet the load and speed requirements. 

Results

Circular dimples were used to conduct high-load 

experiments using stainless steel specimens. The 

load was raised to 93 N on a smaller sample pin so 

that the apparent contact pressure was approximately 

15–20 MPa. To prevent wear of the dimples, a fully 

formulated synthetic lubricant was used. Baseline 

specimens with and without dimples were compared. 

In each case, the dimples caused high friction. 

Grooves also exhibited higher friction when the contact pressure was increased or the speed reduced. 

These results tended to confirm the fluid mechanics 

models. A new design principle was needed to move 

into GPa contact pressure range. 

Considering the basic principles of hydrodynamic lift, we designed a surface feature with a 

sloped bottom to provide an artificial hydrodynamic 

wedge. The idea was that if sufficient contact pressure were exerted, the surface would undergo elastic 

and plastic deformation and the liquid inside the feature would be squeezed, producing reactionary pressure to support the load. Since the pin-on-disk apparatus was limited by load, a four-ball wear tester 

with a configuration of ball-on-three-flat was used to 

generate high contact pressures. A 2.4-mm-diam 

52100 steel ball bearing was used to slide on top of 

three 6-mm-diam disks. The testing conditions were 

0.19–1.9 m/s speed and 2-, 5-, and 8-Kg loads lubricated by a fully formulated engine oil. The disks 

were also 52100 steel without case hardening. The 

surface texture was fabricated using mechanical 

scribing and electrochemical etching to create various sizes and shapes and slopes at the bottoms of the 

features. Since earlier work suggested that orientation with respect to the flow direction was a factor, 

the disks were carefully marked along the major axis 

of the features and mounted in the test apparatus. 

Figures 1 and 3 display the data in a 3-dimensional 

plot showing friction as a function of speed and load 

for steel on copper. 

Figure 1 shows the baseline friction as a function of load and speed. Figure 2 shows the sliding 

direction in a typical hydrodynamic wedge shape; 

the friction coefficients across the speed and load 

range were lower than in the untextured case. 

1000

2000

3000

4000

5000

5

10

15

20

0.00

0.02

0.04

0.06

0.08

Fr ci t oi n coeff ci ei nt

Load, Kg

Rotating speed, rpm

Figure 1. Untextured surface showing friction as a 

function of load and speed. 

Figure 3 shows the sliding direction from shallow to 

deeper slope; the friction showed uniform improvement across the ranges. This finding was surprising, 

but it suggested other mechanisms might be in involved in addition to the classical hydrodynamic 

wedge theory. For these experiments, the contact 

pressure was approaching the GPa range. Instead of 

increasing friction, the hidden wedge geometry was 

able to reduce friction compared with the base case 

under boundary lubrication conditions. The dependence of the orientation direction was also surprising. 

If the major mechanism was to squeeze the fluid 

film along the wedge, then we would expect that 

flow direction along the long axis of the feature 

would give lower friction. For that matter, the cavitation mechanism would also suggest that this geometry would be better. The fact that flow along the 

short axis gives better friction control would seem to 

suggest that back-flow within the feature might play 

a more significant role than previously suspected. 

Conclusions

We have demonstrated a new design principle 

that can reduce friction under boundary lubrication 

conditions. Additional tests on softer materials, in 

which we had more control over shaping the features, showed experimentally that a friction reduction of an order of magnitude was achievable. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

208

Figure 2. Textured surface with the flow toward the long axis of the feature.

Figure 3. Textured surface with the flow toward the short axis of the feature. 

This opens up a whole new avenue—the use of 

surface texturing in energy-intensive applications. 

The lack of an existing theory to adequately explain 

these observations suggests a new theory is needed. 

Therefore, in FY 2005, we will initiate a new modeling effort to increase understanding of the fundamental mechanisms. 

Presentations

S. M. Hsu, “An Integrated Approach to Design 

Intelligent Surfaces for Heavily Loaded Contacts,” 

presented at the International Joint Tribology Conference, Ponte Vedra, FL, October 27–29, 2003. 

S. M. Hsu, “An Integrated Surface Texture Design with Surface Modification Techniques and Thin 

Lubricating Films,” presented at the Smart Surfaces 

in Tribology: Advanced Additives and Structured 

Coatings Conference, Zurich, Switzerland, September 10–12, 2003. 

S. M. Hsu, “An Integrated Surface Technology 

International Program under IEA,” presented at the 

International Energy Agency Executive Committee 

meeting, Oakland, CA, October 20–23, 2003. 

S. M. Hsu, “Surface Texturing under Boundary 

Lubrication for Friction Control,” presented at the 

STLE annual meeting, Toronto, Canada, May 17–

20, 2004. 

Jorn Larson Basse, X. Wang, L. Ives, and S. M. 

Hsu, “Some Friction Experiments with Textured 

Surfaces,” presented at the Nordic Symposium on 

Tribology, Troms, Norway, June 2004. 

1000

2000

3000

4000

5000

5

10

15

20

0.00

0.02

0.04

0.06

0.08

Fr ci t oi n coeff ci ei nt

Load, Kg

Rotating speed, rpm

1000

2000

3000

4000

5000

5

10

15

20

0.00

0.02

0.04

0.06

0.08

Fr ci t oi n coeff ci i ent

Load, Kg

R otating speed , rp m
Heavy Vehicle Propulsion Materials FY 2004 Progress Report 

209

Jorn Larson Basse, X. Wang, L. Ives, and S. M. 

Hsu, “Some Friction Experiments with Dimpled 

Surface Texture,” presented at the Fourth China International Symposium on Tribology, Xian, China, 

November 8–11, 2004. 

S. M. Hsu, “An Integrated Surface Modification 

Technique to Control Friction: A New Paradigm,” 

keynote speech at the Fourth China International 

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,” presented at the First International Conference on Advanced Tribology, Singapore, December 1–3, 2004. 

Publications

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 and S. M. Hsu, “An Integrated Surface 

Modification Technique to Control Friction: A New 

Paradigm,” keynote paper, in Proceeding of the 

Fourth China International Symposium on Tribology, Xian, China, November 8–11, 2004. 
FY 2004 Progress Report Heavy Vehicle Propulsion Materials 

210
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

211

Appendix A. ACRONYMS AND ABBREVIATIONS

2D two-dimensional

3D three-dimensional

3BOB three-ball-on-ball

3BOR three-ball-on-rod

ACEM aberration-corrected electron microscope

ADF annular dark-field

AE acoustic emission

AFM atomic force microscope

Al aluminum

AMCL Advanced Materials Characterization Laboratory

ANL Argonne National Laboratory

ASM American Society for Metals

ASME American Society of Mechanical Engineers

ASTM American Society for Testing and Materials

Ba barium

BET Brunauer-Emmett-Teller theory

BMO base metal oxides

BU Bournemouth University

CARES Ceramic Analysis and Reliability Evaluation of Structures

CC current collector

CCVD combustion chemical vapor deposition

CDF cumulative distribution function

CEN Committee for European Normalization

CERT Committee on Energy Research and Technology

CHIP cold + hot isostatic pressing

CIM chemically interactive material

CP commercially pure

CVD chemical vapor deposition

CF8C an Fe-19Cr-12NiCNb austenitic stainless steel

CN-12 an Fe-25Cr-13NiCNNb austenitic stainless steel

CRADA cooperative research and development agreement

CRT continuously regenerating technology

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

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
FY 2004 Progress Report Heavy Vehicle Propulsion Materials

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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/finite element modeling

FTIR Fourier transform infrared spectroscopy

GPa gigapascals

HA-ADG high-angle annular dark field

HC hydrocarbons

HDI high-density infrared

HFIR High Flux Isotope Reactor

HIP hot isostatic pressing

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

IITS instrumented indentation test system

IEA International Energy Agency

IEST integrated engineered surface technology

INEEL Idaho National Engineering and Environmental Laboratory

ISO International Organization for Standards

JIS Japan Industrial Standards

LANL Los Alamos National Laboratory

LBNL Lawrence Berkeley National Laboratory

LNT lean Nox trap

LS low stress

LSCF La0.6Sr0.4Co0.8Fe0.2O3

LST laser surface texturing

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 of Standards and Technology

NOx oxides of nitrogen; nitrites/nitrates species
Heavy Vehicle Propulsion Materials FY 2004 Progress Report

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NPD neutron powder diffraction

NSF National Science Foundation

NSR NOx stoichiometric ratio

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

PM particulate matter

PM parts per million

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

SAcM scanning acoustic microscopy

SAM scanning auger microprobe

SAXS small angel x-ray scattering

SCG slow crack growth

SCR selective catalytic reduction

SE sensing electrode

SEM scanning electron microscopy

SiC silicon carbide

SIDI spark-ignition direct-injection

SIUC Southern Illinois University–Carbondale

SMPS scanning mobility particle sizer

SOFC solid oxide fuel cell

SOx oxy-sulfur species

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 thru focal series

TGA thermal gravimetric analysis (or analyzer)

Ti titanium

TiAl titanium aluminide
FY 2004 Progress Report Heavy Vehicle Propulsion Materials

214

TiC titanium carbide

TTBCs thick thermal barrier coatings

TTZ transformation-toughened zirconia

TWC three-way catalyst

TWS Third-Wave Systems

UHR-TEM ultra-high resolution TEM

UHMW-PE ultra-high molecular weight polyethylene

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


A Strong Energy Portfolio for a Strong America 

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

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

invests in a diverse portfolio of energy technologies.

For more information contact:

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