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EERE-AVTAEA_Cover'06.indd 2 
U.S. Department of Energy 

Office of FreedomCAR and Vehicle Technologies 

1000 Independence Avenue S.W. 

Washington, DC 20585-0121 

FY 2006 

Progress Report for Heavy Vehicle Propulsion Materials 

Energy Efficiency and Renewable Energy 

Office of FreedomCAR and Vehicle Technologies 

Advanced Materials Technologies 

Edward J. Wall Program Manager, OFCVT 

Rogelio A. Sullivan Advanced Materials Technologies Team Leader 

Jerry L. Gibbs Technology Development Manager 

January 2007 
CONTENTS 

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

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

A. Mechanical Reliability of Piezo Stack Actuators .............................................................................. 5 

3. MATERIALS FOR EXHAUST AFTERTREATMENT.................................................................... 11 

A. Materials for Exhaust Aftertreatment............................................................................................... 11 

B. Characterization of Catalyst Microstructures and Deactivation Mechanisms......................... 27 

C. Development of NOx Sensors for Heavy Vehicle Applications....................................................... 35 

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

(CRADA No. ORNL-02-0659 with Cummins, Inc.) ....................................................................... 39 

E. Catalysis by First Principles.............................................................................................................. 51 

F. Durability of Diesel Particulate Filters (CRADA with Cummins, Inc.) .......................................... 67 

4. MATERIALS FOR AIR HANDLING AND THERMAL MANAGEMENT.................................... 75 

A. Austenitic Stainless Steel Alloys for Exhaust Manifold (CRADA with Caterpillar, Inc.).............. 75 

B. Friction and Wear Reduction in Diesel Engine Valve Trains.................................................. 81 

5. MATERIALS FOR STRUCTURAL COMPONENTS...................................................... 87 

A. Cost-Effective Machining of Titanium and HCCI Fuel Systems Spray Hole Drilling ................. 87 

B. Synergistic Processing-Enhanced Grinding of Advanced Engine Materials........................... 97 

6. MATERIALS FOR HOT SECTION COMPONENTS..................................................... 101 

A. NDE of Diesel Engine Components........................................................................ 101 

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

C. Engineered Surfaces for Diesel Engine Components...................................................................... 121 

D. High-Temperature Advanced Materials for Lightweight Valve Train Components................. 125 

7. CHARACTERIZATION OF PHYSICAL AND MECHANICAL PROPERTIES................... 143 

A. IEA - Implementing Agreement for a Programme of Research and Development on Advanced

Materials for Transportation Applications.................................................................. 143 

B. IEA- Rolling Contact Fatigue............................................... ............................................ 147 

8. EMERGING TECHNOLOGIES............................................................................... 153 

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

Applications ................................................................................................... 153 

B. Powder Processing of Nanostructured Alloys Produced by Machining.......................................... 161 

C. Deformation Processes for the Next-Generation Ceramics....................................................... 169

ii 
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

1. INTRODUCTION 

Heavy Vehicle Propulsion Materials 

Advanced materials are an enabling technology for fuel-efficient commercial vehicle engines. Heavy Vehicle 

Propulsion Materials supports the engines and emissions activities in the FreedomCAR and Vehicle 

Technologies Program to provide enabling materials technologies in a timely fashion. 

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. 

The use of piezoelectric stack actuators as diesel fuel injectors has the potential to reduce injector response 

time, provide greater precision and control of the fuel injection event, and lessen energy consumption (e.g., 

for use in homogeneous charge compression ignition, HCCI, engines). However, the piezoelectric stacks are 

brittle and potentially subject to premature failure. A new R&D effort was begun to apply established brittle 

material probabilistic design and reliability analyses to piezo-stack actuators, generate the required 

mechanical property data for the design and reliability analysis, and identify minimum mechanical 

performance requirements for fuel injector piezo-stack actuators. 

Novel scuffing tests to evaluate fuel injector materials in diesel fuel and biodiesel fuel environments have 

been developed and are in beneficial use for evaluating fuel system materials. The testing system is available 

to High Temperature Materials Lab (HTML) users as well as to the Propulsion Materials Program. 

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 

commercial vehicles are greatly complicated by engine and exhaust aftertreatment technologies designed to 

meet the mandatory U.S. Environmental Protection Agency (EPA) emission regulations for 2010 and 2012. 

Materials and systems research is being conducted to minimize the potentially negative effects of emissionreduction technologies on fuel economy and to result in cleaner and more efficient engines. 

Durability of exhaust aftertreatment systems in commercial 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. 

The characterizations of new materials from various stages of the catalyst’s life cycle were continued this 

year. X-ray diffraction, X-ray photoelectron Spectroscopy and Scanning Transmission Electron Microscopy 

were used to evaluate the thermal degradation of a NOx storage and reduction (NSR) catalyst within the 

catalyst support brick as a function of macroscopic position, catalyst history and operating conditions. Two 

studies characterizing a new catalyst system and a diesel oxidation catalyst in the unused and aged states were 

completed. 

1
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

We have used Density Functional Theory (DFT) methods to investigate formation and adsorption energies as 

well as reaction energy profiles. Formation energies were calculated for Pt1-3Oy on MgO(100), while 

adsorption energies of O and CO on PtxOx and PtxO2x and O, O2, CO, CO2, NO, and NO2 on Ptx were 

calculated. We also calculated reaction energy profiles for CO+1/2O2 →CO2 and NO+1/2O2 →NO2. In 

concert with the first principles calculations we completed the synthesis of a series of supported Pt catalysts 

and evaluated them for CO oxidation. 

The atomic-scale resolution of the aberration-corrected scanning transmission electron microscope (ACEM) 

was used to characterize the nature of platinum catalyst structures on alumina, silica and titania support 

materials, including the imaging of platinum species dispersed as single atoms and in precrystalline clusters 

known as “rafts.” In particular, the different behavior of the Pt atoms, clusters and rafts on the support 

materials was characterized, as a function of electron beam excitation. 

The development of test techniques to characterize the physical and mechanical properties of ceramic diesel 

particulate filters (DPFs) and develop analysis and inspection tools for assessing their reliability and 

durability was continued. Probabilistic design tools were used to model the durability and expected life of 

DPF ceramic substrates. Physical and mechanical property measurements were made on DPF substrates to 

better understand the time dependant fracture mechanics behavior. 

Caterpillar continued to assess durability of various catalyst technologies in terms of phosphorous, sulfur and 

thermal degradation, and to identify materials that have high diesel particulate filtration efficiency, low 

pressure drop, and comply with future emission regulations for commercial diesel engines. The effort was 

focused on identifying the impact of phosphorus (P) exposure on diesel particulate filters (DPF) and diesel 

oxidation catalysts and of P plus sulfur (S) exposure on selective catalytic reduction (SCR) catalyst 

formulations. 

Diesel engine components that require durability and low friction in order to improve propulsion system 

energy-efficiency and reduce emissions were identified. These include wastegate bushings for exhaust gas 

recirculation (EGR) components, fuel injector plungers, and valve and valve seat materials. Test methods, 

data analysis methodology, and models to evaluate and quantify the performance of candidate materials under 

simulated use conditions were developed. 

Caterpillar developed high-velocity oxygen fueled (HVOF) coating processes to apply amorphous steels and 

quasicrystal materials to engine pistons and cylinder heads for use as thermal barrier coatings in homogenous 

charged combustion ignition (HCCI) engine designs. The diffusion stability of the new materials at engine 

operating temperatures was determined by means of diffusion couples and microprobe analysis. 

Caterpillar and ORNL continued the development of CF8C-Plus cast stainless steel for near-term applications 

in diesel engine exhaust manifolds and turbocharger housings. Highlights include demonstrating that CF8CPlus cast stainless steel is far superior to SiMo cast iron at 600°C and above. CF8C-Plus also shows strength, 

creep resistance and aging resistance benefits compared with other commercial cast stainless steels being 

considered for heavy-duty diesel turbocharger housings at 750°C. Significant progress has been made in 

commercializing the alloy. Creep-rupture resistance was determined for CF8C-Plus Cu/W and CF8C-Plus to 

further the interest of major diesel turbocharger OEMs to upgrade housings. ASTM weldability was 

characterized, along with castability of thin-walled CF8C-Plus and CF8C-Plus Cu/W in support of exhaust 

component applications. 

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. The objective is 

valves 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 

2
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

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. 

This year a 500-hr durability test of the light-weight valves was successfully completed. A Caterpillar C15 

ACERTTM engine was donated to the National Transportation Research Center (NTRC) for future component 

testing. 

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

University and Third-Wave Systems. The machining characteristics of difficult materials such as fuel 

injection plungers and wear-resistant engine parts were characterized and modeled in order to develop highspeed, cost-effective manufacturing processes. 

Rogelio A. Sullivan Jerry L. Gibbs 

Team Leader, Materials Technologies Technology Development Manager 

Office of FreedomCAR and Vehicle Technologies Office of FreedomCAR and Vehicle Technologies 

Energy Efficiency and Renewable Energy Energy Efficiency and Renewable Energy 

3
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

2. MATERIALS FOR ADVANCED FUEL SYSTEMS 

A. Mechanical Reliability of PZT Piezo-Stack Actuators for Fuel Injectors 

Andrew A. Wereszczak and Hong Wang 

Ceramic Science and Technology Group (CerSaT) 

Oak Ridge National Laboratory 

P.O. Box 2008, MS 6068 

Oak Ridge, TN 37831-6068 

(865) 576-1169; fax: (865) 574-6098; e-mail: wereszczakaa@ornl.gov 

DOE Technology Development Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

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

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Objectives 

• Apply established structural ceramic probabilistic design and reliability analysis to piezo-stack actuators. 

• Generate the required mechanical property data for the design and reliability analysis. 

• Identify minimum mechanical performance requirements for fuel injector piezo-stack actuators. 

Approach 

• Develop finite element models that estimate service stress states in piezo-stack actuators. 

• Measure strength and fatigue distributions of piezoelectric ceramics that comprise piezo-actuators as a function 

of electric field and loading rate, and combine with FEA modeling to predict performance of whole stacks. 

• Evaluate piezoelectric and mechanical response of whole piezo-stack actuators and compare to predictions. 

Accomplishments 

• Created preliminary FEA models to predict stress states in piezo-stack actuators. 

• Established miniature mechanical test setup to measure strength of thin (~200 micron or 0.008” thick) 

piezoceramics. 

• Designed and procured hardware that will enable mechanical testing of piezoceramics as a function of electric 

field, and cyclic fatigue testing of piezostacks with concurrent confined load. 

Future Direction 

• Refine FEA models. 

• Characterize effect of electric field on piezoceramic mechanical performance. 

• Destructively evaluate piezo-stack actuators to identify performance limiting stresses and flaw types. 

5
Heavy Vehicle Propulsion Materials 

Introduction 

The use of piezoelectric stack actuators as diesel 

fuel injectors has the potential to reduce injector 

response time, provide greater precision and control 

of the fuel injection event, and lessen energy 

consumption (e.g., for use in homogeneous charge 

compression ignition, HCCI, engines). Examples of 

piezo-actuators are shown in Fig. 1. Though 

piezoelectric function is the obvious primary 

function of lead zirconate titanate (PZT) ceramic 

stacks for fuel injectors, their mechanical reliability 

can be a performance and life limiter because PZT is 

brittle, lacks high strength, and may exhibit fatigue 

(i.e., slow crack growth) susceptibility. That 

brittleness and relatively low strength can be 

overcome with proper design though. This project 

combines in-situ micromechanical testing, 

mechanical characterization of full piezo-stack 

actuators, microstructural-scale finite element 

analysis, probabilistic design sensitivity, and 

structural ceramic probabilistic life prediction 

methods to systematically characterize and 

optimally design PZT piezoelectric stack actuators 

that will enable maximized performance and 

operational lifetime. 

Figure 1. Examples of piezo-stack actuators. Left is 

a top view, middle is a side view, and right is a stack 

removed from the protective covering. 

FY 2006 Progress Report 

A significant portion of this project is devoted to 

understanding the link between the microstructural 

and macroscopic responses of a piezo-stack. A 

multilaminate architecture comprises these piezostack with alternating layers of electrode and 

piezoceramic. Such an internal architecture is 

shown in Fig. 2. Cracking can initiate in the piezostacks detrimentally affecting the actuation, so 

concerted efforts are underway to understand that 

cracking from a structural ceramic perspective. 

Figure 2. Cross-section of a piezo-stack actuator. 

Approach 

The strength test setup (see Fig. 3) consists of a 

micromechanical test frame that has integrated load 

and acoustic emission measuring capabilities. Load 

and acoustic histories for a typical test are shown in 

Fig. 4. Interestingly, the specimens do not always 

fracture catastrophically and Fig. 4 can illustrate 

that. The first acoustic event (occurring at the ~ 90s 

6
Heavy Vehicle Propulsion Materials 

mark) is associated with crack initiation and that 

corresponding load was used to calculate strength 

after Shetty et al. [1]. That fracture does 

immediately run through the whole specimen 

causing the specimen to withstand additional 

loading. The second event is undergoing further 

interpretation and the third event is associated with 

complete specimen fracture. 

Figure 3. Micromechanical ball-on-ring biaxial 

flexure strength test setup. Load and acoustic 

activity are continuously measured during each test. 

A thin disk specimen is shown whose thickness is 

approximately 200 microns. 

Figure 4. The upper graph shows acoustic signal 

rise time as a function of time and the lower is a 

double-Y graph that shows acoustic events (red dots) 

and load in Newtons as a function of time. The first 

event is associated with crack initiation and the 

corresponding load is used to calculate strength. 

FY 2006 Progress Report 

In order to confidently estimate strength of the 

PZT films tested with the configuration shown in 

Fig. 3, finite element analysis (FEA) was conducted 

of the ball-on-ring tested films, and then compared 

against the typically used formula for this loading 

configuration. This validation was an important task 

because measured strength is perhaps the most 

prominent dependent parameter being measured. 

To initiate the study of the translation of 

piezoceramic strength and probabilistic life 

prediction and design of piezoactuators, industry 

type 5A4E and 5H4E piezoceramic sheets were 

sectioned into square specimens and flexure tested 

in ball-on-ring testing. The manufacturer’s reported 

d33 for the 5A4E and 5H4E dielectrics were 390 x 

10-12 m/V and 650 x 10-12 m/V, respectively while 

their reported Curie temperatures were 350 and 

230°C, respectively. 

To start to bridge mechanical responses between 

discrete PZT ceramic films and the macroscopic 

piezo-stack actuators, FEA modeling of the latter 

was initiated that will facilitate the effect of applied 

voltage on resulting stress in the piezo-stacks. 

Lastly, hardware (amplifier, loading platens, 

capacitance gages, data acquisition system, dead 

weight loading system, etc.) were acquired in 

preparation for piezostack testing that will enable 

concurrent mechanical loading and piezoactuation. 

This use of this equipment will play key roles in this 

project’s FY07 work. 

Results 

The calculated strength underwent scrutiny 

through comparisons of stresses predicted from 

finite element analysis (Fig. 5). Based on those 

results, it appears that the more classical calculation 

of failure stress [1] was overestimating the stresses 

as shown in Fig. 6. Strength-size-scaling 

relationships were developed and will be used to 

integrate strengths of these thin film PZT test 

coupons into the predicted performance of piezostack actuators comprising them. 

7
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 5. Finite element analysis was performed to 

accurately estimate “ball-on-ring” biaxial bending 

stresses in the thin PZT piezoceramic specimens. 

Figure 7. Weibull strength distributions for the 

5H4E and 5A4E piezoceramics. The same loading 

fixture was used to test both materials. The 5H4E 

specimens were approximately 50% thicker than the 

5A4E. 

Figure 6. Finite element analysis results showed 

that the classical formula for determining biaxial 

flexure stress overestimates actual stress. 

The 5A4E PZT dielectric had a higher strength 

than the 5H4E as illustrated in Figs. 7-8. The 

Weibull moduli of both materials were statistically 

equivalent. Preliminary optical fractography has 

showed failure initiation is at the surface indicating 

that a surface-type flaw or volume-type flaw located 

at the surface was the strength-limiter. 

Figure 8. 95% confidence ratio rings for 

characteristic strength and Weibull modulus 

A one-quarter FEA model was created of a 

piezo-stack to aid in our interpretations of the link 

between voltage and resulting stress. An example of 

that relationship is illustrated in Figs. 9-10. For 

structural ceramic components, the greatest concern 

for reliability compromise is the magnitude of the 

first principal tensile stresses. This project will 

examine the design the existing piezo-stack 

actuators and strive to make educated changes so to 

lessen those tensile stresses and improve their 

reliability. 

8
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 9. Voltage profile in an actuated piezo-stack. 

Figure 10. Resulting first principal stress field from 

the applied voltage field shown in Fig. 9. 

Conclusions 

FY06 was this project’s first year and a great 

deal was accomplished that will enable design 

optimization and reliability enhancements of 

ceramic piezo-stack actuators for heavy duty diesel 

engine fuel injectors. An ultimate goal is to link 

systematic micromechanical property data of the 

piezoceramic materials to mechanical performance 

of piezo-stack actuators. Toward that, in FY06 

micromechanical test facilities were established and 

are being used to measure strength and fatigue 

responses of the piezoceramics. The bridge between 

the micro and macro will be supported through 

modeling, and finite element models were developed 

to both validate stress fields in mechanical testing 

and stress fields in an actuated piezo-stack. 

Hardware was identified and procured that will 

enable the commencement of piezo-stack reliability 

analyses in FY07 and linking to property data of the 

PZT ceramic films. 

References 

[1] D. K. Shetty, A. R. Rosenfield, P. McGuire, 

G. K. Bansal, and W. H. Duckworth, “Biaxial 

Flexure Tests for Ceramics,” Ceramic Bulletin, 

59:1193-97 (1980). 

Presentations and Publications 

H. Wang and A. A. Wereszczak, “Effects of 

Electric Field on the Biaxial Strength of Poled 

PZT,” to be given at 31st International Cocoa Beach 

Conference and Exposition on Advanced Ceramics 

and Composites, Daytona Beach, FL, 21-26 January, 

2007. 

9
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

3. MATERIALS FOR EXHAUST AFTERTREATMENT 

A. Materials for Exhaust Aftertreatment 

Corey Shannon, Herbert DaCosta, Matt Stefanick, Ronald Silver, Tom Paulson, and 

Craig Habeger 

Caterpillar Inc. 

P.O. Box 1875 

Peoria, IL 61656 

(309) 578-6040 ; fax (309) 578-2953; e-mail: Paulson_tom@cat.com 

DOE Technology Development Manager: Jerry L. Gibbs 

 (202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@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. 

• Assess impact of phosphorus on oxidation (DOC), catalyzed DPF (CDPF) and urea-SCR catalysts, 

alone and in combination with sulfur and thermal aging. 

Approach 

• Assess impact of phosphorous (P) and sulfur (S) on aftertreatment systems by comparing fresh 

and aged catalyst performance using fuel burner and bench test system 

• For DOC: Evaluate P impact on CO and HC oxidation as indicated via lightoff temperature. 

• SCR: Evaluate impact of P and sulfur and /or thermal degradation on zeolite-based samples using 

temperature and NO2/ NOx ratio sweeps. 

• CDPF: Evaluate P impact on PM filtration efficiency and soot regeneration. 

Accomplishments 

• Determined the impact of phosphorus (P) exposure on HC and CO lightoff over oxicats to be 

comparable to the loss of NOx reduction over SCRs at similar exposure levels. 

• The deactivation of zeolite-based SCR catalysts due to combinations of phosphorus, sulfur and 

thermal exposure was identified. 

9 At low exposure temperatures (210°C), sulfur initially increases the observed 

deactivation over P alone for both NH3 and NO conversions; the effect is reversible upon 

exposure to sulfur free gas. 

9 At higher temperatures (450°C), the effect of sulfur on deactivation is less reversible. 

9 Sulfur initially adversely affects NH3 conversion more than NOx. Thermal exposure 

alone does not significantly affect SCR activity, but does inhibit SCR in combination 

with sulfur and phosphorus. Phosphorus alone or in combination with sulfur or thermal 

exposure blocks NH3 adsorption sites which in turn prevent NO reduction. 

9 Effects from exposure to phosphorus and sulfur can be mitigated by operating at higher 

NO2/ NOx ratios, or by increasing the catalyst substrate volume 

11
________________________________________________________________________ 

Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

• Determined that the main impact of combusted phosphorus on particulate filters is on the 

regeneration step that occurs over the catalyst.

Future Direction 

• Explore 1 or 2 mitigation strategies in more detail for larger catalysts; blank in front; wash catalyst 

to restore activity. 

• Evaluate the impact of non-combusted P on exhaust catalyst and compare it to the effect of 

combusted P. 

• Correlate results from lab aging studies with real world catalyst samples aged on an engine and or 

in the field. Use correlation to develop rapid aging protocols that accurately predict in-use 

deactivation modes. 

• Continue to identify novel DPFs with improved pressure drop, filtration efficiency, and 

regeneration capabilities.

Introduction 

The objective of this effort is to develop 

and evaluate materials that will be 

utilized in aftertreatment systems for 

diesel engine applications to comply 

with future government emission 

regulations. The materials include 

catalysts for NOx abatement and filtration 

media for particulate control in the 

exhaust system. The project is a part of 

Caterpillar strategy to meet EPA 

requirements for regulated diesel 

emissions in 2007 and beyond. This 

year’s focus is to assess durability of 

various catalyst technologies in terms of 

phosphorous, sulfur and thermal 

degradation, and to identify materials 

that have high diesel particulate filtration 

efficiency, low pressure drop, and 

comply with future emission regulations 

for heavy-duty diesel engines. 

Approach 

Catalyst Deactivation: The effort was 

focused on identifying the impact of 

phosphorus (P) exposure on diesel 

particulate filters (DPF) and diesel 

oxidation catalysts (oxicats) and of P 

plus sulfur (S) exposure on SCR catalyst 

formulations. The performance of fresh 

DPFs and catalysts was tested, and then 

each DPF or catalyst was exposed to P in 

a diesel fuel burner using fuel doped 

with tricresyl phosphate that mimicked 

the effect of P from combusted lube oil. 

In the case of the Selective Catalytic 

Reduction (SCR) catalyst, the diesel fuel 

was additionally doped with S to expose 

the sample to SO2. All aged samples 

were evaluated and compared to the 

analogous fresh samples, while the aged 

SCR catalysts were additionally 

compared to the activity of SCR samples 

exposed to P only. 

Results 

Oxicat durability: 

A set of 10g Pt oxidation catalysts 

(oxicats) were evaluated for CO and HC 

lightoff performance before and after 

aging via P exposure using a diesel fuel 

burner. The P had only a mild effect on 

the oxicat activity, with the worst impact 

being a 20o

C increase in CO lightoff 

after an exposure of 19g/L P. The 

lightoff curves shifted toward higher 

temperature with increasing P exposure 

without changing shape, which is typical 

of a decline in activity due to physical 

blockage of active sites. The results were 

compared to previous tests that 

measured the P effect on NOx reduction 

over Lean NOx Traps (LNTs) and SCR 

catalysts. CO and HC oxidation over the 

tested catalysts are affected by P to 

12
Heavy Vehicle Propulsion Materials 

about the same extent as NOx reduction 

over SCR catalysts. P impacts LNT 

performance more than either oxicats or 

SCRs. 

Results in Figures 1a and 1b respectively 

show the effect of P exposure on the CO 

and HC lightoff performance of a Pt 

catalyst. The lightoff temperature (T50, 

the temperature at which 50% 

conversion is obtained) of the catalyst 

increased by about 8oC for HC and 15oC 

for CO when exposed to 5.5g/L P. The 

lightoff temperatures increase with 

increasing exposure levels of P. A 

second catalyst core was exposed to CO conversion 

100% 

90% 

80% 

70% 

60% 

50% 

40% 

30% 

20% 

10% 

0% 

FY 2006 Progress Report 

3.6g/L P, and results are in line with 

those from the first core. All of the 

samples, which were aged in the lab on 

the diesel fuel burner, have lightoff 

curves that are similar in shape to the 

fresh benchmark. The shift in the curves 

with increasing P is typical of a decline 

in activity due to physical blockage of 

active sites. Most likely, the 

accumulated P blocks some of the sites 

where HC or CO is adsorbed. The actual 

P concentration that a catalyst will see is 

a function of the catalyst volume, the 

concentration of P in the oil, and the 

engine oil consumption rate. 

150 170 190 210 230 250 

Temperature (o

C) 

Fresh 

1.8 g/L P 

3.6 g/L P-alt core 

5.5 g/L P 

FIGURE 1a. Effect of P exposure level on CO lightoff over a Pt oxicat 

13
Heavy Vehicle Propulsion Materials FY 2006 Progress Report HC conversion

100% 

90% 

80% 

70% 

60% 

50% 

40% 

30% 

20% 

10% 

0% 

220 225 230 235 240 245 250 255 260 265 270 

Temperature (o

C) 

Fresh 

1.8 g/L P 

3.6 g/L P-alt core 

5.5 g/L P 

FIGURE 1b. Effect of P exposure level on HC lightoff over a Pt oxicat 

The impact of P on oxicats was 

compared to its effect on LNTs and 

SCR. First the temperature for 90% CO 

conversion over the fresh catalyst was 

obtained, then amount of CO converted 

at that same temperature for a P aged 

catalyst was determined. The procedure 

was repeated for each level of P 

exposure, and then again for the HC. 

Results for LNT and SCR catalysts were 

determined based on fresh vs. aged NOx 

reduction. Oxicat conversion of CO and 

HC is affected by P to about the same 

extent as NOx reduction over vanadia 

based, or zeolite based, SCR catalysts. 

The impact of P on LNTs is greater than 

that for oxicats or SCRs. In all cases, P 

is physically blocking adsorption sites. 

SCR Catalyst durability : 

Zeolite based urea SCR catalysts 

provided by Umicore Catalyst Company 

were previously mildly aged with P 

alone. Using a diesel fuel burner, 

additional core samples were taken from 

this catalyst and exposed to P plus 

significant levels of SO2, such that the 

resulting S weight exposure was equal to 

the P exposure. The activity of these 

aged samples was characterized by 

doing a temperature sweep, an NH3/ NOx 

ratio sweep, and a NO2/NOx ratio sweep 

over the aged catalyst. These results 

indicate the addition of S into the aging 

mix initially accelerates the degradation 

of zeolite SCR performance. 

Results in Figure 2 show the effect of 

exposure to 5g/L P and 5g/L S on the 

NOx and NH3 conversions, calculated by 

dividing the aged catalyst activity by its 

previously tested fresh performance 

under the same conditions. In this plot 

100% means the aged sample has 

activity equal to a fresh sample, not 

complete conversion. The relative 

performance of the sample aged using P 

only is also plotted in this figure. Results 

from mild aging with P showed little or 

no loss of NH3 and NOx conversion 

activity. The addition of S to the aging 

mix drops the relative performance by 

about 10 % for both the NOx and the 

NH3. The relative NOx conversions 

from both agings appear to be slightly 

lower at low temperatures, and then 

14
Heavy Vehicle Propulsion Materials 

increase with temperature. This seems 

to indicate that active sites for low 

temperature NOx reaction are more 

strongly impacted by P and S. An 

analysis of NO and NO2 conversions 

from these agings indicates that NO 

conversion is lost at low temperature. 

The decline in activity is due to a 

decrease in NO oxidation to NO2. In 

contrast, the activity loss for NH3 

conversion due to exposure to P, or P 

and S, does not seem to correlate with 

increasing temperature. 

The effect of core size was also 

investigated. A 1-inch diameter by 3­

inch long core from a zeolite SCR 

FY 2006 Progress Report 

catalyst was exposed to the same level of 

S and P as a 1 inch by 6-inch long core. 

The results are shown in Figure 3 and 

indicate the longer core has better low 

temperature NOx conversion, both fresh 

and after aging. In fact, the aged 6-inch 

core has a NOx lightoff temperature 

similar to the fresh 3-inch core. At high 

temperature, the activity of the 6-inch 

core drops below the performance of the 

3-inch core. This may be due to the 

longer residence time of NH3 in the 6­

inch core, which allows NH3 oxidation 

to occur earlier. The results suggest one 

method of reducing the impact of S and 

P exposure may be to increase the 

volume of the SCR catalyst. 

Aged NOx/ Fresh NOx conversion

100% 

90% 

80% 

70% 

60% 

300 350 400 450 500 550 600 

Temperature (°C) 

NOx- P only NOx - P + S 

NH3 - P only NH3 - P + S 

FIGURE 2. Relative performance of SCR catalysts after exposure to 5g/L P plus S or P 

only. 

15
Heavy Vehicle Propulsion Materials FY 2006 Progress Report NOx conversion

100% 

80% 

60% 

40% 

20% 

0% 

200 250 300 350 400 450 500 

Temperature (°C) 

Fresh 1 x 3 core 

5g/L P and S aged 1 x 3 core 

Fresh 1 x 6 core 

5g/L P and S aged 1x6 core 

FIGURE 3. Comparison of NOx conversion as a function of space velocity, before and 

after exposure to 5g/L P and S. 

The relative activity of zeolite based 

urea SCR NOx reduction catalyst after 

exposure to higher levels of P and S is 

shown in Figure 4. As seen also in 

Figure 2, catalysts that were exposed to 

5g/L P and S at 210oC showed a slight 

degradation in NOx conversion relative to 

the degreened activity. Increasing the 

exposure level at that temperature to a 

total of 10g/L P and S resulted in an 

initial apparent loss of 70% of the 

degreened activity at 300oC. A repeat of 

the activity test showed a significant 

recovery in NOx conversion, such that the 

activity loss at 300o

C was just 20%. For 

comparison, this amount of activity was 

lost from exposure to 10g/L P only. 

Activity testing was performed in the 

absence of S and P, so it appears that 

exposing the catalyst to S-free gas 

removes most of the S that was 

responsible for the initial activity loss. 

The resulting “cleaned” catalyst 

performs similar to a sample that was 

just exposed to 10g/L P. The cleaning 

effect can even be observed during the 

first activity test, where there is 

significant recovery in activity as the 

sample is ramped up in temperature. 

16
Heavy Vehicle Propulsion Materials FY 2006 Progress Report Aged Conversion / Fresh Conversion

100% 

90% 

80% 

70% 

60% 

50% 

40% 

30% 

20% 

10% 

0% 

300 350 400 450 500 550 600 

Temperature (°C) 

5g P+S 

10g P only 

10g P+S 

10g P+S 2nd run 

10g P+ S, thermally aged 

FIGURE 4. Relative NOx performance of SCR catalysts after exposure to 10g/L P, 5g/L 

and 10g/L P and S and after 10g/L P and S plus thermal aging for 24h at 650°C. 

A drop in initial SCR NOx conversion 

activity at 175°C and 400°C as a 

function of exposure level to P and S 

can be seen in Figure 5. The results in 

this figure show the decline in activity 

accelerates with increasing exposure 

amounts. It may be that the Fe-zeolite 

catalysts have a buffering capacity for up 

to about 5g/L P and S exposure, but after 

that the active sites become 

overwhelmed. The figure also indicates 

the difference between the initial activity 

loss due to exposure to 10g/L P and S is 

about 10 to 15% greater than that due to 

exposure to a similar amount of P alone. 

If the NOx activity measured in the 

second and subsequent runs are plotted 

instead of the initial decline in activity 

due to P and S, the resulting curve is 

much flatter and the data point at 10 g/L 

exposure falls on top of the P only data 

point. However, even if the data from 

the repeat runs are utilized, the decline in 

NOx activity between 0 and 5g/L is less 

than the decline in going from 5 to 10g/L 

P and S exposure. 

Similar relative performance results for 

ammonia conversion are observed in 

Figure 6. Typically NH3 is converted 

either by reacting with NOx or with O2 

under SCR reaction conditions. Just as 

with NOx in Figure 4, the initial NH3 

conversion drops significantly when the 

P and S exposure level is increased from 

5g/L to 10g/L. In this case, sweeping 

the catalyst with S free gas recovers 90% 

of the initial activity, or results in a loss 

of just 10% of degreened NH3 

conversion. The recovered activity also 

corresponds to the effect of a P-only 

exposure of 10g/L. 

17
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

100%

90%

SCR NOx conversion 

80% 

70% 

60% 

50% 

40% 

30% 

20% 

10% 

0% 

0 2 4 6 8 10 12 

g/L P and S Exposure 

175°C 

400°C 

175°C, P only 

400°C, P only 

FIGURE 5. Impact of exposure level on initial NOx activity at 175°C and 400°C 

100% 

Aged Conversion / Fresh Conversion

90% 

80% 

70% 

60% 

50% 

40% 

30% 

20% 

10% 

0% 

300 350 400 450 500 550 600 

Temperature (°C) 

5g P+S 

10g P only 

10g P+S 

10g P+S 2nd run 

10g P+ S, thermally aged 

FIGURE 6. Relative NH3 conversion of SCR catalysts after exposure to 10g/L P, 5g/L 

and 10g/L P and S and after 10g/L P and S plus thermal aging for 24h at 650°C 

The 10g/L P and S exposed catalyst was sample to S free gas, or by repeated 

then thermally aged at 650oC for 24 testing. The effect of the thermal aging 

hours in a mixture of 10% steam and air. on NH3 conversion (shown in Figure 6) 

The resulting loss of NOx conversion, is less dramatic, resulting in a total 

shown in Figure 4, was about 40% of the decline of just 20% of degreened 

degreened performance or about the activity. This level is not much different 

same effect as that of the P and S than the effect of P and S exposure 

exposure. This loss of activity did not alone. The difference in temperature 

appear to be recoverable by exposing the effect may be due to sintering of the 

18
Heavy Vehicle Propulsion Materials 

active metal (e.g. iron or copper) due to 

the hydrothermal aging. The active 

metal seems to play an important role in 

activating the NOx, but may not be as 

critical in converting NH3. 

Comparing Figures 4 and 6, the relative 

NOx performance loss appears to be 

slightly higher than the NH3 conversion 

loss for any given P and S exposure. 

This may indicate that the catalyst 

oxidation of NO to NO2 is more 

susceptible to activity loss due to 

thermal aging or P and S exposure than 

is the De NOx reaction, which converts 

equal molar amounts of both NH3 and 

NOx 

FY 2006 Progress Report 

Figure 7 shows the NOx activity as a 

function of temperature after thermally 

aging a fresh sample for 24h at 650°C. 

The results are compared to the repeat 

test of the sample exposed to P and S, 

which was then thermally aged. The 

catalyst that was only thermally aged 

shows very similar activity to the fresh 

sample, indicating little or no loss of 

performance. In contrast, there is about 

a 15% performance decline after the P 

and S exposed sample is thermally aged 

under the same conditions. This 

suggests that exposing the P and S 

contaminants to an elevated temperature 

can cause them to adversely interact with 

other elements of the catalyst surface, 

resulting in an increased loss of activity. 

NOx conversion

100% 

90% 

80% 

70% 

60% 

50% 

40% 

30% 

20% 

10% 

0% 

150 200 250 300 350 400 450 500 550 

Temperature (°C) 

Fresh Baseline 

Thermal sample 6 aged 

P+S sample 3, 10g/L,30K r2 

P+S sample 3, 10g/L +thermal, 30K 

FIGURE 7. Effect of thermal aging on fresh and P + S exposed catalyst. 

The relative activity (aged conversion/ 

fresh conversion) of these same catalysts 

at 190°C and an NH3/ NOx ratio of 

1.1was plotted versus a varying NO2/ 

NOx ratio in Figure 8. In this case, the 

sample that was only thermally aged 

retains >95% of the fresh catalyst’s 

performance. The sample exposed to 

10g/L P and S saw a decline of more 

than 50% off the baseline at low NO2 / 

NOx ratios, but performance improved to 

90% of the fresh sample when the NO2 / 

NOx ratio increased to between 0.25 and 

0.65. A similar improvement in activity 

was seen over the thermally aged, P and 

S exposed sample, which was 65% less 

than the baseline at low NO2 / NOx ratios 

but improved to 90% of the fresh 

performance at a ratio of 0.5. At a NO2 / 

NOx ratio of 0.5, the NO2 level is equal to 

19
Heavy Vehicle Propulsion Materials 

the amount of NO, and is the condition 

where “fast SCR” can occur. The fact 

that there is only a minor decline in 

activity at this point, and despite the 

temperature being just 190°C, indicates 

the P and S primarily impact NO 

oxidation or its interaction with NH3. 

The relatively uniform decline after 

thermally aging the P and S exposed 

sample suggests that some NH3 Aged NOx / Fresh NOx conversion

100% 

90% 

80% 

70% 

60% 

50% 

40% 

30% 

20% 

10% 

0% 

FY 2006 Progress Report 

adsorption might be adversely affected 

by heating these contaminants, possibly 

by physically blocking more of the NH3 

adsorption sites. The relatively mild 

decline in activity at 0.5 NO2 / NOx ratio 

even after significant aging also suggests 

another mitigation strategy may be to try 

to run closer to this condition as the 

system ages. 

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 

NO2 / NOx 

P+S 10g aged sample 

Thermal aged sample 

P+S 10g + thermal aged 

sample 

FIGURE 8. Relative NOx performance of SCR catalysts after exposure to P plus S and 

after exposure to thermal aging only. 

Figure 9 compares the effect of aging the 

catalyst with 5g/L P and S at 450°C 

versus aging at 210°C, in order to see the 

impact of exposure temperature. 

Increasing the exposure temperature 

appears to decrease the relative activity 

by about 8% across the range of the 

temperature sweep. The other effect 

seems to be that the activity loss due to 

sulfur after the higher temperature 

exposure is not as readily reversible as it 

is when the temperature is lower. It is 

also seen that relative NH3 performance 

is better than NOx in most cases. 

However, the gap between the high and 

low temperature exposure activity loss 

seems slightly higher for NH3 than it 

does for NOx. Again, this may be due to 

the higher exposure temperature leading 

to more significant interference of the P 

and S with NH3 adsorption. 

20
Heavy Vehicle Propulsion Materials FY 2006 Progress Report Relative NOx/ NH3 conversion

0% 

10% 

20% 

30% 

40% 

50% 

60% 

70% 

80% 

90% 

100% 

300 350 400 450 500 550 600 

Temperature (°C) 

P+S NOx Aged 210C 

P+S NOx Aged 450C 

P+ S NH3 Aged 210C 

P+ S NH3 Aged 450C 

FIGURE 9. Effect of P and S exposure temperature on NOx and NH3 relative conversion. 

Diesel Particulate Filter (DPF) 

Catalyzed and uncatalyzed diesel 

particulate filters (DPF) obtained from 

suppliers were loaded with particulate 

matter on a diesel fuel-burner bench and 

regenerated. This process was repeated 

after 100 hours and after 200 total hours 

of hydrothermal aging. In all cases, the 

filtration efficiency did not change 

significantly. One trend found in this 

study was that the pressure drop during 

the tests decreased slightly on catalyzed 

samples (Figures 10 and 11). However, 

it could be due to extended degreening. 

Several projects have been executed to 

improve the performance and efficiency 

. 

of the diesel fuel burner on which the 1-

inch diameter DPF cores are tested. One 

such project was a 6-sigma “lean” event. 

A team composed of primary users and 

customers of the test bench examined the 

entire process of loading particulate 

matter onto a DPF. This team was able 

to suggest changes to the procedure as 

well as physical changes on the bench 

itself. Many safety issues were also 

resolved. Another 6 sigma-like project 

was carried out and determined the 

reproducibility of the measurements 

acquired from the fuel burner bench. 

This project ascertained that the 

filtration efficiency of a DPF can be 

measured with less than 2% error 

21
14 

Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

12 

0 5 10 15 20 25 30 35 

Air Flow (L/min) 

Run 1 

Run 2, aged 100 hours 

Run 3, aged 200 hours 

Backpressure (kPa) 10 

8 

6 

4 

2 

0 

FIGURE 10. Pressure drop versus flow rate for catalyzed cordierite before and after 

hydrothermal aging for 100 h and 200 h. 

3 

Backpressure (kPa)2.5 

2 

1.5 

1 

0.5 

0 

0 5 10 15 20 25 30 35 

Air Flow (L/min) 

Run 1 

Run 2, aged 100 hours 

Run 3, aged 200 hours 

FIGURE 11. Pressure drop versus flow rate for catalyzed alternative DPF, before and 

after hydrothermal aging for 100 h and 200 h. 

22
Heavy Vehicle Propulsion Materials 

DPF P durability 

A P deactivation study on catalyzed 

cordierite and the catalyzed alternative 

DPF was done. The catalyzed 

alternative DPFs were exposed to P for 

200 and 400 minutes, corresponding to 

8g/L and 16g/L of P exposure. Figure 12 

presents results for the effect of P 

exposure on the filtration efficiency of 

blank cordierite, catalyzed cordierite, 

100% 

96% 

92% 

88% 

84% 

80% 

0 4 

Filtration efficiency 

FY 2006 Progress Report 

and catalyzed alternative substrate. The 

results showed that the largest decline in 

efficiency occurred over the blank 

cordierite. The two catalyzed samples 

had similar losses up to about 7g/L P 

exposure, but the catalyzed alternative 

substrate declined more slowly than the 

catalyzed cordierite at P levels above 

that. 

8 12 

g/L P exposure 

FIGURE 12. The effect of P exposure level on filtration efficiency over blank cordierite 

(#), catalyzed cordierite (!) and catalyzed alternative substrate (%). Error bars are shown 

at 8 g / L P exposure. 

P exposure slightly affected the filtration 

efficiency but had no significant effect 

on pressure drop. The largest impact of 

the P exposure was on soot regeneration. 

Figure 13 shows the concentration of CO 

+ CO2 released during soot regeneration 

as a function of temperature at different 

levels of P exposure for catalyzed 

cordierite. Figure 14 shows similar 

results for the catalyzed alternative 

substrate and includes additional results 

for hydrothermal aging. . The fresh 

samples each show two peaks, with the 

lower temperature peak being larger than 

the high temperature peak for the 

alternative substrate and the opposite 

result obtained over the cordierite. The 

lower temperature peak is thought to be 

23

16 
Heavy Vehicle Propulsion Materials FY 2006 Progress Report

associated with passive regeneration of levels continue to increase up to 16g/L, 

soot, and the higher temperature peak is the regeneration peak shifts all the way 

likely associated with active to the point where it is the same as the 

regeneration. As the samples are peak temperature found over blank 

exposed to increasing levels of P, at first cordierite, suggesting the P has 

the passive regeneration peak declines in completely deactivated the catalyst 

area and the CO + CO2 emissions shift activity for soot regeneration. 

toward the high temperature peak. As P 

400 

300 

Peak 2

200 Peak 1 

100 

0 

200 300 400 500 600 700 

Temperature (o

C) 

FIGURE 13. Concentration of CO + CO2 from soot regeneration over catalyzed 

cordierite DPF as a function temperature for different levels of P exposure: ()) Fresh, 

uncatalyzed; (') Fresh catalyzed; (() 8 g / L P aged catalyzed; (!) 16 g / L P aged 

catalyzed. ppm CO + CO

2

24 
400 

Heavy Vehicle Propulsion Materials FY 2006 Progress Report

Peak 1 Peak 2

300

200 

100 

0 

200 300 400 500 600 700 

Temperature (o

C) 

FIGURE 14. Concentration of CO + CO2 from soot regeneration over alternative DPF 

substrate material as a function temperature for different aging protocols: ()) Fresh, 

uncatalyzed ; (+) Fresh catalyzed; (& ) Thermally aged catalyzed; (,) 8 g / L P aged ppm CO + CO2 

catalyzed; (%) 16 g / L P aged catalyzed. 

Conclusions 

Results from the catalyst deactivation 

study over oxidation catalysts showed 

that exposure to 5g/L P reduces 

conversion of CO and HC over these 

samples by about 10%, which is about 

the same extent as it reduces NOx

reduction over SCR catalysts. The 

impact of P on LNTs is greater than that 

for oxicats or SCRs. In the case where P 

exposure is combined with S, the initial 

loss of activity proved to be reversible 

simply by repeating the evaluation test in 

S-free gas, suggesting that at an 

exposure temperature of 210°C the 

sulfur is a reversible contaminant. The 

addition of mild thermal aging was 

shown to be detrimental only when P 

and S were already on the surface. 

Exposing the catalyst to P and S at 

higher temperatures also increased the 

loss of performance. Based on previous 

work and work cited in the literature, it 

is thought that P impacts these catalysts 

by physically blocking adsorption sites. 

Regarding DPFs, thermal aging had a 

minor effect on the filtration efficiency 

and pressure drop, while P aging 

drastically affected the regeneration 

capability of the filters. 

Travel

1. Ron Silver traveled to Detroit, 

MI, to attend the 2006 SAE 

World Congress. 

2. Herbert DaCosta traveled to 

Detroit, MI, to attend the 2006 

DEER meeting. 

3. Herbert DaCosta traveled to 

Boston, MA, to attend the 2006 

COMSOL meeting. 

25 
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Presentations/Publications 

H. DaCosta C. Shannon, and R. Silver, 

“Thermal and Chemical Aging of Diesel 

Particulate Filters,” submitted to the 

2007 SAE World Congress. 

H. DaCosta, C. Shannon and R.G. 

Silver, “DPF Durability – Bench 

Studies on Cordierite Filters,” oral 

presentation at the 2006 DEER meeting, 

Chicago, IL (2005) 

26
____________________ 

Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

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

Microstructures and Deactivation Mechanisms 

L. F. Allard, D. A. Blom, C. K. Narula, S. A. Bradley*and M. Jose-Yacaman** 

Oak Ridge National Laboratory 

P.O. Box 2008, MS-6064 

Oak Ridge, TN 37831-6064 

*UOP LLC, Des Plaines, IL 60017 

** University of Texas-Austin, Austin, TX 78712 

DOE Technology Development Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

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

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Objectives 

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

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

emissions in diesel and automotive exhaust systems. 

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

support materials (typically oxides). 

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

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

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

microscope (ACEM). 

Approach 

• Utilize Oak Ridge National Laboratory ACEM to characterize the atomic morphology and behavior of heavymetal species such as platinum and rhenium on oxide support materials (alumina, silica and titania) via 

techniques of annular dark-field (ADF) or “Z-contrast” imaging. 

• Utilize ACEM ADF imaging and techniques of spectrum imaging to characterize the structure and chemistry of 

a series of model bimetallic nanoparticle comprising controlled compositions in the platinum/gold system. 

Accomplishments 

• Characterized the nature of platinum catalyst structures on alumina, silica and titania support materials, 

including the imaging of platinum species dispersed as single atoms and in precrystalline clusters known as 

“rafts.” In particular, the different behavior of the Pt atoms, clusters and rafts on the support materials was 

characterized, as a function of electron beam excitation. 

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

shapes of nanoparticles and clusters. 

• Tested and demonstrated the functionality of a new catalyst reactor specimen rod for before-after studies of 

catalyst reactions. 

27
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Future Direction 

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

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

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

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

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

• Perform precisely controlled thermal experiments using ex-situ catalyst reactor to understanding the interaction 

of Pt atoms with catalyst supports, which is critical for the development of new and improved catalysts. 

Technical Progress 

Imaging of ultrafine atomic clusters of heavy 

metals on low atomic number supports has been 

demonstrated to provide a new level of capability for 

characterizing catalytic materials. The advent of 

research on the HTML’s new aberration-corrected 

electron microscope (JEOL 2200FS-AC “ACEM”), 

has given us the ability to image the locations of 

heavy metal atomic species at the single-atom level, 

and to observe the different behaviors of, for 

example, Pt atoms and clusters dependent on the 

nature of the support materials, under the influence 

of the incident electron beam. Using the technique 

of high-angle annular dark-field imaging (HA-ADF) 

to image the heavy atoms in high contrast also 

allows the characterization of the structures of 

nanoparticles in the 1-3 nm size range, which 

represent structures found on typical “real” catalyst 

materials such as those used in commercial 

automotive systems. In this report, we show 

examples of the structures and behaviors of Pt 

species on alumina, silica and titania supports, 

which gives information on the different metalsupport interactions of these systems. The analysis 

of the structure of discrete bimetallic Au/Pd particles 

suggesting the nature of “active sites” on the 

particles surfaces is also illustrated. The latter 

samples are part of a specimen suite provided by 

Prof. M. Jose-Yacaman of the University of Texas at 

Austin. The former samples, discussed first, 

represent work with colleagues S. A. Bradley and 

W. Sinkler of UOP Co. Pt on titania samples were 

provided originally by colleague C.K. Narula, in 

prior work when Dr. Narula was working at Ford 

Research Laboratory. 

Imaging Pt on oxide supports: Although the 

preparation and properties of heterogeneous 

catalysts have been studied for decades, it is not well 

understood how fine dispersions of heavy metal 

species in “as-prepared” catalysts are arranged on 

common supports such as γ-alumina, titania, or 

amorphous silica. Aging or calcining treatments 

typically cause atomic level dispersions to coalesce 

into discrete particles of sizes >2 nm, which have 

crystal lattices that can be imaged using 

conventional high-resolution bright field imaging 

techniques. The present study, however, illustrates 

the advanced capability of aberration-corrected 

imaging to characterize catalysts comprising Pt on γAl2O3 and amorphous SiO2; systems which have 

different metal-support interactions. To study the 

effect of the supports on metal cluster behavior, 

catalysts were prepared by impregnating the support 

materials with 0.35 wt. % Pt, and oxidizing the 

material at 525oC. The Pt on γ-Al2O3 was then 

reduced at 500oC in H2. Samples were examined in 

ORNL’s JEOL 2200FS FEG ACEM, fitted with a 

hexapole aberration-corrector (CEOS GmbH, 

Heidelberg, Germany) for the probe forming optics, 

which provides a nominal probe diameter of 0.07

nm (0.7Ångström) [1]. Goals of the investigation 

were to determine the number of atoms in a cluster 

and the structure of the cluster (or raft), investigate 

metal-support interactions, and by using an ex-situ 

reactor, to understand the formation of discrete 

metal clusters as a result of reduction of the oxidized 

catalyst. 

Figure 1 shows an example of oxidized Pt on γAl2O3 at increasing magnification. The majority of 

the Pt species are in the form of single atoms, dimers 

and some small clusters. These are seen in greater 

detail in Fig. 2. No discrete particles or rafts were 

seen on this sample. However, after a hydrogen 

reduction treatment at 500°C in the ex-situ reactor, 

the Pt was observed to form rafts of 1-2 layers in 

28
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 1. (a) Oxidized platinum species on γ- Al2O3; no large discrete particles are seen. b) higher 

magnification shows Pt dispersed in single atoms or dimers. 

Figure 2. Detail of the central area on Fig. 1b. Only single atoms or dimers are present. 

thickness, with some single atoms still present. It 

was interesting to note that, under the influence of 

the electron beam (which can be argued as 

representing effects similar to direct heating of the 

sample), the rafts tended to re-arrange their 

structures, but single atoms remained pinned to the 

support. Figure 3 shows a pair of images of the 

reduced sample, taken from a sequential series of 

images of the same area, several scans apart. The 

raft of atoms shows some changes in structure (but 

no dramatic change in size or structural 

rearrangements), while 3 single atoms remain 

clearly pinned in position. The oxidized Pt on silica 

sample was different from the equivalent Pt on γAl2O3 sample, in that it showed rafts similar to the 

reduced Pt on alumina, but also numerous single 

atoms, dimers and small clusters, as seen in Fig. 4. 

These rafts were very mobile under observation, and 

showed a tendency to sinter together into larger 

agglomerates under beam conditions similar to those 

used for the Pt/alumina sample. This suggests a 

stronger metal-support interaction in the Pt/alumina 

system compared to the Pt/silica system. 

29
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 3. HA-ADF image pair from separate scans of Pt species on g alumina, after a reduction at 500°C in 

H2-Ar mixture. The raft represents a collection of 25-30 atoms, but 3 isolated atoms are also seen 

(circled). While the atoms in the raft are somewhat mobile under incident beam irradiation, the 

single atoms remain pinned to sites on the alumina surface. 

Figure 4. HA-ADF image pair from separate scans over rafts of Pt on amorphous silica. 

Insets show areas from which respective colorized intensity perspective views 

were taken. The mobility of the Pt species on the surface of the silica results in 

the conjoining of separate reafts into one larger aggregate of Pt atoms. 

30
Heavy Vehicle Propulsion Materials 

In the Pt/TiO2 (rutile) system, for a sample of 

1% Pt reduced under conditions similar to the 

alumina and silica samples, the Pt species were 

observed to form small clusters with some rafts on 

the order of 1 nm in size, with also a significant 

fraction of individual atoms present. The Pt rafts 

tended to show alignment of atoms within the raft, 

FY 2006 Progress Report 

and in line with the underlying rutile structure. 

Figure 5 is an example of the analyis of a Pt raft. 

The intensity profile over a line of Pt atoms 

suggests that the brighter atoms represent 2 atoms 

aligned directly atop one another, in comparison to 

an adjacent single atom at 1⁄2 the intensity. 

a 

b 

c 

Figure 5. (a) Pt species deposited onto TiO2 (rutile) support material. Note regular underlying lattice image 

contrast from the rutile structure. Inset (b) shows one raft with atoms aligned along rutile planes. 

An intensity profile (c) as indicated in (b) suggests the bright triangle of atoms results from an 

essentially perfect alignment of atoms on a second layer, directly above the first layer atoms, giving 

twice the intensity in the HA-ADF image. 

A final example of our capability to study the about 1 nm in size after reduction. There are small 

behavior of these catalyst systems before and after changes in the underlying alumina structure, but it 

reaction treatments is shown in Fig. 6. This image is easy to recognize that the areas are essentially 

pair shows the same sample area of Pt on γ-Al2O3 equivalent. This remarkable capability will be 

at medium magnification, before and after a used routinely in further work on these catalyst 

reduction treatment. There is almost a single-atom systems. 

level dispersion of Pt species prior to reduction, 

and these atoms have agglomerated into rafts 

31
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 6. a) “Before” HA-ADF image of oxidized Pt species on γ- Al2O3, showing high dispersion of Pt. b) 

Identical area “After” reduction treatment in ex-situ reactor. Pt aggregates into small clusters and 

rafts 0.5-1.0 nm in size. 

Imaging Au/Pd nanocrystals: It has been 

known for several years that the addition of a 

second metal into monometallic nanoparticles 

significantly changes some of their properties, e.g. 

optical properties, catalytic properties, electronic 

properties, etc. [2,3]. When two or more different 

elements compose one nanoparticle, the 

differences in atomic radii are going to influence 

the way the atoms accommodate. However, very 

little is known about the crystal structure of 

bimetallic nanoparticles in the size regime from 

0.5 to 2 nm [4]. For bimetallic nanoparticles 

composed of a noble metal and a transition metal, 

recent experimental results indicate that an 

icosahedral phase structure is preferred over the 

others [5]. In the present work we are conducting a 

complete study of the structure, size, shape, 

optical response and catalytic properties of single, 

bimetallic and multimetallic nanoparticles. 

Figure 7 shows an aggregate of minimal 

Au1Pd5 nanoparticles which show significant 

variations in intensity in atomic columns, e.g. as 

indicated in the inset. The particle is near a <100> 

projection. Intensity profiles over several single 

atoms on the adjacent thin carbon film indicated 

intensity levels of ~900 and ~500 units, suggesting 

that these represented Au and Pd atoms, 

respectively. The profile over the row in the inset 

suggests that the column with the highest intensity 

must contain one or more Au atoms. The observed 

intensity, if considering only Pd atoms in the 

column, would indicate a column thickness at least 

3 atoms greater than surrounding columns, an 

energetically unfavorable situation. Results such 

at these are at the leading edge for analysis of the 

structures of bimetallic nanoparticles, but 

correlation to ADF image and DFT structure 

calculations will provide more reliable 

interpretation as we continue to develop our 

imaging expertise with the sub-Ångström 

resolution of the ACEM. 

32
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 7. a) HA-ADF image of particle with nominal Au1-Pd1 composition. B) Density Functional Theory 

computation of Au1Pd1 nanoparticle, originally created as a perfectly ordered multiply-twinned particle, then 

heated to 1000K and cooled to RT. The latter disordered configuration is a reasonable match to the observed 

particle, with numerous irregular surface sites. 

a b 

Figure 8. (a) Aggregate of Au1Pd5 nanoparticles in 2-4 nm size range. Inset shows variation in intensity of atomic 

columns; b) intensity profile over atom row arrowed in inset. Profiles of individual atoms labeled 1 and 2 

suggest the intensity associated with Au and Pd respectively. Central column intensity in (b) is too great to 

represent only Pd atoms; therefore one or more Au atoms must be present. 

33
Heavy Vehicle Propulsion Materials 

References 

1. M. A. O’Keefe, L. F. Allard, S. J. Pennycook and

D. A. Blom, “Transcending the One-Ångström 

Atomic Resolution Barrier in the TEM,” Microsc 

Microanal 12(Supp 2), Cambridge University Press, 

p. 162-163 (2006) 

2. A.M. Venezia, V. La Parola, V. Nicolì and G. 

Deganello, J. Catalysis 212, 56-62 (2002) 

3. L. Guczi, et al., Surface and Interface Analysis, 

34(1) 72-75 (2002) 

4. J.L. Rodriguez-Lopez J.L., et al., Physical 

Review Letters 92(19) 2004 

5. M. Jose-Yacaman, et al., J.Vac. Sci. & Tech. B, 

19(4) 1091-1103 2001 

Presentations and Publications 

D. A. Blom, S. A. Bradley, W. Sinkler and L. F. 

Allard, “Observation of Pt Atoms, Clusters and 

Rafts on Oxide Supports, by SubÅngström ZContrast Imaging in an Aberration-Corrected 

STEM/TEM,” Microsc Microanal 12(Supp 2), 

Cambridge University Press, p. 50-51 (2006) 

M. José-Yacamán, S. Mejía-Rosales, E. PérezTijerina, D.A. Blom and L.F. Allard, “Imaging AuPd Nanoparticles with the Aberration-Corrected 

STEM/TEM,” Microsc Microanal 12(Supp 2), 

Cambridge University Press, p. 772CD (2006) 

D. A. Blom, L. F. Allard, S. Mishina and M. A. 

O’Keefe, “Early Results from an AberrationCorrected JEOL 2200FS STEM/TEM at Oak Ridge 

National Laboratory,” Microsc Microanal 12, 

Cambridge University Press, p. 483-491 (2006) 

M. José-Yacamán, S. Mejía-Rosales, E. PérezTijerina, D. A. Blom and L. F. Allard, “Catalyst 

Characterization Coupling Cs-Corrected High-Angle 

Annular Dark-Field Imaging with Molecular 

Dynamics Calculations,” ICM16, Sapporo, Japan, p. 

1554 (2006) 

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

“Analysis of Microstructural Changes in Lean NOx 

Trap Materials Isolates Parameters Responsible for 

Activity Deterioration,” SAE 2006-01-3420, 

Dearborn, MI (2006). 

FY 2006 Progress Report 

34
______________________ 

Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

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 6054 

Oak Ridge, TN 37831-6054 

(865) 576-2264; fax: (865) 574-1249; e-mail: westdl@ornl.gov 

*Scientific Research Laboratory, Ford Motor Co. 

DOE Technology Development Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

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

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

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

Subcontractor: Ford Motor Company, Detroit, Michigan 

Objectives 

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

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

from ∼10–200 ppm at oxygen levels from 5 to 20 vol %. 

• Ensure that sensors are selective for either NO (insensitive to NO2), or able to measure “total NOx” (combined 

NO + NO2). 

Approach 

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

yttria stabilized zirconia (YSZ)] substrates. 

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

electrical bias. 

• Characterize sensor in bench setting for NOx sensor response, sensitivity to varying [O2] and the interferents 

propylene (C3H6), CO, and stability under long-term operation. 

Accomplishments 

• Developed sensing elements with improved long-term stability by modifying electrode architecture. 

• Sensing elements displayed near “total NOx” behavior. 

• Sensitivity to C3H6 and CO was observed, indicating an oxidation catalyst will be required in the sensor. 

Future Direction 

• Dynometer testing at Ford Motor Co. 

35
Heavy Vehicle Propulsion Materials 

Introduction 

Environmental stewardship is placing increasing 

demands on emission control of vehicle exhausts. 

For spark ignited, direct injection (SIDI) engine 

exhausts, these demands are largely being met with 

the use of a “three-way catalyst” (TWC), capable of 

simultaneously removing carbon monoxide (CO), 

hydrocarbons (CyHz), and oxides of nitrogen (NOx). 

Unfortunately, the current generation TWC is only 

effective for all three of these pollutants within a 

narrow range of oxygen concentrations ([O2]), losing 

its effectiveness for NOx reduction in the presence of 

excess O2 (Fig. 1). Thus existing TWCs cannot be 

used for NOx remediation of the O2-rich exhausts 

from diesel engines. 

Figure 1. Three-way catalyst (TWC) efficiency.1 As 

mair/mfuel is increased through ~14.6 ("stoichiometry") the 

O2 content of the exhaust will increase by several orders 

of magnitude. This increased O2 content strongly inhibits 

reduction of NOx to N2 and O2. 

Two technologies proposed for NOx remediation 

of diesel exhausts are the lean-NOx trap (LNT), and 

selective catalytic reduction (SCR). Both of these 

approaches require on-board NOx sensors, to control 

trap regeneration in the case of LNT or reagent 

injection in the case of SCR. 

The “NOx” in combustion exhausts is typically a 

mixture of nitric oxide (NO) and nitrogen dioxide 

(NO2), with NO predicted to be the dominant 

equilibrium species above ~500 o

C2. However, 

equilibrium conditions cannot be assumed for the 

exhaust and the relative abundances of NO and NO2 

in the exhaust rarely are known with certainty. In 

FY 2006 Progress Report 

such a situation it will be useful to have a “total 

NOx” sensor that possesses equal response 

characteristics to either NO or NO2. Such a sensor 

would measure [NO] + [NO2].

 In this project ORNL is developing “total NOx” 

sensing elements that can be used in diesel engine 

exhausts. The minimum requirement for this 

application is an operating temperature of 600– 

700oC and the ability to measure 10–200 ppm NOx. 

The project is a cooperative research and 

development agreement (CRADA) with Ford Motor 

Company. 

Approach 

The approach selected is to build a sensing 

platform on yttria-stabilized zirconia (YSZ). YSZ, 

an oxygen-ion conducting material, is widely used in 

oxygen sensors for gasoline engines. Two other 

NOx sensors based on YSZ are also near 

commercialization but these two designs have 

shortfalls; either poor sensitivity at low NOx or 

reliance on conversion of the NOx to NO2 before 

sensing. In addition, both are relatively complex 

with multicavity constructions. This construction is 

required because in the first design3 preliminary O2 

removal from the sampled exhaust is required before 

the NOx is measured amperometrically. In the 

second design4 the sampled exhaust must pass over 

a conversion electrode in order to convert the NOx 

into NO2 before detection in the sensing cavity. 

Most of the work at ORNL has focused on the 

sensing element portion of the sensor (the portion of 

the sensor that responds to the presence of NOx), the 

CRADA partner (Ford) has been focusing on the 

packaging and electronics required to make a 

working sensor. ORNL fabricates sensing elements 

by screen-printing electrodes onto YSZ substrates 

and subsequently firing. The elements are then 

tested for sensing performance using a furnace to 

simulate elevated temperature service and a 

commercial gas mixing unit to deliver blends of N2, 

O2, NOx (either NO or NO2), and other species as 

required. 

Results 

During FY 05 "all-oxide" total NOx sensing 

elements were developed and characterized. As 

indicated in the previous year's report two main 

challenges in bench-top testing remained. First, the 

36
Heavy Vehicle Propulsion Materials 

long-term stability of the elements had to be 

improved (the "background" signal drifted) and the 

response to potential interferents such as CO and 

hydrocarbons needed to be characterized. 

In order to improve the long-term stability of the 

sensing elements the electrode architecture was 

altered to improve the electronic conductivity. This 

alteration involves partially overcoating the 

electrodes with a noble metal. This will cause small 

increase in materials and manufacturing cost, but the 

CRADA partner felt that reduction/elimination of 

sensor drift could readily justify a small cost 

increase in the sensing element. 

Figure 2 illustrates the stability of a sensing 

element with the improved electrode architecture. 

Over approximately one day of continuous 

operation, the output from the sensor only changes 

by about 1%. This is much smaller than the changes 

induced by ~100 ppm NO2 (also shown in Fig. 2) 

and shows that the drift has been reduced to a more 

manageable level. 

Figure 2. Drift of signal with time for sensing element 

with improved electrode architecture 

As mentioned previously, the NOx in diesel 

exhaust is a mixture of NO and NO2, with the 

relative amount of each typically unknown. 

Therefore it is important that the sensing element 

yield approximately the same response irrespective 

of the [NO]:[NO2] ratio present. Some data 

verifying this important characteristic appears in 

Fig. 3, which shows the sensor output as 

[NO]:[NO2] is varied under conditions of 100 ppm 

FY 2006 Progress Report 

NOx. Varying the NO2 content from ~50 to 0%* 

induced an approximately 6% change in the output 

signal. This shows that the sensing element is fairly 

insensitive to the relative abundance of the two NOx 

species. 

Figure 3. Response trace showing how the sensing 

element output changes with [NO]:[NO2]. 

The final two characteristics evaluated for these 

sensing elements with improved electrode 

architecture were their response to changing [O2] 

and the effect of the potential interferents CO and 

hydrocarbons. The effect of changing O2 is 

particularly important as diesel exhaust will 

typically contain between 10 and 15% O2.

5 

The effect of changing [O2] under conditions of 

fixed [NO] and [NO2] is illustrated in Fig. 4. 

Changing [O2] does affect the sensing element 

output, indicating that an independent assessment of 

[O2] will be required during operation of the sensor. 

However, the CRADA partner indicated that this 

may be preferable to the O2 pumping step required 

with the current amperometric NOx sensor

3

. 

Figure 5 shows how the potential interferents 

CO and C3H6 (propylene, a "representative" 

hydrocarbon) affected the sensor output. It can be 

seen these species do affect the sensor output when 

present at concentration levels around 100 ppm. 

Fortunately, these species are easily eliminated with 

an oxidation catalyst in the O2-rich diesel exhaust 

environment (as Fig. 1 suggests). One of the 

*

Increasing the NO2 content from 50 to 100% also caused an 

approximate 6% change in the signal relative to its value at 

[NO]:[NO2] = 1:1. 

37
Heavy Vehicle Propulsion Materials 

commercial sensors mentioned earlier4 incorporates 

such an oxidation catalyst, so the need for a similar 

catalyst with the sensing elements described herein 

does not represent a substantial disadvantage. 

Figure 4. Change in sensor output with oxygen content. 

Figure 5. Comparative response to NOx, C3H6, and CO. 

Conclusions 

Sensing elements that are capable of “total NOx” 

sensing behavior have been further improved during 

FY 06. The drift of these sensing elements was 

dramatically reduced by altering the electrode 

architecture. The sensing elements are sensitive to 

changing [O2] and the potential interferents CO and 

propylene (C3H6). Therefore an independent 

FY 2006 Progress Report 

assessment of [O2] as well as an upstream oxidation 

catalyst will be required. 

References 

1. J. Woestmann and E. Logothetis, "Controlling 

automotive emissions," The Industrial Physicist, 

1, 21, 1995. 

2. K. Schnell and C. Browne, “NOx control,” in 

Air pollution control technology handbook, 

CRC Press, 2001, pp. 241–255. 

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

film ZrO2 sensor for the measurement of low 

NOx concentration,” SAE Technical Paper 

980170, 1998. 

4. A. Kunimoto, M. Hasei, Y. Yan, Y. Gao, T. 

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

sensor based on mixed potential for 

automobiles,” SAE Technical Paper 1999-01-

1280, 1999. 

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

"Automotive catalytic converters: Current status 

and some perspectives," Catalysis Today, 77, 

419, 2003. 

Patents 

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

Armstrong, “NOx sensing devices having 

conductive oxide electrodes,” U.S. Patent App. 

20060073070 (2006). 

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

Armstrong, “NOx sensing devices having 

conductive oxide electrodes,” Int. Pat. App. 

PCT/US2005/034279, 2006. 

Presentations and Publications 

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

Armstrong, “Total NOx sensing elements with 

compositionally identical oxide electrodes,” 

Journal of the Electrochemical Society, 153, 

H23–8, 2006. 

38
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

D. Development of Materials Analysis Tools for Studying NOx Adsorber 

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

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

Oak Ridge National Laboratory 

P.O. Box 2008 

Oak Ridge, TN 37831-6064 

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

DOE Technology Development Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: jerry.gibbs@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 

• The objective of this effort is to produce a quantitative understanding of the processing and in-service effects on 

NOx adsorber catalyst technology leading to an exhaust aftertreatment system with improved catalyst 

performance capable of meeting the 2007 Environmental Protection Agency (EPA) emission requirements in 

the United States. 

Approach 

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

findings with Cummins data and experience. 

Accomplishments 

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

• X-ray diffraction, X-ray photoelectron Spectroscopy and Scanning Transmission Electron Microscopy were 

used to evaluate the thermal degradation of a NOx storage and reduction (NSR) catalyst within the catalyst 

support brick as a function of macroscopic position, catalyst history and operating conditions. 

• Completed two studies characterizing a new catalyst system and a diesel oxidation catalyst in the unused and 

aged states. 

Future Direction 

• Assist Cummins to competitively produce engines which attain the required emission levels for 2010 and 

beyond while maintaining the advantage of the diesel’s inherent energy efficiency. 

• Evaluation of hydrothermal aging of a SCR catalyst. 

• Continue to characterize the soot, coke and ash formed on the catalyst from different fuels, including biodiesel. 

39
Heavy Vehicle Propulsion Materials 

Introduction 

In order to meet the 2007 US Environmental 

Protection Agency (EPA) emission requirements for 

diesel exhaust, aftertreatment in diesel engines may 

be necessary. The technology necessary 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 utilize these catalyst systems. ORNL is 

assisting with the materials characterization effort. 

Prior reports and presentations have covered the 

accomplishments related to the NOx adsorber 

catalyst technology, this report will focus on the 

diesel oxidation catalyst (DOC) study. 

A diesel oxidation catalyst (DOC) is a flow 

through device composed of a porous honeycomb 

structure within which a high surface area ceramic 

washcoat is situated. Small amounts of precious 

metal are dispersed within the washcoat and provide 

the catalytic activity necessary to oxidize carbon 

monoxide, gaseous hydrocarbons and liquid 

hydrocarbon particles (unburned fuel and oil).1 

Regulated emissions are thus controlled. 

Base metal oxides (BMOs) and precious metals 

are major and critical components in current DOCs, 

respectively, that Cummins seeks to use. The 

microstuctural, chemical and phase changes of these 

components are of great interest as DOCs 

experience aging/normal operating conditions. 

Ultimately, it is desired to prevent/control these 

changes such that this technology becomes 

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 transmission 

electron microscopy (TEM), X-ray diffraction 

(XRD), X-ray photoelectron spectroscopy (XPS), 

and Raman Spectroscopy. 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. 

FY 2005 Progress Report 

Samples 

Briefly, Cummins provided various series of 

“core” samples taken from cordierite “bricks” 

containing and an aluminosilicate support. The 

general features of each series will be given below 

in the text as needed. The brick/core samples 

consisted of a new/never-been-used DOC and an 

aged DOC. Here, the aged DOC originate from the 

same unused DOC. Aging was accomplished in a 

retort oven at 700°C for 24 hours, with 10% water 

added. 

Results 

The above samples were examined with XPS, 

Raman, XRD and TEM from October 2005 to 

September 2006, which is summarized below. 

Spectroscopy 

Spectroscopic techniques provide important 

information about the chemical state of and 

identification of adsorbed surface species. Two 

complementary spectroscopic techniques are X-ray 

Photoelectron Spectroscopy (XPS) and Raman 

Spectroscopy which provide low frequency 

vibrational and molecular vibrational information, 

respectively. 

X-ray photoelectron spectroscopy (XPS) was used 

to examine compositional variations along the length 

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

new and aged samples The average surface 

compositions are shown in Table I. Single-trough 

samples were prepared by sectioning out a layer 

containing 8-10 troughs from the interior of the 

honeycomb core. Equidistant positions along 2 

separate troughs were defined; two sets of data were 

taken within each trough (15 points/trough), 

analyzed and averaged. For the new sample, a third 

trough was examined 5 months after the first two, 

confirming the reproducibility of the data (i.e. no 

significant change with time). The compositional 

variation along the troughs is demonstrated in 

Figures 1 and 2 for the new and aged DOCs, 

respectively. Each shows two graphs, one for the 

major elements (O, Al, Si, and C) and one for the 

minor elements (N, Pt, and F) identified. Presently, 

the origin of F is unknown. The figures show a 

small decrease in the amount of O within in the 

probe depth and a slight increase in the amounts of 

Pt, Al and C in the aged relative to the new DOC. 

40
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

The amount of Si seems unchanged, and the 

variability of N and F is inconclusive. Examination 

of the individual runs (not shown), made prior to 

adding/averaging, show the O, C, and Al data were 

more scattered, which could be due to the 

atmospheric contamination for O and C. The origin 

of the changes in Al is unclear, as absence of 

magnesium indicates that the signal originates from 

41
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Table I. Surface composition (atom %) of new and aged DOC samples averaged from data taken 

along the length of a single trough using XPS. 

42
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 1. Typical XPS results showing the major and minor elements in the new/unused DOC sample along 

two troughs (1&2) of the cordierite honeycomb. 

43
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 2. Typical XPS results showing the major and minor elements in the aged DOC sample along two troughs 

(1&2) of the cordierite honeycomb. 

44
Heavy Vehicle Propulsion Materials 

the washcoat only. The washcoat likely contains 

some sort of alumino-silicate, based on the majority 

elements. In contrast, the Pt scans were more 

consistent, trough-to-trough and run-to-run. Based 

on this, two major conclusions can be drawn from 

these data. First, the aging process does not 

influence the nominal composition of the major 

elements or change the uniform distribution along 

the trough. Second, the Pt signal is nearly doubled 

for the aged sample and implies that Pt atoms have 

migrated/redistributed to the near surface region 

(~5-10 nm, i.e. the probe depth for XPS). 

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. Raman 

spectroscopy confirmed the presence of γ-alumina in 

the washcoat (see Figure 3). There is evidence to 

suggest a small a mount of θ-alumina may also be 

present. 

X-ray 

X-ray diffraction (XRD) was employed to 

understand the crystalline nature of the samples. For 

the unfamiliar, the following analogy can be applied: 

As a fingerprint identifies a person, so a diffraction 

pattern identifies a crystalline material. 

In Figure 4, XRD indicates that the sample 

contains orthorhombic cordierite 

(2MgO•5SiO2•2Al2O3) and platinum. The broad 

peak at low angle at ~20° 2Θ is due to an amorphous 

phase, possibly part of the washcoat. Previously, γalumina was tentatively identified within the 

washcoat of the new DOC. Figure 5 compares the 

XRD patterns of the new/unused and aged DOC 

samples showing the marked crystallite size growth 

of the Pt (<1, i.e., not observable, to ~30 nm). 

Microscopy 

Scanning Transmission Electron Microscopy 

(STEM) techniques were used to characterization of 

the microstructure of the on both the new and aged 

DOCs. In the new sample, a majority of the 

washcoat/support is found to be amorphous (see 

above), and energy dispersive spectroscopy (EDS) 

indicates the presence of Si, Al and O (perhaps 

alumina and silica). The washcoat contains Pt metal 

crystals, which are 1 - 4 nm in size (see Figure 6). 

Much of the Pt is atomically dispersed as seen with 

Aberration-corrected STEM. This is the first time 

FY 2005 Progress Report 

such a non-ideal sample has been examined in this 

fashion. Small clusters of Pt (1-3 atoms thick) are 

also observed. For thicker areas, the more limited 

depth of field in aberration-corrected STEM 

provides richer, 3D data to be extracted by using a 

focusing/defocusing technique over ~5nm. Again, 

this is the first time that this has been applied to 

large cluster. The aged sample demonstrated 

evidence of massive Pt sintering (see Figure 7), 

which is consistent with the XRD results above. 

Small nanocrystals, atomic clusters, and single 

atoms are also still present (see Figure 8). In Figure 

9, the moderately sized nanocrystals are surrounded 

by “cloud” of Pt clusters, which suggests a 

redispersion of Pt, supporting the XPS result above 

The mechanism is presently unknown but the 

electron beam from the STEM has been eliminated 

as a potential cause due to lack of change with time. 

Finally, the Pt atoms are observed to preferentially 

attach to alumina planes in the support. 

References 

[1] Washington State University Extension Energy 

Program, Diesel Oxidation Catalyst, 

http://www.energy.wsu.edu/pubs/. 

Acronyms 

BMOs Base metal oxides 

DOC diesel oxidation catalyst 

EDS Energy dispersive spectroscopy 

EPA Environmental Protection Agency 

NOx Nitrogen and Oxygen containing 

compounds 

ORNL Oak Ridge National Laboratory 

SCR selective catalytic reduction 

STEM Scanning Transmission Electron 

Microscopy 

XPS X-ray photoelectron spectroscopy 

XRD X-Ray Diffraction 

45
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 3. Raman spectrum of the DOC washcoat. 

Figure 4. XRD pattern of the aged DOC sample. 

46
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 5. XRD patterns of the new (black curve) and aged (blue curve) DOC samples and highlights the large Pt 

peaks present in the aged relative to the new DOC. 

47
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 6. Images show the near atomically dispersed Pt particles and clusters in the new DOC taken with

Aberration-corrected STEM. 

48
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 7. STEM images showing the large Pt crystals in the aged DOC. 

Figure 8. STEM images showing the near atomically dispersed Pt cluster and particles still present within the aged

DOC. 

49
Heavy Vehicle Propulsion Materials FY 2005 Progress Report 

Figure 9. STEM images showing nanocrystals (white arrows) surrounded by “cloud” (hazy “milky-way”) of Pt 

clusters within the aged DOC. 

50
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

E. Catalysis by First Principles 

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

 Oak Ridge National Laboratory 

 P.O. Box 2008, MS 6116 

(865)574-8445, narulack@ornl.govl 

DOE Technology Development Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@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 

The objective of this work is to search for durable emission treatment catalysts (LNT, TWC, OC) from a 

protocol based on 

 an integrated approach between computational modeling and experimental development, 

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

necessary for improving the catalytic efficiency of these materials. 

Approach 

 Theoretical Modeling 

o Density functional theory-based first-principles calculations 

o Optimization of Pt clusters supported on alumina 

o Interaction of CO, NOx, and HC with Pt clusters supported on alumina 

 Experimental System 

o Synthesize Pt Nanoclusters on morphologically diverse alumina supports 

o Evaluate systems for CO, NOx, HC oxidation 

o We are employing Ex-situ reactor to expose Pt cluster in 2%Pt//γ-Al2O3 (1nm Pt) catalyst to 

CO oxidation conditions in order to monitor microstructural changes by atomic resolution 

electron microscopy. This approach permits us to repeatedly image the exact same areas of 

the fresh supported Pt catalyst and after multiple CO oxidation reaction conditions. 

o We carried out In-situ FTIR studies to determine the CO absorption modes on fresh 2%/γAl2O3 and after CO oxidation testing. Catalysts with 1 nm and 25 nm Pt clusters/particles 

were compared at various stages of exposure to CO oxidation conditions to understand the 

effect of Pt microstructure on CO absorption. 

Accomplishments 

 We completed the synthesis of a series of supported Pt catalysts [2%Pt/MgO, 2%Pt/γ-Al2O3 (1nm and 

25nm Pt), 2%Pt/θ-Al2O3 (1nm Pt)] and evaluated them for CO oxidation. 

o We have established that the nanostructure of 2%Pt/γ-Al2O3 is quite different form that 

reported in literature and found that it undergoes changes even on exposure to CO oxidation 

initiation conditions. 

o We monitored the light-off temperatures of these catalysts on a bench reactor (%CO 

oxidation) and initiated the efforts to correlate the activity with catalysts’ microstructure. 

51
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

o The preliminary studies on 1nm 2%Pt/θ-Al2O3 suggest that the structure of substrate has a 

significant impact on CO oxidation initiation and light-off. The efforts to determine the 

impact of substrate structure on nano-structure of catalysts are in progress. 

 We have used DFT methods to investigate formation and adsorption energies as well as reaction 

energy profiles. 

o Formation energies Pt1-3Oy on MgO(100) 

o Adsorption energies of O and CO on PtxOx and PtxO2x 

o Adsorption energies of O, O2, CO, CO2, NO, and NO2 on Ptx 

o Reaction energy profiles for CO+1/2O2 →CO2 and NO+1/2O2 →NO2 

Future Direction 

 Understanding of the structures of nano-clusters on support. Our results show that nano-clusters are 

10-15 atoms 

o Theoretical models to understand cluster oxidation state (oxidized, reduced, in equilibrium), 

dependence on cluster size, and the kinetics of oxidation 

o Experimental studies on the structures of these systems using EXAFS, XANES 

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

o Cationic or zero-valent metals or both 

o Cationic metals at the metal–support interface 

 Reactivity of the clusters 

o Theoretical and experimental studies on structure-reactivity correlation towards NOx, and HC 

oxidation studies 

o The impact of cluster size on catalytic activity 

 Translate these results into design of durable supported catalysts for lean NOx traps (and other 

systems such as TWC, OC for diesel etc.) 

Introduction 

This research focuses on an integrated 

approach between computational modeling and 

experimental development, design and testing of 

new catalyst materials, that we believe will 

rapidly identify the key physiochemical 

parameters necessary for improving the catalytic 

efficiency of these materials. 

The typical solid catalyst consists of 

nano-particles on porous supports. The 

development of new catalytic materials is still 

dominated by trial and error methods, even 

though the experimental and theoretical bases 

for their characterization have improved 

dramatically in recent years. Although it has 

been successful, the empirical development of 

catalytic materials is time consuming and 

expensive and brings no guarantees of success. 

Part of the difficulty is that most catalytic 

materials are highly non-uniform and complex, 

and most characterization methods provide only 

average structural data. Now, with improved 

capabilities for synthesis of nearly uniform 

catalysts, which offer the prospects of high 

selectivity as well as susceptibility to incisive 

characterization combined with state-of-the 

science characterization methods, including 

those that allow imaging of individual catalytic 

sites, we have compelling opportunity to 

markedly accelerate the advancement of the 

science and technology of catalysis. 

Computational approaches, on the other 

hand, have been limited to examining processes 

and phenomena using models that had been 

much simplified in comparison to real materials. 

This limitation was mainly a consequence of 

limitations in computer hardware and in the 

development of sophisticated algorithms that are 

computationally efficient. In particular, 

experimental catalysis has not benefited from 

the recent advances in high performance 

computing that enables more realistic 

simulations (empirical and first-principles) of 

large ensemble atoms including the local 

environment of a catalyst site in heterogeneous 

52
Heavy Vehicle Propulsion Materials 

catalysis. These types of simulations, when 

combined with incisive microscopic and 

spectroscopic characterization of catalysts, can 

lead to a much deeper understanding of the 

reaction chemistry that is difficult to decipher 

from experimental work alone. 

Ptx Pt/Al Pt 2O3 x Pt/Al

2 3 TWC, LNT

Theoretical ModelsExperimental Models

Although it is beyond doubt 

computationally challenging, the study of 

surface, nanometer-sized, metal clusters may be 

accomplished by merging state-of-the-art, 

density-functional-based, electronic-structure 

techniques and molecular-dynamic techniques. 

These techniques provide accurate energetics, 

force, and electronic information. Theoretical 

work must be based on electronic-structure 

methods, as opposed to more empirical-based 

techniques, so as to provide realistic energetics 

and direct electronic information. 

A computationally complex system, in 

principle, will be a model of a simple catalyst 

that can be synthesized and evaluated in the 

laboratory. It is important to point out that such 

a system for experimentalist will be an idealized 

simple model catalyst system that will probably 

model a “real-world” catalyst. 

Thus it is conceivable that 

“computationally complex but experimentally 

simple” system can be examined by both 

theoretical models and experimental work to 

forecast improvements to obtain optimum 

catalyst systems. 

Approach 

The theoretical modeling is based on 

DFT studies of Pt clusters to understand the 

relationship between cluster size, structure, 

composition, and reactivity. This coupled with 

first-principles thermodynamics provides 

insights into the effects of oxidizing atmosphere 

(O2) – finite (T, pO2

), structure, composition, 

FY 2006 Progress Report 

Thus, a protocol to systematically find 

the optimum catalyst can be developed that 

combines the power of theory and experiment 

for atomistic design of catalytically active sites 

and can translate the fundamental insights 

gained directly to a complete catalyst system 

that can be technically deployed. 

Candidate 

Materials 

Candidate

Materials

CombinatorialDurable 

Catalyst

Screening Materials ning Materials

redox potential on particle size. The results 

provide guidance for investigations of larger / 

supported clusters. The next steps of this work 

will model supported clusters, based on input 

from experimental results, and their interactions 

with CO, NOx, and HC. 

Experimentally, we have synthesized a 

series of Pt clusters and particles supported on 

morphologically different aluminas. The 

microstructural characterization reveals the 

presence of 1-3 atom and 10-20 atom clusters. 

We have carried our CO oxidation studies on 

these catalysts and we plan to study NOx and HC 

oxidation activity. 

Guided by the theoretical models and 

experimental results of our study, we will 

synthesize new set of catalyst materials with 

higher durability under operating conditions. 

Results 

Over the last year, we have carried out 

theoretical studies and experimental studies of Pt 

nanoclusters and nanoparticles and their CO 

oxidation activity employing a bench-top reactor 

connected to Agilent GC for gas analysis. We 

monitored CO adsorption and desorption on 

catalysts by in-situ FTIR studies using a 1%CO 

gas flow (He balance) at 25 mL/min across a 

pressed wafer of the catalyst. We employed our 

ex-situ reactor6

 and ACEM to study the 

nanostructural changes in catalysts upon 

exposure to various reaction condition of CO 

oxidation. It is important to note that the ex-situ 

reactor permits duplication of CO oxidation 

conditions of bench-top reactor. 

53
Heavy Vehicle Propulsion Materials 

The results show the correlation between 

nanostructural changes and their impact on the 

catalytic activity and are summarized in the 

following subsections: 

OXIDATION CATALYSTS - Pt/MgO 

Magnesia (MgO) is a model basic oxide surface 

that is relatively inert toward many adsorbates 

and its (100) facet is thermodynamically the 

most stable. Our DFT investigation for Pt 

clusters supported on MgO(100) is on-going. 

Preliminary results are summarized in Figure 1. 

The oxides of the Pt dimer and trimer exhibit 

very similar formation energies on MgO(100) 

compared to those of the free Pt2 and Pt3 

clusters, while Pt1Oy clusters have markedly 

lower formation energies. Furthermore, the 

morphologies of the various Pt oxide clusters on 

MgO(100) closely resemble their gas-phase 

counterparts. These results support our 

hypothesis that the oxidation energetics and 

morphologies of the free Pt nanoclusters should 

by and large be preserved on a relatively inert 

surface and therefore our approach of 

investigating the isolated Pt nanoclusters first. 

While efforts to experimentally synthesize Pt 

nanoclusters supported on MgO are in progress 

to validate theoretical models, we have 

synthesized Pt/MgO with Pt particle size 

predominately in the 15-20 nm range supported 

on single crystal MgO particles. The results of 

CO oxidation studies are shown in Figure 2. The 

CO oxidation gradually reaches ~10% in 170­

230oC range but rapidly becomes quantitative 

just above 230oC. 

OXIDATION CATALYSTS - Pt/γ-Al2O3 

1 nm 2%Pt/γ-Al2O3: We synthesized 1 nm 

2%Pt/γ-Al2O3 by impregnation of γ-Al2O3 with 

H2PtCl6•6H2O followed by calcination in air at 

450°C. STEM analysis shows that platinum 

particle size ranges from 0.6-1.4 nm and the 

distribution is centered at 1.1 nm (Figure 3). In 

general, the models for such nanoclusters 

assume perfect spherical shape to calculate 

number of particles in such cluster. Atomic 

resolution of the particles by Aberration 

 FY 2006 Progress Report 

Corrected Electron Microscopy (ACEM) in the 

HAADF-STEM mode clearly shows that such 

models are incorrect and the atomic make-up of 

these platinum nano-clusters involves on 

average between 10-20 atoms (Figure 4). In 

addition, single atoms and 2-3 atom clusters can 

also be observed that are not visible in TEM 

images obtained from non-aberration correct 

STEM. Structurally, such clusters can be 

expected to be quite different from bulk 

platinum and our previously reported theoretical 

studies also suggest that these clusters could 

incorporate oxygen. 

We tested the 1 nm 2%Pt/γ-Al2O3 catalyst for 

CO oxidation activity on a bench-top reactor. 

The conversion of CO to CO2 was monitored by 

the analysis of the off gases by gas 

chromatography (GC) as a function of 

temperature. The reaction was stopped after CO 

oxidation had initiated (4.4% CO conversion) at 

210 °C (Figure 5) and the sample was collected 

for analysis. A fresh sample was placed in the 

reactor and the temperature was increased till 

CO conversion became quantitative. The 

reaction was stopped and the reactor was 

allowed to attain room temperature. The CO 

oxidation reaction was repeated 2 more times 

(Figure 6). In the first cycle, CO oxidation began 

at 210 °C while the third cycle required only 180 

°C for the CO oxidation initiation. Analysis of 

these test samples by X-ray diffraction (XRD) 

gave broad or unobservable Pt diffraction 

patterns due to the small Pt particle sizes. The 

images obtained on an electron microscope 

suggest that Pt nanoparticle have increase in size 

to some degree. 

The fresh 1nm 2%Pt/γ-Al2O3 catalyst showed no 

CO adsorption at 25°C during in-situ IR studies. 

However, at 85°C, one CO peak was seen by IR 

at 2116 cm-1 which has a small shoulder (ca. 

2094 cm-1). When the adsorbed sample was 

subject to desorption by flowing only helium the 

peak shifted to 2114 cm-1 but did not desorb 

(Figure 7). The 1 nm 2%Pt/γ-Al2O3 catalyst was 

also studied by the in-situ FTIR process after it 

was tested for 100% CO oxidation 3 times. This 

catalyst had three CO vibrations at 25°C (2175, 

2117, 2063 cm-1) under adsorption conditions 

54
Heavy Vehicle Propulsion Materials 

(Figure 7). The two lower energy vibrations 

desorbed in the absence of CO while the higher 

energy vibration shifted to 2058 cm-1. When the 

sample was tested at 85°C three similar CO 

vibrations were seen but the high and the low 

energy vibrations were shifted (2177, 2116, 

2069 cm-1). While the two lower energy 

vibrations mostly desorbed, vibrations were still 

discernable with the lowest shifting in 

comparison to the adsorption data (2248 and 

2116 cm-1). The major vibration seen under 

desorption conditions was the highest energy 

vibration which was shifted (2050 cm-1). 

Interestingly, the vibration around 2116 cm-1 is 

present in all spectra and is most intense in the 

fresh sample. Furthermore, it does not change 

after exposing fresh sample to desorbing 

conditions. This suggests that this vibration is 

caused by CO adsorption on monatomic Pt. 

Thus TEM and IR studies suggest that even 

exposure to CO oxidation initiation conditions 

results in changes in nanostructure of 1nm 

2%Pt/γ-Al2O3 catalyst. In order to ascertain the 

extent of nanostructure changes, we have 

initiated the study of nanostructural changes in 

1nm 2%Pt/γ-Al2O3 catalyst on our unique ex-situ 

reactor at ORNL in conjunction with 

microscopy using the ACEM. This study allows 

us to monitor changes in selected areas 

repeatedly after exposure to various reaction 

conditions such as (1)fresh samples, (2)CO 

oxidation initiation conditions, (3) degreening 

process (100% conversion 2x), and (4)100% 

CO conversion after sample degreening. The 

fresh sample shows small Pt clusters and single 

Pt atoms (Figures 8a, 9a and 10a). After 

exposure to CO oxidation initiation conditions, 

large clusters (Figure 8b and 9b) and small Pt 

particles (2nm) (Figure 10) can also be seen. 

The ex-situ reactor study confirms our 

observations that nanostructural changes in 1nm 

2%Pt/γ-Al2O3 catalyst occur even on exposure to 

CO oxidation initiation conditions (4.4% CO 

oxidation). This study is continuing and will be 

completed in the next quarter. 

These results are also supported by theoretical 

modeling. Preliminary DFT calculations of the 

oxidation of CO and NO on Pt clusters indicate 

that the reaction energetics are strongly 

 FY 2006 Progress Report 

dependent on the size of the clusters as well as 

the extent of oxidation on the clusters. We first 

look at the adsorption of the relevant atomic and 

molecular species on Pt nanoclusters, including 

O, O2, CO, CO2, NO, and NO2. Figure 11 shows 

the lowest adsorption energies (i.e., strongest 

adsorption) for the six species on Ptx. Adsorption 

geometry, including adsorbate and the cluster, 

are fully relaxed. Overall, the adsorption of all 

six species is highly size-dependent in the 

angstrom regime and substantially stronger than 

on the bulk Pt surface (Pt(111)). As particle size 

increases, the adsorption energies trend toward 

the Pt(111) level (“∞” in Figure 11). That the 

adsorption of these species is enhanced 

compared to Pt(111) means that the energy 

profiles of CO and NO oxidation on Ptx also 

differ from their bulk-surface counterparts 

(Figure 12). CO oxidation on the bulk Pt 

surface is exothermic every step along the 

reaction path, but because of the enhanced 

adsorption of CO, there is a minimum in energy 

when CO adsorbs on the clusters. The same 

phenomenon can be seen in NO oxidation. This 

indicates that the ability of the Pt clusters to 

catalyze CO and NO oxidation may be inhibited 

by the strong adsorption of CO and NO. 

Our latest results indicate that the adsorption of 

O, CO (Figure 13) and NO (not shown) are 

weakened on the PtxOx and PtxO2x clusters, 

which suggests that the energetics of CO and 

NO oxidation may be more favorable on the 

small oxidized Pt clusters than on the original, 

metallic clusters. 

25 nm 2%Pt/γ-Al2O3: The sintering of 1 nm 

2%Pt/γ-Al2O3 at 650C for 4h led to the 

formation of 25 nm 2%Pt/γ-Al2O3 with Pt 

particles of an average size of 25 nm (Figure 14) 

as calculated from analyzing Pt diffraction 

peaks in XRD pattern using Scherrer’s formula. 

The CO oxidation studies mirrored the CO 

oxidation protocol described for 1 nm 2%Pt/γAl2O3 catalyst (Figures 5 & 6). 100% CO 

oxidation occurred at a lower temperature for the 

25 nm Pt catalysts than that seen for the 1 nm Pt 

catalysts. The larger Pt particles also exhibit a 

shift in the temperature required to achieve 

100% oxidation between the 2nd and 3rd cycle 

(from 250 to 242 °C). Pt particle growth was 

55
Heavy Vehicle Propulsion Materials 

noted by XRD analysis after both the initiation 

test and the 100% oxidation 3 times, see Figure 

14. The analysis of diffraction peaks in XRD by 

Scherrer formula showed that Pt particles grew 

from 25 nm to 28 nm after initiation which was 

stopped at 8.1% conversion and to 34 nm after 

cycling to 100% conversion 3 times (Figure 14). 

No significant difference in CO oxidation 

activity was observed when 30 nm instead of 25 

nm Pt particles were tested. 

The in-situ FTIR studies of CO adsorptiondesorption on fresh 25 nm 2%Pt/γ-Al2O3 catalyst 

at 25 and 85°C showed bands at 2170, 2117, 

2069 cm-1 and 2179, 2116, 2060 cm-1 

respectively (Figure 15). At both temperatures, 

only the high energy vibration remained during 

desorption conditions and were shifted to 5 cm-1 

higher energy. The catalyst, after going through 

three CO oxidation cycle, exhibits bands at 

2173, 2116, 2073 cm-1 and 2173, 2117, 2063 cm1

 in its IR spectra at 25 and 85°C respectively. It 

is important to note that the CO adsorption 

behavior does not significantly change for this 

catalyst after exposure to CO oxidation 

conditions. Desorption, on the other hand, 

results in a strong band at 2075 and 2056 cm-1 

and weak band at 2173 and 2139 cm-1 at 25 and 

85°C respectively (Figure 15). All samples also 

exhibit a weak vibration at 2116 cm-1 also. 

OXIDATION CATALYSTS - Pt/θ-Al2O3 

The θ–alumina is the first structurally diverse 

alumina that we selected to understand the 

impact of surface properties and structure of 

substrate on the behavior of nanocatalysts. 

Compared with γ–alumina, θ-alumina (BET 

surface area ~100 m2

/g) has a simple structure 

and its surface properties while less attractive 

than sol-gel derived γ–alumina (BET surface are 

~200m2

/g) are still comparable to commercial γalumina (BET surface are ~ 100m2

/g). It can be 

easily synthesized by sintering sol-gel derived γAl2O3 at 900°C for 34 hours. The Pt was loaded 

by impregnating it with H2PtCl6•6H2O and 

subsequent calcination in air at 450°C. The 

electron microscopic analysis shows that the Pt 

clusters sizes ranges from 0.6-1.3 nm and are 

centered at 1.1 nm (Figure 16). Thus, the Pt 

 FY 2006 Progress Report 

particles distribution is similar to that for fresh 

2%Pt/γ-Al2O3, prepared from commercial γalumina. The nanostructure analysis of this 

sample by ACEM is in progress and will be 

reported in a future report. 

We subjected this catalyst to the test protocol 

described for 1 nm 2%Pt/γ-Al2O3 and the results 

and comparison with 1 nm 2%Pt/γ-Al2O3 are 

presented in Figure 17 and Figure 18. The 

2%Pt/θ-alumina samples show lower CO 

oxidation initiation and light-off temperatures. 

Further nanostructural studies will reveal the 

differences between Pt/θ-alumina and Pt/γ-Al2O3 

catalyst that could explain their CO oxidation 

behavior. 

Conclusions 

The most important result of this study is that 

the nanostructure of Pt supported on alumina is 

quite different from that reported in several 

literature studies. Furthermore, we have shown 

that even exposure to CO oxidation initiation 

conditions results in dramatic changes in the 

nano-structure of Pt. The increase in Pt size from 

1 to 25 nm results in lower CO oxidation 

initiation and completion temperatures. We 

have also shown the changes in CO absorption 

supported Pt catalysts after exposure to various 

CO oxidation conditions by FTIR. 

Computational calculations and simulation have 

shown that the stronger adsorption of reactants 

on small Pt clusters may inhibit there oxidation. 

Adsorption of oxidation reactants on oxidized Pt 

clusters is weaker, suggesting that CO and NO 

oxidation may be more favorable than on nonoxidized Pt clusters. 

The impact of support on the nanostructure of 

supported Pt is not fully understood. 

Experimentally, we have observed that Pt 

supported on θ-Al2O3 exhibits lower CO 

oxidation initiation and light-off temperatures as 

compared with Pt supported on commercial γAl2O3. Our next goals are to complete 

microstructural studies of the 2%Pt/γ-Al2O3 

catalysts used for CO oxidation by microscopy 

and start to initiate the study of NO oxidation on 

these catalysts. 

56
Heavy Vehicle Propulsion Materials 

Presentations and Publications 

C.K. Narula presented this work at DEER 

conference. M. Moses presented a paper at SAE 

Powertrain and Fluid Systems Meeting. 

1. Narula, C.K.; Blom, D,A,; Allard, L.F.; 

Moses, M.J.; Xu, Y.; Shelton, W.A.; 

Schneider, W.F.; “Catalysis by Design – 

Theoretical and Experimental studies of 

model catalysts for Lean NOx Treatment.” 

DEER 2006, Marriott, Detroit, Aug. 21-25, 

2006. 

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

“Analysis of Microstructural Changes In 

Lean NOx Trap Materials Isolates 

Parameters Responsible for Activity 

Deterioration” SAE no. 2006-01-3420, 2006. 

3. C.K. Narula, M. Moses(speaker), L.F. 

Allard, SAE Powertrain and Fluid Systems 

Meeting 2006, Toronto, Ontario, October 

18-20, 2006. 

4. Narula, C. K., Moses, M.J., Xu, Y., Blom, 

D.A., Allard, L.F., Shelton, W.A. “Catalysis 

by Design – Theoretical and Experimental 

Studies of Model Catalysts.” SAE no. 

2007M-418, 2007, abstract accepted-invited. 

5. Y. Xu, W. A. Shelton, W. F. Schneider, 

Theoretical Aspects of Oxide Particle 

Stability and Chemical Reactivity, in 

Synthesis and Application of Oxide 

Nanoparticles and Nanostructures; Wiley, 

in press. 

6. Y. Xu (speaker), W. A. Shelton, and W. F. 

Schneider, “Nano-scale effects in the 

 FY 2006 Progress Report 

oxidation and the reactivity of platinum 

clusters,” AIChE Annual Meeting, San 

Francisco, CA, 11/2006, abstract accepted. 

7. Y. Xu (speaker), W. A. Shelton, and W. F. 

Schneider, “Platinum clusters for oxidation 

catalysis: Nano-scale effects in 

thermodynamics and reactivity,” 232nd ACS 

National Meeting, San Francisco, CA, 231st 

ACS National Meeting, Atlanta, GA, 

3/2006. 

8. Y. Xu, W.A. Shelton, and W.F. Schneider, 

“The thermodynamic equilibrium 

compositions, structures, and reaction 

energies of PtxOy (x = 1-3) clusters predicted 

from first principles,” in press, Journal of 

Physical Chemistry B, 110 (2006) 16591. 

9. Y. Xu, W. A. Shelton, and W. F. Schneider, 

“Effect of particle size on the oxidizability 

of platinum clusters,” Journal of Physical 

Chemistry A, 110 (2006) 5839. 

References 

1. Deng, W.; Toepke, M.W., Shanks, B.H., 

Adv. Func. Mater., 2003, 13, 61. 

2. Gates, B.C., Chem. Rev., 1995, 95, 511 

3. Narula, C.K.; Nakouzi, S.R., Wu, R., 

Goralski, C.T., Allard, L.F., AIChE Journal, 

2001, 47, 744. 

4. Andrews, M. A., Kirtley, S. W., Kaesz, H. 

D., Inorg. Synth. 1977, 17, 16. 

5. Fung, A. S., Tooley, P. A., Kelley, M. J., 

Koningsberger, D. C., Gates, B. C. J. Phys. 

Chem. 1991, 95, 225. 

6. Narula, C.K., Moses, M.J., Allard, L.F. SAE 

no. 2006-01-3420, 2006. 

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Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

nO 

free 

on MgO 

nO 

free 

on MgO 

free 

on MgO 

Pt1Oy Pt2Oy 

Pt3Oy 

nO 

Figure 1. The formation energies of Pt1-3Oy on MgO(100). 

150 160 170 180 190 200 210 220 230 240 250 

Temperature (C) 

Figure 2. Complete catalytic conversion of CO on 2% Pt/MgO as a function of temperature, with initial 

sampling and after 15 min. at the given temperature. 

0 

20 

40 

60 

80 

100 

%

C

O conversion 

Pt/MgO 

0 minutes 15 minutes 

58
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

160 

# of particles 120 

80 

40 

0 

0.5 1.0 1.5 2.0 

Particle size (nm) 

Figure 3. STEM image of fresh 2%Pt/γ-Al2O3 (Left), the Pt particles size distribution for this area 

(Right). 

Figure 4. ACEM HAADF-STEM images of 2%Ptγ-Al2O3. Bright white dot are Pt atoms, grayish are is 

alumina support. 

59
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 5. Initiation of catalyst conversion as a function of temperature with sampling at 0 (blue) and 15 

(pink) minutes at the indicated temperatures. The solid lines are results for the 2% Pt/ γ-Al2O3 annealed 

at 650°C (ca. 25 nm Pt) and the dotted lines are results for 2% Pt/ γ-Al2O3 (ca. 1 nm Pt). 

Figure 6. Complete catalyst conversion of CO to CO2 as a function of temperature with sampling after 

15mins at the given temperature. The solid lines are results for the 2% Pt/ γ-Al2O3 annealed at 650°C (ca. 

25 nm Pt) and the dotted lines are results for 2% Pt/ γ-Al2O3 (ca. 1 nm Pt) catalysts. 

60
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 7. In-situ FTIR CO adsorption data for fresh (1 nm Pt) 2%Pt/γ-Al2O3 at 85°C (Ads_F_85) and 

adsorption data at 25 and 85°C for the catalyst used for 100% CO oxidation tests 3 times (Ads_U_25 and 

Ads_U_85). 

1.2 

(A) (B) 

Figure 8. ACEM HAADF-STEM images where the white dots are Pt atoms and grey area is the alumina 

support. (A) Image of fresh 2%Pt/γ-Al2O3 showing both single Pt atoms and small Pt cluster. (B) Image of 

the exact same spot (for reference squares indicate same area) after exposure to CO oxidation initiation 

conditions. 

61
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

1.2 nm 

1.6 nm 

5 nm 5 nm 

(A) (B) 

Figure 9. ACEM HAADF-STEM images where the white dots are Pt atoms and grey area is the alumina 

support. (A) Image of fresh 2%Pt/γ-Al2O3 showing mostly small Pt clusters of uniform size. (B) Image of 

the exact same spot (for reference circles indicate same area) after exposure to CO oxidation initiation 

conditions shows a small increase in some of the Pt clusters. 

Metallic Pt 

2 nm 2 nm 

(A) (B) 

Figure 10. ACEM HAADF-STEM images where the white dots are Pt atoms which form clusters and 

the grey area is the alumina support. (A) Image of fresh 2%Pt/γ-Al2O3 showing both single Pt atoms and 

small Pt cluster. (B) Image of the exact same spot (for reference circles indicate same area), after 

exposure to CO oxidation initiation conditions, shows an increase in some of the Pt clusters to metallic Pt 

particles about 2 nm in size. 

62
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 11. The adsorption energies of O, O2, CO, CO2, NO, and NO2 on Ptx nanoclusters plotted against 

the number of Pt atoms in the Ptx clusters. 

Figure 12. The reaction energy profiles for CO + 1⁄2O2 Æ CO2 and NO + 1⁄2O2 Æ NO2 on Ptx (x = 1-5 and 

10). The profiles on the bulk Pt surface (represented by the Pt(111) surface) are included. 

Number of Pt atoms (x) 

... ∞ 

O CO 

Number of Pt atoms (x) 

... ∞ 

Figure 13. The adsorption energies of O and CO on Pt (Ptx) and Pt oxide (PtxOx and PtxO2x) clusters, 

plotted against the number of Pt atoms in the clusters. 

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Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 14. X-ray diffraction pattern of fresh 2% Pt/γ-A2O3 annealed at 650 °C (a) fresh and after CO 

oxidation tests: CO conversion (b) initiation (8.1%) (c) full range (3 cycles at 100%).* denotes a sample 

holder impurity. 

Figure 15. In-situ FTIR CO adsorption data for fresh (25 nm Pt) 2%Pt/γ-Al2O3 at 85°C (Ads_F_25 and 

Ads_F_85) and adsorption data at 25 and 85°C for the catalyst used for 100% CO oxidation tests 3 times 

(Ads_U_25 and Ads_U_85). 

64 
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 16. STEM image of fresh 2%Pt/θ-Al2O3 (Left), the Pt particles size distribution for this area 

(Right). 

Figure 17. Initiation of catalyst conversion as a function of temperature comparing the different 

initiation temperature required for 1 nm Pt on θ-Al2O3, 1 nm Pt on γ-Al2O3 and 25 nm Pt on γ-Al2O3. 

65
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 18. 100% CO oxidation cycled 3 times presented as a function of temperature comparing the 

different temperature ramps required for 1 nm Pt on θ-Al2O3, 1 nm Pt on γ-Al2O3 and 25 nm Pt on γAl2O3. 

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Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

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

Edgar Lara-Curzio, Amit Shyam and Thomas R. Watkins 

Oak Ridge National Laboratory 

P.O. Box 2008, MS-6069 

Oak Ridge, TN 37831-6069 

(865) 574-1749; fax: (865) 574-6098; e-mail: laracurzioe@ornl.gov 

DOE Technology Development Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@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 

• Implement test techniques to characterize the physical and mechanical properties of ceramic diesel particulate 

filters (DPFs) and develop analysis and inspection tools for assessing their reliability and durability. 

Approach 

• Implement test techniques to determine the physical and mechanical properties of DPF ceramic substrates. 

• Application of probabilistic design tools to DPF ceramic substrates. 

Accomplishments 

• Characterized the fracture mechanical behavior of diesel particulate filter substrates in relevant conditions. 

• Carried out physical and mechanical property measurements to better understand the time dependant and 

fracture mechanical behavior of DPF substrates. 

Future Direction 

• Implementation of characterization techniques to assess the integrity of DPF ceramic substrates. Inspection 

schedules would be determined based on the developed probabilistic analysis tools. 

Introduction 

The pollution emitted by diesel engines 

contributes greatly to the nation’s air quality 

problems. Even with more stringent environmental 

regulations set to take effect in 2007, existing trucks 

and buses will continue to emit nitrogen oxides 

(NOx) and particulate matter (PM), both of which 

contribute to serious public health problems

1

. There 

are several technologies designed to reduce pollution 

from existing trucks and buses; prominent among 

these technologies are 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, hydrocarbons, and 

CO by 60 to 90%. 

Most DPFs consist of a ceramic honeycomb 

with hundreds of cell passages partitioned by walls 

(Figure 1). Each cell passage has a square cell 

67
Heavy Vehicle Propulsion Materials 

opening at one end and is closed at the other end so 

that the cell passages are alternately closed at each 

end. This structure forces the exhaust gases through 

the porous, thin ceramic honeycomb walls. When 

the gases carrying the carbon particles flow through 

the fine pores of the walls, the carbon particles are 

filtered out. High porosity values, in the range of 

60%, heighten filtration efficiency to more than 90% 

while reducing gas-flow resistance to avoid affecting 

the engine performance. 

The process of diesel PM collection continues 

while the engine is operating. Carbon particles are 

Figure 1. Corning’s cordierite-based DPFs. 

collected on the ceramic walls and as a result, the 

backpressure of the system increases. This problem 

is alleviated by burning the trapped PM, through a 

catalytic reaction using exhaust gas heat at 400°C or 

more, into CO2 and water vapor. This process, called 

regeneration, results in a cleaner filter. The 

regeneration process is dependent upon exhaust 

temperature, oxygen, NOx content, time, and PM 

levels. 

FY 2006 Progress Report 

The key to the successful application of DPFs is 

to reliably regenerate the filter (e.g., to burn the PM 

that the filter continues to trap or collect). 

Traditionally, combustion of soot is done in an 

oxygen atmosphere (air). In air, soot will burn at 

about 500°C. However, this is not a typical 

operating temperature for diesel engine exhaust. 

Therefore, to burn soot in air, an active system—i.e., 

one that increases the temperature of the exhaust 

using some external heat source—is required. But if 

an active system is not carefully controlled, or if too 

much PM collects on the filter walls, the filters can 

experience an “uncontrolled burn” where the 

temperature increases to 600°C or more, resulting in 

damage to the filter element. 

The objective of this project is to develop and 

implement methodologies to predict the reliability of 

DPFs. A useful conceptual model for this purpose is 

the reliability bathtub curve, which describes 

reliability-related phenomena of a component over 

its life cycle.2

 A schematic of the reliability bathtub 

curve is depicted in Figure 2. It consists of three 

stages: the infant mortality phase is characterized by 

premature failures due to improper manufacturing or 

assembly, poor workmanship, or defects introduced 

during processing. The second stage of the curve 

corresponds to the useful life of the component and 

Figure 2. Reliability bathtub curves. 

is characterized by a constant failure rate. In this 

regime, failures are typically associated with random, excessive loads. If sufficiently high safety 

factors are used during the design process, the 

magnitude of this failure rate should be negligible. 

The third stage of the bathtub curve is known as the 

wear-out phase, where the failure rate increases with 

time as a result of aging phenomena. Aging phenomena include thermal and mechanical fatigue, 

68
Heavy Vehicle Propulsion Materials 

corrosion, creep deformation and environmentally 

assisted crack growth. The time at the onset of wearout is often regarded as the useful life of the 

component. 

The reliability bathtub curve can be used as a 

descriptor of how the failure rate of DPFs evolves 

over time. Infancy failures of DPFs could be related 

to manufacturing or process defects (large pores, 

inclusions, cracking) or defects introduced during 

assembly. Failures of DPFs during their useful life 

will be dictated by the intersection between the 

spectrum of thermomechanical loads and the distribution of DPF strengths. Such failures could result 

from excessive vibration, for example, or unwanted 

thermal excursions during transients associated with 

regeneration. Wear-out and degradation failures of 

DPFs could be associated with the growth of microcracks assisted by thermal fatigue, and/or by the 

chemistry of the environment, and/or by chemical 

and microstructural changes in the material due to 

long-term exposure to elevated temperatures in the 

exhaust environment. 

The objective of this project is to implement test 

techniques to characterize the physical and 

mechanical properties of ceramic DPF substrates 

and to develop analysis and inspection tools for 

assessing the reliability and durability of DPFs. The 

developed tools and methods would allow the design 

of 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 and stiffness. 

Designing mechanically reliable DPFs is important 

because these components will experience 

demanding thermomechanical conditions during 

service. These include, for example, thermal shock 

resulting from rapid heating/cooling and thermal 

stresses that arise from temperature gradients. The 

approach that will be followed in this project 

includes identifying and implementing test 

techniques for the 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 will also be determined 

and identified. These properties include thermal 

FY 2006 Progress Report 

expansion, thermal conductivity, heat capacity, 

density, porosity, elastic properties, strength, 

fracture toughness, and resistance to crack growth at 

ambient and elevated temperatures, in air and in 

relevant environments. 

The information generated will be used in turn 

to implement probabilistic design tools. In particular, the applicability of the CARES code3

 to predict 

the reliability of DPF ceramic substrates will be 

investigated. Such probabilistic design methodologies are based on a combination of experimentally 

determined strength data, stress analyses of the 

component using a finite-element analysis, and 

selection of appropriate failure criteria. The durability (service life) of the component can also be 

predicted using this framework by considering the 

mechanisms that are responsible for the degradation 

of material strength, such as slow crack growth or 

creep. 

Results 

During FY 2006, the procedure to prepare 

double-torsion (DT) test specimens was further 

developed and this testing technique was applied 

determine the fracture toughness and slow-crackgrowth behavior of porous cordierite. The doubletorsion test configuration is particularly attractive for 

determination of the fracture mechanical properties 

of porous cordierite used for fabrication of DPFs. 

This is due to the following reasons (a) crack length 

measurements are difficult in porous cordierite and 

DT characterization does not require crack length 

measurements, (b) the DT specimen geometry is 

ideally suitable for plate-like specimens that can be 

fabricated from DPF walls and (c) the DT 

configuration is also attractive for slow-crackgrowth studies because of the relative stability of 

crack extension, in contrast to other testing 

configurations (e.g., SENB, CT). 

The preparation of double-torsion test specimens 

of porous cordierite has been reported earlier.4 The 

sample preparation technique was improved further 

by development of a vacuum fixture to grip 

specimens while grinding and notching. The fixture 

increased the parallelism of the prepared double 

torsion test specimens to 5 μm. The mean thickness 

of the DT test specimens was equal to 0.375 mm and 

reproducible to within 10 μm. The consistency of 

the thickness (and hence volume) of the doubletorsion specimens allowed the measurement of the 

69
Heavy Vehicle Propulsion Materials 

porosity via high precision mass measurements prior 

to notching the specimens. The porosity of the 

specimens determined by this means was 

50.00+0.34%. Resonant ultrasound spectroscopy 

(RUS) of the rectangular plates was also performed 

prior to notching, in order to obtain an elastic 

property signature of all the test specimens. Two 

independent parameter elastic constant fits led to an 

average elastic modulus of 12.33+0.26 GPa and an 

average Poisson’s ratio of 0.24+0.04. 

The fracture toughness of porous cordierite test 

specimens was determined by loading pre-cracked 

test specimens to failure at a constant crosshead displacement rate of 0.01 mm/s. The fracture toughness 

was calculated from the peak load using Equation 1: 

⎡3 1+ ν ⎤

1 / 2 ( ) KIC = PSm⎢ ⎥ , (1)

⎣ St4

ξ ⎦ 

where P is the peak load, S is the width of the test 

specimens, Sm is the moment arm, t is the specimen 

thickness, ν is Poisson’s ratio, and ξ is a finite beam 

thickness correction factor given by 

⎛ πS ⎞ ⎛ t ⎞ ⎛ t ⎞ ⎝

⎜− 2t ⎠

⎟

ξ = 1 − 1.26 ⎜ ⎟ + 2.4 ⎜ ⎟e .

 (2) ⎝ S ⎠ ⎝ S ⎠ 

Figure 3. Fracture toughness of porous cordierite as a 

function of temperature. 

In addition, the slow crack growth curves and 

the slow crack growth exponent were determined by 

application of the double-torsion load-relaxation 

method. The fracture toughness and the slow crack 

FY 2006 Progress Report 

growth behavior were determined at temperatures up 

to 900o

C in ambient, water-vapor and simulated 

diesel engine exhaust environments. The effect of 

temperature on the fracture toughness of porous 

cordierite is illustrated in Figure 3. The fracture 

toughness versus temperature variation of porous 

cordierite is unusual, in that the fracture toughness 

decreases from 0.38 ± 0.02 MPa m at 20°C to 0.30 

± 0.06 MPa m at 500°C and subsequently 

increases with temperature to 0.37 ± 0.03 MPa m 

at 800°C. The fracture toughness value at 900°C 

was obtained as 0.35 ± 0.03 MPa m . As expected, 

porosity plays an important role in determination of 

the fracture toughness. Cracks extend by joining 

pores that are favorably oriented with respect to the 

loading direction. The increase of fracture toughness 

with temperature above 500o

C was attributed to the 

presence of a glassy phase at the grain boundary. 

Figure 4 compares the fracture surface of doubletorsion specimens tested at 20 and 900o

C, 

respectively. At temperatures above 500o

C, periodic 

undulations of the fracture surface such as that 

shown in Figure 4(b) were observed. These features 

are formed due to the activation of viscoplastic 

processes at elevated temperatures, possibly due to 

the presence of a glassy phase. 

Direct evidence of the amorphous phase at the 

grain boundary is illustrated in the high resolution 

transmission electron micrograph (HRTEM) shown 

in Figure 5 (region between arrows). The thickness 

of the grain boundary phase varied between 5-20 

nm. Electron dispersive spectroscopy (EDS) 

revealed that the amorphous grain boundary phase 

was rich in silica when compared to the crystalline 

cordierite phase. The amorphous phase is beneficial 

for the elevated temperature fracture resistance due 

to the softening of the grain-boundary phase. It is, 

however, well known that the amorphous phase 

decreases the slow crack growth resistance of 

silicate ceramics. The grain boundary phase 

characteristics, therefore, need to be carefully 

tailored in order to obtain optimum fracture 

resistance over the range of operating conditions. 

The double-torsion test specimens in Figure 3 

came from the same region of the filter. It is not 

necessary that the characteristics of the substrate 

will remain unchanged in different locations of the 

filter. During FY2006, studies were conducted that 

determined the properties of diesel particulate filter 

substrates as a function of location in the filter. 

70
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 6 contains a series of micrographs that 

illustrate the influence of the location of the filter on 

the porosity of the walls. The porosity of the 

Figure 5. HRTEM image of the silica-rich amorphous 

phase at the grain boundary of cordierite. 

Figure 4. High resolution SEM images comparing the 

fracture surfaces of failed DT specimens at (a) 

20 and (b) 900o

C. 

71
Heavy Vehicle Propulsion Materials 

specimens was determined by image analysis and 

helium pycnometry. The porosity remained close to 

50% in all cases, except for tangential specimens 

close to the edge of the filter (number 13 in Figure 

6) where the porosity dropped to about 36%. Both 

image analysis on SEM micrographs and helium 

pycnometry showed similar trends except that the 

helium density measurements yielded slightly lower 

porosity (by 3-4%) values. This could be due to the 

fact that the closed, unconnected pores are not 

accounted for in the helium density measurements. 

During FY2006, the effect of thermal cycling on 

the elastic and thermal properties was characterized. 

The effect of thermal cycling on the elastic 

properties of 2x2 cordierite rectangular bars is 

presented in Figure 7. The modulus measurements 

were performed in three-point bending in a dynamic 

mechanical analyzer (DMA). The results indicate 

that the elastic properties degradation is strongly 

related to the maximum temperature in the cycle. 

Figure 7 demonstrates that damage in the 

specimen is higher for a maximum temperature of 

1000o

C compared to 800oC. Moreover, at the 

thermal loading rate (approximately 40o

C/min) 

applied, most of the damage occurs in the first few 

FY 2006 Progress Report 

cycles and subsequent degradation in elastic 

properties is minimal. Thermal conductivity was 

also measured for the same specimens with the hot 

disk method and it showed a similar variation with 

thermal cycling. 

 The thermal expansion and phase composition 

of the porous cordierite was characterized during 

FY2006. The thermal expansion behavior of porous 

cordierite, parallel to the extrusion direction, is 

illustrated in Figure 8. It can be seen that cordierite 

has a negative thermal expansion coefficient at 

lower temperatures and above 300o

C; the thermal 

expansion coefficient becomes positive. The net 

thermal expansion coefficient is low and this results 

in excellent resistance to thermal cycling as 

illustrated in Figure 7. 

Cordierite exists in more than one allotropic 

form. The phase composition cannot be resolved 

with typical laboratory X-ray diffraction (XRD) due 

to the low peak resolution and to the fact that the 

closely spaced overlapping XRD peaks of the 

orthorhombic crystal structure, making it 

indistinguishable from the hexagonal structure 

cordierite. During FY2006, synchrotron XRD 

72
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 6. Scanning electron micrographs illustrating the variation of porosity with location in the filter. 5 represents 

the center of the filter and 13 represents the edge. R and T stand for radial and tangential locations, 

respectively. 

Figure 8. The thermal expansion behavior of porous 

Figure 7. Effect of thermal cycles on elastic property cordierite for a test specimen parallel to the 

degradation in porous cordierite. extrusion direction. 

73
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

revealed that the porous cordierite material studied 

has the orthorhombic (ordered) crystal structure. The 

existence of porous cordierite in the ordered 

orthorhombic form under ambient conditions rules 

out the possibility that the mechanical property 

variations (e.g., Figure 3) are due to phase 

transitions in the crystal structure. The 

orthorhombic phase is thermodynamically stable up 

to temperatures much higher than the operating 

service temperatures of DPFs. 

Summary 

The fracture mechanical behavior of porous 

cordierite was characterized in relevant conditions. 

It was observed that the fracture toughness variation 

with temperature was unusual, in that the fracture 

toughness values increased with temperature above 

500o

C. The amorphous phase at the grain boundary 

was characterized and found to increase the fracture 

toughness of porous cordierite at elevated 

temperature. Additional physical and mechanical 

property measurements were carried out to better 

understand the time dependant, spatial and fracture 

mechanical behavior of DPF substrates. 

References 

1. United States Environmental Protection 

Agency, Office of Transportation and Air Quality 

EPA420-F-03-017, June 2003. “Questions and 

Answers on Using a Diesel Particulate Matter Filter 

in Heavy-Duty Trucks and Buses.” 

2. G. S. Wasserman, Reliability Verification, 

Testing and Analysis in Engineering Design, Marcel 

Dekker, New York, 2003. 

3. N. N. Nemeth, L. P. Powers, L. A. Janosik, 

and J. P. Gyekenyesi, “Ceramics Analysis and 

Reliability Evaluation of Structures Life Prediction 

Program: Users and Programmers Manual,” NASAGRC, Life Prediction Branch (1993). 

4. A. Shyam, E. Lara-Curzio, H.-T. Lin, and 

R.J. Parten, “Fracture Toughness of Porous 

Cordierite,” in Proceedings of 30th International 

Symposium on Advanced Ceramics and 

Composites, Cocoa Beach, 2006 (to be published). 

Collaborators 

H-T. Lin, Randy Parten, and Claire Luttrell. 

Presentations and Publications 

A. Shyam and E. Lara-Curzio, “The doubletorsion testing technique for determination of 

fracture toughness and slow crack growth behavior 

of materials: a review,” Journal of Materials 

Science, vol. 41, no. 13, pp. 4093-4104, July 2006. 

A. Shyam, E. Lara-Curzio, H.-T. Lin and R. J. 

Parten, “Fracture toughness of porous cordierite,” to 

be published in the proceedings of the Cocoa Beach 

Meeting of the American Ceramic Society, 2006. 

74
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

4. MATERIALS FOR AIR HANDLING AND THERMAL 

MANAGEMENT 

A. Austenitic Stainless Steels Alloys for Exhaust Manifold 

P. J. Maziasz, J. P. Shingledecker, and N. D. Evans 

Materials Science and Technology 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: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

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

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Objectives 

• Commercialize new CF8C-Plus cast austenitic stainless steel. 

• Complete the high-temperature properties database for test specimens from new commercial heats of CF8CPlus Cu/W and determine the properties of prototype components of both CF8C-Plus and CF8C-Plus Cu/W. 

Approach 

• Support Caterpillar (CAT) evaluation efforts to upgrade turbochargers from cast iron to cast austenitic stainless 

steel. 

• Perform and evaluate direct comparisons between commercial CF8C-Plus steel and other alloys used by 

component suppliers/end-users to establish performance benefits at 700–850°C. 

Accomplishments 

• Provided creep-rupture resistance data on CF8C-Plus Cu/W and CF8C-Plus to further the interest of major 

diesel turbocharger OEMs to upgrade housings. 

• Contracted trial-licensee Stainless Foundry & Engineering to produce ASTM weldability data, and to cast thinwalled step castings of CF8C-Plus and CF8C-Plus Cu/W in support of exhaust component applications. 

Future Direction 

• Finish all necessary properties measurements and microcharacterization before CRADA ends in July, 2006. 

• Support commercial licensing for diesel turbocharger housings, and other applications. 

75
Heavy Vehicle Propulsion Materials 

Introduction 

Advanced heavy truck diesel engines are 

increasingly required to have higher fuel efficiency 

and reduced exhaust emissions without sacrificing 

durability and reliability. The most demanding 

normal duty cycles require exhaust manifolds and 

turbocharger housing materials to withstand temperatures ranging from 70 to nearly 800°C. Such 

materials must withstand prolonged, steady hightemperature exposure (which requires tensile 

strength and creep and oxidation resistance) as well 

as more rapid and severe thermal cycling (which 

requires resistance to aging and thermal fatigue). 

New emissions reduction technology and transient 

power excursions can push temperatures in these 

critical components even higher. 

Higher diesel engine exhaust temperatures push 

components made from SiMo ductile cast iron well 

beyond their current strength and corrosion limits. 

The first Oak Ridge National Laboratory 

(ORNL)/CAT cooperative research and development agreement (CRADA) (3y) developed a new, 

modified CF8C cast austenitic stainless steel (CF8CPlus) as an upgrade alternative to SiMo cast iron and 

performed the initial mechanical properties testing 

on a lab-scale heat. The new CF8C-Plus steel was 

found to have outstanding creep resistance at 850°C. 

The second ORNL/CAT CRADA project (3y) 

capitalized on the initial success and began 

commercial scale-up of the new CF8C-Plus. The 

new material was found to have good castability, as 

well as much better mechanical properties at 600– 

850°C than standard CF8C steel, including 

resistance to aging, creep, and fatigue/thermal 

fatigue. The new CF8C-Plus cast steel won a 2003 

R&D 100 Award, and in 2004 testing continued of 

different commercial casting process effects (static, 

centrifugal, and larger heats). 

 In 2005, efforts to commercialize CF8C-Plus 

expanded dramatically, so this CRADA project was 

extended for 2 more years. Commercial interest to 

upgrade diesel turbocharger housing intensified in 

2006. In 2005, successful casting trials of CF8CPlus were carried out for components ranging from 

exhaust manifolds to industrial turbine casing. In 

addition, the first commercial heats of CF8C-Plus 

Cu/W were made in 2005 to boost strength and 

creep resistance at 750–850°C, and were tested in 

2006. Three companies have taken trial licenses and, 

FY 2006 Progress Report 

to date, over 31,000 lb of the new steel has been 

produced commercially. 

Approach 

Three U.S. foundries obtained trial licenses 

for CF8C-Plus steel in 2005: MetalTek International, Wollaston, and Stainless Foundry and Engineering. CAT and ORNL both ordered additional 

heats of CF8C-Plus and CF8C-Plus Cu/W steels in 

2005, and cast the pieces in various sizes, using both 

static casting in sand molds and centrifugal casting 

of rings of varying thickness, to enable more 

comprehensive mechanical properties testing. Creep 

evaluate for further improvements in strength and 

creep-resistance in the CF8C-Plus Cu/W steel at 

750o

C and above began in 2006. Finally, in 2006, 

interest in CF8C-Plus steel as an upgrade 

replacement for SiMo cast iron for turbocharger 

housing intensified by the turbocharger OEMs. 

Upgrading various critical exhaust components to 

stainless steel should enable the increased durability 

at the increased engine temperatures needed to allow 

a 3% decrease in fuel consumption for on-highway 

trucks. Both CAT and ORNL are performing 

specific mechanical properties testing and 

microcharacterization in support of these various 

applications or efforts. 

Technical Progress 

Caterpillar 

The Caterpillar Technical Center completed 

characterization and properties measurements on a 

large industrial turbine end-cover casting of CF8CPlus steel (8000-lb heat to yield 6700-lb finished 

component), made by MetalTek International for 

Solar Turbines (a Caterpillar Company). No defects 

were obvious visually, and sectioning revealed no 

cracking or hot tearing at the side flange. This is significant because fully austenitic alloys such as 

CF8C-Plus are usually more susceptible to hot tearing than similar steel with a much higher δ-ferrite 

content (standard CF8C steel has 15–25% such 

phase content). Mechanical properties testing 

revealed properties generally consistent with other 

data on CF8C-Plus. This particular demonstration of 

good properties and castability in making thicksection components (slow cooling rates) compared 

with smaller test castings (faster cooling rates) is 

76
Heavy Vehicle Propulsion Materials 

stimulating commercial interest for larger gasturbine and advanced steam turbine components. 

CF8C-Plus

 steel 

SiMo 

cast-iron 

Figure 1. The first commercial application of CF8CPlus is the Caterpillar Regeneration System 

(CRS) housing, announced in July, 2006. This 

combustor ignites exhaust gas and injected 

diesel fuel to clean filters, and must withstand 

high temperatures and rapid thermal cycling. 

The 21st Century Truck Partnership technical 

goal of developing emission-compliant engine systems for Class 7–8 highway trucks is an improvement in engine efficiency from 42 to 50% by 2010. 

CAT has determined that replacing various SiMo 

cast iron exhaust components with CF8C-Plus cast 

stainless steels allows a 90°C increase in exhaust gas 

temperatures, which translates directly into a 3% 

increase in fuel economy while also significantly 

improving performance and durability. CF8C-Plus 

stainless steel was chosen for Caterpillar 

Regeneration System (CRS) housing in July, 2006. 

This component is showing in Fig.1. The CSR will 

be standard equipment on all heavy-duty highway 

truck diesel engines starting in 2007. 

Efforts also continued at CAT to support and 

advance CF8C-Plus steel upgrades for other exhaust 

components. Interest by the turbocharger OEMs 

intensified in 2006 for upgrading turbocharger 

housings from SiMo cast iron to CF8C-Plus steel, 

and CAT and ORNL continue their technical efforts 

to support such new applications. 

FY 2006 Progress Report 

ORNL 

In 2005, ORNL began the longer term creep 

testing to prove the overwhelming advantage of 

CF8C-Plus steel over other stainless steels and 

alloys at 700-850°C, which continued during 2006. 

ORNL also began similar creep-testing of the new 

CF8C-Plus Cu/W alloy. The creep-rupture strength 

advantage of CF8C-Plus over standard CF8C steel 

and Ni-resist austenitic cast-iron (an upgrade for 

some applications relative to SiMo cast-iron) is 

shown in the Larson-Miller parameter (LMP) plot in 

Figure 2. CF8C-Plus cast steel is much stronger than 

either SiMo or Ni-resist cast-irons, and it is almost 

twice as strong as standard CF8C steel. Figure 3 

shows the new creep-rupture data on CF8C-Plus 

Cu/W at 750o

C and 100 MPa. The CF8C-Plus 

Cu/W has an even larger improvement in creeprutpure resistance compared to standard CF8C steel, 

and has about 8 times longer rupture life, but almost 

20 times longer secondary creep regime to about 2% 

strain. Despite the increased creep strength, the 

CF8C-Plus Cu/W also still has good rupture 

ductility. 

Finally, comparison of CF8C-Plus and CF8CPlus Cu/W cast stainless steels on an LMP plot 

(Figure 4) with alloy 617 shows that these new cast 

stainless steels come close to the creep-strength of 

alloy 617, a Ni-Cr-Co superalloy (which is about 10 

times more expensive than the cast stainless steel), 

particularly at the highest temperatures. Such 

properties comparisons are very important to 

considering CF8C-Plus for potential gas- and steamturbine applications, as well as many other 

chemical/petrochemical, energy, or metals 

processing applications 

77
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Temperature ( oC) for Rupture in 10,000 Hours 

500 550 600 650 700 750 800 850 

Stress (

MPa) - Log 

Scale

100 

10 

18,000 20,000 22,000 24,000 26,000 28,000 

SiMo 

Cast Iron 

Ni-Resist 

Std. CF8C 

CF8C-Plus 

D5S* 

D5B** 

*NiResist D5S Mo from Waukasha Manifold 

**NiResist D5B Mo from Stainless Foundry 

Larson-Miller Parameter (K, C=20) 

Figure 2. Creep-rupture stress plotted vs LMP for various ORNL creep-rupture 

tests at 500–850°C in air of various commercial heats of SiMo ductile and 

Ni-resist austenitic cast irons and standard CF8C and CF8C-Plus. CF8CPlus shows significant creep-strength advantage over the standard steel or 

the cast irons. 

Figure 3. Plots of creep-strain vs time for heats of standard CF8C, CF8C-Plus and 

the new CF8C-Plus Cu/W, all tested in air at 750°C. Clearly the CF8CPlus Cu/W has much more creep-resistance, with an extended secondary 

creep-regime and a very low steady-state creep rate. 

78
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

TEMPERATURE (oC) FOR RUPTURE IN 100,000 HOURS

600 640 680 720 760 800 840

300

100

80

60

40

20

21,000 22,000 23,000 24,000 25,000 26,000 27,000 28,000 

Alloy 617 

CF8C-Plus 

CF8C-Plus (in-test) 

CF8C-Plus+Cu/W 

CF8C-Plus+Cu/W (in test) 

LMP (C=20) 

Alloy 617 

ST

RES

S (

M

Pa) 

Figure 4. Creep-rupture stress plotted vs LMP for various ORNL creep-rupture tests 

at 700–850°C in air of CF8C-Plus and CF8C-Plus Cu/W. CF8C-Plus has 

better creep-strength than most wrought heat-resistant austenitic stainless 

steels and alloys and compares well with alloy 617, a Ni-Cr-Co superalloy. 

79
Heavy Vehicle Propulsion Materials 

In FY 2006, ORNL completed some initial 

testing of the new CF8C-Plus steel with the Diado 

KN2 cast stainless steel and other steels and alloys 

that are being considered as turbocharger housing 

upgrades relative to SiMo cast iron. The hightemperature creep-strength, fatigue and thermal 

fatigue properties of CF8C-Plus steel has stimulated 

considerable interest from several different 

turbocharger OEMs. At the end of FY2006, ORNL 

had Stainless Foundry & Engineering make stepcasting of various thicknesses to provide thin-walled 

test material of both CF8C-Plus and CF8C-Plus 

Cu/W to provide data to support the diesel 

turbocharger housing application. Testing will 

begin next year. 

Conclusions 

ORNL and CAT in 2006 supported and 

expanded commercialization of the CF8C-Plus cast 

stainless steel. To date, more than 31,000 lb of 

commercial CF8C-Plus stainless steel has been cast 

for various applications, and all prototype 

component castings have been successful on the first 

trial, as a result of good castability. Caterpillar 

chose CF8C-Plus steel as the new CRS housing 

material in July, 2006. CRS components will be 

standard equipment on heavy-duty highway truck 

diesels beginning in 2007. In 2006, interest 

intensified by several turbocharger OEMs in CF8CPlus stainless steel as an upgrade to SiMo cast iron 

for the housing. Testing of CF8C-Plus for diesel 

exhaust manifold applications also continues. 

In 2006, creep-rupture testing of a new 

commercial heat of CF8C-Plus Cu/W produced by 

MetalTek International showed significant 

improvement compared to CF8C-Plus steel. In 2006 

thin-walled castings of CF8C-Plus and CF8C-Plus 

Cu/W were made by Stainless Foundry & 

Engineering, and will be tested year in support of the 

turbocharger housing applications. 

Publications/Presentations 

J.P. Shingledecker, P.J. Maziasz, N.D. Evans, M.L. 

Santella, and M.J. Pollard, “CF8C-Plus: A New 

High Temperature Austenitic Casting Alloy for 

Advanced Power Systems,” Energy Materials 1 

(2006) 25-32. 

J.P. Shingledecker, P.J. Maziasz, N.D. Evans, and 

M.J. Pollard, “Creep Behavior of a New Cast 

FY 2006 Progress Report 

Austenitic Alloy,” to be published in International 

Journal of Pressure Vessels and Piping (Special 

issue on New Steels) in 2006. 

Special Recognitions and Awards/Patents Issued 

A new, second U.S. Patent Application was filed 

in July 31, 2006 by Caterpillar, Inc. for 11/495,671 

“Heat and Corrosion Resistant Cast Austenitic 

Stainless Steel Alloy With Improved High 

Temperature Strength,” by P.J. Maziasz, M.J. 

Pollard and J.P. Shingledecker. 

80
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

B. Friction and Wear Reduction in Diesel Engine Valve Trains 

Peter J. Blau with 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 Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@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 

• To enable the application of next-generation diesel engine materials by developing test 

methods which simulate the mechanical, thermal, and chemical environments of wear-prone 

engine parts, like exhaust valves and exhaust gas recirculation (EGR) systems. 

• To conduct friction and wear characterizations of promising new engine component 

materials, surface treatments, composites, and coatings with a goal of understanding the 

scientific basis for their performance in severe environments. 

Approach 

• Identify diesel engine components that require durability and low friction in order to improve 

propulsion system energy-efficiency and reduce emissions. These include wastegate 

bushings for exhaust gas recirculation (EGR) components, fuel injector plungers, and valve 

and valve seat materials. 

• Develop test methods, data analysis methodology, and models to evaluate and quantify the 

performance of candidate materials under simulated use conditions. 

• Investigate and understand the characteristics of friction and wear transitions that affect the 

lifetimes of materials in wear-critical applications like valve trains. 

Accomplishments 

• Designed, built, and used a high-temperature oscillatory scuffing test system that operates at 

wastegate bushing temperatures (~ 600 – 700o

C). 

• Published results of tests on a range of metallic alloys, ceramics, and coatings to determine 

which of these had the best durability under high-temperature conditions. 

81
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

• Developed a novel, ‘pin-on-twin’ scuffing test to evaluate fuel injector materials in diesel fuel 

and low sulfur fuel environments. 

• Developed criteria for the on-set of scuffing damage and evaluated traditional steel fuel 

injector materials as well as ceramics, advanced cermets, and hard coatings in diesel fuel and 

low sulfur fuel. Represented results in terms of ‘scuffing maps’ and transition diagrams. 

• Developed and published a model for scuffing tendency that considers both lubricant 

characteristics and solid material characteristics. 

• Designed a high-temperature, controlled environment exhaust valve wear testing system to 

investigate the combined influences of repetitive impacts, elevated temperatures, and gaseous 

environments on the durability of candidate valve and seat materials. 

Future Directions 

• Use the test data from the new instrument to assess implications of material and surface 

treatment selection for maintaining engine efficiency and reducing diesel engine emissions. 

Introduction 

The diesel engine industry faces the 

important challenge of improving fuel 

efficiency while meeting increasingly strict 

emissions regulations. These challenges are 

being addressed by modifying engine 

designs, introducing homogeneous 

combustion, designing advanced control 

systems, and using exhaust gas aftertreatments. Such modifications affect the 

mechanical, thermal and chemical 

environments in which the engine materials 

must operate, and currently-used materials 

may fall short of meeting these needs. New 

materials or engineered surfaces are needed 

to enable new technological goals. 

The objective of this effort is to enable 

the selection and use of durable, lowerfriction 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 underlying 

approach involves developing simulative 

test methods, analyzing microstructures of 

candidate materials, developing design maps 

with variables like surface finish and 

frictional behavior, and modeling the 

damage process itself. 

In prior years, the focus of this effort 

was on wastegate bushings for exhaust gas 

recirculation (EGR) systems and fuel 

injector plungers that must operate in low 

sulfur fuels. A variety of materials was 

evaluated using specially-designed tests to 

simulate the operating environment of the 

components, and data that compared 

traditional materials with potential 

replacement materials was analyzed and 

shared with industry. For example, 

computer-based techniques for measuring 

the onset of scuffing damage were 

developed, as was a model to describe the 

progression of surface damage under 

boundary-lubricated sliding. This work 

resulted in seven journal publications and 

reports. 

In FY 2006, the focus of this effort was 

redirected to address the selection and use of 

materials for the demanding environment of 

diesel engine exhaust valves, where 

temperatures can reach over 850o

 C and 

materials are subjected to both contact 

fatigue and deformation in the presence of 

gaseous environments. The challenge was 

to investigate candidate materials under 

82
Heavy Vehicle Propulsion Materials 

operating conditions that provided a 

reasonable simulation of their performance 

in engines but without the high cost and 

complications of full-scale engine tests. 

Approach 

The four-pronged approach involves (a) 

defining the tribosystem operating 

environment, (b) designing and building an 

exhaust valve testing system, (c) conducting 

tests to compare current valve materials with 

candidate alternative materials and surface 

treatments, and (d) providing an assessment 

of the candidate materials’ response to 

engine manufacturers and designers. This 

annual report describes background 

considerations and recent progress in 

designing a high-temperature, controlledatmosphere, repetitive impact testing 

apparatus with which to study valve 

materials. It considers issues such as contact 

geometry, motion control, thermal and 

environmental considerations, and it 

indicates the progress to-date. 

1.6 

1.5 

1.4 

1.3 

( N / P ) 

1.2 

1.1 

1.0 

0.9 

0.8 

FY 2006 Progress Report 

When designing a high-temperature, 

repetitive impact (HTRI) testing system, it 

was necessary to consider the damage 

modes that exhaust valves and valve seats 

experience. Valve face abrasion is caused 

by fine particles from residual deposits. 

Adhesive wear and micro-welding can occur 

to produce pitting and localized pull-out of 

material. Oxidation and sulfidization can 

attack the mating surfaces, introducing a 

tribochemical aspect into the HTRI design. 

The ratio of the normal force on a valve 

seat (N) to the closing force collinear with 

the stem (P) was discussed in a paper by 

Giles [N. Giles (1966) “Fundamentals of 

Valve Design and Material Selection,” SAE 

Paper 660471]. From the contact geometry, 

(N/P) = 1 / (cos θ + μ sin θ) (1) 

where θ is the seat angle and μ is the 

friction coefficient between the valve and 

valve seat. Figure 1 depicts the relationship 

expressed in Equation (1) for three common 

valve seat angles. 

θ = 20 deg 

θ = 45 deg 

θ = 30 deg 

0 0.1 0.2 0.3 0.4 0.5 0.6 

μ 

Figure 1. Relationship between the normal force on the seat (N) and the axial force (P) as a function of friction 

coefficient. 

Larger seat angles and lower friction more wedging between the valve face and 

coefficients cause higher normal forces and insert. While higher normal forces tend to 

83
Heavy Vehicle Propulsion Materials 

crush deposits that could otherwise build-up 

on the mating surfaces, but they also 

generate more wear. In cases of low-sulfur 

fuel where the friction coefficients are 

somewhat higher, higher normal force can 

be problematical. 

On the other hand, shallower seat angles 

reduce the shear stress and wear, but they 

allow deposits to form more readily. 

In designing the HTRI apparatus, it was 

desirable to be able to test not only simple 

test coupons, but also actual diesel engine 

FY 2006 Progress Report 

valves. Therefore, the design included a 

measure of flexibility. High-temperature 

alloys were used for all fixtures in the hot 

zone and the use of screws was minimized 

to avoid corrosion that could cause binding 

of the threads. Figure 2 shows the holder 

arrangement. A selection of production 

diesel engine exhaust valves was provided 

by Caterpillar Corporation to aid in the 

design of the holders, and the holder was 

made large enough to accommodate a range 

of valve diameters and bevel sizes. 

Figure 2. Close-up view of a diesel exhaust valve mounted (inverted) in the holder. When in operation, the 

valve will repetitively impact the two test blocks that can be fabricated from a variety of candidate seat 

materials. 

One of the challenges in the design of the two ‘seat specimens’ to rotate into place 

HTRI apparatus was to enable specimens to below the valve tulip. Additional features 

be installed and replaced with minimum were engineered into the fixtures to make 

difficulty. This was accomplished by the apparatus ‘user-friendly’. 

incorporating a specimen rotation feature In order to reduce the mechanical forces 

into the seat specimens (Figure 2). After the that had to be supplied by the vertical drive 

valve specimen is installed, a small shaft, it was decided to support the load only 

displacement of the upper holder causes the at two locations on the valve seat and use a 

84
Heavy Vehicle Propulsion Materials 

non-conformal contact, rather than seating 

around the entire circumference. This load 

concentrating design also had the advantage 

of enabling several tests to be run on the 

same valve simply by rotating it to a new 

location. The typical compression pressure 

for a diesel engine is about 2800 psi. If the 

valve head is 40 mm in diameter and the 

angle on the seat is 30 degrees, then the 

contact pressure is 61 MPa. To achieve this 

contact stress, based on the Hertz equations 

of elastic contact, a down force of 

approximately 22.8 kg-f would be required. 

If the seat angle were increased to 45 

degree, the required down force is 33 kg-f. 

FY 2006 Progress Report 

We will attempt to achieve contact stresses 

of this level. 

The main features of the complete HTRI 

system are schematically shown in Figure 3. 

Certain details, like the loading system, have 

been omitted for clarity. Symbol “A” 

represents the location of the test specimen 

holder inside an environmental chamber 

“B”, which is surrounded by a custom-built 

tube furnace “C.” The lower specimen is 

actuated by a cam drive “D” and a variablespeed motor “E.” Gas feed-throughs, 

labeled “F” and “G,” control the 

environment in the test chamber. Water 

cooling is supplied to critical components at 

the top and bottom of the furnace chamber. 

Figure 3. Schematic representation of the HTRI system. 

Completion of the HTRI system is expected 

by the end of calendar year 2006. The 

system will then be calibrated, and the test 

specimen wear modes for baseline material 

combinations will be compared to those on 

worn valves from actual fired diesel engines 

to ensure that the test system adequately 

simulates the mechanical, thermal, and 

chemical attributes of engine use. 

85
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Summary and Conclusions 

• Design and construction in under way

for a repetitive-impact valve material 

testing system that operates at 

temperatures in excess of 800o

 C in 

controlled environments. . 

• Once baseline tests are concluded, it will 

be possible to test a wide variety of 

candidate exhaust valve materials, 

including novel coatings and surface 

treatments. 

Publications/Presentations 

None this period. 

86
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

5. MATERIALS FOR STRUCTURAL COMPONENTS 

A. Cost-Effective Machining of Titanium and HCCI Fuel Systems Spray Hole 

Drilling 

Albert J. Shih 

Mechanical Engineering 

University of Michigan 

Ann Arbor, Michigan 48109-9739 

(734) 647-176; fax: (734) 936-0363; e-mail: shiha@umich.edu 

DOE Technology Development Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

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

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: University of Michigan, Ann Arbor, Michigan 

Objective 

• Investigate new technologies for cost-effective machining of titanium alloys with particular emphasis on highthroughput drilling technology 

• Evaluate the microstructure and mechanical properties of the workpiece and chip in high-throughput drilling of 

titanium alloys 

• Provide accurate prediction of temperature and stress of cutting tool in high-throughput machining 

• Explore the needs and conduct joint research with Caterpillar on the advanced material and manufacturing 

issues for the future homogeneous charge compression ignition (HCCI) fuel systems 

Approach 

• Conduct high-throughput drilling test and 3D finite-element modeling using various process parameters. 

• Study the Ti-6Al-4V machined surface and the chips using the metallurgical analyses, including Scanning 

Electron Microscopy (SEM), X-Ray Diffraction (XRD), electron microprobe, and nanoindentation. 

• Investigate the spray hole and reverse taper generation for HCCI fuel systems. 

Accomplishments 

• Achieved 156 mm3/s material removal rate in titanium drilling using 4 mm diameter tool. 

• Analyzed the phase transformation and hardness change in high-throughput drilling of Ti-6Al-4V. 

• Predicted drilling process behavior including cutting forces, tool and workpiece temperatures, stresses, etc. 

Future Direction 

• Explore new manufacturing technologies to enable advanced materials in HCCI fuel systems. 

• Finish the 3D finite element modeling of high-throughput drilling and advance the thermal analysis and 

management of the cutting tool. 

87
Heavy Vehicle Propulsion Materials 

Introduction 

This report presents the research work to on 

cost-effective machining of Ti and the exploratory 

work in spray hole drilling for HCCI fuel systems. 

Ti has two crystal structures: the hexagonal 

closed-pack (HCP) α phase, and body centered 

cubic (BCC) β phase1

. Pure Ti is 100% α at room 

temperature. The allotropic transformation from α

to β phase takes place at the β transus temperature 

883°C. For Ti-6Al-4V, Vanadium is added to pure 

Ti to stabilize the β phase by lowering the β transus 

temperature1

. Aluminum is added to increase the β

transus temperature1

. With 6 wt% of Al and 4 wt% 

V, the β transus temperature of Ti-6Al-4V is 980°C. 

An example of the microstructure of the Ti-6Al-4V 

used in this study is shown in Figure 1. The dark 

region is α phase and the light region is β phase. 

Figure 1. SEM micrograph of an etched microstructure 

of Ti-6Al-4V with dark region for HCP α

phase and light region for BCC β phase. 

Drilling is a widely used machining process and 

has considerable economical importance because it 

is usually among final steps in the fabrication of 

mechanical components. In drilling, high material 

removal rate (MRR) and long drill life are essential 

to increase productivity and reduce cost. To achieve 

the desired MRR, high cutting speed and feed per 

revolution of drill are required2

. The drilling of Ti6Al-4V at high cutting speed has been reported in 

2005 Annual Report. In this report, the drill life and 

metallurgical analysis, including SEM, XRD, 

electron microprobe, and nanoindentation tests, of 

workpiece and chip in high-throughput drilling of 

Ti-6Al-4V are studied extensively. The Ti 

FY 2006 Progress Report 

workpiece and chips undergo large strain and high 

strain rate deformation and subject to high 

temperature, which are expected to alter the 

microstructure and material properties. The balance 

of cutting speed and feed in high throughput drilling 

to achieve long drill life and good hole surface 

roughness is investigated. 

To achieve the cost-effective machining of Ti, a 

better understanding of the underlying mechanism 

through better modeling is necessary3

. In this report, 

the 3D FEM modeling of Ti drilling is conducted. 

The model includes fully adaptive unstructured 

mesh generation, thermo-mechanically coupling, 

deformable tool-chip-workpiece contact, interfacial 

heat transfer across the tool-chip boundary, 

momentum effects at high speeds and constitutive 

models appropriate for high strain rate, large 

deformation analyses. 

The geometry of injector spray hole is important 

for the fuel efficiency and emission of the HCCI fuel 

systems. To produce consistent quality and control 

the shape of spray holes, the monitoring and control 

of electrical discharge machining (EDM) drilling 

process and the measurement of geometrical features 

of spray holes are performed. 

Approach 

High Throughput Drilling of Ti-6Al-4V 

High throughput drilling test was conducted at 

High Temperature Materials Lab at ORNL using a 

Cincinnati Saber machine with a 4 mm diameter 

Kennametal K285A01563 WC-Co spiral point drill 

(WC-Co Spiral). The workpiece was a 6.35 mm 

thick Ti-6Al-4V plate. This machine has the high 

speed spindle and through-the-spindle high pressure 

coolant system. The cutting fluid was 5% 

CIMTECH 500, a synthetic metalworking fluid. 

Table 1 summarizes the cutting fluid supply, 

peripheral cutting speed, and feed ub four sets of 

experiments, marked as Exps. I-VI. Exp. I was dry 

drilling at 183 m/min peripheral cutting speed 

and 0.051 mm/rev feed. Exp. II was conducted 

with external cutting fluid supply. Exp. III-IV 

was conducted with cutting fluid supplied 

internally through the drill body. The feed was 

gradually increased from 0.051 to 0.076, 0.102, and 

0.152 mm/rev at corresponding drill rotational speed 

to maintain the same 0.75 m/min feed rate and 156 

mm3

/s MRR. 

88
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

The holes were cut in half to measure the 

roughness of hole surface using the Taylor Hobson 

Talysurf model 120 stylus surface profilometer with 

0.8 mm cutoff length and 100:1 bandwidth ratio. 

For each drilling test, the first three holes were 

measured. 

Table 1. Experimental conditions for high throughput 

drilling tests. 

Cutting 

Exp. Cutting speed Feed 

fluid (m/min) (mm/rev) 

I No 183 0.051 

II Ext. 183 0.051 

III Int. 183 0.051 

IV Int. 122 0.076 

V Int. 91 0.102 

VI Int. 61 0.152 

Metallurgical Analysis 

After drilling, one of the sectioned specimens 

was metallographically mounted in epoxy potting 

and then polished with fine diamond compound for 

nanoindentation testing and chemical composition 

analysis. After nanoindentation, the specimens were 

etched with a solution containing 6 ml HNO3 and 3 

ml HF in 100 ml H2O to distinguish the α and β

phase structure of Ti-6Al-4V in the SEM. The other 

sectioned specimen was used for XRD. An original 

Ti-6Al-4V bulk material was also polished and 

etched to extract the initial α−β phase information in 

SEM and XRD testing. The chips were cleaned by 

acetone, metallographically mounted in epoxy, and 

sectioned and polished for nanoindentation testing 

and SEM work. 

The polished and etched hole and chip cross 

sections were examined by a Hitachi S-4700 SEM 

machine to identify the changes in the 

microstructure. The phase transformation of α−β

phase was examined by a PANalytical XPERT-PRO 

MPD X-ray diffractometer with Cu tube (generator 

setting: 40 mA, 45 kV). The chemical composition 

was investigated by a JEOL 8200 electron 

microprobe. The nanoindentation tests were 

conducted on the subsurface adjacent to the hole 

surface by using a Hysitron triboindenterTM with 2 

mN applied load. A MTS Nano-indenterTM II was 

used for the nanoindentation of drilling chips. 

3D FEM Drilling Modeling 

The AdvantEdgeTM 3D machining simulation 

software by Third Wave Systems with updatedLagrangian finite element method, continuous 

remeshing, and adaptive meshing techniques4

 was 

used to model Ti machining process. The workmaterial model contains the power-law strain 

hardening, thermal softening and rate sensitivity4

. 

Figure 2 shows the 3D finite element mesh using a 

conventional twist drill. The workpiece is modeled 

with a predrilled cone-shape blind hole to enable the 

early full-engagement of the whole drill point region 

to reduce the simulation time. 

The model is validated by comparison with 

experimentally measured thrust force and torque. 

The drilling test was conducted on a Mori Seiki 

computer numerical control machining center using 

a M2 high speed steel twist drill with 3.97 mm 

diameter, 0.77 mm web thickness, 118° point angle, 

and 30° helix angle. Workpiece is a 6.35 mm thick 

Ti-6Al-4V plate. Drilling tests were conducted at 

three drill rotational speeds, 734, 1100, and 1468 

rpm, which corresponds to 9.1, 13.7, and 18.3 m/min 

drill peripheral cutting speed, respectively. A 

constant feed of 0.051 mm/rev was used. 

Figure 2. Finite element mesh for 3D drilling model. 

Spray Hole EDM 

An Agilent Infiniium 54833A digital 

oscilloscope with 2 GHz sampling rate was used for 

sub-ns real-time measurement of the gap voltage and 

current in spray hole EDM. The roundness, 

straightness, and taper of 0.15 mm diameter fuel 

injector spray hole was measured using a Werth 

VideoCheck HA 400 coordinate measurement 

machine (CMM) using a glass-fiber probe with a 

89
0 

3 

1 

2 

3

Torque (N-m) 

(9783 rpm) 

Heavy Vehicle Propulsion Materials 

ball end tip and a charge coupling device (CCD) 

optical sensor to measure the ball contact position. 

Results

High Throughput Drilling of Ti-6Al-4V 

Supplying the cutting fluid via through-the-drill 

holes is essential to drill life in high throughput 

drilling. At 183 m/min cutting speed and 0.051 

mm/rev feed, the internal cutting fluid supply (Exp.

III) significantly increased the drilling life to 101 

holes from 10 holes in dry drilling (Exp. I). 

Adjusting the feed also has a significant impact

FY 2006 Progress Report

The drill which had good drill life also produced 

holes with lower surface roughness. Figure 4 shows 

the hole surface roughness vs. drill life (number of 

holes) for high throughput drilling tests. All tests 

have the same MRR, 156 mm3

/s. The range of

upper and lower surface roughness of the first three

holes is presented. In Exp. IV, the internal cutting 

fluid supply helps to improve the surface finish. The 

long drill life positively correlates to the good initial

hole surface finish. 

250 

200 

Average roughness Drill life 

R a (μm) (number of holes)

on drill life. Feed is a process parameter which has 

been overlooked by drilling researchers in the past. 

Advanced spiral point drill geometry design and

new tool material enable the drill to take a higher 

cutting force and the use of larger feed, which 

corresponds to lower cutting speed. As shown in

Figure 3, at 0.076 mm/rev feed (Exp. IV), the thrust

force was at about 200 N and the drill life was 153 

holes. Increasing the feed to 0.102 mm/rev (Exp. V) 

only raised the thrust force to about 260 N for the 

150 

100 

50 

0 

3 

2.5 

2 

1.5 

1 

0.5 

0 

new drill and 330 N at the end of drill life (204 

holes). The lower cutting speed at higher feed helps 

to reduce the drill temperature, which is beneficial to

the drill life. There was a limit of drill life 

associated with the increase of feed. As the feed 

increased to 0.152 mm/rev (Exp. VI), the drill life 

(164 holes) started to deteriorate. 

Exp. I II III IV V VI 

Figure 4. Hole surface roughness under high throughput

drilling with the same material removal rate 

(156 mm3/s) using WC-Co Spiral drill. 

Metallurgical Analysis 

Figure 5 shows the SEM micrographs of the 

hole cross-sections of Exps. I and III, which have Hole A BC A BC 

400 4 number the same drilling peripheral cutting speed and feed 122 m/min 122 m/min

Thrust force (N) Thrust force (N) Thrust force (N) 300 per revoluation. In Exp. I (dry drilling), the hole (9783 rpm) 

0.076 mm/rev 0.076 mm/rev 1st surface is rough with many debris welded on the 0.57 s 0.57 s

200 

75th 

100 hole surface and 10-15 μm thick subsurface layer. 152nd 

0 In this layer, there is a gradation of distinct grain 

400 91 m/min 4 91 m/min boundaries between α and β phases to indiscernable Torque (N-m) Torque (N-m)

(7338 rpm) 

0.102 mm/rev

(7338 rpm) 3 300 0.076 mm/rev 1st ones from the bulk toward the hole surface. The 

200 2 100th high magnification micograph of the subsurface 

100 204th layer shows some tiny acicular or elongated grains. 1 

0 0 This phenomenon is likely due to two factors: one is

400 4 61 m/min 61 m/min the phase transformtion from β to α as a result of 

(4892 rpm) 

0.152 mm/rev

(4892 rpm) 

0.152 mm/rev 1st high temperature and the other is the decrease of 2 80th grain size due to large plastic deformation. 

163rd 1 In Exp. III, the β structure is obseravable at 

0 0

05 10 15 05 10 15 Drilling depth (mm) Drilling depth (mm) 

Figure 3. Thrust force and torque of WC-Co Spiral drill 

for high throughput drilling at 0.75 m/min feed

rate and 156 mm3/s material removal rate with

internal cutting fluid supply.

90

regions very close to the hole surface, suggesting 

that the extent of the β→α phase transformation is 

significantly reduced due to the cutting fluid cooling

effect. The subsurface layer is about only 3 μm. As

shown in the high magnification micograph, the

100 

200 

300 
Heavy Vehicle Propulsion Materials 

grain structure within the subsurface layer is 

elongated parallel with the hole edge, indicating the 

severe plastic deformation. The micrographs of

Exps. V and VI showed similar features and

dimensions of the features as those in Exp. III. 

Figure 5. SEM micrographs of the polished and etched

cross-sections of drilled holes. 

As polished bulk material

FY 2006 Progress Report

that the residual β peaks are submerged by 

broadened neighboring α peaks and become

indiscernible. The amount of β phase increases to 

13.6, 7.8, and 11.4% for Exps. III, V, and VI, 

respectively, with the introduction of cutting fluid.

Figure 7 compares the nanoindentation hardness 

vs. distance from the hole surface in Exps. I and III. 

The layer of high hardness exists in all conditions,

but the extent and the depth of high hardness layer 

are different from one condtion to another. The 

highest hardness exists under the dry drilling 

condition with 183 m/min cutting speed and 0.051 

mm/rev feed (Exp. I). At the reigon less than 5 μm 

from the hole surface, the hardness is over 9 GPa, 

which is about twice the hardness measured away

from the hole surface. The high hardness layer is 

about 15-20 μm wide. The existence of this 

hardened layer is the result of severe plastic

deformation outweighing the thermal softening 

effect. The phase transformation from β to α due to

high temperature may also contributes to the 

formation of this layer [5,6]. Beyond this layer, the

hardness is stablized around 4.1-5 GPa. When the

cutting fluid is supplied at the same cutting speed 

and feed (Exp. III), the peak hardness is 5-5.5 GPa 

in the subsurface layer close to the hole edge, which 

183 m/min, 0.051 mm/rev, dry is just slightly higher than the hardness of bulk 

material. The thickness of hardened layer is reduced 

to between 5-15 μm. 

12 183 m/min 

Square root of counts

α (100)

α (002)

Exp. I β (110) β (200) β (211) β (220) 0.051 mm/rev 

Dry

Hardness, H (GPa) Hardness,

H (GPa) 

α (101)

α (102) 

α (110)

α (103)

α (200)α (112)α (201)α (004)α (202) 

8 

4 

30 40 50 60 70 80 90 

Diffraction angle, 2θ (°) 

Figure 6. X-ray diffraction patterns of the as-polished

bulk material and hole surface of Exp. I.

Quantitative analysis by Rietveld method shows 0 

183 m/min 7.3% β in the bulk material. In Exp. I, the β is 12 

0.051 mm/rev 

reduced to 0%. The X-ray penetration depth is Exp. III Wet

estimated to vary from 4 to 12 μm at 30 to 90° 2θ, 

respectively, assuming 95% of the total diffraction 

intensity. According to the micrographs in Figure 1, 

small amount of β grains still exists in the 

subsurface layer very close to the hole edge. This is 

8 

4 

0

slightly different with the XRD observation. There 0 20 40 60 

are two possible causes for the disappearance of β Distance from hole wall (μ m) 

peaks. The main cause is the transformation of a Figure 7. Nanoindentation profile of hole subsurface 

large portion of the β to α phase. The other cause is hardness. 

91
0 

) ( q

Hardness (GPa) 

Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Nanoindentation on the chip cross-section of

Tor ue N-m

4.1 3.9 

4.1 4.1 

4.0 

4.0 

4.2 

3.9 

4.2 

4.2 

4.3 

3.8 

4.3 

4.3 

4.2 

4.4 

4.2 

4.1 

4.1 

4.0 

For Ti chip in drilling, the saw tooth chip 

formation is not a common feature. Both saw teeth 

and small lamellae were observed in the drilling 

chips, depending on the locations where crosssections were taken. Saw teeth are obvious at the 

outmost edge of the chip which is generated at the 

outmost point of the drill cutting edge, as shown in 

the SEM micrographs of chip cross-sections of Exp. 

I in Figure 8. Even in dry drilling condition, the 

light β phase can be observed, as well as in the 

condition with cutting fluid supply. It means the 

Figure 9. temperature does not reach the β transus temperature Exp. I. across the entire primary shear zone during the 

drilling process. In chips, shear bands are observed 

to initiate from the valley of saw-tooth chips. The 3D FEM Drilling Modeling 

grain structures are elongated along the both sides of Figure 10 shows the finite element modeling and

shear bands, clearly identifying the severe plastic 

shear deformation. 

averages of experimentally measured thrust forces 

and torques. The FEM model provides a reasonable 

estimation of thrust force. The thrust force is less 

than 20% underestimation, the torque is overestimated by 50%. 

Force-Z (N) 

TToorr

qquue e (N-m(Mmod) 

Thrust force (Mod) 

Thrust force (Exp) 

Torque (Exp) 

9.1 m/min 

(1100 rpm) 

60600 0 1.5 

Thrust force (N) Thrust force (N)

400 400 1.0 

Figure 8. SEM micrograph of chip cross-section in Exp. I. 

In the chip formation process at the outmost part 

of drill cutting edge, high temperature and strain are 

concentrated in the narrow bands dividing the chips 

into segments. High temperature tends to soften the 

material and possibly change the crystal structure, 

202000 0.5 

0 0

0.002 0.004 0.006 0.008 

Time (s) while high strain is expected to increase the hardness. 

The influence of these competing factors on the Force-Z (N) 

Torque (N-mm) 

Thrust force (Mod

Torque (Mod) 

200 

600 

0.5 

Thrust force (Exp) 

Torque (Exp) 

18.3 m/min 

(1470 rpm) 

1.5 

mechanical properties was examined by comparing 

the hardness of shear bands with that of other 

regions in nanoindentation tests. Figure 9 shows an 

example of indents on a chip sample. The spacing 400 1.0 Torque (N-m) 

between each indent is 3-6 μm. No significant 

difference of hardness can be found between the 

indents on the shear bands and those on the bulk of 

chips. Over 140 and 60 indents were made in a 

similar configuration on saw-tooth chip crosssections under dry and wet conditions, respectively. 0 0 The experimental results conclude that the formation 0 0.002 0.004 0.006 0.008 

Time (s) of saw tooth feature and shear bands does not 

change the mechanical properties of chips severely. Figure 10. Comparison of modeling and experimental 

thrust forces and torques. 

92 
-150 

-100 

Voltage (V) 

Heavy Vehicle Propulsion Materials 

In experiments, a large amount of heat is 

generated from the initial contact of the drill and 

workpiece to the full engagement of drill point 

region with workpiece. It results in high 

temperature and softens the workpiece. In 

comparison, the drill temperature is at room 

temperature when the cutting edge fully engages the

workpiece in FEM. As a result, the workpiece is 

harder and results in higher torque. Another 

possible reason for the discrepancy is the inaccuracy 

of material and friction modeling. A constant 

coefficient of friction was used. 

Temperature (°C) 

(a) 

Temperature (°C) 

(b) 

Temperature (°C) 

FY 2006 Progress Report

13.7 m/min. The chip and the engagement of the

drill and workpiece are shown in Figure 11(a). 

Figure 11(b) shows the temperature distribution of

the chip and workpiece. The temperature of Ti chip

is higher than the cutter surface on the workpiece.

The highest temperature, about 453°C, is

concentrated in regions contacting the drill cutting 

edge near the drill margin, where the cutting speed is

the highest along the cutting edge. Figure 11(c)

shows the temperature distribution on the drill flank 

face. The peak temperature reaches 405°C. Such

high temperature is detrimental to the tool life. 

Spray Hole EDM Process Development 

Figure 12 shows the gap voltage and current of

sample spark and arc pulses and the input and output 

process parameters in spray hole EDM. The use of

thin, 0.075 to 0.15 mm diameter W electrode and 

short pulse on and off time in spray hole EDM 

requires the use of sub-ns fast sampling rate to 

monitor the plasma generation in spark erosion.

Figure 13 shows a sample spark using 0.125 mm 

diameter W wire for EDM AISI 52100. The spark 

generation events can be divided into six periods, 

marked as Periods I to IV. A close-up view of the 

voltage discharge between Period II and III is also

show in Figure 13. Applying this monitoring 

technology, the spark, arc, and short pulses can be 

identified, the quality of EDM process can be 

quantified, and the process setup can be adjusted for 

EDM process optimization. 

Spark Arc 250 Pre-discharging time, T p

60 

200 

Open circuit

Pre-discharging voltage, Vo

current, Ip

Interval

voltage,Vi

150 

100 

50 Voltage 

0 35 Ringing effect 

Current (A) 

-50 30 

0.5 μs 25 

20 

-200 Discharging 15

Discharging

current, Id

Reverse

current, Ir

-250 time, T d Ringing time, T r

(c) -300 Figure 11. FEM temperature results: (a) drill, chip, and Current -350 

workpiece, (b) chip and workpiece, and (c) Pulse off time, Toff

-400 -5 Pulse duration time, T D

drill. -450 t o t 1 Time 

In drilling, the chip formation has a significant 12. EDM process input and output parameters of

sample spark and arc pulses. influence on the thrust force and torque as well as

the temperature and tool performance. An example

of the chip formation process is shown in Figure 11. 

The cutting speed for the selected example is

93

-10 

0 

5 

10 

40 

45 

50 

55 
-200 

-150 

50 

100 

-50 

Heavy Vehicle Propulsion Materials FY 2006 Progress Report

Spark Figure 14 by the 3D wire mesh representation. The 300 55 230 ns 330 ns effect of AFM to enlarge the size of spray hole on 250 T p = 30 ns 100 ns 50 

200 62 ns 45 one side on the flow entrance end can be recognized. A 150 40 Using this measurement technology the level of Vo

35 reverse taper and its effect on engine performance 

can be correlated. 

Conclusions 

Vi

Voltage 

Current (A) 

Voltage (V) Voltage (V) 

30 

0 25 

-50 Ringing effect 20 

-100 15 

10 

Ip

Id

Experiments of high throughput drilling of Ti­ Current Ir

5

-250 0 6Al-4V with 156 mm3

/s material removal rate using I I II II IV V VI -300 -5 a 4 mm diameter WC-Co spiral point drill Time demonstrated that supplying the cutting fluid via

250 

Spark 

30 through-the-drill holes was a critical factor for drill Rate of voltage drop: 

–37.6 V/ns 10 ns life. The balance of cutting speed and feed is 200 25 

Voltage essential to achieve long drill life and good hole 

150 20 surface roughness. 

100 Vo

15 

50 10 

Current (A) 

Metallurgical studies, including SEM, XRD, 

electron microprobe, and nanoindentation tests were 

conducted on the hole surface and subsurface and 

the chips in high-throughput drillling of Ti-6Al-4V. 0 Current 5 

In dry drilling, the transformation of BCC β phase 0 

Ip

into HCP α phase was identified in a 10 to 15 μm 

-100 II III -5 wide subsurface layer adjacent to the hole surface by

Time SEM and XRD analysis. High hardness was found 

Figure 13. Close up view of the voltage and current in this layer by nanoindentation testing. No obvious 

change in spark. β→α phase transformation occurrd with the supply

of cutting fluid. The high hardness layer was also 

narrower than that of dry drilling. Unlike chips 0 0 

formed in orthogonal cutting, drilling chips had -0.1 -0.1 

complicated morphology. The saw-tooth feature

-0.2 -0.2 only formed at the outmost region of the drill cutting 

edge, mainly due to the variable cutting speed along -0.3 -0.3 

the cutting edge. Narrrow shear bands could be -0.4 observed in the cross-sections of chips with distinct 

-0.5 saw teeth, but no significant change of mechanical

-0.6 properties along the chips was found using Hole depth (mm)

Hole depth (mm) -0.4 

-0.5 

-0.6 

nanoindentation tests. 

A 3D Finite Element Drilling Modeling was 

validated by comparing the experimentally

measured and modeling thrust force and torque. The 

model provides good estimation of thrust force but

the torque is overestimated by 50%. Further

improvement of the FEM drilling model is currently

ongoing. This study will be extended to highthroughput drilling of Ti. 

The future research will move to the HCCI fuel 

systems advanced material manufacturing 

development. Spray hole will remain to be a critical 

engine technology that will be further developed in 

this research to achieve the efficiency and quality

Flow -0.7 -0.7 

direction 

-0.8 -0.8 

-0.9 0.1 -0.9 0.1 

-0.1 

0

Unit (mm) -0.1 

0Unit (mm) 

-0.1 0 -0.1 0 0.1 

Unit (mm) 0.1 Unit (mm) 

 (a) (b) 

Figure 14. 3D wireframe representation of a spray hole: 

(a) before AFM and (b) after AFM.

The geometrical features, including 

roundness, straightness, and taper, of a 0.15 mm 

nominal diameter spray hole before and after the 

abrasive flow machining (AFM) is illustrated in 

94
Heavy Vehicle Propulsion Materials 

control of spray hole size smaller than the current 

standard size spray hole and the associated process 

improvement and geometrical measurements. 

References 

1. M. J. Donachie, Jr., Titanium: A Technical Guide, 

ASM International, Material Park, OH, 1988. 

2. D. A. Stephenson, J. S. Agapiou, Metal Cutting 

Theory and Practice, Marcel Dekker, New York, 

2005. 

3. X. Yang and C. R. Liu, “Machining Titanium and 

Its Alloys,” Machining Science and Technology, 

3(1), 107–139 (1999). 

4. T.D. Marusich, M. Ortiz, “Modeling and 

Simulation of High-Speed Machining,” 

International Journal for Numerical Methods in 

Engineering, 38(21), 3675-3694 (1995). 

5. A.R. Machado, J. Wallbank, “Machining of 

Titanium and Its Alloys-A Review,” Proceedings 

of Institution of Mechanical Engineers, Part B: 

Management and Engineering Manufacture, 

204(1), 53-60)1990). 

6. J. Cantero, M. Tardío, J. Canteli, M. Marcos, M. 

Miguélez, “Dry Drilling of Alloy Ti-6Al-4V”, 

International Journal of Machine Tool s and 

Manufacture, 45(11), 1246-1255(2005). 

7. C.A. van Luttervelt, T.H.C. Childs, I.S. Jawahir, 

F. Klocke, P.K. Venuvinod, “Present Situation 

and Future Trends in Modelling of Machining 

Operations -- Progress Report of the CIRP 

Working Group `Modelling of Machining 

Operations',” CIRP Annals - Manufacturing 

Technology, 47, 587-626 (1998). 

Presentations and Publications 

R. Li, P. Hegde, A.J. Shih, "High Throughput 

Drilling of Titanium Alloys," International Journal 

of Machine Tools and Manufacture, 43, 63-74 

(2007). 

R. Li and A. J. Shih, “Finite Element Modeling 

of 3D Turning of Titanium,” International Journal 

of Advanced Manufacturing Technology 29(3-4), 

253-261 (2006). 

T. D. Marusich, S. Usui, R. Aphale, N. Saini, R. 

Li, A. J. Shih, “Three-Dimensional Finite Element 

Modeling of Drilling Processes,” presented at 2006 

ASME International Conference on Manufacturing 

Science and Engineering, October 8-11, 2006, 

Ypsilanti, Michigan. 

FY 2006 Progress Report 

R. Li, A. J. Shih, “High Throughput Drilling of 

Titanium Alloys,” presented at the 2006 

International Conference on Frontiers of Design and 

Manufacturing, Guangzhou, China, June 19-22, 

2006. 

R. Li and A. J. Shih, “Inverse Heat Transfer 

Solution of Tool Temperature in Titanium Drilling,” 

Journal of Manufacturing Science and Engineering 

(submitted). 

R. Li, L. Riester, T. R. Watkins, P. J. Blau, A. J. 

Shih, “Metallurgical Analysis and Nanoindentation 

Characterization of Ti-6Al-4V Workpiece and Chips 

in High Throughput Drilling,” Materials Science 

and Engineering: A (submitted). 

95
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

B. Synergistic Process-Enhanced Grinding of Advanced Engine Materials 

Ronald H. Chand 

Oak Ridge National Laboratory 

P.O. Box 2008, Mail Stop 6063 

Oak Ridge, TN 37831-6063 

(865) 574-3123; fax: (865) 574-6918; e-mail: chandrh@ornl.gov 

DOE Technology Development Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@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 

• The objective of this effort is to understand, develop, and apply synergistic process-enhanced grinding (SPEG) 

to improve the efficiency of grinding advanced materials while ensuring surface quality. 

Approach 

• Conduct a comprehensive investigation of the current state of the art in SPEG methods. This study is to focus 

on advanced laser machining and its application to machining of advanced engine components. 

• Modify one or more existing machine tools at ORNL to conduct instrumented, SPEG experiments on ceramics, 

cermets, and/or metal matrix composites that are candidates for diesel engine components. 

• Develop a high-speed, instrumented scratch testing system to study the microstructural response of advanced 

materials like ceramics and composites to the high deformation rates they experience during grinding. 

Accomplishments 

• A comprehensive search and evaluation of 2000 recent worldwide publications in the area of laser assisted 

machining, treatment, and structuring was conducted. Water jet guided laser machining was identified as a 

potential technology for machining automotive injector nozzle. 

• Specifically designed laser machining technology was successfully tested on machining 3mm diameter holes in 

5mm thick silicon carbide. 

• Milling machine was specifically fitted and instrumented to test high speed boring of holes to improve the cycle 

time of finishing ID of silicon nitride roller followers. 

• A new type of grinding wheel with replaceable diamond inserts was designed for a dynamic scratch test. 

Future Direction 

• Laser machining of small holes shows great promise and will be further developed 

• Hole boring technology will be tested on actual silicon nitride roller follower. 

• Dynamic scratch test will be performed using the new design replaceable diamond insert wheel. 

97
USA 

Germany

Japan 

UK

China 

France 

Italy 

Canada

Russia

Spain 

Taiwan 

Sweden 

Switzerland 

South Korea 

Singapore

Poland

Austria 

Netherlands

Greece 

India 

Heavy Vehicle Propulsion Materials 

Introduction 

The advanced ceramic manufacturers have 

continually faced the challenge of making components for 

the automotive industry at a cost that the industry could 

afford. Major cost of an advanced ceramic component 

continues to be the machining cost. This cost can no 

longer be lowered any further by arbitrary manipulations 

of the processing variables. Application of lower cost 

labor is not a solution either for it results in more out 

sourcing of jobs. New synergistic process-enhanced 

machining methodologies are needed to meet the goals. 

The objective of this effort is to survey and identify 

technologies that could help lower the cost of machining 

FY 2006 Progress Report 

Results 

The results presented here are an abbreviated version of 

more detailed quarterly reports. The comprehensive 

search of published literature on laser machining resulted 

in the evaluation of 2000 most recent publications. 

The findings are summarized below: 

Relative Publication Activity on Laser Machining, Treatment, or Structuring of Materials 

25.0% 

Laser assited machining only 

USA 

67% 

Greece 

4% Germany 

4% 

France 

17% 

China 

4% 

Russia 

4% 

Creare Res & 

Dev Inc 

7% 

Univ. 

Wisconsin 

4% Purdue Univ. Univ. Notre 56% Dame 

33% 

only 24 publications on 

Laser Assisted Machining 

from 1982 - 2005 

no international collaborations 

advanced engine materials through the systematic 

evaluation and development of new machining 

technologies. The current approach involves a 

comprehensive investigation of the current state of the art 

in SPEG methods, machining tests on specially modified 

and instrumented machine tool, especially as related to 

the machining of advanced ceramics for automotive 

20.0% 

15.0% 

10.0% 

Percentage of 2000 publications analyzed

from 1997 - 2005 

5.0% 

0.0% 

engine components. The approach also includes design of 

a high-speed, dynamic scratch testing system to study the 

microstructural response of advanced ceramics to the high 

deformation rates they experience during grinding. The 

focus on roller followers and fuel injector nozzles is based Countries with 30 or more publications 

on discussions with manufacturers. Laser drilling and 

diamond boring tests were conducted on silicon carbide 

material and a new replaceable diamond insert wheel was 

developed for a dynamic scratch test on advanced ceramic 

materials. 

Approach 

A comprehensive search and evaluation of most 

recent publications in the area of laser assisted machining, 

treatment, and structuring was conducted. A water jet 

guided laser technology was identified as a potential 

machining methodology for advanced ceramics. 

Preliminary experiments were successfully conducted 

with specifically developed laser machining technology to 

drill small (3mm diameter) holes in 5mm thick silicon 

carbide. 

For cost-effective machining of the inside diameter 

of roller followers, a unique hole boring approach was 

studied. A milling machine was instrumented with 

Kistler dynamometer and modified with a through the 

tool coolant delivery system and hole boring experiments 

were conducted. 

To study the microstructural response of advanced 

ceramics to the high deformation rates they experience 

during grinding, a scratch test that emulates the dynamics 

of the grinding process is required. As a first step towards 

developing a dynamic scratch test a wheel with 

replaceable diamond inserts was designed. 

Figure 1 Publication Activity on Laser Machining 

Analysis of the literature and trend suggested that 

1. Materials’ processing with lasers is an active 

area of research. 

2. Laser as a tool is making conventional 

manufacturing more accurate and cost-effective. 

3. Laser technology is also opening up new 

manufacturing technologies that go beyond 

conventional methodologies. 

4. Laser technology can be arranged in four 

categories. Namely, Laser Machining, Laser 

Assisted Machining, Laser Deposition, and Laser 

Treatment. 

As a result of the extensive search and evaluation a 

breakthrough technology called the water jet guided laser 

system (Figure 2) was identified. This system utilizes a 

water jet to guide the laser beam over a much longer (10 

cm) distance as compared to conventional focused laser 

beam that has a working distance of just a few 

millimeters. An additional benefit of the water jet is 

thermal management of the work zone and flushing away 

of the molten material. 

98
Heavy Vehicle Propulsion Materials 

Figure 2 Illustration of Conventional & Water Jet Guided 

Laser 

Preliminary trials of material removal via laser 

drilling operation in CVD SiC has yielded unprecedented 

removal rates at an order of magnitude higher than the 

highest rate obtainable with existing technologies. The 

experiment was conducted using fiber optic delivered NdYAG laser (1.06 μm wavelength) at the operating 

parameters of 0.5ms pulse duration, 50 Hz pulse 

frequency and 300W. The high material removal rates and 

associated high dimensional tolerances of holes (Figure 3) 

300 μm 

Figure 3 Optical micrograph of 

photonic drilled holes in silicon carbide. 

pointed 

to the 

possibilit 

y of the 

extension of laser-based machining technique into a 

hybrid water jet-assisted laser machining technique for 

FY 2006 Progress Report 

rapid and cost effective machining of ceramics. 

Furthermore, the employment of water jet-assisted laser 

beam delivery is expected to eliminate generation of the 

thermal stresses during machining that are often 

detrimental to the integrity of ceramic component. 

As a continued effort to test SPEG methodologies a 

milling machine at HTML was modified (Figure 4) to do 

test high speed hole finishing of diesel engine roller 

followers. 

Figure 4 Modified Milling Machine for Hole Boring 

Preliminary drilling tests conducted on silicon 

carbide using the modified machine resulted in 3.5 times 

the conventional material removal rate. 

To study the microstructural response of advanced 

materials like ceramics and composites to the high 

deformation rates they experience during grinding a test 

system comprising of a special wheel with a replaceable 

diamond insert was developed. The wheel is illustrated in 

Figure5. 

Figure 5 Dynamic Scratch Test Diamond Wheel 

Balancing 

Studs 

Wheel 

Core 

In Line Multiple 

Steel Inserts (2) Diamond Grains

99
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Summary and Conclusions 

• The comprehensive search and evaluation of SPEG 

literature resulted in the discovery of water jet guided 

laser technology. Preliminary results obtained on 

conventional laser machining of small holes in 

silicon carbide showed great promise. It is 

anticipated that water jet guided laser system would 

enhance laser machining and should be further 

evaluated as an important tool for machining 

advanced ceramics. 

• Diesel engine roller followers continue to suffer from 

high cost of machining. The bore finishing test 

conducted on modified milling machine showed 

much promise in further developing the high speed 

bore finishing technology. 

• Fundamental understanding of the interaction 

between the diamond grains and the material during 

grinding is critical for making the grinding process 

more efficient. A dynamic scratch test that emulates 

the grinding process is necessary for studying such 

interaction. To this end the replaceable diamond 

insert wheel is a valuable development for it allows 

the diamond insert to be removed from the wheel for 

characterization and be inserted back for continuation 

of the test. 

100
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

6. MATERIALS FOR HOT SECTION COMPONENTS 

A. NDE of Diesel Engine Components 

Jiangang Sun 

Argonne National Laboratory 

9700 South Cass Avenue 

Argonne, IL 60439 

630-252-5169; fax: 630-252-2785; e-mail: sun@anl.gov 

DOE Technology Development Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

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

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: Argonne National Laboratory, Argonne, Illinois 

Objectives 

• Develop laser-scattering and other nondestructive evaluation (NDE) methods for detection and characterization 

of surface and subsurface defect/damage in diesel-engine valve-train components made from advanced 

materials such as ceramics and intermetallics. 

• Develop quantitative NDE analysis tools to identify strength-limiting flaws and failure mechanisms, evaluate 

damage evolution and growth, and predict mechanical properties for advanced materials. 

Approach 

• Develop fast and reliable automated laser-scattering NDE systems for surface/subsurface inspection of Si3N4 

ceramic and TiAl intermetallic valves. 

• Inspect valve surfaces using developed NDE methods to determine the level of machining damage and 

accumulated damages from bench and engine tests. 

• Correlate NDE data with microstructure and strength of silicon-nitride specimens for prediction of fracture 

initiation defects and mechanical properties. 

Accomplishments 

• Established laser-scatter NDE detection sensitivity for surface damage in TiAl specimens. 

• Evaluated 25 SN235P Si3N4 and 11 TiAl valves tested by a cyclic-impact bench rig for >20 h. Identified levels 

of impact damage and verified effectiveness of oxidation treatment for TiAl surface to improve wear resistance. 

• Inspected 10 SN235P Si3N4 and 4 TiAl valves after 100-h engine test. While minor surface wear was found for 

Si3N4 valves, considerable surface damage was detected for the TiAl valves. 

• Developed a high-resolution cross-polarization confocal microscopy technology. For the first time, subsurface 

C-crack trajectory within translucent Si3N4 materials can be directly imaged in 3D. 

101
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Future Direction 

• Complete NDE evaluation for damage accumulation in SN235P and TiAl valves from 500- and 1000-h engine 

tests. 

• Perform NDE analysis for prediction of strength-limiting defects/damage in engine-tested valves. 

• Continue development of the cross-polarization confocal technology to improve detection sensitivity, spatial 

resolution, and penetration depth for translucent materials. 

Introduction 

Advanced ceramics and intermetallics are leading candidates for high-temperature engine applications that offer improved fuel efficiency and reduced 

emissions. Among them, silicon nitrides (Si3N4) and 

titanium aluminide (TiAl) are being evaluated as 

valve-train materials for diesel and natural gas 

engines because of their lighter weight, high strength 

and corrosion resistance at elevated temperatures. 

However, these materials are brittle so microscopic 

defect/damage near the surface may significantly 

degrade their mechanical properties. Material 

damage in the subsurface may be induced by 

machining or from service due to impact, wear, and 

corrosion. For silicon nitrides, material strength may 

also be limited by inherent defects such as voids and 

porosities within subsurface. To detect and 

characterize these defects, Argonne National 

Laboratory (ANL) developed a laser scattering NDE 

method to measure detailed surface/subsurface 

microstructure for these materials. The objective of 

this research is to demonstrate that this method can 

be used to assess/evaluate the cost-effectiveness and 

reliability of valve-train components manufactured 

from advanced ceramic and intermetallic materials 

for diesel engines. The primary effort in FY 2006 

was to utilize a fast NDE valve-scan system 

developed in FY2005 to evaluate SN235P Si3N4 and 

TiAl valves after cyclic-impact rig test and 100-hour 

engine duration test. Another effort was directed on 

developing a cross-polarization confocal microscopy 

technology for high-resolution 3D subsurface 

imaging of translucent materials such as ceramics. 

This research is collaborated with Caterpillar, Inc. 

Approach 

The critical region for brittle ceramic and intermetallic components in structural applications is 

near surface. The common types of defects in this 

region are mechanical, such as cracks, spalls, inclusions, and voids. The size of the defects that limits 

the component strength is generally small, say 

~100 μm. To detect these defects, a laser-scattering 

method based on cross-polarization detection of 

optical scattering originated from surface and subsurface microstructure can be used for noncontact, 

NDE of these materials. Because ceramics are partially translucent to light, optical penetration can 

reach the subsurface to directly interact with the 

subsurface defects. For metallic or intermetallic 

materials, optical penetration is not possible. 

However, cross-polarization detection eliminates 

optical reflection from smooth surface and measures 

only optical scattering from rough or cracked 

surface. By scanning the entire surface (flat or 

curved) of a component and constructing a twodimensional (2D) scatter image, surface/subsurface 

defects can be readily identified as they exhibit 

excessive scattering over the background, and their 

type and severity may be analyzed. To apply this 

technology for NDE of Si3N4 and TiAl valves, an 

automated laser-scatter system was developed by 

ANL for scanning the entire valve surface at high 

spatial resolutions (5-10 μm). The resulting 2D 

scattering image data are used to identify the 

location, size, and relative severity of subsurface 

defects/damage. Analysis of the NDE data can 

determine the statistical properties of the component 

surface, which can be correlated with machining or 

service conditions. Development of quantitative 

image-processing methods may lead to automated 

identification of strength-limiting defect/damage for 

prediction of mechanical properties. 

Results 

Evaluation of Cyclic-Impact-Tested Si3N4 Valves 

Twenty-five SN235P Si3N4 valves, consisting of 

13 intake and 12 exhaust natural-gas valves finishmachined in FY2005, were proof tested in a cyclicimpact rig at Caterpillar for >20 hours. These valves 

were evaluated by ANL’s automated laser-scatter 

NDE system. Only valve-head surface was scanned 

102
Heavy Vehicle Propulsion Materials 

with a scan pixel size of 10 μm. NDE data were 

analyzed and compared with photomicrographic 

surface images to identify damages. 

For all 25 SN235P valves, laser-scatter data did 

not reveal any subsurface damage in the valve head 

surfaces. This observation is consistent with the 

results for the 10 NT551 Si3N4 valves that exhibited 

no subsurface damage after testing in the same 

cyclic-impact rig for up to 1000 hours

1

. A few 

valves, however, showed apparent circumferential 

“wear scars” within the contact surface. A “wear 

scar” is seen in the laser-scatter image (Fig. 1) of 

exhaust valve #29 after 20-h rig test. The scarred 

surface displays lower scatter intensity (darker) in 

the contact surface. Because defect/damage 

normally generates higher scatter intensities, the 

reduction of scatter intensity within the scarred area 

does not indicate damage, but likely due to surface 

contamination of embedded metal particles from the 

metal seat insert. This is verified from 

photomicrograph of the contact surface (in Fig. 1). 

The higher scatter-intensity spots in the enlarged 

NDE image represent material inherent defects such 

as porosities with sizes typically smaller than 50 μm. 

Figure 1. Laser-scatter scan image of head surface of 

SN235P valve #29 after 20-h rig test. 

Evaluation of Cyclic-Impact-Tested TiAl Valves 

Eleven TiAl natural-gas exhaust valves were 

NDE evaluated after cyclic-impact rig test at 

Caterpillar for >20 hours. Because as-processed 

TiAl surfaces are susceptible to wear/impact 

damage, 5 TiAl valves were oxidation treated to 

improve their surface wear resistance. The treatment 

caused slight curvature along valve axis which, 

however, did not significantly affect the NDE 

scanning. Both head and stem surfaces of these 

valves were scanned because of apparent damage on 

these surfaces. NDE data were analyzed and 

compared with photomicrographic surface images. 

For the 6 untreated TiAl valves, significant 

surface damage from the rig test was observed in the 

contact and stem surfaces, especially for TiAl valves 

FY 2006 Progress Report 

#1 and #2. The laser-scatter image of the valve-head 

surface of TiAl valve #1 is shown in Fig. 2. A gross 

variation of scatter intensity along the axial direction 

(in fillet surface) is noticeable, which is due to slight 

angle or focus-distance mismatch during the scan. 

The impact wear damage within the contact surface 

is represented by higher scatter intensities (brighter). 

The photomicrograph and the enlarged NDE image 

in Fig. 2 show clearly the wear scar. It is evident 

that the laser-scatter intensity is very sensitivity to 

the surface condition: surface damage, which may 

result in higher surface roughness, causes strong 

laser scatter from the surface. The wear damage in 

the contact surface of other 4 TiAl valves was less 

severe than that in TiAl valves #1, as seen in Figs. 

3a and 3b. 

For the 5 oxidation-treated TiAl valves, very 

minor surface wear was observed from the NDE 

data (Fig. 3c). Therefore, NDE data verified that 

oxidation treatment of the TiAl surface can 

significantly improve its wear resistance. 

Figure 2. Laser-scatter scan image of head surface of 

TiAl valve #1 after 20-h cyclic-impact rig test. 

(a) (b) (c) 

Figure 3. Detailed laser-scatter images in contact 

surfaces of TiAl valves (a) #1, (b) #3, and (c) 

#6 after cyclic-impact rig test. 

Figure 4. Laser-scatter scan image of stem surface of 

TiAl valve #1 after 20-h rig test. 

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Heavy Vehicle Propulsion Materials 

NDE scans were also performed for the stem 

surfaces of all TiAl valves. Figure 4 is the laserscatter scan image for the stem surface near the 

keeper groove of TiAl valve #1. In addition to the 

axial and circumferential wear stripes of higher 

scatter intensities, three features with distinctly 

higher scattering intensities are identified. These 

features were examined under optical microscope, 

shown in Fig. 5. Correlation of NDE data with 

photomicrographs concluded that these are surface 

damages probably induced by scratch or indent. 

Figure 5. Photomicrographs of features detected in the 

NDE data in Fig. 4. 

Figure 6. Photograph of 100-h engine-tested Si3N4 and 

TiAl valves. 

Figure 7. Location of impact area on valve surface. 

NDE Characterization of Si3N4 and TiAl Valves 

after 100 Hours of Engine Test 

Ten Si3N4 and four TiAl valves were run 

successfully for 100 hours in a Caterpillar naturalgas G3406 generator set at the National 

Transportation Research Center (NTRC). Figure 6 

shows a photograph of these valves. It is seen that 

all valve-head surfaces were covered by a coat of 

combustion byproduct2

. However, the impact area of 

the valve seat surface is clean without any deposit. 

FY 2006 Progress Report 

The impact area is not entirely within the conical 

contact surface; it lies around the edge between the 

contact and fillet surface, as illustrated in Fig. 7. 

NDE scans of the valve-head surface for all 

valves were conducted. Typical results for Si3N4 

intake, Si3N4 exhaust, and TiAl exhaust valves are 

presented and discussed below. Figure 8 shows the 

laser-scatter image of the valve-head surface of 

Si3N4 intake valve #20 that was positioned at I2 in 

the engine. The combustion coating has some effect 

on scatter intensity in the contact and fillet surfaces. 

Within the impact area, shown in the enlarged NDE 

image, the scatter intensity in the conical contact 

surface is higher that that in the fillet surface (see 

edge location), which indicates a difference of wear 

levels in the two surfaces. Nevertheless, no apparent 

damage was detected within the entire impact region 

of all Si3N4 intake valves. 

Figure 9 shows the laser-scatter image of the 

valve-head surface of Si3N4 exhaust valve #11 

positioned at E1 in the engine. The combustion 

coating on this valve is thin, so the valve surface 

underneath the coating is clearly imaged. The impact 

region for this valve shows relatively higher scatter 

intensity than the surrounding surface. The NDE 

data within the impact region also exhibit multiple 

bands of scatter-intensity variation; which could be 

an artifact due to slight laser-incident-angle 

variation. Despite these NDE indications of surface 

wear from engine test, no subsurface damage such 

as cracks was detected for all Si3N4 exhaust valves. 

Figure 8. Laser-scatter scan image of Si3N4 intake valve 

#20 (at I2 in engine) after 100-h engine test. 

Figure 9. Laser-scatter scan image of Si3N4 exhaust 

valve #11 (at E1 in engine) after 100-h engine 

test. 

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Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 10 shows the laser-scatter image of the 

valve-head surface of TiAl exhaust valve #8 

positioned at E15 in the engine. The combustion 

coating on this and other TiAl valves caused very 

high scatter than that from the bare TiAl surface. In 

order to display scatter image for the entire surface, 

the scatter intensity is scaled logarithmically in this 

image (this is a common practice in presenting Si3N4 intake #20 Si3N4 exhaust #11 TiAl exhaust #8 

optical data). For this and other TiAl valves, two 

types of damage indications were observed. First, Figure 11. Photomicrographs of impact surfaces of Si3N4 

several high scatter-intensity spots with sizes up to and TiAl valves after 100-h engine test. 

~1 mm were detected within the impact region (see 

enlarged image in Fig. 10). Second, the impact 

region exhibits considerable wear damage Cross-Polarization Confocal Microscopy of 3D 

represented by distributed scatter-intensity variations Subsurface Microstructure in Silicon Nitrides 

over the entire impact region. These NDE results Measurement of detailed 3D subsurface 

indicate that the impact damage in TiAl valves microstructure is critical in determining material 

appears to be progressing at a faster rate than those properties. However, the laser-scatter method used 

observed in the Si3N4 valves. for detecting subsurface damage in ceramic valves 

Figure 11 shows photomicrographs of impact can only image the 2D microstructure in the 

surfaces in Si3N4 intake valve #20, Si3N4 exhaust subsurface without the depth resolution. To resolve 

valve #11, and TiAl exhaust valve #8. For Si3N4 depth and improve spatial resolution, a crossintake valve #20, the original machining roughness polarization confocal microscopy technology was 

is still present, with little or no wear is observed. developed at ANL. A schematic diagram of the 

For Si3N4 exhaust valve #11, the machining system is shown in Fig. 12. In a typical test, the 

roughness has been worn off, only little remaining sample is scanned in the x-y-z directions in a raster 

along the fillet surface. The worn surface appears fashion, and the resulting 3D scattering data can be 

smooth, with distributed porosity spots and few viewed as lateral or cross-sectional images 

subsurface damage marks (bright line) that are representing the subsurface microstructure of the 

probably remnants of the original machining sample. Preliminary tests demonstrated that, with a 

damage. For the TiAl exhaust valve #8, however, a moderate 40X objective lens at an optical 

large spot with clear boundary is observed. Material wavelength of 633 nm, the system has an axial 

corrosion is suspected within this area, which (depth) resolution of 2 μm and a lateral resolution of 

showed higher scattering in the NDE data (this spot ~0.6 μm, which are significantly higher than other 

is seen in the enlarged image in Fig. 10). In available optical methods used for direct subsurface 

addition, the impact surface of the TiAl valve structure measurement. 

becomes very rough, with significant wear damages. 

In general, the micrographic observations of these 

valves are consistent with those detected by the laser 

scatter NDE. 

Figure 10. Laser-scatter scan image (in log-intensity 

scale) of TiAl valve #8 (at E15 in engine) after 

100-h engine test. 

Figure 12. Schematic of cross-polarization confocal 

microscopy system. 

A NT551 silicon nitride specimen with known 

subsurface damage was tested. The damage is 

subsurface C cracks induced by indentations with 

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Heavy Vehicle Propulsion Materials 

different loading forces. Figure 13 shows a lateral 

scan image of a damaged region with two C cracks 

generated by 2800 and 3000N loads. Strong laser 

scattering intensity is observed from both C cracks as 

well as from many individual material defects likely 

porous spots. Figure 14 is a cross-sectional scan 

image at Line A marked in Fig. 13, which shows 

clearly the direction and depth of the C-crack 

trajectories in the subsurface. The depth of optical 

penetration is ~40 μm, which can be increased using 

longer laser wavelengths. 

Figure 13. Lateral confocal-scan image of a NT551 

silicon nitride specimen with two C cracks. 

Figure 14. Cross-sectional confocal-scan image along 

line A as indicated in Fig. 13. 

Conclusions 

Twenty-five finish-machined SN235P Si3N4 

valves and eleven as-processed TiAl valves were 

tested in a cyclic-impact bench at Caterpillar for >20 

hours and inspected using the laser-scatter valvescan system at ANL. The NDE data indicated no 

detectable damage in SN235P valves. However, 

surface impact/wear damages ranging from severe to 

minor were found for TiAl valves. NDE data also 

verified that oxidation treatment of the TiAl surface 

can significantly improve its wear resistance. 

Ten Si3N4 (6 inlet and 4 exhaust) and four TiAl 

exhaust valves successfully run for 100 hours in a 

Caterpillar natural-gas G3406 generator set at NTRC 

were examined by the laser-scatter NDE system. For 

Si3N4 inlet valves, no apparent damage was detected 

within the entire impact region. For Si3N4 exhaust 

valves, NDE data showed general surface wear 

within the seat surface; however, no subsurface 

damage such as cracks was detected. For TiAl 

FY 2006 Progress Report 

exhaust valves, NDE test detected suspected 

corrosion spots as well as considerable wear damage 

over the entire impact surface. These results 

indicate that the impact damage in TiAl valves 

appears to be progressing at a faster rate than those 

observed in the Si3N4 valves. 

A cross-polarization confocal microscopy 

technology was developed in FY2006. The system 

resolution was found to be excellent: 2 μm in the 

axial (depth) direction and 0.6 μm in the lateral 

direction, which are likely the best resolutions 

achieved by optical methods. For the first time, 

subsurface C-crack trajectory within Si3N4 materials 

can be directly imaged in 3D. 

References 

1. J. G. Sun, “NDE Development for Ceramic 

Valves for Diesel Engines,” Annual Report for 

FY 2004, Heavy Vehicle Propulsion Materials 

Program. 

2. J. S. Trethewey, “High-Temperature 

Advanced Materials for Lightweight Valve Train 

Components,” Quarterly Report for April-June 

2006, Heavy Vehicle Propulsion Materials 

Program. 

Presentations and Publications 

J. M. Zhang, J. G. Sun, M. J. Andrews, A. 

Ramesh, J. S. Trethewey, and D. M. Longanbach, 

“Characterization of Subsurface Defects in Ceramic 

Rods by Laser Scattering and Fractography,” in 

Review of Quantitative Nondestructive Evaluation, 

Vol. 25, eds. D.O. Thompson and D.E. Chimenti, 

pp. 1209-1216, 2006. 

J. G. Sun, J. M. Zhang, and J. S. Trethewey, 

“Nondestructive Evaluation of Machining and 

Engine-Test Damage in Silicon-Nitride Ceramic 

Valves,” paper presented at the 30th International 

Conference & Exposition on Advanced Ceramics & 

Composites, Cocoa Beach, FL, Jan. 22-27, 2006. 

J. G. Sun presented an invited talk on “Optical 

NDE Methods for Characterization of Subsurface 

Microstructure in Silicon Nitride Ceramics” at the 

30th Annual Conference on Composites, Materials, 

and Structures (ITAR Restricted Sessions), Cocoa 

Beach/Cape Canaveral, FL, Jan. 23-26, 2006. 

106
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Patent 

J. G. Sun, “Device and Nondestructive Method 

to Determine Subsurface Micro-Structure in Dense 

Materials,” U.S. Patent No. 7,042,556, issued May 

9, 2006. 

107
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

B. Life Prediction of Diesel Engine Components 

H. T. Lin, T. P. Kirkland, A. A. Wereszczak, and H. Wang 

Oak Ridge National Laboratory 

P.O. Box 2008, MS-6068 

Oak Ridge, TN 37831-6068 

(865) 576-8857; fax: (865) 475-6098; e-mail: linh@ornl.gov 

Jeff Jensen and Nan Yang 

Caterpillar, Inc. 

Peoria, IL 61656-1875 

(309) 578-2969; fax: (309) 578-2953; e-mail: Jensen_Jeff_A@cat.com 

DOE Technology Development Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

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

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: Caterpillar, Inc., Peoria, Illinois 

Objective 

• Generate a database and characterize damage mechanisms of candidate advanced ceramics and intermetallic 

alloys. 

• Apply and verify probabilistic life prediction and component design and verification methods for advanced 

diesel engine components. 

Approach 

• Evaluate the dynamic fatigue, rotary bending fatigue, and high-temperature fatigue and creep performance of 

candidate silicon nitride ceramics and TiAl alloys at elevated temperatures in air before and after long-term 

exposure to simulated engine environments, as well as engine field tests. 

• Characterize the evolution and role of damage mechanisms, and changes in microstructure and chemistry, 

linked to the long-term mechanical performance and reliability of ceramics and intermetallic alloys. 

• Predict the failure probability and reliability of complex-shaped diesel engine components subjected to 

application conditions via the use of life prediction codes. 

Accomplishments 

• Completed 500-h engine field testing for advanced TiAl and silicon nitride valves using Caterpillar G3406 (inline 6 cylinders) nature gas engine. 

• Completed mechanical reliability evaluation of TiAl turbo wheel component. 

• Completed dynamic scratch studies for both Daido and Howmet TiAl alloys. 

• Completed initial generation of long-term tensile creep database for Daido and Howmet TiAl alloys. 

109
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Future Direction 

• Characterize the retained mechanical properties and microstructure of Kyocera SN235P silicon nitride and TiAl 

valves as a function of engine test time, and verify the probabilistic component design and life prediction. 

• Develop a high-temperature tensile creep and step-stress rupture database of commercially available TiAl alloys 

for probabilistic component design and life prediction for diesel exhaust valve and turbine wheel component 

application. 

• Develop mechanical database for complex-shaped TiAl components for verification of probabilistic component 

design and life prediction, and also process optimization. 

Introduction 

There has been considerable interest in the 

potential for extensive use of advanced ceramics and 

intermetallic alloys in advanced diesel engine 

systems because of their superior thermomechanical 

properties at elevated temperatures. The 

implementation of components fabricated from these 

advanced materials would lead to significant 

improvement in engine efficiency and long-term 

durability and reduced nitrogen oxides (NOx) and 

CO exhaust emissions as required in the 21st 

Century Truck Program. Interest has focused 

primarily on research into characterization and 

design methodology development (life prediction) 

for advanced silicon nitride ceramics and TiAl 

alloys to enable the manufacture of consistent, 

reliable complex-shaped components for diesel 

engine. The valid prediction of mechanical 

reliability and service life is a prerequisite for 

successful use of these materials in internal 

combustion engine components. 

This research project has three primary goals: 

the generation of a mechanical engineering database, 

from ambient to high temperatures, of candidate 

advanced materials before and after exposure to 

simulated engine environments; the microstructural 

characterization of failure phenomena in these 

advanced materials and in components fabricated 

from them; and the application and verification of 

probabilistic life prediction methods using diesel 

engine components as test cases. For all three stages, 

results will be provided to both material suppliers 

and component end-users for use in refining and 

optimizing processing parameters to achieve 

consistent mechanical reliability, and in validating 

the probabilistic design and life prediction of engine 

components made from these advanced materials. 

Approach 

The biaxial flexure strength [1,2] of TiAl turbo 

wheel component, as shown in Fig. 1, was measured 

using a ball-on-ring arrangement. Disk-shaped 

specimens were machined from both the airfoil 

surfaces and hub region by first diamond core 

drilling small cylinders having a nominal diameter 

of 10.0 mm. Each cylinder was then machined on 

one face only until the thickness was 0.5 mm. In 

this way, one face of each specimen always 

consisted of the as-manufactured surface of the 

airfoils. Discs machined from hub region were all 

with machined surfaces. Thus during testing, the asmanufactured surfaces of airfoil discs were always 

loaded in tension. 

The test fixture consisted of a 1-mm-diameter 

WC ball, which was mounted on a miniature load 

cell, as shown in Fig. 2. The lower support ring, 

which was 5.0 mm in diameter, was fabricated from 

a high-strength polymer. The test fixture itself was 

mounted on a vertical stepper motor (Z stage), 

which was affixed to the X-Y stages for positioning 

in the horizontal plane. A personal computer 

controlled all three stages using the LabView 

program (National Instruments, Austin, TX). After 

placing a specimen on the lower support ring, the XY stages were used to position the assembly directly 

under the upper load ball. A calibration procedure 

was carried out to ensure the WC ball was centered 

on the lower support ring prior to the testing. The Zstage was then lowered at a rate of 0.1 mm/s until 

the WC ball made a light contact with the specimen. 

The specimen was subsequently loaded to failure at 

a displacement rate of 0.01 mm/s. The computer 

monitored and recorded the displacement, load, and 

time. The disc strength, Sb, was calculated from the 

equation 

Sb = 3P(1+ν)/(4ν t2

)•[1 + 2ln(a/b) + ((1­

ν)/(1+ν))(1 – b2

/2a2)(a2/R2)] (1) 

110
Heavy Vehicle Propulsion Materials 

where P is the ultimate sustained load, a is the radius 

of the support ring, b is the effective radius of 

contact of the loading ball on the specimen, R is the 

specimen radius, t is the specimen thickness, and ν is 

Poisson's ratio. 

The principle of the experimental scratch-testing 

technique carried out on TiAl alloys has been 

described elsewhere [3-5]. The single-grit scratch 

testing apparatus (developed by Michigan 

Technological University Houghton, MI) consists of 

a rigid platform on which the specimen fixture is 

mounted and that has translation capability for ease 

of specimen mounting and positioning. A 

micrometer is mounted on the top of the pendulum 

such that its shaft is set directly on the top of a 

diamond tool. Thus, one can measure the tool 

position with the micrometer, while adjusting the 

tool in the center of the pendulum. The signals from 

load cells are outputted to a digitizer (NI5102) via 

charge amplifiers and acquired by a LabVIEW 

program. Conical diamond tools used (J & L 

Industrial Supply, Pittsburgh, PA) had a 90° 

included angle and shank sizes of 6.35 mm diameter 

x 50.8 mm length. The conical diamond tips 

generally had a radius of ~ 5 μm (based on the SEM 

image analysis). 

Results 

Characterization of TiAl Components 

Studies of biaxial strength of disc specimens 

machined from TiAl turbo wheel airfoils aimed to 

provide database directly from component for 

verification of probabilistic component design and 

life prediction task, which are critical to end users. 

Also, the database would provide guidelines to 

material suppliers for further optimization of 

component manufacturing processes to achieve 

consistent mechanical performance and reliability. 

Figure 3 compares the Weibull strength distribution 

of TiAl discs machined from airfoils with asprocessed and as-machined surface, respectively. 

Results show that the discs machined from airfoils 

with as-processed surface exhibit characteristic 

strength that is 34% lower than those obtained with 

as-machined surface. Also, these as-processed 

airfoil discs exhibit relative low Weibull modulus 

(m ~ 3.5) with respect to the as-machined specimens 

(m ~ 5.2). Note that huge variation in mechanical 

properties between as-processed and as-machined 

airfoils has also been observed in many silicon 

FY 2006 Progress Report 

nitride components; i.e., turbine blade, nozzle, and 

rotor [6]. The significant difference in Weibull 

strength distribution as found in previous studies of 

silicon nitride components was attributed to 

differences in the microstructure as well as strengthlimiting flaws present near the surface region of 

airfoils. Results suggest that the probabilistic life 

prediction task that was carried out for TiAl 

components based on the mechanical data generated 

from the simple shaped specimen geometry could 

possibly over-estimate the lifetime performance of 

turbo wheel under the application conditions. Thus, 

it is critical to re-verify these tasks by using the 

database directly generated from complex-shaped 

components in order to achieve reliable lifetime 

prediction and performance. 

Dynamic Scratch Test of TiAl Alloys 

TiAl alloys have also been considered for 

exhaust/intake valve application of heavy duty 

vehicles. In general, these exhaust/intake valves are 

always subject to the high-speed reciprocal 

movement during services. The responses of TiAl 

materials in the corresponded application condition 

need to be properly addressed before they can be 

successfully implemented in the engine systems. 

Thus, studies of dynamic scratch test using a singlegrit pendulum (rotational) scratch technique were 

carried out and completed during FY2006. The 

objective of this study was to provide insight into 

the dynamic abrasion response of TiAl alloys in 

sliding contact. This dynamic scratch approach is 

employed because of its advantages over the 

traditional sliding scratch technique in 

characterizing the materials removal mechanisms. 

TiAl alloys evaluated include Daido HIP and nonHIP TiAl, and Howmet TiAl material, which exhibit 

differences in grain domain and lamellar size, and 

thus could lead to differences in mechanical and 

wear response. 

Figure 5 shows the typical scratch surface 

feature of Daido HIP and Howmet TiAl alloy. The 

groove surfaces of scratches of two TiAls appeared 

to have maintained a certain degree of integrity with 

clear surface contours, which is distinctly different 

from scratches on brittle materials; i.e., alumina and 

silicon nitride (where the side ridges were not 

preserved due to the lateral cracks). The side ridges 

are developed on banks with appreciable pile-up, 

especially over the wide middle region of the scratch 

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Heavy Vehicle Propulsion Materials 

that corresponds to the suppressed wear rate. 

Frequent fractures take place in the transverse 

direction over the middle region of the scratch, some 

of which are sizable and run into the bottom of the 

groove. For the Daido TiAl, no shallow lateral 

fractures were developed that are observed to occur 

in pure titanium. However, the lateral fractures were 

developed for the Howmet TiAl on the sides of the 

scratch as a result of cleavage; thus the material 

removal has been enhanced with a predicted high 

wear rate. 

The scratch forces measured by load cells in the 

time domain were transformed and expressed in 

terms of the tool position using a linear relationship. 

The force signals of representative scratches for the 

two TiAls are given in the top row of Fig. 6. The 

scratch speeds obtained were ~ 1 m/s, provided that 

these scratches were generated within 0.003 to 0.004 

s with a length of 3 to 4 mm. The normal and 

tangential forces of the two scratches show a bell 

shape; they rise as the tool tip penetrates into the 

specimen, reach the peak in the middle, and then fall 

as the tool returns to the surface. The force signals 

are not necessarily symmetrical, especially for the 

tangential force that usually skews toward the right 

direction (on upright portion of the stroke). As a 

result of that, the values of OFC (overall frictional 

coefficient, defined as the ratio of tangential force to 

normal force) generally exhibit an ascending trend 

over the wide middle region of the scratch. 

The geometries of the scratch groove section 

were extracted at every sampling point along the 

longitudinal axis X. The relevant quantities include 

the depth of groove, groove section area, plowing 

section area, and removed section area. Among 

those parameters, the groove area is the sectional 

area of the scratch below the original level, the 

plowing area is the sectional area in contact with the 

tool face during scratching, and the removed area is 

the removed groove area (i.e. the groove area minus 

the total area of two side ridges). Variations of these 

quantities along the axis X, therefore, result in the 

derived scratch profiles with respect to the groove 

depth and the sectional areas. As expected, the 

profile of the depth of groove shows an arch-shaped 

curve in response to the imposed circular motion of 

the tool tip. The derived profiles of sectional areas 

appear with the similar curves, but they resemble 

more the shape of scratch forces than the profiles of 

groove depth. The bottom row of Fig. 6 gives the 

corresponding profiles of sectional areas of the 

FY 2006 Progress Report 

respective scratches. As expected, the level of the 

overall groove area is higher than that of the 

removed area but lower than that of the plowing 

area. In addition, the profiles of both forces and 

sectional areas are generally seen with appreciable 

oscillations whose amplitude increases 

proportionally with the depth of penetration. 

The differences between the defined sectional 

areas for the shown scratch in Fig. 5 imply the 

susceptibility of material loss or wear. If looking at 

the wear rate (defined as the ratio of the removed 

sectional area to the groove sectional area), then one 

can find that the wear rates of the tested TiAls vary 

slightly over the whole scratch process. Over the 

wide middle region of the scratch, the wear rate 

profiles illustrate a relatively steady decrease with a 

moderate suppression around the middle; this 

becomes pronounced in the case of Daido TiAl HIP. 

Overall, the wear rate of the Daido TiAl HIP is 

lower than other TiAl alloy evaluated due to its 

stronger capability of energy dissipation or 

resistance against material loss or removal from 

dynamic scratching. 

Characterization of Engine-Tested Si3N4 and TiAl 

Valves 

To validate the valve component design and life 

prediction tasks carried out for advanced heavy duty 

engines, the engine field tests using Caterpillar 

G3406 (in-line 6 cylinders), as shown in the top 

photo of Fig. 7, have been performed during 

FY2006. The 500-h engine field test for silicon 

nitride and TiAl exhaust, as well as intake valves, 

has been successfully accomplished. The metallic 

valves currently employed in the engine were also 

included for comparison. The valve position of each 

material in the engine is also illustrated in the 

bottom photo of Fig. 7. Following engine testing 

visual and optical examinations were carried out to 

inspect the tested valves for any apparent surface 

damages (chipping and cracking). Figure 8 shows 

the typical surface features for silicon nitride and 

TiAl exhaust valve after engine test. Similar 

features were also observed for the silicon nitride 

intake valves as well. Results of visual and optical 

inspection showed that there was not apparent 

surface damage introduced during the 500-h engine 

test, except apparent changes in surface morphology 

(roughness) due to the reciprocating contacts 

between valve guide and stem region. 

112
Heavy Vehicle Propulsion Materials 

A Taylor-Hobson surface profiling system 

(Model # Form Talysurf 120) was then employed to 

measurement the surface roughness (and thus degree 

of wear) of these tested valves. Three sections; i.e., 

upper valve guide, lower valve guide, and valve seat 

insert contact region, were evaluated, and three 

scans at each location were taken at 0, 120, and 

240°, respectively. Results showed that there was a 

significant increase (1.5-5 folds) in roughness (Ra 

values) in the valve seat region for both silicon 

nitride and TiAl valves. As for the valve guide 

region, the Ra values (roughness) of silicon nitride 

valves decreased after 100-h engine test probably 

due to the polishing effect, and exhibited a slightly 

increased roughness after 500-h test, as shown in 

Fig. 9. Nonetheless, the Ra values of 500-h tested 

silicon nitride valves are comparable to those 

measured for the as-received valves. On the other 

hand, the TiAl exhaust valves showed a much more 

increase in Ra value (2-10 folds) as compared to the 

silicon nitride valves, probably due to the loss of 

oxide layer formed during the heat treatment 

performed prior to the engine testing. The effect of 

the surface roughness introduced after 500-h engine 

test on the mechanical performance will be 

evaluated via destructive test using half-cylindrical 

bend bars after the completion of NDA task 

performed at ANL. Also, 1000-h engine field test 

will be continued in FY2007 to provide important 

insight into the long-term effect of engine 

environment on the mechanical reliability and 

component durability of TiAl and silicon nitride 

valves. 

References 

1. R. Thiruvengadaswamy and R. O. 

Scattergood, "Biaxial Flexure Testing of Brittle 

Materials," Scripta Metall. et Mater., Vol 25, 1991, 

pp2529-2532. 

2. D. K. Shetty, A. R. Rosenfield, P. McGuire, 

B. K. Bansal, and W. H. Duckworth, "Biaxial 

Flexure Tests for Ceramics," Ceramic Bulletin, Vol 

59 [12], 1193-97 (1980). 

FY 2006 Progress Report 

3. H. Wang, G. Subhash, A. Chandra, 

“Characteristics of Single-grit Rotating Scratch With 

a Conical Tool on Pure Titanium,” Wear, 249: 566­

581, 2001 

4. G. Subhash, J. E. Loukus, S. M. Pandit, 

“Application of Data Dependent Systems Approach 

for Evaluation of Fracture Modes During a SingleGrit Scratch,” Mech. Mater. 34: 25-42, 2002 

5. H. Wang, A. A. Wereszczak, M. J. Lance, 

“Effect of Grain Size on Dynamic Scratch Response 

in Alumina,” Ceram. Eng. Sci. Proc. (2006), in 

print. 

6. H. T. Lin and M. K. Ferber, “Characterization 

of Mechanical reliability of Silicon Nitride 

Microturbine Rotors,” Key Engineering Materials 

Vol. 287 (2005), pp. 393-403. 

Presentations and Publications 

Presentation 

H. Wang, H. T. Lin, A. A. Wereszczak, J. A. 

Jensen, and N. Yang, “Specific Energy and Scratch 

Hardness of Gamma Titanium Aluminides 

Subjected to Single-Grit Pendulum Scratching,” 

presented at ASME Internal Combustion Engine 

Division 2006 Fall Technical Conference, 

November 5-8, Sacramento, CA. 

Publications 

1. H. Wang, H. T. Lin, A. A. Wereszczak, J. A. 

Jensen, and N. Yang, “Specific Energy and Scratch 

Hardness of Gamma Titanium Aluminides 

Subjected to Single-Grit Pendulum Scratching,” 

ICEF2006-1532, published in the Proceedings of 

ASME Internal Combustion Engine Division 2006 

Fall Technical Conference, November 5-8, 

Sacramento, CA. 

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,” 

J. Mater. Sci. (in press). 

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Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 1. Photo of TiAl turbo wheel acquired from Caterpillar. 

Figure 2. Photos show the biaxial test facility used for testing biaxial discs machined from the TiAl turbo 

wheel airfoils. 

114
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

(a) (b) 

Fig. 3 Components of the single-grit pendulum scratch tester: (a) specimen fixture mounted on a rigid 

platform that is operated by a handle (not shown); (b) pendulum, equipped with a positioning micrometer on 

the top, is driven by a piston that is pneumatically operated by high pressure gas 

115
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

TiAl Turbo Wheel Airfoils

Uncensored Biaxial Strength Distribution

20°C - 0.1 mm/s - As-processed Surface

Max. Lik. Fit 

Strength Data 

m = 3.48 (2.49, 4.66)

σθ

 = 795 MPa (704, 892) 

n = 27 specimens 

95% Confidence 

Bands Shown 

Ball-on-Ring

Loading rate: 0.1mm/s 

2 

1 

0 

-1 

-2 

-3 

-4 

-5 

-6 

30 100 1000 

Failure Stress (MPa) 

TiAl Turbo Wheel Airfoils 

99.9 

99.0 

90.0 

50.0 

20.0 

10.0 

5.0 

2.0 

1.0 

0.5 

4000 

ln ln ( 1 / ( 1 - P ) ) ln ln ( 1 / ( 1 - Pf ) ) f

Probability of Failure, P , (%) Probability of Failure, Pf , (%) f

Uncensored Biaxial Strength Distribution

20°C - 0.1 mm/s - As-machined Surface 

2 

Max. Lik. Fit 

Strength Data 

m = 5.20 (3.80, 6.86)

σθ

 = 1200 MPa (1107, 1294) 

n = 28 specimens 

95% Confidence 

Bands Shown 

Ball-on-Ring

Loading rate: 0.1mm/s 

99.9 

99.0 

1 90.0 

0 

50.0 

-1 

20.0 

-2 10.0 

-3 5.0 

-4 2.0

1.0 

-5 

0.5 

-6 

100 1000 4000 

Failure Stress (MPa) 

Figure 4. Uncensored Weibull strength distribution of TiAl biaxial discs machined from airfoil with asprocessed surface (top) and as-machined surface (bottom). 

116
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

 (a)

 (b) 

Fig. 5 SEM micrographs of scratch on: (a) Daido TiAl HIP and (b) HOWMET TiAl. Scratch ran from left to 

right. The micrographs were taken after chemical etching. 

(a) (b) 

Fig. 6 Variations of scratch forces (top row) and sectional areas (bottom row) for a typical scratch on (a) 

Daido TiAl HIP and (b) HOWMET TiAl. 

117
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Si3N4 

Valves 

TiAl 

Valves 

Original Steel Valves 

4 3 2 1 6 5

Figure 7. Photos show Caterpillar G3406 (in-line 6 cylinders) nature gas engine (top) and valve position of 

each material in the engine compartment (bottom). 

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Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

After 100h engine 

test 

After 500h engine 

test 

After 100h engine 

test 

After 500h engine 

test 

After 400h engine 

test 

Figure 8. Comparisons of surface features of silicon nitride valve, TiAl valve, and steel valve as a function of 

engine test time. 

119
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

0 

0.05 

0.1 

0.15 

0.2 

0.25 

0.3 

0.35 

E1 E3 E5 E7 E9 E11 E13 E15 I2 I4 I6 I8 I10 I12 

Valve 

0 h 

100 h 

500 h 

0 

0.2 

0.4 

0.6 

0.8 

1 

1.2 

E1 E3 E5 E7 E9 E11 E13 E15 I2 I4 I6 I8 I10 I12 

Valve 

0 h 

100 h 

500 h 

Figure 9. Roughness Ra valves measured for valves in the as-received, 100-h test, and 500-h test condition. 

Note that Btop and Ctop are two locations above and below the valve guide contact region, respectively. 

120
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

C. Engineered Surfaces for Diesel Engine Components 

M. Brad Beardsley 

Caterpillar, Inc. 

Technical Center Bldg. E/854 

P.O. Box 1875 

Peoria, IL 61656-1875 

(309) 578-8514; fax: (309) 578-2953; e-mail:Beardsley_M_Brad@cat.com 

DOE Technology Development Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

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

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: Caterpillar, Inc., Peoria, Illinois 

Objectives 

• Develop coating processes to apply amorphous steels and quasicrystal materials to engine pistons and cylinder 

heads for use as thermal barrier coatings in homogenous charged combustion ignition (HCCI) engine designs. 

• Quantify the diffusion stability of new quasicrystalline material at engine operating temperatures. 

Approach 

• Use high-velocity oxygen fueled (HVOF) processes to apply amorphous and quasicrystalline materials to 

engine piston and cylinder head. 

• Use microprobe analysis of diffusion couples to determine the stability of quasicrystalline and bond coating 

couples at engine operation temperatures. 

Accomplishments 

• A HCCI piston coated with amorphous steel demonstrated improved coating durability over ceramic thermal 

barrier coatings in single cylinder engine tested. 

• Thermal conductivity measured for the amorphous steel coating was higher than that reported for this type of 

material and resulted in higher than desired heat rejection through the piston. 

Future Direction 

• Develop graded coating structures for the quasicrystalline material based upon the diffusion profiles developed. 

Introduction 

Two new materials, amorphous steels and 

quasicrystalline materials, are investigated as 

thermal barrier coatings. Both materials are 

capable of being sprayed using high-velocity 

oxygen fuel (HVOF) processes that produce 

denser coatings with higher strengths than plasma 

spraying. The denser coating structure of HVOF 

coatings eliminates the need for sealing of the 

coating porosity and the higher strength enhances 

the coating durability. Amorphous steel coatings 

are reported to have low thermal conductivity as 

121 
Heavy Vehicle Propulsion Materials 

well as high hardness and toughness, making them 

attractive as thermal barrier coatings.1,2

 The 

quasicrystal being studied, Al71Co13Cr8Fe8, has 

been shown to be capable of 1100°C operation in a 

turbine engine.3 

Approach 

Amorphous Steel Coatings 

A piston was coated with the SHS717 

amorphous steel produced by The NanoSteel 

Company, Maitland, Florida. The coating was 

nominally 3.5 mm in thickness. The piston was 

run for approximately 24 hours in a single cylinder 

test during which component temperatures were 

measured. 

Quasicrystalline Materials 

Analysis of the chemical mobility of the 

species in the quasicrystalline coating has 

continued. Interaction zones of up to 100 micon 

in thickness at the interface between the 

quasicrystal and NiCrAlY bond coating are being 

analyzed. 

Results 

Amorphous Steel Coatings 

The piston temperatures measured were higher 

than expected for the reported thermal 

conductivity of the amorphous steel coatings. A 

measurement of the HVOF produced coating 

showed that the thermal conductivity was 1.5 to 2 

times higher than the 5 W/m-K reported, Figure 1. 

Lower thermal conductivity amorphous materials 

FY 2006 Progress Report 

will be needed if they are to be of use for 

applications to diesel engine components. 

Conclusions 

Amorphous Steel Coatings 

The durability of the amorphous steel coating 

showed significant improvement over plasma 

sprayed ceramic thermal barrier coatings. 

However, lower thermal conductivity 

compositions will need to be identified in order to 

provide the required thermal resistance needed for 

diesel engine component applications. 

References 

1. D. J. Branagan, W. D. Swank, D. C. 

Haggard, and J. R. Fincke, “Wear Resistant 

Amorphous and Nanocomposite Steel Coatings,” 

Met. & Mat. Transactions A, 32A, 2615–2621 

(2001). 

2. D. Shin, F. Gitzhofer, and C. Moreau, 

“Development of Metal Based Thermal Barrier 

Coatings (MBTBCs) for Low Heat Rejection 

Diesel Engines,” Proceedings of the 2005 

International Thermal Spray Conference, 915–919 

(2005). 

3. A. Sanchez, F. J. Garcia, J. M. Algaba, 

J. Alvarez, P. Valles, M. C. Garcia-Poggio, and 

A. Aguero, “Application of Quasicrystalline 

Materials as Thermal Barriers in Aeronautics and 

Future Perspectives of Use for These Materials,” 

Mat. Res. Soc. Symp. Proc., 553, 447–458 (1999). 

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Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Thermal Conductivity 

5.000 

10.000 

15.000 

20.000 

25.000 

0 200 400 600 800 1000 1200 

Temperature, C 

Ther

m

al Conductivity, W/cm-K 

AS-SPRAYED 

ANNEALED 

` 

Figure 1. Exposure to temperatures above glass transition temperature (700°C) causes an increase in thermal diffusivity 

of amorphous steel. 

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Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

D. High-Temperature Advanced Materials for Lightweight Valve Train 

Components 

Nathaniel S. L. Phillips, Principal Investigator 

Jeff Jensen, Jeremy Trethewey, Natalie Vanderspiegel, David Longanbach, and Chad Logan 

Caterpillar Inc. 

Tech Center 854-22 TC-E 

14009 Old Galena Rd. 

Mossville IL, 61552 

(309) 578- 5788 (voice) 

(309) 578-2953 (fax) 

E-mail: Phillips_Nate@cat.com 

DOE Technology Development Area Specialist: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: jerry.gibbs@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

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

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract Number: DE-AC05-00OR22725 

Subcontractor: Caterpillar Inc., Peoria, Illinois 

Objective 

• Evaluate prospective lightweight valve materials and determine feasibility of implementation in 

heavy-duty diesel and natural gas applications. 

Approach 

• Design, procure, and evaluate γ-phase titanium aluminide (TiAl) and silicon nitride (Si3N4) heavyduty engine valves. 

• Demonstrate performance and durability advantages by performing 1000-h durability test on 

valves in a Caterpillar G3406 natural gas engine. 

• Develop finite-element analysis (FEA)-based life prediction model (NASA CARES/Life) that will 

generate life prediction curves for advanced valve materials. 

• Evaluate the thermal, chemical and mechanical properties of TiAl and Si3N4 that are critical to onengine valve performance and durability. 

• Investigate manufacturing considerations (i.e., machinability and quality inspection technique) of 

these novel materials. 

Accomplishments 

• Completed 500 hour engine testing. 

• Completed 500-hour post-test non-destructive analysis (NDE). 

125
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Future Direction 

• 2/3 of the valves will be re-installed in the engine while the remaining 1/3 will be replaced by 

fresh valves and reserved for destructive testing. 

• After the engine has run for another 500 hours, all valves will be removed, inspected, and prepared 

for destructive testing. By performing fast fracture tests on valves that have accumulated 0, 500, 

and 1000 hours of engine time, the effect of engine aging on the materials’ retained strength will 

be examined. 

• Along with combustion dynamics and fuel consumption measurements, this data will enable the 

team to quantify performance gains and establish useful life estimates of Si3N4 and TiAl valves in 

a heavy-duty engine environment. 

Introduction 

Valve train components in heavyduty engines operate under high 

stresses, at elevated temperatures, and in 

severely corrosive environments. 

Structural ceramics and emerging 

intermetallic materials are highly 

corrosion and oxidation resistant and 

possess high strength and hardness at 

elevated temperatures. These properties 

are expected to allow higher engine 

operating temperatures, lower wear, and 

enhanced reliability. In addition, the 

lighter weight of these materials (~50% 

of production alloys) will lead to lower 

reciprocating valve train mass that could 

improve fuel efficiency. 

Over the past decade, the 

automotive engine industry has 

demonstrated moderate success with 

TiAl as a valve material. Eylon et al.1 

demonstrated a 2% fuel savings on a 

Chevrolet Corvette that ran for 25,000 

km with TiAl valves. This fuel 

efficiency increase is attributed to the 

decrease in valve train parasitic 

frictional losses. In another study with 

TiAl valves, Maki et al.2

 demonstrated a 

1000-rpm increase in over-speed 

performance on a Nissan VRT35 

engine. This capability would allow 

more efficient engines to operate in 

more demanding regimes. The current 

research and development program at 

Caterpillar, Inc. extends this body of 

knowledge from the automotive 

community and examines TiAl and 

Si3N4 in the context of heavy-duty diesel 

and natural gas engine environments. 

The valve train material 

development effort will provide the 

materials, design, manufacturing, and 

economic information necessary to 

bring these new materials and 

technologies to commercial realization. 

With this information, component 

designs will be optimized using 

probabilistic lifetime prediction models, 

and validated in rig bench tests and 

short- and long-term engine tests. After 

establishing proof-of-concept with 

valves, this design approach will be 

applied to other components made from 

high-temperature materials. 

Approach 

Many perspectives must be 

considered when evaluating TiAl and 

Si3N4 materials for use in valve 

applications. As such, there are several 

facets to this project. Validation of these 

components begins with the fabrication 

of TiAl and Si3N4 valves and testing 

them on a Caterpillar G3406 (in-line 6 

cylinder) natural gas engine. Previous 

annual reports documented the design 

process used to modify the current 

baseline metallic valve to accommodate 

the limited ductility of ceramic and 

intermetallic materials.3

 For example, 

the head dimensions were changed to 

reduce the stress concentration 

126
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

experienced in the fillet radius. Silicon 

nitride and TiAl valve blanks were 

formed, rough machined, friction 

welded to a Ti-6Al-4V shaft (TiAl 

only), and finish machined. 

A parallel effort to develop a design 

tool for high-strength, brittle materials is 

also being pursued. A FEA-based life 

prediction code (NASA CARES/Life) is 

used to evaluate component design and 

generate life prediction curves. This 

approach uses material properties and 

thermo-mechanical boundary conditions 

as the primary inputs. A probabilitybased code then calculates the accumulated damage through the life of the 

valve and creates a life prediction curve. 

A variety of other studies contribute 

to the understanding of the behavior of 

TiAl and Si3N4 and how they may 

perform in an engine environment. 

Several studies examine the 

fundamental properties of the materials: 

tensile strength, creep strength, thermophysical properties, friction coefficient, 

wear resistance, oxidation resistance, 

and corrosion resistance. Two other 

studies examine the manufacturability of 

the materials: the effect of machining 

damage on fatigue resistance and a NDE 

technique to evaluate surface quality.

4 

Results 

Pre-engine test valve characterization 

Prior to subjecting the TiAl and 

Si3N4 valves to the full-scale engine test, 

three characterization techniques were 

used to gather baseline information 

about the valve dimensions, material 

surface quality and structural integrity. 

A coordinate measuring machine 

(CMM) profile scan was performed 

along the axis of each valve to verify 

that they met the dimensional 

specifications. 

In order to establish the pre-engine 

test surface condition, profilometry 

scans were performed on 5 key regions: 

keeper groove, Ti-6-4 stem, TiAl stem, 

fillet and the seat contact region. 

Figure 1 shows a series of roughness 

measurements for a single TiAl valve. 

It was confirmed that the surface 

roughness specification (Ra) was met 

for nearly all of the measured regions on 

all of the valves. 

The most sophisticated set of data 

collected was the laser scatter scans 

performed by Dr. JG Sun at Argonne 

National Laboratory. The fundamental 

physics and experimental setup are 

detailed in the “NDE of Diesel Engine 

Components” quarterly reports. Scans 

were performed at the fillet and keeper 

groove locations on all TiAl and Si3N4 

valves. An additional scan of the weld 

region was also performed to determine 

if any cracks formed during the friction 

welding process. Figure 2 shows two 

scans of the fillet region of the same 

Si3N4 valve. Figure 2a shows the scan 

after poor finish machining. During the 

last year’s efforts these valves were remachined to meet the desired surface 

quality. Figure 2b shows the same valve 

after corrective finish machining. 

Valves with little or no detectable 

damage were used for the engine test. 

Engine testing of Titanium 

Aluminide, Silicon Nitride and 

production valves 

One of the primary goals of this 

study is to evaluate the performance of 

ceramic (Si3N4) and intermetallic (TiAl) 

valves in an engine environment. As 

outlined in previous quarterly reports 

(see Oct-Dec 2004), arrangements have 

been made to perform parallel 1000-

hour endurance tests on three valve 

materials: γ-TiAl, Si3N4 and the 

austenitic steel production valve. These 

tests were performed at the National 

Transportation Research Center (NTRC) 

on a natural gas G3406 genset. 

Since the design of these “advanced 

material” valves was optimized for the 

less ductile nature of these materials, 

several additional modifications to the 

G3406 valve train components had to be 

127
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

made, including seat inserts, valve 

guides, keeper locks, and bridges. 

Caterpillar and NTRC personnel 

oversaw the installation of 4 TiAl and 

10 Si3N4 valves at Stowers Industrial 

Power, the East Tennessee Regional 

Caterpillar dealership. During the 

installation process, it was requested 

that the mechanic handle the valves with 

the same degree of care in which he 

would handle a standard steel valve. 

Figure 3 shows the valves as they are 

being installed. It is more realistic to 

require that the new material is 

sufficiently robust to survive the current 

standard installation procedure than to 

modify the procedure to conform to the 

limitations of the new material. The fact 

that all 14 of the “advanced valves” 

survived installation is an indication of 

the handling robustness of the materials. 

Figure 4 shows the valve array 

immediately after installation. It is 

worth noting that all of the customdesigned peripheral valve train 

components (seat inserts, guides, keeper 

locks, and bridges) were successfully 

integrated in the “advanced valve” head 

re-build. 

In addition to valve train hardware, 

instrumentation specific to this engine 

test was also installed. Calibrated 

Kistler 6125B pressure pick-ups were 

installed in each of the six cylinders. 

Pressure traces from these transducers 

enabled the evaluation of the valve 

material’s effect on the combustion 

performance (i.e. heat release rate). 

Also, a fuel consumption meter was 

installed to estimate the valves’ 

aggregate effect on fuel efficiency. 

Finally, as a safety precaution, a fuel 

line shut-off valve was installed 

upstream of the genset intake and was 

calibrated to trigger if abnormal engine 

vibration occurred. 

The table below (Table 1) outlines 

the progress of this engine durability 

test. 

Engine operating conditions 

The valves were installed in the 

following arrangement (Table 2) to 

ensure that the performance of the Si3N4 

and TiAl valves could be isolated during 

post-test analysis. 

Table 1 Engine durability test progress 

Phase Action Status 

Initial Installation Valve Installation Completed in March 

Phase #1 Engine Exposure Completed in May 

(0 h – 100 h) Post-test Analysis Completed in June 

Phase #2 Engine Exposure Completed in July 

(100 h – 500 h) Post-test Analysis Completed in November 

Phase #3 Engine Exposure Not yet started 

(500 h – 1000 h) Post-test Analysis Not yet started 

Table 2 Valve installation arrangement 

Cylinder Intake Exhaust 

1 Silicon Nitride Silicon Nitride 

2 Silicon Nitride Silicon Nitride 

3 Silicon Nitride Titanium Aluminide 

4 Production Steel Titanium Aluminide 

5 Production Steel Production Steel 

6 Production Steel Production Steel 

128
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Within the first few hours of engine 

operation, Cylinder 1, outfitted 

exclusively with Si3N4 valves, began to 

exhibit the signs of knocking. This preignition condition is a symptom of 

altered combustion dynamics and is 

most likely caused by the protrusion of 

the seated valve ~1.5 mm farther into 

the combustion chamber than originally 

designed. This protrusion may cause 

pre-ignition because it changes the 

compression ratio and provokes 

autoignition prior to spark-controlled 

ignition. It is also possible that local 

hotspots are generated on the valve face 

due to the increased thermal mass in the 

hottest region of the combustion 

chamber. The protruding portion of the 

valve head exacerbates this hotspot 

formation, a potential cause of preignition. Because this portion of the 

valve head does not contact the valve 

seat, there is no direct path for thermal 

conductivity and the local material 

temperature increases. It is important to 

note that this protrusion is an artifact of 

the valve geometry re-design and is 

correctable in future valve design 

iterations. To circumvent knocking, an 

exceedingly harsh operating 

environment, the genset was derated 

from 235 kW to 125 kW for the 

remainder of the test. 

Valve Inspection – 

Visual & Profilometry 

In a heavy-duty engine, the first 100 

hours of service time is considered the 

“break-in” period. During this initial 

phase the valve and seat insert conform 

to each other by a scuffing wear 

mechanism. Valve/guide contact may 

also cause wear on the stem while 

valve/bridge contact may cause wear on 

the stem tip. After 100 hours, the Si3N4 

and TiAl valves showed minimal stem 

tip wear, acceptable valve/guide wear, 

and significant seat/insert wear. Figures 

5a-c show a representative Si3N4 

exhaust valve (Si3N4E-32) after the first 

100 hours of testing. As expected, there 

was an even coating of combustion 

byproduct on the valve’s head face, fillet 

and lower valve stem regions. The stem 

wear was evident in the discoloration 

ending halfway up the valve. Figure 5d 

shows the wear profile of the valve seat 

region. Figures 6a-d present the same 

results for a representative TiAl valve 

(TiAl-5). 

In addition to wear scar depth, 

surface roughness, or RA, also allows 

direct comparison of worn contact 

regions. Table 3 displays the surface 

roughness of each Si3N4 and TiAl valve 

before and after the first 100 hours of 

engine exposure. 

Several preliminary conclusions 

may be drawn from these photographs, 

profilometry scans and surface 

roughness measurements. 

Valve/seat wear: The 2x increase in 

valve/seat surface roughness for 

Si3N4E-32 is evidence of break-in 

scuffing wear on the Si3N4 valve. 

Figure 5d shows a relatively constant 

wear scar depth of ~2μm. However, the 

5x increase in valve/seat surface 

roughness for TiAl-5 is an indication 

that wear has exceeded standard breakin levels. This is confirmed by Figure 

6d which shows that the valve/seat wear 

scar depth was ~15μm, a significant 

magnitude for 100 hours of engine 

exposure. This accelerated wear in the 

valve/seat region of the TiAl valve 

suggests that an optimized wearresistant surface treatment may be 

necessary to ensure acceptable wear 

behavior. When the 500 hour and 1000 

hour wear data has been captured firm 

conclusions will be drawn from these 

observations. 

129
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Table 3. Surface roughness measurement of Si3N4 and TiAl valves at 0 hours and 100 hours 

Valve Time Valve Guide Valve Seat 

Valve ID # Location (hrs) (RA) (RA) 

Si3N4 Intake 

Si3N4I-20 I2-SI3N4 0 0.067 0.152 

100 0.095 0.226 

Si3N4I-21 I4-SI3N4 0 0.061 0.15 

100 0.079 0.779 

Si3N4I-25 I6-SI3N4 0 0.069 0.143 

100 0.069 0.199 

Si3N4I-28 I8-SI3N4 0 0.063 0.178 

100 0.093 0.435 

Si3N4I-33 I10-SI3N4 0 0.067 0.174 

100 0.089 0.185 

Si3N4I-41 I12-SI3N4 0 0.069 0.174 

100 0.096 0.394 

Si3N4 Exhaust 

Si3N4E-11 E1-SI3N4 0 0.063 0.232 

100 0.059 0.436 

Si3N4E-17 E3-SI3N4 0 0.06 0.211 

100 0.027 0.888 

Si3N4E-9 E5-SI3N4 0 0.063 0.222 

100 0.037 0.382 

Si3N4E-32 E7-SI3N4 0 0.066 0.206 

100 0.034 0.389 

Ti TiAl 

TiAl Exhaust 

TiAl-5 E9-TiAl 0 0.121 0.121 0.187 

100 0.239 0.316 0.970 

TiAl-6 E11-TiAl 0 0.102 0.102 0.221 

100 0.316 1.054 1.157 

TiAl-7 E13-TiAl 0 0.121 0.121 0.211 

100 0.269 0.661 0.762 

TiAl-8 E15-TiAl 0 0.12 0.12 0.226 

100 0.318 1.097 1.143 

Valve/guide wear: The surface 

roughness data suggests that minor 

scuffing wear moderately increased the 

RA of the Si3N4 intake valves. In 

contrast, the RA of the Si3N4 exhaust 

valves decreased because the higher 

stem temperature enabled the 

stem/guide to undergo a polishing wear 

mode. The TiAl valves exhibited 

inferior wear performance compared to 

the Si3N4 valves in the valve/guide 

region. However, the 2x-3x RA increase 

in the Ti-6-4 segment of the stem 

indicates that the thermal oxidation heat 

treatment was more effective on the Ti6-4 than on the TiAl, which showed a 

3x-10x RA increase. 

Valve integrity: Although the wear 

mechanisms and rates must be better 

understood before Si3N4 and TiAl can be 

successfully implemented into 

production, the fact that these valves 

have survived 500 hours of service is an 

indication that these valve materials may 

be capable of withstanding a heavy duty 

engine environment. 

Valve Inspection – Laser Scatter NDE 

At the end of the first 100 hour 

“break-in” test phase each valve was 

scanned with the laser-based NDE 

(Non-Destructive Evaluation) 

instrument outlined in the project titled 

“NDE of Diesel Engine Components.” 

Figure 7 presents the NDE scans for 

exhaust valves Si3N4E-32 and TiAl-5 

(the same valves in Figures 5 and 6). In 

general, these valves scans are similar to 

the pre-test scans with a couple of 

exceptions. The region of interest in 

these images was the valve seat band, 

the location of highest accumulated 

damage. The white stripes above and 

below the valve seat band were most 

likely caused by a combustion 

byproduct and are not significant flaws. 

The continuous grayscale transition in 

the Si3N4 valve seat band (Figure 7a 

130
Heavy Vehicle Propulsion Materials 

insert) confirms that there was no 

notable surface damage in its valve/seat 

wear region. However, the valve seat 

band in the TiAl valve (Figure 7b insert) 

showed a significant level of surface 

damage. More detail may be found 

regarding NDE data interpretation in Dr. 

JG Sun’s Apr-Jun 2006 quarterly report. 

Development of a valve life prediction 

tool using CARES/Life 

The ongoing effort to predict the life 

expectancy of a heavy-duty valve 

subjected to typical engine loads had 

made significant progress in this past 

reporting period. The results of the first 

generation model were presented in a 

previous quarterly report (Jan-Mar 

2005). Since then, the most significant 

accomplishment has been the shift from 

an axisymmetric model to a threedimensional model. The additional 

degrees of freedom permit a more 

accurate representation of the valves 

deformation modes. In addition to the 

axisymmetric loads, the valve stem is 

permitted to bend and the valve head is 

permitted to deflect in a “pringle” 

deformation mode. Figure 8 presents 

the resultant stress field superimposed 

on a deformed valve. It is interesting to 

note that the maximum principal stresses 

observed on the 3D model is a factor of 

two greater than that observed on the 

preliminary axisymmetric model. 

After applying the thermal and 

mechanical boundary conditions to the 

Si3N4-based 3D model the 

NASA/CARES probabilistic code was 

executed to examine the slow crack 

growth failure mode as a function of 

service time. Figure 9 presents the 

resulting life prediction curve for a 

Si3N4 valve experiencing slow crack 

growth under continuous G3406 genset 

operating conditions. Note that two 

curves are displayed; one represents a 

safety factor of one (i.e. loads are 

identical to those specified by the 

simulation-derived boundary 

FY 2006 Progress Report 

conditions), the other represents a safety 

factor of two (i.e. applied loads were 

doubled). This 3D model is 

considerably more conservative than the 

preliminary axisymmetric model. The 

valve’s lifetime predicted with this 3D 

model is substantially shorter than the 

lifetime predicted by the axisymmetric 

model (see Jan-Mar 2005 quarterly 

report). Even so, a probability of failure 

corresponding to ~0.01 ppm failed 

valves after one year of continuous 

operation in a heavy-duty engine is 

exceptional! 

A couple of cautionary notes must 

be presented. This model relies on 

boundary conditions acquired from FEA 

simulations. Thus, they may not be 

exactly representative of field-tested 

engines. Also, this model only captures 

the probability of failure as a function of 

slow crack growth damage. As in any 

engineering system, several failure 

modes are acting simultaneously. This 

model focuses on the slow crack growth 

failure mode (prominent in ceramics). 

Other failure modes such as wear, 

impact resistance, and environmental 

resistance must be independently 

considered to gain a full perspective on 

the life prediction of a valve component. 

As discussed in previous reports 

(FY2005 Annual Report), Jadaan 

Engineering & Consulting has been 

developing a probabilistic lifetime 

prediction model for TiAl and Si3N4 

valves as they experience engine-like 

thermo-mechanical loads. Progress in 

this quarter expands upon the 

establishment of a three-dimensional 

FEA model for a Si3N4 valve (Oct-Dec 

2005) by incorporating TiAl material 

properties. Also, a sensitivity study 

using the Probabilistic Design System 

(PDS) function in ANSYS was 

performed. Below are three significant 

conclusions that may be drawn from this 

set of analyses: 

131
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Si3N4 Valve Performance: Life 

prediction estimations for Si3N4 valves 

can be construed as a cautious “vote of 

confidence” in the material’s reliability. 

For example, after 15,000 hours, a valve 

will have a probability of failure of 

8x10-11. Assuming an engine will 

contain 12 Si3N4 valves, 1 out of one 

billion engines would experience a 

catastrophic valve failure. However, 

doubling the stress field, an exceedingly 

conservative approach, would increase 

the probability of failure to 2x10-4 (1 

out of 450 engines would experience a 

catastrophic valve failure). 

TiAl Valve Performance: Results 

for TiAl valves were inconclusive. Two 

factors suggest that the TiAl valve 

predictions are of questionable accuracy. 

First, there was insufficient data to 

develop reliable statistical material 

characteristics. Second, the maximum 

stress, shown in Figure 10, is atypical of 

valve behavior. This may be an artifact 

of the insufficient material properties 

and/or boundary conditions. Caterpillar 

will review the FEA model for 

discontinuities. However, if this 

analysis is an accurate representation of 

the TiAl/Ti-6-4 stress field, this atypical 

stress “hot spot” may suggest that a redesign may be required to mitigate the 

valve’s maximum stresses. In other 

words, it is likely that TiAl may not be 

regarded as a “drop in” replacement for 

either the “current steel” or “advanced 

Si3N4” valve design. 

PDS Results: The sensitivity study 

examined the importance of six 

variables (valve face pressure, Weibull 

modulus, scale parameter, Young’s 

modulus, coefficient of thermal 

expansion and thermal conductivity) on 

the probability of failure and maximum 

stress of TiAl and Si3N4 valves. Figure 

11 summarizes these results. Young’s 

modulus, valve face pressure, and 

coefficient of thermal expansion are 

found to be the critical variables for 

defining the maximum stress found in 

the valve. However, Weibull modulus, 

Young’s modulus, and the scale 

parameter are the three strongest 

influences on the ultimate probability of 

failure. Of course, the accuracy of the 

PDS analysis is only as good as the 

accuracy of the FEA model (discussed 

above). With this in mind, the PDS 

analysis suggests that controlling the 

latter set of variables would have the 

largest impact on improving the valves’ 

durability. 

Conclusions 

Lightweight, high-strength, hightemperature materials are enablers for 

advanced combustion concepts and fuelefficient engine designs. This effort 

aims to examine two prospective 

materials, TiAl and Si3N4, in the context 

of heavy-duty diesel and natural gas 

engine environments. Progress has been 

made on several fronts this year, 

including: (1) engine tests on these 

valve materials, (2) development of a 

probabilistic life prediction model, and 

(3) compilation of material properties of 

TiAl and Si3N4 that are critical to 

implementation of these materials in the 

extreme environment of a heavy-duty 

engine. 

Twelve TiAl and forty-five Si3N4 

valves have completed the fabrication 

process. Their incorrect surface finish 

provided by the machining supplier was 

rectified. The valves have been 

subjected to several pre-engine test 

characterization assessments: CMM, 

profilometry and laser-scatter NDE. The 

first set of valves has accumulated 20 h 

on an impact rig and has demonstrated 

acceptable performance. Upon 

successful completion of the impact rig 

test, the valves will be tested on a 

Caterpillar G3406 engine at the NTRC. 

In addition to fabricating custom 

components that will enable this engine 

to accept the new modified valve design, 

132
Heavy Vehicle Propulsion Materials 

the engine has been instrumented to 

capture the performance of each 

cylinder. 

Significant progress was made 

toward developing a NASA 

CARES/Life prediction model. A 

preliminary model was developed, and a 

life prediction curve was generated. It 

was speculated that this preliminary 

curve significantly under predicted the 

probability of failure due to incomplete 

representation of the boundary 

conditions. Model enhancements are 

currently underway and updated results 

will be correlated to the engine 

performance of the valve hardware. 

Thermo-physical properties (e.g., 

coefficient of thermal expansion, 

thermal conductivity, density) and 

mechanical properties (e.g., ultimate 

tensile strength, Young’s modulus, 

ductility) of TiAl and Si3N4 have been 

collected. Properties specific to 

implementation of these advanced 

materials in a valve environment are 

also under investigation (e.g., 

machinability, frictional coefficients, 

corrosion resistance). 

This diverse effort will help to equip 

design engineers with the information 

required to implement advanced 

lightweight materials in a valve train 

application. This program has made 

significant progress toward this goal in 

FY 2006 and is poised to significantly 

advance the case for TiAl and/or Si3N4 

as a future valvetrain material in FY 

2007. 

References 

1. D. Eylon et al., “Development of 

permanent-mold cast TiAl automotive 

valves,” Intermetallics, 6, 1998. 

2. K. Maki et al., “Development of 

a high-performance TiAl exhaust 

valve,” SAE Paper 960303, Warrendale, 

Pennsylvania, 1996. 

3. M. J. Andrews and J. ChapaCabrera, “High-temperature advanced 

materials for lightweight valve train 

FY 2006 Progress Report 

components,” Heavy Vehicle Propulsion 

Materials 2003 Annual Progress Report, 

Energy Efficiency and Renewable 

Energy Office of FreedomCAR and 

Vehicle Technologies, March 2004. 

4. J. G. Sun, “NDE development 

for ceramic valves for diesel engines,” 

Heavy Vehicle Propulsion Materials 

2004 Annual Progress Report, Energy 

Efficiency and Renewable Energy 

Office of FreedomCAR and Vehicle 

Technologies, May 2005. 

5. J. S. Trethewey et al., “Hightemperature advanced materials for 

lightweight valve train components,” 

Heavy Vehicle Propulsion Materials 

2004 Annual Progress Report, Energy 

Efficiency and Renewable Energy 

Office of FreedomCAR and Vehicle 

Technologies, May 2005. 

6. Biery et al., “Use of Weibull 

statistics to quantify property variability 

in TiAl alloys,” Metallurgical and 

Materials Transactions A, 33A, 

3127–3136 (October 2002). 

Presentations and Publications 

O. Jadaan, J. S. Trethewey 

“Reliability of high temperature 

lightweight valve train components in a 

total probabilistic design environment,” 

paper presented at the American 

Ceramic Society’s 30th Ann. Intl. Cocoa 

Beach Conf. & Exposition on Advanced 

Ceramics & Composites, Cocoa Beach, 

Florida, January 21-26, 2006, and to be 

published in conference proceedings. 

J. G. Sun, J. M. Zhang, J. S. 

Trethewey “Nondestructive evaluation 

of machining damage in silicon nitride 

ceramic valves,” paper presented at the 

American Ceramic Society’s 30th Ann. 

Intl. Cocoa Beach Conf. & Exposition 

on Advanced Ceramics & Composites, 

Cocoa Beach, Florida, January 21-26, 

2006, and to be published in conference 

proceedings. 

133
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 1 Surface roughness measurements on a TiAl valve. Note that scan fluxuation may not 

appear to correlate with the stated Ra value – this is because the plotted range for each scan is 

different 

134
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 2 Laser scatter NDE scans of a Si3N4 valve (a) after incorrect finish machining and (b) after 

corrective finish machining. Note that the white streaks in (a-inset) that indicate subsurface damage 

have been removed by the corrective finish machining (b-inset) 

135
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 3 Si3N4 valve under compression from an installation tool 

Figure 4 Caterpillar G3406 cylinder head after “advanced valve” installation. From left to right 

there are 10 production steel valves (2 temporarily removed), 4 TiAl exhaust valves and 10 Si3N4 

valves 

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Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 5 (A) An image of Cylinder ‘B’ immediately after the completion of the first 100 hour breakin phase. (B,C) Valve Si3N4E-32, a representative Si3N4 exhaust valve. (D) Profilometry scan of the 

valve/seat wear scar showing a wear scar depth of ~2 μm 

137
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 6 (A) An image of Cylinder ‘C’ immediately after the completion of the first 100 hour breakin phase. (B,C) Valve TiAl-5, a representative TiAl exhaust valve. (D) Profilometry scan of the 

valve/seat wear scar showing a wear scar depth of ~15 μm 

138
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 7 Laser-based NDE scans of (A) Si3N4E-32 and (B) TiAl-5 indicating that the Si3N4 valve has 

not sustained major wear damage, but the TiAl valve is showing signs of significant valve/seat wear 

139
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 8 Principal stress field of the three-dimensional valve FEA model. Note the non-axisymmetric 

stem bending and “pringled” seat warping. The maximum stress observed in this 3D model is 

approximately a factor of two greater than the maximum stress observed in the axisymmetric model 

140
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 9 Life prediction curve for Si3N4 valve experiencing slow crack growth under continuous

G3406 genset operating conditions using the 3D model. Note that the safety factor corresponds to a 

multiplication factor on the applied loads. As such, a load factor of two would generate more severe 

loading conditions resulting in a higher probability of failure 

Figure 10 Principal stress field of TiAl/Ti-6-4 valve near peak cylinder pressure. The location of the 

maximum stress value (431 MPa) is not at the center of the fillet radius, which may indicate that the 

valve design is not optimized for the TiAl’s material properties 

141
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Figure 11 Results of the PDS sensitivity study highlighting the relative effect of the input variables on 

the probability of failure and maximum principal stress for Si3N4 and TiAl valves 

142
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

7. CHARACTERIZATION OF PHYSICAL AND MECHANICAL 

PROPERTIES 

A. Implementing Agreement for a Programme of Research and Development on 

Advanced Materials for Transportation Applications 

Andrew A. Wereszczak and Mattison K. Ferber 

Ceramic Science and Technology Group 

Oak Ridge National Laboratory 

P.O. Box 2008, MS: 6068 

Oak Ridge, TN 37831-6068 

(865) 576-1169; fax: (865) 574-6098; e-mail: wereszczakaa@ornl.gov 

DOE Technology Development Manager: Jerry L. Gibbs 

 (202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@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 

• Integrate new technologies into the transportation sector by implementing research that improves material 

properties and performance while maintaining acceptable life-cycle costs. 

• Promote commercialization of new materials technologies by developing standard testing and characterization 

methods with domestic and international standards communities. 

Approach 

• Define and implement research in conjunction with International Energy Agency’s (IEA) Implementing 

Agreement Advanced Materials For Transportation (IA-AMT) Applications. 

• Promote candidate annex on coating performance and evaluation and assess alternative coating test methods. 

Accomplishments 

• Defended proposed new annex on Coating Performance and Evaluation to IA-AMT Executive Committee. 

• Assessed laser spallation and large-scale deflection test methods to characterize coating spallation and bonding 

adherence. 

Future Direction 

• Project concluded in FY06. 

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Heavy Vehicle Propulsion Materials 

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 new 

material technologies are required to increase 

efficiency and reduce harmful emissions in these 

systems. Examples of these technologies include 

lightweighting to improve fuel efficiency; surface 

engineering to improve the resistance to wear and 

contact damage; development of durable coating 

systems for thermal, wear, and environmental 

management; and development of revolutionary 

materials (structural ceramics and ceramic matrix 

composites) for operation at much higher 

temperatures and pressures. 

In FY06, the primary effort of this project was 

to explore and develop innovative test methods to 

characterize coat performance. Coatings are 

technologically important to the transportation 

sector for applications needing improved 

tribological, thermal, and environmental protection 

of metallic (substrate) components. However, for 

coatings to enhance those enhancements, they need 

to be mechanically robust and remain adherent to the 

metallic component. 

Approach 

Three methods were considered this year to test 

coating adherence and spallation resistance. 

The first method involved using a laser-induced 

shock to cause coating spallation, and collaboration 

with Laser Shock Peening Technologies, Inc. 

(Dublin, OH). A schematic of the system is shown 

in Fig. 1. A high-energy laser is pulsed onto an 

ablative tape that is adhered to a metal substrate. 

The coating under evaluation is positioned on the 

opposite end to where the laser will strike. Water 

floods the impact side, is transparent to the laser, 

and serves as an impedance barrier. The laser goes 

through the water, impacts the ablative tape, and 

creates a plasma that lasts only up to a couple tens of 

nanoseconds. The water’s impedance causes the 

resulting shock wave to transmit into the substrate. 

The wave transmits into the substrate as a 

compressive wave. It continues into the specimen 

and reflects off the back surface becoming a tensile 

wave. The created tensile stress causes a portion of 

FY 2006 Progress Report 

the coating to spall off. The energy of the laser, 

dimensions and properties of the substrate, and the 

produced spall are all known or are measurable, so 

there is the potential to calculate the spall strength of 

the coating using this technique. 

Figure 1. Schematic of the laser shock test that 

was very effective at inducing coating spallation. 

The second method involved using a pendulumbased scratch tester. A picture of the system is 

shown in Fig. 2. A gas gun pushes on the pendulum 

horizontally. That event coupled with the 

pendulum’s radius causes high speed scratching with 

a conical diamond tool that is positioned at the 

bottom of the pendulum. Two high frequency load 

cells, positioned to measure normal and tangential 

forces, continuously capture load as a function of 

time. Those signals are then combined with 

postmortem of the produced scratch geometry to 

interpret scratch resistance. This system has been 

primarily used with monolithic materials; however, 

interest existed to assess its viability for 

characterization coating adhesion properties. 

The third considered method involved 

subjecting a coated metallic substrate to ball-on-ring 

equibiaxial flexure. Essentially, this is an 

axisymmetric three-point-bend test. Examples of 

tested coupons and a side-view schematic are shown 

in Fig. 3. The coated side of the coupon is loaded 

against with a ball using a mechanical testing frame. 

The coated specimen is deflected by several times 

the thickness of the specimen causing (at least) 

permanent deformation of the metallic substrate. 

During that deflection, the coating will spall. Its 

diameter is then linked to the radius of curvature of 

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Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

the permanently deformed specimen. That radius of 

curvature can then be used to calculate the imposed 

strain at that location, or effectively, the coating 

strain to failure or spallation. If spallation is more 

apt to occur in tension or compression, then the 

coating can simply be placed on the top or bottom to 

promote. This technique relies on the deformation 

of the substrate to impose a strain on the coating, 

and then a critical strain of spallation or cracking 

can be estimated. 

Figure 2. High-speed pendulum scratch tester. The 

horizontal piston shown at the top strikes the 

vertical steel bar resulting in a pendulum motion at 

its bottom where a conical diamond tool is mounted. 

Figure 3. Examples of coated specimens tested in 

ball-on-ring biaxial flexure (top). Sectional view 

of ball-on-ring biaxial flexure test configuration. 

Results 

The laser shock spallation test method was quite 

effective at producing coating spallation. An 

example of a test is shown in Fig. 4. Spall diameter is 

the dependent parameter, whereas laser energy or 

pressure, pressure waveform, substrate and coating 

physical properties and dimensions, and coating 

strength are independent parameters. Modeling of the 

shock event is needed to link those independent 

parameters to the produced spall diameter so that 

coating strength can be back calculated out. 

Polished cross-sections of the spall were 

examined with scanning electron microscopy (SEM). 

In one example (Fig. 5), the laser shock caused 

spallation at the interface between a bond coat and 

topcoat in a thermal barrier coating. Because the 

laser pressure could be controlled and varied, 

spallation was able to be initiated at different 

locations within the substrate - interface - bond coat - 

interface - top coat system. This is attractive because 

it enables the adherence characterization of the 

discrete layers within this multilaminate coating 

system (and presumably other multilaminate systems 

too). 

Figure 4. Examples of coating spallation 

produced by laser shock test system. Diameter of 

this coated coupon was 1 inch (25.4 mm). 

Sponsorship of the project terminated before 

scratch adhesion testing of coatings could be 

interrogated using the pendulum tester; however, 

strong potential for it to be an effective method for 

such exists. A schematic of the produced groove 

from the scratch testing is shown in Fig. 6. The 

height of the swinging pendulum can be adjusted to 

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Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

produce different depths of plowing or adjusted to 

best match the depth corresponding to coating 

thickness. Work is ongoing in other projects to 

continue to assess the viability of using pendulumbased scratch testing to characterize coating 

mechanical properties. 

Strain to failure is estimated using the spall size 

(or the location where cracking initiates if no 

spallation has occurred), radial deflection profile 

(measured with a profilometer), and finite element 

analysis (FEA). In one example (Fig. 7), the coating 

spall diameter was approximately 6 mm; that value Figure 7. Stress profile predicted by FEA that is on the predicted FEA curve was found to correspond used to estimate strain-to-failure. to a strain to failure of ~4%. This method is a simple 

and inexpensive test to conduct to estimate coating 

failure strains. Conclusions 

Three innovated test methods were considered 

for use to characterize coating strain to failure. The 

demonstration of two of the techniques (laser 

spallation and biaxial deflection) is very effective at 

enabling coating adherence characterization. The 

third technique (scratch adhesion testing) has 

potential to also be an effect test in characterizing 

coating adherence. Coatings (tribological, 

environmental, and thermal) can enhance or enable 

the performance of transportation-related 

components, so developing techniques to understand 

their adhesion characteristics is very relevant. 

Presentations and Publications 

M. K. Ferber, “Coating Evaluation: New 

Annex,” IEA IA-AMT Executive Committee 

Meeting, Porto, Portugal, 11 October 2005. 

A. A. Wereszczak, “Coating Evaluation & 

Performance: New Annex?” IEA IA-AMT 

Executive Committee Meeting, Calgary, Canada, 

07 May 2006. 

Figure 5. Polished cross-section of spall. 

Figure 6. Schematic of groove produced by 

pendulum scratch tester. 

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Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

B. 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 Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@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 

Objectives 

• Identify and study surface and sub-surface flaws in ceramic balls. 

• Summarize rolling contact fatigue (RCF) test methods and result interpretations used in Germany, Japan, 

United Kingdom, and the United States. 

Approach 

• Develop a robust test method that enables the exploitation and study of surface and sub-surface flaws in 

ceramic balls used in bearings. 

Accomplishments 

• Developed “c-sphere” strength (i.e., hoop tensile strength) test method to characterize ceramic balls. 

• Measured and compared c-sphere strength of three commercial grades of silicon nitride ball bearings. 

• Published ORNL Technical Report summarizing internationally used RCF test methods. 

Future Direction 

• Project concluded in FY06. 

147
Heavy Vehicle Propulsion Materials 

Introduction 

Silicon nitride (Si3N4) balls for use in hybrid 

bearings have been shown to extend rolling contact 

fatigue (RCF) lifetime, reduce the dynamic loading 

in high-speed roller element applications because of 

their lower density, compared to steel balls, and 

enable bearing use in more corrosive and higher 

temperature environments. During their 

development, many parameters such as density, 

hardness, elastic modulus, fracture toughness, 3- and 

4-pt flexure strength and Weibull modulus, crushing 

strength, grain size and secondary phase 

morphology, porosity, and surface finish have been 

extensively studied in an effort to improve or predict 

RCF performance of Si3N4 balls. As an outcome of 

many of those studies, ASTM F2094 was developed 

to specify allowable minimums and maximums of 

many of those properties and finish conditions. 

In spite of the above work, little or no study has 

been devoted to the quantification of sub-surface 

inhomogeneities and what effect they could have on 

RCF performance. Sub-surface factors such as 

damage due to machining and remnants from hotisostatic-pressing glass encapsulation that are 

sometimes used in processing are examples of this. 

Such sub-surface inhomogeneities usually play a 

dominant role in the performance (strength, fatigue, 

etc.) of structural ceramics. Their study in structural 

ceramics is typically facilitated using simple tensile 

or flexure strength specimen geometries that enable 

the concurrent measurement of strength and 

strength-limiting flaw size (latter identified through 

fractographical practices). 

However, for ceramic spheres, two historical 

barriers to the study of strength-flaw-size 

relationships within them are the (1) the inherent 

geometry of the sphere itself, and (2) sphere size 

which inhibits the size of extractable strength 

specimens (e.g., bend bars). Regarding the former, 

it is nearly impossible to mechanically load a 

ceramic ball in a manner to produce a predictable 

tensile stress field at the sphere’s surface sufficient 

to cause its fracture. Compressively loading a 

ceramic sphere diametrally produces two Hertzian 

contacts where ring- and cone-crack initiation 

occurs, and ultimate fracture can result if the contact 

stresses are sufficiently high. Hence, the evolution 

of ultimate fracture with this test method makes the 

study of strength limitation from surface- and subsurface-located intrinsic and extrinsic flaws nearly 

FY 2006 Progress Report 

impossible. The fracture of brittle spheres from 

diametral compression occurs at such high loads 

(i.e., a great deal of stored energy exists in the 

sphere at fracture) that the test coupon disintegrates 

into fine rubble rendering fractography nearly 

impossible, and thus, inhibiting the study of the 

relationship between strength and flaw size. Other 

complicating issues can result if one opts to crush 

small ceramic spheres; for example, exorbitantly 

high contact stresses are created causing the 

(typically very strong) sphere, in effect, to act as a 

spherical indenter against the loading platens before 

it is crushed. 

Because the desire exists to identify and study 

surface-located flaws and sub-surface damage in 

ceramic spheres, and crush testing was deemed to be 

impractical, a new test coupon called the “C-sphere” 

flexure strength specimen was conceived and 

developed to enable the study of and measurement 

of a strength and associated linked flaw size. The 

C-sphere flexure strength specimen fractures at low 

to modest loads, and fractographical analysis can 

readily be performed and strength-limiting-flaw 

studies can ensue in a practical manner because 

excessive fragmentation is not produced. Enabling 

the identification of a flaw type and measurement of 

its size in finished ceramic balls is indeed important 

for the study of strength, but more importantly, for 

the study and predictability of RCF performance (a 

response limited by surface- or near-surface-located 

flaws or in changes thereof). Hadfield reported that 

a critical depth of ring cracking in RCF is 5-20 

microns below the surface. The C-sphere specimen 

facilitates the study of flaws that can exist at those 

depths in ceramic ball bearings and ceramic spheres 

in general. 

Approach 

A schematic of the C-sphere specimen is shown 

in Fig. 1. It is produced through the controlled 

slotting of a ceramic sphere. It is then diametrally 

loaded perpendicular to that slot (or flexed) to 

initiate hoop tensile fracture at the sphere’s surface 

or “outer fiber” as shown in Fig. 2. This specimen is 

analogous to a “C-ring” flexure specimen, which is 

produced through controlled slotting of a ceramic 

ring and enables strength evaluations of tubular 

components. 

148 
Heavy Vehicle Propulsion Materials 

Figure 1. Side and front views of the C-sphere 

flexure strength specimen. 

FY 2006 Progress Report 

depth (related to this offset) in the FEA model was 

varied until a slot-width and slot-depth combination 

produced at least a 10-times higher outer-fiber 

surface or hoop tensile stress than any other surface 

tensile stress located elsewhere on the coupon. The 

final offset size (0.635 mm) resulted in a C-sphere 

geometry that, when diametrally compressed, 

produced an outer-fiber or hoop tensile stress much 

larger than that on the ground interior and a 

(desirably) high likelihood that fracture would occur 

at the outer fiber. 

The nodal stress distribution plots for the 

C-sphere model are shown in Fig. 3 for an applied 

compressive force of 200N. One observes from 

these plots how the maximum tensile stress of 

132.1 MPa is located at the outer fiber area of the 

C-sphere. This model shown was used to compute 

the effective area and effective volume for the 

C-sphere specimen, and those results are shown in 

Fig. 4(a). 

Area and volume efficiencies were computed by 

dividing the effective area and volume by the 

original sphere’s total area and volume. Their 

dependencies as a function of Weibull modulus are 

shown in Fig. 4(b). The use of the concepts of area 

and volume efficiencies is useful to consider when 

evaluating how efficient a particular specimen is in 

sampling surface-type and volume-type strengthlimiting flaws. If a ceramic has a Weibull modulus 

of 10, then the C-sphere geometry described herein 

would sample flaw populations on about 3% of the a 

sphere’s surface and in about 0.1% of its volume. 

As a comparison to an established test configuration, 

the area efficiency is about 24% for an ASTM 

C1161B flexure configuration with a 3 x 4 x 50 mm 

bend bar of a ceramic with a Weibull modulus of 10. 

Figure 2. Diametral compression of the C-sphere 

flexure specimen causes fracture initiation from a 

hoop stress at the outer fiber. 

An important aspect of the geometry is the 

purposeful non-coincidal positioning of the sphere’s 

center and the center of curvature of the slot’s 

innermost depth. The introduction of this offset acts 

to increase the outer-fiber tensile stress through the 

introduction of an applied bending moment; if it 

were not introduced, then too much tensile stress 

could develop on the ground surface of the notch 

and fracture would undesirably initiate there. Slot 

While the C-sphere’s area efficiency is modestly 

good, its volume efficiency is not. However, its 

design was never intended to sample volume flaw 

populations in ceramic spheres, so its poor volume 

efficiency is not viewed as a limitation. 

The inherent geometry of the C-sphere produces 

a desirable artifact associated with the stress state at 

the edges that intersect the original sphere surface 

and the ground surfaces from the slotting procedure. 

Additional machining steps (i.e., edge chamfering 

after C-sphere slotting) are not necessary because 

significant tensile stresses never develop at the 

edges during diametral compression. Therefore, 

rudimentary chipping that can potentially result at 

149 
Heavy Vehicle Propulsion Materials 

these edges from the slotting procedure will not 

initiate fracture and not influence measured strength. 

Figure 3. Nodal first principal stress distribution at 

three different views for an applied compressive 

force of 200 N. Shown values have units of MPa. 

Figure 4. (a) Effective area and effective volume 

and (b) area and volume efficiency as a function of 

Weibull modulus for the C-sphere flexure strength 

specimen. The latter is the effective area and 

volume divided by the sphere’s total surface area 

and volume, respectively. 

FY 2006 Progress Report 

NBD200, SN101C, and TSN-03NH balls with 

12.7 mm diameter were machined into C-sphere 

flexure strength specimens. Examples of machined 

C-spheres are shown in Fig. 5. 

After ligament thickness was measured in all Csphere specimens, they were monotonically loaded 

in compression to failure using an electromechanical 

universal testing machine. A special alignment jig 

was used to horizontally align the C-sphere slot prior 

to loading (see Fig. 6). A slight compressive 

preload (~ 25 N) was applied to the jigged C-sphere 

specimen enabling the jig to be pulled out without 

disturbing the C-sphere alignment. The specimens 

were then tested using a constant crosshead 

displacement rate. Load to fracture was recorded 

and used to determine C-sphere flexure strength 

using interpretations of Fig. 3. Approximately 30 

C-sphere specimens each of NBD200, SN101C and 

TSN-03NH were tested. Weibull strength 

distributions were determined using commercially 

available software. Optical fractography was 

conducted on all specimens to identify failure 

location and the fracture surfaces of a select few 

specimens were examined with scanning electron 

microscopy. 

Figure 5. Machined Si3N4 C-sphere flexure strength 

specimens. 

Results 

The C-sphere characteristic strength of SN101C 

was 27% and 17% larger than the NBD200 and 

TSN-03NH, respectively. The Weibull moduli of 

SN101C and NBD200 were equivalent, whereas that 

of the TSN-03NH was slightly larger. The 95% 

confidence ratio rings for the two-parameter Weibull 

distributions of each are illustrated in Fig. 7. 

150 
Heavy Vehicle Propulsion Materials 

Figure 6. A plastic jig was used to align the 

C-sphere specimens prior to their diametral 

compressive loading. 

Figure 7. 95% confidence ratio rings for the three 

tested silicon nitrides. 

Optical fractography showed that failure always 

initiated at the surface but not necessarily always 

exactly at the apex of the outer fiber. Failureinitiation locations for all the tested C-sphere 

specimens are mapped in Fig. 8. The scatter in the 

failure locations matches the region where there is 

tensile stress on the C-sphere’s surface as shown in 

Fig. 3. No failures initiated at the edges, which is 

consistent with there being little or no stress in that 

region during loading. 

A representative example of a failure location is 

shown in Fig. 9. The fracture of all specimens 

initiated at the C-sphere surface, which is an 

occurrence that was sought with the design of the 

C-sphere flexure specimen. For a strength of 

770 MPa, a fracture toughness of 6 MPa√m, and a 

crack geometry factor of 1.5, the estimated Griffith 

flaw size for the specimen shown in Fig. 9 would be 

FY 2006 Progress Report 

approximately 30 microns. That is of the size scale 

that Hadfield had indicated to be a critical size for 

crack penetration during RCF. 

Conclusions 

A “C-sphere” flexure strength specimen 

geometry was developed to measure the hoop tensile 

strength of bearing-grade silicon nitride balls and 

enable the study of surface-located strength-limiting 

flaws in these ceramic spheres. The C-sphere has 

the advantage of possessing a simple geometry that 

is easy to test. A simple, monotonically increasing 

uniaxial compressive force produces a hoop tensile 

stress at the C-sphere’s outer surface that ultimately 

initiates fracture. The C-sphere strength is 

determined using the combination of measured 

failure load, geometry, and finite element analysis. 

Representations of C-sphere effective area, effective 

volumes, area efficiency, and volume efficiency as a 

function of Weibull modulus were developed. 

C-sphere flexure strength distributions were 

determined for NBD200, SN101C, and TSN-03NH 

(all commercial grade Si3N4 balls). SN101C was 

found to have the highest characteristic strength and 

TSN-03NH the highest Weibull modulus. The 

C-sphere test configuration was successful 100% of 

the time in causing fracture initiation at the surface, 

and therefore is effective at enabling the study of 

surface-located strength-limiting flaws. 

Figure 8. Map of locations of failure initiation for 

all tested specimens. Data projected onto a onequarter domain of a C-sphere. 

151 
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

ORNL/TM-2006/521 Report: “Rolling Contact 

Fatigue of Ceramics,” A. A. Wereszczak, W. Wang, 

Y. Wang, M. Hadfield, W. Kanematsu, 

T. P. Kirkland, and O. M. Jadaan, 2006. 

A. A. Wereszczak, W. Wang, O. M. Jadaan, and 

M. J. Lance, “Strength of a C-Sphere Flexure 

Specimen,” Ceramic Engineering and Science 

Proceedings, 2006. 

Figure 9. Example of a surface-located strengthlimiting flaw in a SN101C C-sphere flexure strength 

specimen. This specimen had a strength of 

770 MPa. 

Presentations and Publications 

A. A. Wereszczak, “IEA Annex III: Rolling 

Contact Fatigue,” IEA IA-AMT Executive 

Committee Meeting, Porto, Portugal, 11 October 

2005. 

A. A. Wereszczak, “Strength of a C-Sphere 

Flexure Specimen,” 30th Annual Conference on 

Composites and Advanced Ceramics, Cocoa Beach, 

FL, Jan. 24, 2006. 

A. A. Wereszczak, “IEA Annex III: Rolling 

Contact Fatigue,” IEA IA-AMT Executive 

Committee Meeting, Calgary, Canada, 07 May 

2006. A. A. Wereszczak, O. M. Jadaan, and 

T. P. Kirkland, “Hoop Tensile Strength of Ceramic 

Spheres Using a Diametrally Compressed 

‘C-Sphere’ Specimen,” in review, Journal of the 

American Ceramic Society, 2006. 

A. A. Wereszczak, O. M. Jadaan, H. -T. Lin, 

G. J. Champoux, and D. P. Ryan, “Hoop Tensile 

Strength Testing of Small Diameter Ceramic 

Particles,” in press, Journal of Nuclear Materials, 

2006. 

152 
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

8. EMERGING TECHNOLOGIES 

A. Processing and Characterization of Structural and Functional Materials for 

Heavy-Vehicle Applications 

J. Sankar, Z. Xu, and S. Yarmolenko 

North Carolina A &T State University 

1601 East Market Street 

Greensboro, NC 27411 

(336) 256-1151, extension 2282; fax: (336) 256-1153; e-mail: sankar@ncat.edu 

DOE Technology Development Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

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

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Subcontractor: North Carolina A&T State University, Greensboro, North Carolina 

Objectives 

• Produce doped strontium cerate thin film electrolytes with flame-assisted chemical vapor deposition (FA-CVD) for 

solid oxide fuel cells (SOFCs). 

• Study the effect of the processing parameters on the porosity of the NiO/YSZ anode substrates. 

Approach 

• Study the experimental parameters of the FA-CVD process and deposit yttria-doped strontium cerate thin film 

electrolytes thin film coatings using this technique. 

• Characterize the coatings using scanning electron microscopy (SEM) and X-ray diffraction (XRD). 

• Study the composition of raw materials, pressing, and sintering parameters for the anode substrate fabrication. 

• Use optimization technique to study the properties of the anode materials. 

Accomplishments 

• Obtained dense and thin yttria-doped strontium cerate thin film electrolyte coatings. 

• A design matrix was developed and successfully used to develop the anode material for SOFC application. 

• The effect of the processing variables and their levels on porosity and conductivity of the anodes were 

identified. 

Future Direction 

• Optimize the FA-CVD process to prepare protonic conductive electrolyte with designed stoichiometry. 

• Characterize the electrical conductivity of the protonic conductive electrolytes. 

• Standardize the fabrication process of the anode materials. 

• Fabricate single SOFCs and conduct performance tests. 

153 
Heavy Vehicle Propulsion Materials 

Introduction 

Proton-conductive perovskite materials have 

been receiving more and more interests since it was, 

for the first time, reported by H. Iwahara et al. over 

twenty years ago [1]. These materials have already 

found applications in fuel cells [2,3], hydrogen 

separation [4-6], and hydrogen sensors [7]. State-ofthe-art solid oxide fuel cells (SOFCs), one of the 

most well known types of fuel cells, use solid state 

oxygen ion conductive electrolytes, which permit 

only oxygen ions to transfer through it from the 

cathode to anode. When proton-conductive materials 

are used as electrolyte for SOFCs, it permits only 

hydrogen ions, i.e. protons, to migrate from the 

anode to the cathode. Proton-conductive perovskite 

electrolytes have many advantages over oxygen ion 

conductor electrolytes. When a proton conductor is 

used instead of an oxygen ion conductor, fuel 

circulation is unnecessary, because no water vapor is 

generated at the hydrogen electrode. Moreover, the 

operating temperature can be reduced to about 600­

700°C, where reactions such as methane reformation 

are kinetically favorable without the need for costly 

catalysts. Therefore, fuel cells could be constructed 

from less expensive materials. Recently, acceptordoped ABO3 perovskite oxides (A2+=Sr, Ba, and 

B4+=Ce, Zr) which exhibit proton conduction, have 

received much attention as solid oxide electrolytes 

of the hydrogen–air fuel cells. Among many 

different proton-conductive perovskites, Barium 

cerates (BaCexM1-xO3-δ) and strontium cerates 

(SrCexM1-xO3-δ) doped with acceptors M (trivalent 

element, such as, Y or rare earth elements, Yb, Gd, 

Nd, Dy, etc.) found to have the highest 

conductivities. 

Up-to-date, almost all the research efforts 

towards proton-conductive perovskites are focused 

on materials processing, microstructure, phase and 

electrochemical properties characterization [8-10]. 

In addition to widely used powder sintering methods 

to prepare electrolytes [11], thin film technologies 

are also favorable in applying proton-conductive 

perovskites as electrolytes in SOFCs. Thin films 

were essentially prepared with sol-gel techniques 

[12, 13], screen printing [14], ultrasonically 

atomized spray of powder colloid [15] and RFmagnetron sputtering techniques [16]. In this report, 

preparation of yttria-doped strontium cerate thin 

films with flame-assisted chemical vapor deposition 

(FA-CVD) technique will be presented for the first 

FY 2006 Progress Report 

time. FA-CVD technique had been used to process 

yttrium-stabilized zirconia thin films on both dense 

and porous substrates [17, 18], ceria [19] and 

alumina [20], and so on. It is a cost-effective 

technique to prepare metal oxides. It is operated in 

ambient, so that it is made possible to coat on large 

area. It is especially important to scaling up this 

technique for industrial applications. In the present 

work, we demonstrate again that the FA-CVD 

technique is effective in preparing thin films of 

perovskite materials. The purpose of the study was 

to understand the effects of the operating conditions 

on the microstructure, phase and deposition rate in 

order to optimize the process. 

The two most common designs of SOFCs are 

the tubular and the planar [21]. Planar cells are 

capable of generating high power densities in 

comparison with the tubular cells. Anode supported 

planar cells provide very high cycle efficiency and 

pollution free energy. However, the microstructure 

of the cell degrades with respect to time due to the 

high temperature operating conditions. This leads 

to the deterioration of the fuel cell [22] and 

subsequent decomposition. This problem causes the 

major concern towards the commercialization of 

the fuel cells. The current study was to improve the 

life and performance of the anode cell for the 

SOFCs in a long run. In the earlier studies, Muller 

et al. [23] and Lee et al. [24] did extensive studies 

on the porosity and the electrical measurements for 

anode material with varying proportions of size and 

mol. % of the raw materials and the sintering 

temperature. Takehisa et al. [25] have used varying 

sizes of NiO-YSZ composite particles to fabricate a 

porous anode structure. Clemmer et al. [26] have 

observed dramatic improvements in electrical 

conductivity with increase in Ni content. Lee et al. 

[27] have investigated the effect of microstructure 

on anode substrate on the unit cell performance. 

However, the earlier investigators have studied the 

effect of individual process variables on the 

performance of the anode material. It is imperative 

to consider all the contributing factors that affect 

the anode material performance. So, it was decided 

to use the statistical method to study the individual 

effects and the combined effects of the different 

processing parameters on the output variable. 

154 
Heavy Vehicle Propulsion Materials 

Approach 

A liquid fuel FA-CVD system similar to the one 

introduced in previous work [17] was used for 

strontium cerate thin film processing. The metalorganic reagents used in all the experiments were 

strontium 2-ethylhexanoate (Sr-2EH, 40%, Strem 

Chemicals, Inc.), Cerium 2-ethylhexanoate (Ce2EH, 49%, Alfa Aesar) and yttrium 2­

ethylhexanoate (Y-2EH, 99.8%, Alfa Aesar). High 

purity oxygen (99.999%, National Welders) was 

used as oxidant. In order to vary the substrate 

temperature in a relatively large range, air-blow 

cooling and oxy-acetylene flame heating under the 

substrate support were developed. They were used 

alternatively according to the requirement of the 

substrate temperature. The ranges and values of the 

parameters employed in the experiments were 

designed as shown in Table 1. Si(100) was used as 

substrates. 

To investigate the effect of the total-metal 

concentration and substrate temperature on the 

microstructure and phase development in the 

coatings, a full arrangement of 18 experiments for 

the two parameters with three levels was designed 

with a duplicate. The experiments were also 

randomized in order to eliminate the temporal effect 

on the results. 

Table 1. Experimental parameter design 

Parameters Ranges and values 

Total flow rate of solution 

(cm3

/min) 2.0 

Flow rate of oxygen (cm3/min) 1600 

Substrate temperature (°C) 900, 1050, 1200 

Total-metal concentration (M) 2×10-3, 4×10-3, 6×10-3 

Atomizer-to-substrate distance 

(mm) 51 

Deposition time (min) 20 

The microstructures of the deposited films and 

composition of the coatings were characterized with 

SEM (Hitachi 3000N) and EDS (Oxford 

Instruments). The phases of the sintered materials 

were determined with X-ray diffraction (XRD, D8 

Discover, Bruker AXS, Germany) with Cu Kα1 

radiation at λ=1.5406 Å. 

The anode material was prepared by mixing the 

raw materials (graphite, NiO and YSZ) in different 

proportions, compacted uniaxially, sintered in air 

and finally reduced in H2/Ar atmosphere. The 

FY 2006 Progress Report 

graphite powder ratios and sintering temperatures 

used for the fabrication is shown in the Table 2. 

The NiO/YSZ ratio was 57.5/42.5 and the 

compaction pressure was 150 MPa. The sintered 

samples were reduced at 1000 °C. The different 

powder ratios, compaction pressure, the sintering 

temperatures and the reduction temperature were 

decided based on the literature. The different 

powders used were mixed together using ball 

milling technique. The die used for powder 

compaction was half-inch steel die. The measured 

output variables were porosity and the electrical 

conductivity. 

Table 2. Composition and process parameters used 

for the experimental design 

Independent 

Treatment level 

variable 0 

Low 

1 

Medium 

2 

High 

Graphite volume 

fraction (vol %) 30 35 40 

Sintering temp 

(°C) 1200 1250 1350 

Results 

Phase Characterization of the Cerate Thin Films 

Fig. 1 shows the XRD spectra of the coatings 

of experiments SCY-1, 3, 5, 7, 9, 11, 13, 15, and 17 

with the characteristic peaks of SrCe0.85Y0.15O2.925 

(Pattern 01-083-1157) and CeO2 (Pattern 01-081­

0792) at the bottom for reference. The diffraction 

spectra of the three coatings of the experiments 

SCY-1, 3, and 5 show a tendency of increased 

intensity of the perovskite characteristic peaks at 2θ

= 29.728o with the increased substrate temperature 

at the lowest total-metal concentration of 2.0×10-3 

M. At the total-metal concentration of level 1 and 2 

(4.0×10-3M and 6.0×10-3 M), the highest intensities 

of the perovskite peaks were found from the 

coatings deposited at the intermediate substrate 

temperature, i.e., 1050oC. All the coatings show 

(131) and (122) peaks at about 2θ = 37.78o and 2θ

= 39.47o (not labeled in the graph). Comparing all 

the spectra, it seems that the coating of SCY-15 

deposited at 6.0×10-3 M and 1050oC represent the 

best result of the perovskite phases. However, the 

characteristic peaks of fluorite CeO2 phases can be 

155 
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

SCY-17 

SCY-15 

SCY-13 

SCY-11 

SCY-9 

SCY-7 

SCY-5 

SCY-3 

SCY-1 

Intensity (arb. unit) 10 15 20 25 30 35 40 45 50 55 60 

Yttria-doped Strontium Cerate 

CeO2 

2 θ

Figure 1 XRD spectra of yttria-doped strontium 

cerate coatings. 

noticed in all the coatings. It shows the possibility of 

CeO2 phase precipitation in the coatings caused by 

nonstoichiometry between Sr and Ce. Examination 

of the elemental ratios between Sr and Ce for the 

coatings SCY-2, 4 and 6 with EDS demonstrated the 

deficiencies of Sr. The atomic ratios of Sr to Ce 

were about 0.8. This implies that to achieve the 

stoichiometry of the SrCe0.85Y0.15O2.925, more starting 

chemical of Sr is needed in the solution. High 

temperature sintering of BaCe0.85Y0.15O3-δ had 

experienced Ba evaporation and CeO2 precipitation 

[28]. 

Microstructure of the Cerate Thin Films 

 The surface morphologies of all the coatings were 

examined with SEM. These micrographs shown in 

Figure 2 depict that the intermediate substrate 

temperature of 1050oC produced smoothest surfaces. 

The grain sizes are not proportional to the substrate 

temperature. Both 900 and 1200oC substrate 

temperatures produced large grains, rough surfaces 

consequently, while the substrate temperature of 

1050oC produced smallest grains. It is noteworthy to 

discern that the combination of substrate 

temperature of 1050 oC and the total-metal 

concentration of 4.0×10-3M generated the best 

quality coating in terms of morphology uniformity 

and surface smoothness. Cracks were apparently 

perceived in all the coatings, which demonstrate that 

there is a large difference in the coefficients of 

thermal expansion between the coating and 

substrate. The coefficient of thermal expansion for 

the coatings needs to be determined by further 

experiments. 

Porosity and Conductivity of the Anodes 

The graphite was used to obtain the desired 

porosity. The porosity calculations were based on 

the weigh and size measurements of the sample 

followed by the actual and the theoretical volume 

calculations. The four-point probe method was used 

for the conductivity measurements. The porosity 

and the conductivity results are shown in the Table 

3. For each sample, three replicates were used and 

the results shown were averaged to one value. 

The results show that the sintering temperature was 

the major variable that controls the porosity and the 

conductivity of the cell. Graphite content also had a 

role in controlling the porosity but it mostly 

depends on the sintering temperature values. Both 

graphite and the sintering temperature almost had a 

linear effect on the porosity and the conductivity 

values (Fig 3). For the samples sintered at 1200°C, 

with 30 vol. % of graphite content, the porosity 

value was 56.06 and for 40 vol. % of graphite 

content, the porosity was 58.96. But the porosity 

and the conductivity values had an inverse effect 

for the given sintering temperature. The reason 

was, the graphite oxidizes at lower temperature (~ 

1000°C) compared to the maximum temperatures 

used and it directly leads to the formation of pores. 

So the difference in the graphite vol. % had some 

effect on the pore volume at the lower temperature 

sintering (i.e., at 1200°C). But at 1350 °C, the 

difference in the graphite volume did not make 

much of a difference in the porosity volume. This 

was because of the fact that at higher sintering 

temperature, the particle growth and particle fusion 

started taking place and it leads to reduced porosity 

and improved inter-particle connectivity. This 

replicates on the conductivity of the cell. At lower 

sintering temperatures, the inter particle 

connectivity was poor and that leads to poor 

electrical charge carrying paths. Whereas, at higher 

sintering temperature, the conductivity increases 

drastically due to the fact that it has more electron 

conducting path. Thus, the conductivity and the 

porosity values obtained for the samples sintered at 

1350 °C were more suitable for the SOFC 

application. 

156 
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

(a) (b) (c) 

(d) (e) (f) 

(g) (h) (i) 

1 μm 1 μm 1 μm 

2 μm 2 μm 2 μm 

5 μm 2 μm 5 μm 

Figure 2. Micrographs of the SrCe0.85Y0.15O2.925 coatings deposited at different conditions. The total-metal 

concentrations for the coatings shown in the first, second and the third rows are 2.0×10-3, 4.0×10-3 and 

6.0×10-3 M, respectively. Substrate temperatures for the coatings shown in the first, second and the third 

columns are 900, 1050 and 1200oC, respectively. 

Table 3. Porosity and Conductivity results for 

different processing conditions 

Sintering 

Temp / °C 

Graphite 

Vol % 

Porosity 

Vol % 

Conductivity 

Ω-1 cm-1 

1200 

30 56.06 69.59 

35 56.85 69.27 

40 58.96 58.53 

1250 

30 52.53 147.99 

35 55.41 94.39 

40 57.66 81.70 

1350 

30 41.35 744.87 

35 37.08 686.19 

40 41.01 733.46 

Conclusions 

Yttria-doped strontium cerate coatings were 

developed with FA-CVD technique. XRD 

characterization revealed combined phase of the 

perovskite and fluorite in the coatings. EDS study 

found strontium depletion in the coatings as 

compared with the composition in the solution. The 

studies of the micrographs of the surface of the 

coatings also uncovered that both intermediate 

substrate temperature and total-metal concentration 

produced coating with smoothest surface. 

The design matrix was developed and 

successfully used to develop the anode material for 

SOFC application. The results showed that the 

processing variables and their levels had a strong 

impact on the output variables. The porosity values 

had a controlling effect on the conductivity values. 

It indicates that the electrode performance can be 

optimized by controlling the porosity in order to 

achieve desired conductivity. The current work 

sparkle the idea of changing the ratios of NiO in the 

mixture to study the changes in conductivity values 

and also in the microstructure of the electrode. 

157 
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

Graphite vol. % Vs Porosity vol. % Graphite vol. % Vs Conductivity 

70.00 800.00

1200 °C 

1250 °C 

1350 °C 

1200 °C 

1250 °C 

1350 °C

0.00 20.00 

25 30 35 40 45 25 30 35 40 45 

Graphite (vol. %) Graphite (vol. %) 

(a) (b) 

Figure 3. (a) Comparison of porosity values for different sintering temperatures (b) Comparison of 

Conductivity values for different sintering temperatures 

Conductivity (

Ω-1 cm -1) 

60.00 600.00

Porosity (vol. %) 50.00 

400.00 

40.00 

200.00 

30.00 

References 

1. H. Iwahara, H. Uchida and S. Tanaka, Solid State 

Ionics, 9-10(1983) 1021-1025. 

2. W. G. Coors, J. Power Sources, 118(2003) 150­

156. 

3. S. Yamaguchi, H. Yugami, and S. Ikeda, J. 

Alloys and Compounds, 330-332(2002) 911-915. 

4. U. Balachandran, T. H Lee, S . Wang and S. E. 

Dorris, Inter. J. Hydrogen Energy 29 (2004) 291. 

5. U. Balachandran, B. Mar, P . S . Maiya, R. L. 

Mieville, J. T. Dusek, J. Picciole, J. Guan, S. E. 

Dorris and M. LIU, Solid State Ionics 108 (1998) 

363. 

6. R. V. Siriwardane, J . A. Poster, E. P . Fisher, T. 

H. Lee, S . E. Dorris and U. Balachandran, App. 

Surf. Sci. 167 (2000) 34. 

7. H. Iwahara, Y. Asakura, K. Katahira; M. Tanaka, 

Solid State Ionics, 168 (2004), 299-310. 

8. F. Chen, O. T. Sørensen, G. Menga

 and D. Peng, 

Solid State Ionics, 100(1997) 63-72. 

9. S. D. Flint, M. Hartmanová, J. S. Jones and R. C. 

T. Slade, Solid State Ionics, 86-88(1996) 679­

683. 

10. R. Peng, Y. Wu, L. Yang and Z. Mao, Solid 

State Ionics, 177(2006) 389-393. 

11. W. Grover Coors and Dennis W. Readey, J. 

Am. Ceram. Soc., 85 (2002) 2637-40. 

12. V. Agarwal, M. Liu, J. Mater Sci, 32 (1997) 

619-625. 

13. S. Higgins, N. Sammes and A. Smirnova, 

Proceedings of the 26th Risø International 

Symposium on Materials Science: Solid State 

Electrochemistry, Editors: S. Linderoth, A. 

Smith, N. Bonanos, A. Hagen, L. Mikkelsen, K. 

Kammer,D. Lybye, P.V. Hendriksen, F.W. 

Poulsen, M. Mogensen, W.G. Wang Risø 

National Laboratory, Roskilde, Denmark 2005 

14. G. Meng, P. Wang, Y.Gu and D. Peng, Solid 

State Ionics, 136-137(2000)209-213. 

15. T.H. Lee, S.E. Dorris and U. Balachandran, 

Solid State Ionics, 176( 2005) 1479-1484. 

16. T. He, C. L. Jia, P. Ehrhart and P. Meuffels, 

Solid State Ionics, 89(1996) 9-12. 

17. Z. Xu, J. Sankar and S. Yarmolenko, Surf. Coat. 

Technol. 177–178 (2004) 52-59. 

18. Z. Xu, S. Yarmolenko, J. Sankar, NATO 

Science Series, II: Mathematics, Physics and 

Chemistry (2005), 202 (Fuel Cell Technologies), 

49-57. 

19. W. B. Carter, G.W. Book, T.A. Polley, D.W. 

Stollberg, J,M. Hampikian, Thin Solid Films 

(1999), 347(1999), 25-30. 

20. B. Naoufal, Surf. Coat. Technol, 200(2006) 

4097-4103. 

21. S.C. Singhal, Solid State Ionics, 152-153(2002), 

405-410. 

22. M.J. Jorgensen, P. Holtappels, and C.C. Appel,. 

Journal of Applied Electrochemistry, 30(2000), 

411-418. 

23. A.C Muller, D. Herbstritt, and E.I. Tiffee, Solid 

State Ionics. 152-153(2002): 537-542. 

24. J.-H Lee, H. Moon, H.-W. Lee, J. Kim, J-D. 

Kim, and K-H. Yoon, Solid State Ionics, 

148(2002), 15-26. 

158 
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

25. F. Takehisa, K. Murata, S. Ohara, H. Abe, M. 

Naito, and K. Nogi, Journal of Power 

Sources, 125(2004), 17-21. 

26. R. M. C. Clemmer, S.F. Corbin, Solid State 

Ionics, 166(2004), 251-259 

27. J. H. Lee, J.W. Heo, D.S. Lee, J. Kim, G.H. 

Kim, H.W. Lee, H.S. Song, and J.M. Moon, 

Solid State Ionics, 158(2003), 225-232. 

28. J. Wu, L.P. Li, W.T.P. Espinosa, and S.M. 

Haile, J. Mater. Res., 19(2004) 2366-2376 

Publications and Presentations 

1. Zhigang Xu, Gregory Young, Gukan Rajaram, 

Jag Sankar, “Proton Conductive Strontium 

Cerate Thin Films Processed Using FlameAssisted Chemical Vapor Deposition”, 

IMECE2006-15022, Proceedings of 

IMECE2006 

2006 ASME International Mechanical 

Engineering Congress and Exposition 

November 5-10, 2006, Chicago, Illinois, USA 

2. G. Young, Z. Xu, G. Ramjram, J. Sankar, 

“Deposition of Yttrium Doped Strontium 

Cerate Thin Films by Combustion Chemical 

Vapor Deposition”, presented in the 30th 

International Conference and Exposition on 

Advanced Ceramics and Composites, Cocoa 

Beach, FL, January 22-27, 2006 

3. Zhigang Xu,

, Gukan Rajaram, Jag Sankar and 

Devdas Pai, “Electrophoretic deposition of 

YSZ electrolyte coatings for solid oxide fuel 

cells”, Surface and Coatings Technology, 

201(2006), 4484-4488. 

4. Zhigang Xu, Jag, Sankar, “ Preparation and 

Characterization of Nano-crystallite 

Electrolyte for SOFCs”, presented in the 2006 

MRS Fall Meetings, November 27 - 

December 1, 2006 

Boston, MA. 

159 
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

B. 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 Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

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

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

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

Subcontractor: Purdue University, West Lafayette, Indiana 

Objective 

• Develop manufacturing technologies for producing bulk nanostructured alloys with enhanced mechanical 

properties for transportation applications. 

Approach 

• Create nanostructured chips and particulate by machining alloys of aluminum, titanium, and iron under 

controlled machining conditions. 

• Develop machining methods for direct production of bulk forms (e.g, sheet, bar) with nano-scale 

microstructure. 

• Characterize microstructure, structure and mechanical properties of nanostructured chip particulate and bulk 

forms. 

• Develop low-temperature consolidation routes for converting the chip particulate into monolithic and composite 

materials without significant coarsening of the nanocrystalline microstructures. 

• Characterize microstructure, mechanical properties, and performance of the bulk forms after the consolidation. 

Accomplishments 

• Produced nanostructured chips of a variety of alloys and metals—including Al6061-T6, commercially pure 

titanium, stainless steel, steels, and Inconel—by machining. The nanostructured chips are up to three times as 

strong as the bulk material. The grain size in most cases is ~ 100 nm. 

• Developed a new process, large-strain extrusion-machining (LSEM), to directly produce bulk forms (e.g. sheet, 

bar) with ultra-fine grained (UFG) microstructure and enhanced mechanical properties. 

• Developed a new Modulation –Assisted Machining (MAM) method for making nanostructured particulate of 

controlled size (30-500 micrometers) and shape (e.g., equi-axed, platelet, needle) directly by chip formation. 

• Developed innovative, low-temperature processing routes for producing bulk nanostructured alloys from chip 

particulate using polymeric binders. These polymer-bonded alloys show promising enhancements in hardness 

and wear properties. 

161
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

• Developed low-temperature consolidation routes that limit coarsening in Al6061-T6. Achieved a high degree of 

densification and interparticle bonding in room temperature powder extrusion of nanostructured Al6061-T6. 

• Developed high temperature, high-strength nanostructured alloys, for example, Inconel 718, Waspaloy, with 

extraordinary thermal stability. The process involved a combination of nanostructure refinement by machining 

and creation of ultra-fine fine-scale precipitates in the nanostructured matrix by thermal ageing after the 

machining. 

Future Direction 

• Continue development of the extrusion machining process to make sheets of nanostructured Al and Mg alloys 

for potential automotive applications. 

• Continue development of low-temperature consolidation routes to produce bulk monolithic and/or composite 

alloys with nanocrystalline microstructure. 

• Characterize microstructure and properties of the bulk forms to demonstrate retention of nanostructures and 

enhanced mechanical properties. 

• Determine tensile and wear properties of bulk forms to demonstrate integrity of the consolidated alloy samples 

and guide application. 

• Initiate intensive efforts to scale-up these processes for prototype applications in the transportation and energy 

sectors. 

Introduction 

This program builds on the recent discovery of a 

low-cost means of producing nanostructured 

materials in essentially any alloy via machining. A 

notable achievement has been the development of a 

large-strain extrusion machining process (LSEM) 

for direct production of bulk forms, such as sheet 

and bar, with an ultra-fine grained (UFG) 

microstructure. These bulk forms, with or without 

additional thermo-mechanical processing, can serve 

as pre-cursor materials for component applications. 

Particulate having controlled size and shape can be 

produced directly by a unique modulation-assisted 

machining (MAM) process. The chips and 

particulates may be used as continuous or 

discontinuous reinforcements in metal or polymer 

matrices to create advanced composites; or 

consolidated into monolithic bulk forms using lowtemperature shear-based consolidation processes 

with a binding agent (e.g., polymer). 

Progress from FY 2006 in four main areas is 

highlighted in this report: (1) LSEM for direct 

production of sheet and bar with an UFG 

microstructure (2) particulate production by MAM 

(3) production of bulk alloys by low-temperature 

pressing of particulate using a very thin layer of 

polymer as binder and (4) continued development of 

high-strength nanostructured alloys with 

extraordinary thermal stability. 

Approach 

Two routes were pursued for the creation of bulk 

forms with nanocrystalline and/or UFG 

microstructures. The first involved a modified 

version of conventional chip-formation machining 

process called LSEM. In LSEM, the chip is 

produced with a controlled geometry by constraining 

its shape and size at the point of its formation. This 

method combines inherent advantages afforded by 

large strain deformation in chip formation by 

machining, with simultaneous dimensional control 

of extrusion in a single step of deformation. The 

resulting sheet offer enormous promise for ground 

transportation applications. Second, by incorporating 

controlled superimposed modulation with a 

conventional machining process, nanostructured 

chips could be directly created in micron-sized 

particulate form. The particulate have controlled size 

and shape, a tight particle size-distribution, and 

macroscopic dimensions resembling those of 

powders used in P/M processes. This process called 

MAM has provided a cost-effective solution to the 

problem of making nanostructured particulate 

suitable for consolidation. The resulting particulate 

162
Heavy Vehicle Propulsion Materials 

has served as pre-cursor material for shear-based 

consolidation processes such as powder extrusion 

and pressing with a polymer binder, and 

reinforcements for metal and polymer matrices in 

composites processing. 

Bulk Al 6061T6 alloys have been have been 

produced using a low-temperature consolidation 

pressing process wherein a very thin layer of 

polymer is used as a binder. The polymer bonding is 

a new approach in which a thermosetting polymer 

resin is used, together with high-pressure 

compression molding in a steel die to yield a 

structure consisting of low-volume fractions of 

polymer bonding the nanostructured metal particles. 

The goal of this approach is to bond the machining 

chip particulate with as fine a polymer bond layer as 

possible. Previous work has established that 

premixing the particulate and liquid thermoset and 

pressing in a steel die is preferred to pressing the 

powder alone followed by infiltration. Although the 

metal volume fraction that can be achieved is 

comparable, a significant fraction of the residual 

pores in the preform in the infiltration route are 

closed or otherwise inaccessible to the infiltrant 

leading to more residual porosity. Although less 

residual porosity occurs pressing a mixture of the 

liquid epoxy and alloy particulate, there still is 

residual porosity, mostly in large epoxy regions 

larger alloy particles exhibit did not pack well. In an 

attempt to overcome the packing limitation, 

experiments were conducted using the equi-axed 

particles now being produced by the modulationassisted machining method. 

Another avenue of investigation has explored 

creation of high-strength, high-temperature 

nanostructured materials such as titanium and 

nickel-based alloys, with exceptional thermal 

stability of the microstructure. We have 

demonstrated a means of extending the temperature 

range for structural applications of nanomaterials 

through the introduction of a dispersion of secondphase precipitates, whose thermal stability 

determines the temperature at which the strength 

deteriorates. Precipitate-stabilized nanostructured 

materials synthesized from solution-treated Inconel 

718 were found to be extremely stable even after 

prolonged heat treatment for 500 h at temperatures 

as high as one-half of the melting point. The 

enhanced thermal stability of the material’s strength 

results from retention of both the fine microstructure 

and the fine precipitate phase coexisting at the high 

FY 2006 Progress Report 

temperatures. This novel microstructure was created 

by first subjecting solution-treated Inconel 718 to 

severe plastic deformation (SPD) by machining, 

followed by thermal aging of the chip to create finescale precipitates. The principles developed herein 

appear to be equally well-applicable to a variety of 

precipitation-treatable materials systems such as P 

13 stainless steels and Waspaloy. This class of 

materials is designated - Highly-Stable 

Nanostructured Materials (HSN). 

Results 

Bulk Nanostructured Materials by LSEM 

Figure 1 illustrates a plane strain LSEM process 

wherein material (chip) of controlled thickness and 

an UFG microstructure is produced by machining 

and extrusion imposed in a single step using a 

specially designed tool. In this configuration, an 

LSEM tool moves into a disk-shaped workpiece 

rotating at a constant peripheral speed, V. The tool 

consists of two components – a bottom section with 

a sharp cutting edge inclined at a rake angle (α) 

(Figure 1); and a wedge-shaped top section that acts 

as a constraining edge. Both sections are made of a 

hard material. The machined material is 

simultaneously forced through an ‘extrusion’ die 

formed by the bottom rake face and the top 

constraining edge, thereby, effecting dimension and 

shape control. It is important to note that LSEM, 

unlike conventional machining, involves constrained 

chip formation where the deformed chip thickness is 

controlled a priori. LSEM can be considered a 

generalization of ECAP with the exemption that the 

thickness of the material at the tool exit can be larger 

or smaller than that at the inlet. This exemption 

ensures the imposition of levels of strain that are 

V

Workpiece (bulk) 

t

Cutting edge 

α tc

D

w

Rake face 

t

Constraining edge Workpiece (bulk)

t

Cutting edge

α tc

D

w

Rake face

t

Constraining edge

Figure 1. Schematic of LSEM process 

163
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

ABC

Figure 2. Bulk forms made by LSEM 

D: λ = 11 γ = 11 157± 3

C: λ = 7.4 γ = 7.4 158± 4

B: λ = 4.2 γ = 4.3 157± 5

A: λ = 1.8 γ = 2.2 148± 6

97± 3 Bu * lk copper 

Hardness 

(kg/mm2) 

Deformation condition 

D: λ = 11 γ = 11 157± 3

C: λ = 7.4 γ = 7.4 158± 4

B: λ = 4.2 γ = 4.3 157± 5

A: λ = 1.8 γ = 2.2 148± 6

97± 3 Bu * lk copper

Hardness 

(kg/mm2)

Deformation condition

Table 1. Vickers hardness of copper 

samples created with different levels of 

strain. 

AB

CD

Figure 3. TEM images and SAD patterns 

(insets) of copper subjected to different levels 

of strain in LSEM. A: γ = 2.2 (elongated), B: γ

= 4.3 (mix of elongated and equi-axed), C: γ = 

7.4 (equi-axed) and D: γ = 11 (equi-axed). 

sufficiently high to induce formation of an UFG 

microstructure in a single-step deformation process. 

Figures 2A) – 2C) show examples of bulk sheet and 

bar samples of copper created by the LSEM process. 

Figure 3 shows TEM images of microstructure of 

bulk copper samples with different levels of strain 

made by the LSEM, and the corresponding selected 

area diffraction (SAD) patterns (inset). The sequence 

of images A, B, C and D, illustrates the progressive 

refinement of microstructure and formation of UFG 

structures with increasing strain; these are consistent 

with previously observed dependence of 

microstructure on strain. An elongated UFG 

microstructure is seen at a relatively small strain 

value (Figure 3A, γ = 2.2). With increasing strain, an 

increased level of grain sub-division and formation 

of equi-axed grains are observed (Figure 3B, γ =4.3). 

Larger strains lead to essentially equi-axed UFG 

structures with grain sizes of ~ 250 nm (Figures 3C 

and 3D). In consonance with this progressive 

refinement of the microstructure, the diffraction 

pattern can also be seen to progressively evolve 

from what is practically a smeared single-crystal 

pattern (Figure 3A) to one that is indicative of larger 

misorientation angles (Figures 3C and 3D). The 

TEM analysis also showed the microstructure to be 

reasonably uniform over much of the chip thickness, 

with the exception of a small region near the edges, 

consistent with the distribution of the strain. 

Table 1 gives hardness (including standard 

deviation) of the bulk copper created with the 

different levels of strain by the LSEM. In all cases, 

the hardness values are substantially greater than 

that of the bulk copper. The highest hardness of ~ 

160 kg/mm2 is essentially indistinguishable from that 

of UFG copper microstructures created by 

machining or ECAE. The standard deviations on the 

hardness values are small, less than 4% of the mean, 

indicating relative uniformity of the material 

strength across most of the sample thickness. 

Recent tests have shown that this LSEM can be 

applied to a variety of metals (e.g., copper, titanium, 

Al 6061-T6, tantalum) resulting in bulk sheet 

samples with UFG microstructures. Furthermore, by 

combining a scaled-up version of the LSEM tool 

used here, with an appropriate machine, large scale 

production of nanostructured sheet and bar material 

appears very feasible. These will be taken up in the 

coming year. The process offers exciting 

opportunities for creating bulk sheets of Al, Mg and 

Ti alloys for ground transportation applications. 

Nanostructured Particulate by ModulationAssisted Machining (MAM) 

The machining process imposes large shear 

strains on the order of 1-10 in the chip resulting in 

significant microstructure refinement. While the 

chip has an UFG microstructure, the macroscopic 

shape and size of the chip are generally uncontrolled 

in conventional machining (note LSEM is an 

exception to this as shown earlier). Recently, a new 

machining method with superimposed lowfrequency vibration (0-500Hz), referred to as 

modulation assisted machining (MAM), has 

demonstrated the ability to control the chip 

formation process. This control of chip formation in 

164
Heavy Vehicle Propulsion Materials 

MAM has been exploited to develop a novel process 

for producing particulate metals with controlled 

sizes and shapes, and having a nanocrystalline or 

UFG microstructure. The basis of this process is the 

observation that the chip dimensions and aspectratios are established geometrically by the discrete 

cutting events, which in turn are determined by 

machining and modulation parameters. Control of 

machining and modulation parameters provides the 

capability to produce particulate of different shapes 

and sizes in a deterministic way. The effectiveness 

of this process for making particulate with sizes in 

the range of 20 μm to 2,000 μm and with shapes as 

diverse as platelet, needle and equi-axed, has been 

demonstrated for a wide variety of materials (e.g. 

titanium, stainless steel, tantalum, Al alloys). Both 

the particle shape and the particle size distribution 

can be tightly controlled by adjusting the modulation 

parameters. 

Figure 4 shows nanostructured Al6061-T6 alloy 

particulate (chips) in the form of needle and platelet 

produced by changing the modulation frequency in 

Figure 4. Al6061-T6 alloy particulate (chips) in 

the form of needle and platelet. 

identical machining conditions. The needle and 

platelet particles were produced from a d=6mm 

diameter bar rotating at frequency fw= 360rpm. 

Additional MAM tests have shown production of 

equi-axed particles and filaments during similar 

turning operations. The needle and platelet particles 

demonstrate the shape control and related effects of 

MAM on the particle surface area. For comparison, 

the needle particles in Figure 4 were produced with 

MAM conditions which produce a particle volume 

that is identical to the volume of the platelet particle. 

The needle and platelet particles develop high 

surface-areas compared to spherical shaped particles 

of the same volume. The sphericity (ratio of surface 

area of equivalent volume sphere to surface area of 

the particle) is ~0.77 for the needle particulate and 

~0.65 for the platelet particulate. This comparison 

demonstrates the extraordinary control of particle 

FY 2006 Progress Report 

size and shape control by MAM and the ability to 

create very high-aspect ratio, high surface-area 

particulates. Typical particle sizes are in the range of 

20 μm to 2,000 μm. Preliminary TEM analysis has 

shown the particles to have UFG microstructures. 

The Vickers hardness of the particulate was found to 

be substantially higher than that of the bulk material 

consistent with the UFG microstructure. 

The MAM process offers considerable scope 

for low-cost production of nanostructured particulate 

of micron-sized dimensions, extremely suitable for 

powder consolidation and spraying processes. 

Bulk Al Alloys using a Polymer Bonding Agent 

Die press bonding has been used to produce 

bulk Al-6061-T6 samples from nanostructured 6061 

T6 particulate using a small amount (<10%) of 

epoxy as binder. Figure 5A) shows the results of 

pressing the nominal 50-150 μm Al alloy particulate, 

produced by MAM, with Buehler Epoxide resin at 

800 MPa pressure in a steel die, followed by curing 

at 100°C for 1 hour. A thin "glue line" of epoxy is 

present at all the boundaries but the high strength of 

the particles seems to preclude any significant 

component of their densification due to particle 

plasticity. On the other hand, the fraction of epoxy 

+ porosity is similar to composites produced by the 

same process, but using higher aspect ratio 

particulate produced by milling machining chips. A 

potential property advantage of the equiaxed particle 

composites may be less anisotropy compared to 

those made with anisometric particles; this aspect is 

yet to be quantified. 

The effect of removing the native surface oxide 

from the particles before bonding was also 

investigated. The same procedure as above was 

followed except that the particles were stripped in a 

10% NaOH solution for 10 minutes under 

ultrasonication and then rinsed in distilled water and 

dried in air. The resulting composite is shown in 

Fig. 5B). Overall, more uniform packing was 

observed, but the overall packing density (and 

porosity) levels were comparable to the composites 

made without pre-cleaning the particles in this way. 

An intriguing observation is that the interparticle 

bond ("glue line") appears to be thicker with the precleaned particles. 

Epoxy and porosity volumes as low as 6% have 

been achieved so far, and preliminary nanohardness 

measurements show little or no loss of hardening 

165
Heavy Vehicle Propulsion Materials 

imparted in producing the chips, consistent with 

200 μm 

200 μm 

A 

B 

200 μm

200 μm

A

B

Figure 5. Polymer-bonded Al alloy systems. 

separate measurements of coarsening-softening 

kinetics in this alloy. 

Pressing experiments soon to be carried out 

in an autoclave are likely to reduce the porosity to 

minimal levels. Work in the next year will proceed 

to tensile testing in order to evaluate the degree of 

bonding. Although high strength and ductility are 

not predicted, the high hardness should impart 

improved wear resistance together with the shaping 

advantages of powder processing. Applications for 

lightweight, wear-resistant components are 

envisioned. 

Highly Stable Nanostructured Materials (HSN) 

The utility of the extraordinary strengths of 

nanostructured materials is typically limited by their 

low thermal stability. We have demonstrated a 

means of extending the temperature range for 

structural applications of nanomaterials through the 

introduction of a dispersion of second phase 

precipitates, whose thermal stability determines the 

temperature at which the strength deteriorates. This 

FY 2006 Progress Report 

demonstration has been realized in a model hightemperature alloy system-Inconel 718; the resulting 

material is called HSN Inconel. 

Figure 6 illustrates the nanostructured material 

composed of grains typically 100 nm in size 

produced from solution-treated microcrystalline 

Inconel 718 (grain size ~50 μm in the initial state) 

due to plastic shear strains of ~6 imposed using a ­

20° rake angle tool. The microstructure is composed 

primarily of equi-axed grains, along with some 

elongated structures. The diffraction pattern inset in 

Figure 6 indicates the presence of large angle grain 

boundaries in this nanostructured material. The 

Vickers microhardness value of this microstructure 

was 550 kg/mm2

 (~5.4 GPa), nearly twice that of the 

undeformed bulk, microcrystalline solution-treated 

material (300 kg/mm2

). Figure 7 illustrates the much 

improved ultimate tensile strength of this 

nanostructured material compared to the bulk 

microcrystalline material. Nanostructured Inconel is 

seen to possess a peak tensile strength of ~1.3 GPa, 

more than twice the yield strength of the bulk 

material (~600MPa). As expected its ductility is 

limited, but the tensile fracture surface showed veinlike features typical of micro-scale plasticity during 

crack propagation and fracture. Interestingly, the 

micrometer-scale features on this fracture surface 

are not characteristic of fracture surfaces of 

conventional nanocrystalline materials that usually 

show sub-micrometer scale dimpled features. 

Figure 6. TEM micrograph of nanostructured 

Inconel 718 (shear strain ~6). Grain size 

~100 nm. 

166
Heavy Vehicle Propulsion Materials 

When this high-strength nanostructured material 

was heat treated at 600°C (~ one-half the melting 

point), a rapid strengthening was found that can be 

associated with the formation of second-phase 

precipitates. A peak hardness value of 630 kg/mm2 

(~ 6.2 GPa) is reached after about 6 hours of 

annealing. In contrast, precipitation of the 

intermetallic γ" and γ’ phases to achieve peak 

precipitation strengthening in the conventional 

microcrystalline structure usually involves a 

lengthier heat treatment process due to the 

characteristically sluggish precipitation kinetics. 

Additional evidence for the extraordinary 

strengthening and thermal stability comes from 

measurements of stress-strain curve of HSN Inconel 

after 6 hours of heat treatment at 700°C. The 

material has a failure stress of 1.4 GPa as illustrated 

FY 2006 Progress Report 

distributed in the matrix. In the nanostructured 

matrix, this phase is more likely to precipitate on the 

grain boundaries. This preferential precipitation can 

aid efficient pinning of the grain boundaries and 

enhance thermal stability. Thermal stability is also 

aided by the stability of the fine precipitate phase, 

which has been reported to have large activation 

energy for coarsening of ~1000 kJ/mol. This ensures 

the stability of the fine precipitate phase and 

consequent retention of strength in the 

nanostructured Inconel even after prolonged 

annealing at 600°C, resulting in its designation as a 

Highly-Stable Nanostructured (HSN) material. 

The materials design principle propounded 

above can be generalized to other precipitationtreatable materials and offer a new line of 

development for this important class of structural 

materials. Other material systems that form stable, 

Engineering stress (MPa) 1600 Nanostructured Inconel 718

(after SPD)

1400 

Inconel 718

HSN Inconel 718 o after 6 hours at 700 C 1200 Bulk microcrystalline

1000 

800 

600 

400 

200 

0 

-200 0 0.01 0.02 0.03 0.04 0.05 

Engineering strain

Figure 7. Stress-strain curves for Inconel-718 

in Figure 7. The measurements also indicate an 

expected, albeit, modest improvement in the 

ductility of the HSN Inconel. Much more 

noteworthy, however, is the observation that this 

tensile strength following the heat treatment at 

700°C is very comparable to, if not markedly greater 

than, that of the HSN Inconel prior to the heat 

treatment. This reinforces our observation of the 

extraordinary stability of the tensile strength of HSN 

Inconel. 

In the nanostructured state, the large diffusion 

coefficient and the large number of nucleation sites 

resulting from SPD of the solution-treated material 

allow a rapid precipitation of a fine dispersion. 

Furthermore, the Ni-Nb intermetallic γ" phase, that 

can significantly contribute to material strength, is 

characterized by a very fine and dense dispersion of 

coherent, disc shaped precipitates, that are 

fine precipitate phases, for example, special Al 

alloys, P-series stainless steels, Waspaloy, can be 

expected to behave in a similar manner in the 

nanostructured state. Exceptional thermal stability 

and retention of the material strength are ensured by 

the stable, fine precipitate phase and the fine 

microstructure coexisting at high temperatures for 

significant lengths of time, resulting in a class of 

HSN materials with enormous potential for highstrength and/or high-temperature applications. These 

avenues will be investigated in detail in the coming 

year. 

Publications 

P. Iglesias, M. D. Bermudez, F. J. Carrion, S. 

Chandrasekar, and W.D. Compton, A study of the 

wear behavior of polymer-matrix composite 

containing discontinuous nanocrystalline alloy 

reinforcements, Tribology International, 40, 479 

(2006). 

M. R. Shankar, B. C. Rao, S. Lee, S. 

Chandrasekar, A. H. King, and W. D. Compton, 

Severe plastic deformation of titanium at near 

ambient temperature, Acta Materiala, 54, 3691 

(2006). 

Y. Uluca, B. C. Rao, M. Ravi Shankar, T. L. 

Brown, J. B. Mann, S. Chandrasekar, and W. D. 

Compton, Low-cost ultra-fine grained materials 

from aerospace machining chips, SAE Journal of 

Aerospace, (2006). 

S. Lee, J. Hwang, M. Ravi Shankar, S. 

Chandrasekar and W. Dale Compton, Large strain 

167
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

deformation field in machining, Metallurgical and 

Materials Transactions A, 37, 1633 (2006). 

W. Moscoso, M. R. Shankar, J. B. Mann, W. D. 

Compton, and S. Chandrasekar, Bulk nanostructured 

materials by large-strain extrusion machining, 

Journal of Materials Research, in press, to appear in 

January 2007. 

S. Swaminathan, M. R. Shankar, B. C. Rao, W. 

D. Compton, S. Chandrasekar, A. H. King, and K. P. 

Trumble, Severe plastic deformation and 

nanostructured materials by machining, Journal of 

Materials Science, in press, to appear in January 

2007. 

Presentations 

S. Chandrasekar et al., Severe Plastic 

Deformation and Nanostructured Materials by 

Machining, Sandia National Labs, Albuquerque, 

NM, June 2006. 

S. Chandrasekar et al., Severe Plastic 

Deformation and Nanostructured Materials, Indian 

Inst. Sci., Bangalore, India, March 2006. 

S. Chandrasekar, J. B. Mann and M. R. Shankar, 

Nanostructured Materials by Machining, Timken 

Company, Canton, OH, April 2006. 

S. Chandrasekar, Manufacturing Innovations, 

State of Indiana Advancing Manufacturing Summit, 

Purdue University, West Lafayette, IN, May 2006. 

J. B. Mann et al., Modulation-Assisted 

Machining for Particulate Production, Presentation 

before Army Research Labs panel, Purdue 

University, West Lafayette, IN, July 2006. 

S. Chandrasekar, Severe Plastic Deformation 

and Nanostructured Materials, Naval Post Graduate 

School, Monterrey, CA, June 2006. 

S. Chandrasekar et al., Advanced Manufacturing 

Technologies, State of India Senate Sub-Committee 

on Science and Technology, Indianapolis, IN, 2006. 

J. B. Mann et al., Modulation-Assisted 

Machining for Particulate Production, US Army 

Benet Labs, Watervliet, NY, August 2006. 

J. B. Mann et al., Modulation-Assisted 

Machining for Particulate Production, US Army 

Aberdeen Arsenal, MD, September 2006. 

S. Chandrasekar et al., Measurements of Strain 

in Severe Plastic Deformation, UFG 2006, 

Kloitsersee, Germany, September 2006. 

K. P. Trumble et al., Severe Plastic Deformation 

of Difficult-to-Deforma Materials at Ambinet 

Temperature, UFG 2006, Kloistersee, Germany, 

2006. 

K. P. Trumble et al., Consolidation of 

Nanostructured Chip Particulate by Powder 

Extrusion, TMS meeting, Orlando, FL, March 2006. 

A. H. King et al., Large Strain Deformation 

Maps for Copper, TMS meeting, March 2006. 

K. P. Trumble, Severe Plastic Deformation and 

Nanostructured Materials, Max Planck Institute, 

Stuttgart, Germany, September 2006. 

Related Patents 

Four patent applications pertaining to the 

machining technology for making bulk and 

particulate materials have been filed in May 2006. 

Two of these technologies have already been 

licensed. 

168
Heavy Vehicle Propulsion Materials FY 2006 Progress Report 

C. Deformation Processes for the Next-Generation Ceramics 

P. F. Becher 

Oak Ridge National Laboratory 

P.O. 2008, MS 6068 

Oak Ridge, Tennessee 

(865) 574-5157; fax: (865) 574-6098; e-mail: becherpf@ornl.gov 

DOE Technology Development Manager: Jerry L. Gibbs 

(202) 586-1182; fax: (202) 586-1600; e-mail: Jerry.gibbs@ee.doe.gov 

ORNL Technical Advisor: D. Ray Johnson 

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

Contractor: Oak Ridge National Laboratory, Oak Ridge, Tennessee 

Contract No.: DE-AC05-00OR22725 

Objectives 

• Promote plastic deformation of oxide ceramics by the application of electric fields for application to 

densification and shape forging of nanocrystalline ceramics. 

• Characterize thermal and mechanical properties to assess influence of nanocrystalline microstructures for 

gaining improved performance for wear and other components. 

Approach 

• Collaborate with members of both the Pennsylvania State University and the Korea Institute of Machinery and 

Materials to develop fabrication technology to be able to fabricate dense monolithic nanocrystalline (i.e., grain 

sizes <100-nm) bodies. 

• Characterize the potential for enhanced plasticity in ceramics resulting from sub-micron grain size on materials 

and from the application of applied electric fields with North Carolina State University researchers. 

• Characterize the effects of nanometer scale microstructures on the mechanical and wear performance of 

ceramics. 

Accomplishments 

• In collaboration with Korea Institute of Machinery and Materials, it was demonstrated that dense 

nanocrystalline (grain size < 100 nm) could be produced through the use of a two-step sintering cycle and 

control of the composition. 

• Demonstrated large reduction in tensile flow stress in a dense submicron grain sized zirconia ceramic 

containing 2.5 mol % yttria with the application of an applied D. C. electric field and showed that flow stress 

reduction increased to over 90% with increase in both permanent strain and strain rate. 

Future Direction 

• The fracture and wear properties of nanocrystalline ceramics will be examined to seek approaches for 

developing ceramics with substantially improved wear resistance. 

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Heavy Vehicle Propulsion Materials 

Introduction 

In addition to lower specific weight, ceramics 

have a number of attractive properties for application as components in various heavy-vehicle 

systems. However, many heavy-vehicle components 

require ceramics with complex shapes as well as 

function. Furthermore the demanding service conditions necessitate improved mechanical reliability 

together with fracture and/or wear resistance. 

Underlying the performance and use of ceramics is 

their tendency to fail in a brittle fashion except at 

temperatures generally well above 1000°C. Complex 

shape forming processes for ceramics are relegated 

to approaches based on powder processing as a 

result of the excessive temperatures for hot forming. 

On the other hand, metallic alloys can be hot worked 

at temperatures below 1000°C to not only form the 

desired shape but also to enhance their properties. 

In ceramic systems, one needs to combine processing to tailor microstructures with novel concepts 

that offer the potential to enhance deformation. One 

such concept that has seen considerable interest is 

that of superplastic deformation of submicron-grainsized ceramics above 1000°C.1

 Recent studies 

suggest that reducing the microstructure to achieve 

grain sizes <100 nm could lead to permanent/ 

plastic deformation in ceramics at quite modest 

temperatures.

2,3

 Recent studies have shown that the 

application of electric fields can reduce the yield 

stress or increase the creep rates of ceramic 

systems.4

 To date, this has not been exploited to 

determine the effects on the deformation of 

nanocrystalline oxides. The present study seeks to 

examine the influence of true nanocrystalline microstructures on the mechanical properties of ceramics, 

as well as approaches to economically fabricate 

nanocrystalline ceramics monoliths. 

Hardness values generally obey the Petch-Hall 

behavior with the hardness rising as the grain size is 

reduced. However, results show that when the grain 

size is ~20 nm, the room temperature hardness then 

decreases as the grain size continues to be 

decreased.1,2

 This behavior has two important 

implications for applications of interest here. One is 

that some permanent deformation might be achieved 

in ceramics with grain sizes <30 nm, and this might 

offer improved reliability. More likely of interest is 

the fact that ceramics with grain sizes below 100 nm 

but above the softening point will be extremely hard, 

and that offers substantial advantages in overcoming 

FY 2006 Progress Report 

wear problems in rolling stock. In addition, surface 

roughness of as-sintered ceramics decreases as the 

grain size decreases. In fact, the as-sintered surfaces 

of alumina substrates with grain sizes <500 nm 

exhibit a mirror-like finish without any surface 

preparation. This feature should be applicable to the 

fabrication of low-friction components. 

Approach 

This project involves two major thrusts. The 

first is a collaborative effort with researchers at the 

Korea Institute of Machinery and Materials on the 

sintering of nanocrystalline silicon nitride ceramics. 

In addition, collaborations continue with researchers 

at Pennsylvania State University, who are 

fabricating nanocrystalline oxides. These efforts are 

aimed at finding approaches to produce sizeable 

monolithic bodies for subsequent evaluation. The 

second thrust focuses on the mechanical and friction 

and wear behavior of nanocrystalline ceramics and 

the influence of external parameters [e.g., stressing 

rate, temperature, environment (e.g., air, vacuum)], 

and electric fields on the plastic deformation 

behavior. 

Fabrication of Nanostructured Ceramics: 

Densification of nanocrystalline ceramics has 

typically involved the use of applied pressure to 

achieve dense sub-100-nm grain sized samples. In 

those few cases where pressureless sintering has 

been employed successfully, long hold times at the 

sintering temperatures were required to achieve densities ≥95% of theoretical, which allows undesired 

grain growth to occur. On the other hand, a two step 

sintering cycle offers a viable alternative approach.5 

This involves first heating to a temperature 

sufficient to achieve densities of ≥75% of theoretical 

followed by rapidly lowering the temperature to a 

level to obtain sufficient boundary diffusion. Using 

an extend hold time at the lower temperature results 

in the removal of pores trapped on boundaries while 

minimizing grain growth during the final stage of 

densification. 

Deformation Behavior in Nanostructured 

Materials: Baseline deformation data will be established through studies of high purity, dense submicron grain size zirconia ceramics in order to 

evaluate the influence of nanocrystalline microstructures. These studies will then be extended to 

<100-nm grain sized ceramics. This will include 

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Heavy Vehicle Propulsion Materials 

(a) nanostructured ceramics to explore for the 

softening effects observed in nanocrystalline metals 

by evaluating both the creep response over a range 

of temperatures and applied stresses, and the yield 

stress as a function of strain rate and temperature, 

and (b) the influence of external factors (e.g., 

electric fields) that have been shown to enhance 

deformation processes in ceramics with micrometer 

grain sizes. Modeling studies will be conducted to 

describe the deformation behavior and to address 

such phenomenon as space charge/solute 

segregation coupled with electric field effects. 

Additionally, the fracture, hardness, and wear 

response of nanocrystalline oxides will be evaluated 

to assess their application in wear components. 

Results 

Fabrication of nanocrystalline ceramics 

The ability to fabricate dense bulk 

nanocrystalline ceramics represents a substantial 

barrier to their eventual application. In the present 

project, the fabrication of bulk nanocrystalline 

zirconia has been achieved using a combination of 

colloidal and granulation processing methods in 

conjunction with novel two-step sintering cycles 

working with researchers at Pennsylvania State 

University. In addition, bulk nanocrystalline silicon 

nitride ceramics have been fabricated using a similar 

approach to sintering working with researchers at 

the Korea Institute for Machinery and Materials.6 

In both cases the sintering cycle involves 

heating to a peak temperature with a very 

abbreviated hold there followed by lowering the 

temperature by ≤ 200°C and holding there for 10 to 

20 hours. The first step is designed to allow the 

ceramic body to reach a density in excess of 75% of 

the theoretical density. The second step allows 

further densification to occur by grain boundary 

diffusion. The key is to inhibit grain coarsening 

such that pores are retained at grain boundaries 

where they can pin boundaries, as well as be 

removed by diffusion. These collaborative studies 

make it possible to fabricate materials with a range 

of microstructures with grain sizes less than 100 nm 

where nanocrystalline effects are expected to have 

the greatest impact. 

Electric field effects on plastic deformation 

Collaborations with Prof. Hans Conrad and 

Dr. Kang Jung of North Carolina State University 

FY 2006 Progress Report 

Figure 1. Dense nanocrystalline silicon nitride 

ceramic produced by two-step sintering.6 

have assessed the reduction in flow stress in zirconia 

ceramics due to the application of an applied direct 

current (dc) field. Earlier studies of the 

enhancement of deformation by the application of 

electric fields to dense nanocrystalline zirconia 

ceramics subjected to compressive stresses 

suggested that the field effects were small. On the 

other hand, previous observations on oxides4 

exhibited substantial effects in oxides subjected to 

tensile loading. Our current studies seek to 

understand how superplastic behavior might be 

promoted at lower temperatures. These studies seek 

to determine how to enhance plasticity in that can 

also be employed to form nanocrystalline ceramics. 

Re-evaluation of the electric field effects in the 

compression test results suggested that the 

magnitude of the electric field was much lower than 

expected due to the higher resistivity of the dense 

SiC push rods required to support the compression 

loads as compared to the resistivity of the lower 

density electric heating element grade of SiC used as 

pull rods in prior tensile tests. If this resulted in a 

much lower applied electric field in the compression 

tests, the reduction in the flow stress with the field 

applied versus not applied would be expected to be 

quite small. To test this, the current tensile tests at 

1450°C included to electrode configurations: (1) the 

voltage applied directly through a circuit consisting 

of the upper and lower low resistivity SiC pull rods 

and the zirconia tensile sample and (2) the circuit 

path of a SiC sleeve with an air gap separating the 

sleeve from the ZrO2 sample that was connected to 

the lower low resistivity SiC pull rod. In 

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Heavy Vehicle Propulsion Materials 

configuration 2, the maximum applied field was 

~3.3 V/cm while in configuration 1 fields of up to 

~46 V/cm were achieved. Using configuration 2 

with an applied field of ~3.3 V/cm, the reduction in 

flow stress induced by the field was ≤10% of that 

obtained with configuration 1 where an applied field 

of 46 V/cm was achieved. Thus, one can attribute 

the minimal electric field effects on the flow stress 

reduction in the earlier compression tests to a 

substantial reduction in the actual applied electric 

field achieved. 

With configuration 1, the tensile flow stress at 

1450°C was reduced by upwards of 90% with 

increase in strain and strain rate with the application 

of an electric field of 46 V/cm. In order to 

accurately assess the impact of the electric field on 

the flow stress, one must account for the increase in 

specimen temperature due to Joule heating, which 

additional test results indicate minimal (~25°C) rise 

in temperature due to Joule heating. 

The flow of vacancies needed for material 

transport and dislocation motion by applying stress 

and an electric field. The direction of the vacancy 

flux is a function of the sign of the applied stress 

(i.e., vacancies flow to regions under compression). 

This applies locally to vacancy flow in the 

deformation of a polycrystalline ceramic where 

vacancies move from grain boundaries under tension 

to those in compression while ions/atoms move in 

the reverse direction. The presence of space charge 

zones at grain boundaries also impacts the 

deformation.7

 Such zones are due to the migration 

of solute atoms (e.g., Y) that concentrate at the 

boundaries. As a result, the reduction in the flow 

stress (ΔσE) is a function of the electric field (E): 

ΔσE = ZeExSC Ω , (1) 

where Z is the valence of the ion (Zr+4 in this case); 

e is the electronic charge (1.6 x 10-19 coulomb)s; Ω

is the atomic volume of Zr; ΧSC is the width of the 

space charge zone at the grain boundaries. 

Based on the flow stress reduction (~6 MPa) at 

1450°C due to an applied field of 46 V/cm, the 

space charge zone width is calculated to be 4 to 5 

nm. The width of the space charge zone is also 

related to the Debye length of the material: 

XSC = βλ (2) 

FY 2006 Progress Report 

where it is a proportionality factor in the range of 2 

to 5. The value of XSC calculated from the current 

results are consistent with values obtained from 

studies of the ionic conductivity of zirconia 

ceramics.8 However, theoretical treatments show 

that the Debye length (δ) in ionic solids increases as 

both the bulk solute content and temperature are 

decreased.9,10 The increases in the Debye length due 

to decreases in temperature and solute content will 

be reflected as increases in the space charge width. 

One should then expect that the decrease in 

flow/yield stress brought about by application of an 

electric field would increase as the temperature 

decreases. 

Conclusions 

Until now, the ability to fabricate dense bulk 

nanocrystalline ceramics has been a major barrier to 

their evaluation for heavy vehicle applications. 

With the advances in the fabrication technology, one 

can now assess the wear resistance and 

electrochemical performance of nanocrystalline 

nitrides and oxides, respectively. The demonstration 

of electric field effects on deformation offer 

opportunities for deformation forming of 

nanocrystalline ceramics to reduce the temperatures 

required for densifying sub-100 nm grain sized 

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). Ibid, “Superplasticity-Like Deformation of 

Nanocrystalline Monoclinic Zirconia at Elevated 

Temperatures,” J. Am. Ceram. Soc., 87(6), 1122–25 

(2004). 

4. H. Conrad, “Thermally Activated Plastic 

Flow of Metals and Ceramics with an Electric Field 

or Current,” Mater. Sci. Eng. A322: 100-107 (2002). 

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Heavy Vehicle Propulsion Materials 

5. I.-W. Chen and X.-H. Wang, “Sintering Dense 

Nanocrystalline Ceramics Without Final-Stage 

Grain Growth,” Nature, 404, 168–71 (2000). 

6. H.-D. Kim, Y.-J. Park, B.-D. Han, M.-W. 

Park, W.-T. Bae, Y.-W. Kim, H.-T. Lin and P. F. 

Becher, Scripta Mater., 54: 615-19 (2006). 

7. H. Conrad, “Space Charge and the Dependence of the Flow Stress of Ceramics on an Applied 

Electric Field,” Scripta Mater., 44, 211–16 (2001). 

8, X. Guo, J. Maier, “Grain Boundary Blocking 

in Zirconia: A Schottky Barrier Analysis,” J. 

Electrochem. Soc. 148 (3) E121– E126 (2001). 

9. K. L. Kliewer and J. S. Koehler, “Space 

Charge in Ionic Crystals. I. General Approach with 

Application to NaCl,” Phys. Rev. A, 140 (4A), 

1226–40 (1965). 

10. M. F. Yan, R. M. Cannon, H. K. Bowen, 

and R. L. Coble, “Space-Charge Contribution to 

Grain-Boundary Diffusion,” J. Am. Ceram. Soc., 60 

(3–4),120–27(1977). 

FY 2006 Progress Report 

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