The Development of Alumina-Forming Austenitic Stainless Steels for High-Temperature Structural Use

By Brady, M P Yamamoto, Y; Santella, M L; Maziasz, P J; Pint, B A; Liu, C T; Lu, Z P; Bei, H

A new family of alumina-forming au.slenitic stainless steels is under development at Oak Ridge National Laboratory for structural use in aggressive oxidizing environments at 600-900[degrees]C. Data obtained to date indicate the potential to achieve superior oxidation resistance compared to conventional Cr^sub 2^O^sub 3^- forming iron- and nickel-based heat-resistant alloys, with creep strength comparable to state-of-the-art advanced austenitic stainless steels. A preliminary assessment also indicated that the newly developed alloys are amenable to welding. Details of the alloy design approach and compositionmicrostructure-property relationships are presented. INTRODUCTION

Oxidation resistance is one of the primary considerations that determine the durability of heat-resistant alloys. The key to good oxidation resistance is to establish an external, continuous layer of a slow-growing, thermodynamically stable oxide phase such that subsequent oxidation is limited by diffusion of metal or oxygen species across this layer. For high-temperature applications (i.e., >600[degrees]C), Cr^sub 2^O^sub 3^ and Al^sub 2^O^sub 3^ are the principal oxides used for the protection of metallic alloys.

In many high-temperature environments Al^sub 2^O^sub 3^ scales offer a superior degree of protection to Cr^sub 2^O^sub 3^.8-15 Alumina scales grow at a rate that is 1 to 2 orders of magnitude lower than that of Cr^sub 2^O^sub 3^ (Figure 1a). Alumina is also significantly more thermodynamically stable than is Cr^sub 2^O^sub 3^ (Figure 1b). Alumina scales have proven to be particularly beneficial in the presence of aggressive carbon- or sulfur-species encountered in combustion and chemical process industry applications.8,9 Further, in combustion environments, a key advantage of Al^sub 2^O^sub 3^ over Cr^sub 2^O^sub 3^ is greater stability in the presence of water vapor.16 With both oxygen and water vapor, volatile chromium oxy-hydroxide species can form and significantly reduce oxidation lifetime, in part by constantly thinning the protective Cr^sub 2^O^sub 3^ scale that is established, resulting in linear oxidation kinetics.16,17 Such attack is particularly relevant for thin-walled components such as heat exchangers.18 However, despite the many advantages of protective Al^sub 2^O^sub 3^ scales, virtually all iron-based, heat-resistant structural alloys for use above ~600[degrees]C utilize Cr^sub 2^O^sub 3^-based scales for protection.8’10 This is due to the extensive solid solubility and excellent metallurgical compatibility of chromium in Fe/Fe(Ni), which permits ready formation of a protective Cr^sub 2^O^sub 3^-based scale with ample alloy design flexibility to co-optimize oxidation resistance with other needed properties such as creep resistance, weldability, etc.

Ferritic Fe-Cr-Al-based alloys capable of Al^sub 2^O^sub 3^ formation are widely used in specialty applications such as heating elements and furnace liners. Despite their outstanding oxidation resistance, they are not suitable for structural applications above ~500-600[degrees]C due to poor creep resistance resulting from their open body-centered cubic structure. To obtain creep resistance above ~600[degrees]C in conventional cast or wrought iron-based alloys, an austenitic face-centered cubic structure is needed. Oxide dispersion strengthened (ODS) ferritic Fe-Cr-Al-based alloys10 and nickel- based alloys8-10 capable of alumina scale formation and with excellent high-temperature creep resistance are also available, but their high cost limits their use.

The interest to create Al2O3-forming austenitic (AFA) stainless steels for use as heat-resistant structural alloys dates back at least 30 years (e.g., References 17, 20-23). A major complication for developing a successful AFA stainless steel is that aluminum is a strong ferrite stabilizer. Further, the alloys also require the addition of significant quantities of chromium, and a ferrite stabilizer, to help promote protective Al^sub 2^O^sub 3^ scale formation. (Additions of chromium reduce the critical level of aluminum in an alloy needed to form a protective Al^sub 2^O^sub 3^ scale, often referred to as the third-element effect.10,24,25) Typically explored alloying addition levels of ~4-6 wt.% aluminum and ~10-25 wt.% chromium can destabilize the parent austenitic matrix structure, resulting in duplex ferritic/austenitic microstructures and a loss of creep resistance (e.g.. References 2, 17). The desired austenitic matrix structure can be stabilized in these alloys by large additions of nickel, but the levels needed usually result in a nickel-based alloy rather than an iron-based alloy, and the cost advantages are lost. This paper overviews recent efforts at Oak Ridge National Laboratory (ORNL)1-7 devoted to developing creep-resistant AFA stainless steels with relatively low levels of nickel, comparable to existing advanced austenitic stainless steels and alloys.

ORNL AFA DEVELOPMENT APPROACH

Alumina-forming austenitic alloy development efforts at ORNL used the high-temperature ultrafine-precipitation-strengthened (HTUPS) family of austenitic stainless steel alloys26 as a starting point for alloy modification.2 These alloys exhibit among the highest creep strengths ever reported for austenitic stainless steel alloys. However, they were originally developed for advanced liquid-metal nuclear reactor environments where gas-metal oxidation was not a major consideration. Consequently their oxidation resistance is relatively poor. A typical HTUPS composition is Fe-14Cr-16Ni-2.5Mo- 2Mn-0.5Ti0.3V-0.15Nb wt.% base with additions of B, C, and P to form stable nanoscale precipitates such as MC (M = Nb, Ti, V) carbides for strengthening.26 Levels of -2.5 (alloy HTUPS-I) and 3.8 wt.% aluminum (alloy HTUPS-2) were added initially to a baseline HTUPS alloy composition, with the nickel level increased to 20 wt.% in order to stabilize austenite, in an attempt to promote protective Al^sub 2^O^sub 3^ scale formation (Table I).2

Evaluation of creep-rupture life at 750[degrees]C and 100 MPa in air indicated that the HTUPS-1 alloy with 2.5 wt.% aluminum exhibited excellent creep resistance (Figure 2a). For comparative purposes, conventional austenitic stainless steels such as type 347 stainless steel alloys (-Fe-18Cr-I INi wt.% base) ruptured in less than -100-300 h under these conditions. State-of-the-art austenitic stainless steel alloys such as alloy 709 (~Fe-25Ni-20Cr wt.% base) can exhibit creep rupture lives at 750[degrees]C and 100 MPa from – 2.000-6,000 h.

Increasing the level of aluminum to 3.8 wt.% (HTUPS-2) resulted in the formation of delta-ferrite in the microstructure, which converted to sigma phase on exposure at 750[degrees]C and degraded creep resistance.2 However, a short-term oxidation screening at 800[degrees]C in air indicated that neither the 2.5 wt.% nor 3.8 wt.% level of aluminum addition was sufficient to enable Al^sub 2^O^sub 3^ scale formation (Figure 3). Instead, the aluminum was internally oxidized in HTUPS-1 and HTUPS-2, and the external scale consisted of faster-grow ing, less-protective mixed iron- and chromium-rich oxide phases.2 Such behavior is the reason previous alloy development efforts generally utilized higher levels of aluminum (4-5 wt.%) and chromium (>15 wt.%) with the drawback of destabilizing austenite and a loss of creep resistance without large nickel additions.

Protective Al^sub 2^O^sub 3^ scale formation was, however, achieved in an HTUPS type alloy with only -2.5 wt.% aluminum (AFA 2- 1) (Table I and Figure 3c). (All Al^sub 2^O^sub 3^-forming alloys are referred to as AFA, with 2, 3, or 4 designating 2.5, 3, or 4 wt.% aluminum series). This was accomplished by eliminating the titanium and vanadium additions to the alloy and increasing the niobium level to 0.9 wt.% niobium2,4 (Table I). Despite the absence of titanium and vanadium, the AFA 2-1 alloy still exhibited excellent creep resistance at 75O0C and 100 MPa in air with a rupture lifetime of -2,000 h (Figure 2a). Creep-rupture life data for AFA 2-1 and some comparable commercial Cr,O3-forming alloys are presented in the Larson-Miller plot shown in Figure 2b.

Transmission-electron microscopy (TEM) analysis indicated that the creep resistance in AFA 2-1 resulted from nano MC precipitates, similar to those observed in alloy 709 (Figure 4). It should be noted that as with the original HTUPS alloys,26 AFA 2-1 utilized 10% cold work2 to enhance MC precipitation and creep resistance. Subsequent studies of related developmental AFA alloys have indicated that good creep resistance can also be obtained in solution-treated material without prior cold work.

ALLOY DESIGN DIRECTIONS

Subsequent studies of AFA HTUPS alloys at ORNL have explored the effects of alloying additions on oxidation and creep behavior, with an emphasis on the variation of Nb, Ti, and V content as a function of Ni and Al levels (Table I).3-7 The alloys were manufactured by arc-casting (200 g or 1 kg castings) which were then solution heat treated in the 1,200-1,250[degrees]C range, cold worked 40-70%, and recrystallized at 1,200-1,250[degrees]C to control alloy grain size (nominally -50 [mu]m to 150 [mu]m size range). Surprisingly, niobium additions appeared to be the key to oxidation resistance,4 particularly in water-vapor containing environments. Water-vapor environments can be highly aggressive, due not only to volatilization of chromium-oxide scales, but also because of increased tendency for internal oxidation as compared to dry air exposure.30 This latter factor is of great relevance to the ORNL AFA alloy composition range, which falls close to the critical minimum levels of aluminum and chromium needed for external Al^sub 2^O^sub 3^ scale formation.2,4,6,31 Figure 5a shows oxidation kinetics at 800[degrees]C in air with 10% water vapor. AFA 4-1 (4AI/0.6Nb/ 0.1Ti) was able initially to establish an Al^sub 2^O^sub 3^ scale but was not able to maintain it despite the relatively high level of 4 wt.% aluminum. Rather, a transition to iron-rich oxide nodule formation and subsequent scale spallation was observed (mass loss in Figure 5a). In contrast, better oxidation resistance was exhibited by lower aluminum- but higher niobium-containing alloys, AFA 2-4 (2.5A1/3NW0.2V) and AFA 3-8 (3Al/2.5Nb/0.1Ti), illustrating the beneficial effects of niobium on oxidation resistance of AFA alloys in water-vapor-containing environments. These alloys were able to form protective Al^sub 2^O^sub 1^ scales under these aggressive test conditions despite the presence of small additions of vanadium or titanium, which were observed to degrade Al2O3-forming ability at higher levels (0.5Ti and 0.3V) in HTUPS-1 and HT-UPS-2 (Figure 3). Some tolerance for vanadium and titanium is important because vanadium and titanium can be used to enhance MC carbide formation for improved creep resistance, and because vanadium and titanium also are common impurities present in commercial stainless steel grades. The ability to tolerate 0.2 wt.% vanadium and form Al^sub 2^O^sub 3^ was lost at 800[degrees]C in air with 10% water vapor when the niobium content was decreased to I wt.% in alloy 3-5 (3A1/ 1Nb/0.2V) (Figure 5a).

Findings to date indicate that increasing niobium, aluminum, and/ or nickel content all favor the establishment and maintenance of protective Al^sub 2^O^sub 3^ scale formation in these alloys.4,6 The mechanisms behind these trends are not fully understood and are the subject of ongoing investigation. It is speculated that the key factor controlling if and how long Al^sub 2^O^sub 3^ scale formation occurs in AFA alloys is oxygen solubility in the alloys. Increased niobium and nickel levels may reduce alloy oxygen solubility, which favors external protective Al^sub 2^O^sub 3^ scale formation, while the addition of titanium and vanadium potentially increases oxygen solubility, favoring internal oxidation of aluminum.

The capacity of the AFA alloys to form Al^sub 2^O^sub 3^ scales was degraded with increasing oxidation temperature,46 with the lower aluminum-, niobium-, and nickel-containing alloys exhibiting a transition to internal oxidation of aluminum (i.e., losing the capacity to form protective Al^sub 2^O^sub 3^ scales) between -800 and 900[degrees]C, and the higher aluminum, niobium, and nickel alloys losing the capacity to form Al^sub 2^O^sub 3^ between ~900 and 1,000[degrees]C. This behavior is illustrated in Figure 5b, which shows oxidation behavior (500 h cycles) at 900[degrees]C in air. AFA 3-2 (3 l/0.6Nb/0. ITi) and AFA 2-1 4(2.5Al/0.9Nb) (not shown in Figure 5b) showed large mass gains after the initial 500 h cycle, with cross sections indicating many locations with ironrich oxide nodules overlying regions of internally oxidized aluminum, rather than external protective Al^sub 2^O^sub 3^.

The AFA 2-4 alloy (2.5Al/3Nb/ 0.2V) which exhibited excellent oxidation resistance at 800[degrees]C in air with 10% water vapor (Figure 5a) lost the capacity to maintain Al^sub 2^O^sub 3^ scale formation after 500-1,500 h of total exposure at 900[degrees]C. Analysis of a local intact region of continuous Al^sub 2^O^sub 3^ indicated significant aluminum depletion in the alloy underneath the scale to a level less than 1 wt.% aluminum.6 The solubility of aluminum in austenite is on the order of -2 wt.% to 2.5 wt.% aluminum, with higher bulk alloy aluminum levels resulting in second- phase dispersions of (Fe,Ni)Al-type B2 phase. In the 3 wt.% and 4 wt.% aluminum-containing alloys, the B2 precipitates act as aluminum reservoirs for the scale,6 resulting in a B2-denuded zone directly underneath the scale (e.g., Figure 6). Within this B2-denuded region, the aluminum profile remained constant at the 2-2.5 wt.% Al range solubility limit of aluminum in austenite (i.e., no aluminum depletion was observed). Such a microstructure was observed, for example, in AFA 3-7 after 5,000 h at 900[degrees]C in air.

Figure 5b also illustrates the strong dependence of oxidation behavior on alloying additions other than aluminum.6 AFA 3-3 (3Al/ 0.6Nb/26Ni) was more oxidation resistant at 900[degrees]C in air than AFA 3-2 (3Al/0.6Nb/0.1Ti/20Ni), with comparable aluminum and niobium levels, but increased nickel at 26 wt.% vs. 20 wt.%. As at 800[degrees]C in air with 10% water vapor, niobium also played a strong role, with AFA 3-7 at 3Al and 1.5Nb exhibiting greater oxidation resistance than AFA 4-1, with 4Al but only 0.6Nb.

The niobium content also played a significant role in creep resistance. Figure 7 shows a compilation of creeprupture life for various alloys with 50-150 [mu]m grain sizes under a highload screening condition of 750[degrees]C and 170 MPa in air.7 (This condition was selected to provide rapid feedback for the effects of composition on creep rupture life behavior.) A maximum in creep rupture life was observed at ~1 wt.% niobium for AFA alloys, independent of the aluminum level in the 2.5-4 wt.% aluminum range. This latter observation suggests little contribution of the B2 precipitates to the creep strengthening at 750[degrees]C and 170 MPa. The higher niobium levels also resulted in significant quantities of Fe^sub 2^(Mo,Nb) type Laves phase precipitates, but did not correlate with improved creep resistance, consistent with the supposition that creep resistance primarily results from MC carbide formation.3 Preliminary computational thermodynamic calculations suggest that in the 4 wt.% aluminum alloys, MC carbide precipitation may reach an optimum at the 1.5 wt.% niobium level.7 Based on the trends observed, niobium ranges of -1-3 wt.% are of primary interest for applications requiring a balance of creep and oxidation resistance, with the higher niobium levels in this range favoring oxidation and the lower levels favoring creep resistance.

INITIAL SCALE-UP EFFORTS

The studies reported1-7 were made on small laboratory-scale arc- melted castings. Efforts are being initiated to determine if similar properties can be achieved in AFA alloys made by more commercially viable methods in larger scale castings. Results shown in Figures 8- 11 are from a vacuum-melted and hot-rolled 50 lb test heat of alloy AFA 4-1. Screening of tensile properties in the solution-treated condition indicated room-temperature elongation of greater than 50% and yield strength of -250 MPa (Figure 8). Tensile evaluations at 650[degrees]C, 750[degrees]C, and 850[degrees]C indicated a decrease in elongation with increasing temperature (Figure 8b). It is speculated that this ductility reduction results from precipitation of B2, Laves, and MC phases, although necking instability may also be playing a major role. Further study of this issue is planned.

Creep-rupture life measurements were made for the as- hot-rolled material (perpendicular and parallel to the rolling direction) and solution heattreated conditions, with and without cold work, under the accelerated test conditions of 170 MPa and 750[degrees]C in air (Figure 9). Creep-rupture lifetimes ranged from 130 h for as hot- rolled, oriented perpendicular to the rolling direction, to 350 h for solution-treated and cold-worked material, which is in the range observed for arc-cast material. Elongations were on the order of 10- 20%, despite the B2 and Laves phase precipitation. The microstructure after creep testing is shown in Figure 10 for the solution-treated and coldworked sample (Figure 9). The alloy grain boundaries were decorated with coarse B2 and Laves phase precipitates, with some primary undissolved MC also observed (Figure 10a). The intragrain regions contained a high density of submicrometer B2 and Laves phase precipitates (Figure 10b), for which TEM analysis indicated well-distributed, nanoscale MC carbide precipitates (Figure 10c).

The weldability of the AFA 4-1 plate was screened using gas tungsten arc welding with AFA 4-1 material as weld filler (Figure 11). Detailed microstructural analysis has not yet been completed, but preliminary evaluation indicated that no cracking had occurred. A similar weldability screening conducted for AFA 2-1(1) also showed no cracking after gas tungsten arc welding. Therefore, it appears that the AFA alloys may be amenable to conventional welding processes.

CONCLUSION

The AFA alloys show a promising combination of oxidation resistance, creep resistance, tensile properties, and potential for good welding behavior. The results to date suggest that the optimum alloy composition is in the range of Fe-(20-25)Ni-(12-15)Cr-(3- 4)Al(1-3)Nb wt. % base, with an upper usetemperature limit from an oxidation standpoint approaching ~900[degrees]C for alloy compositions at the higher aluminum, niobium, and nickel contents of the range. The upper temperature limit of ~900[degrees]C should siill permit widespread applicability of AFA alloys to various energy production and chemical and process industries, although the tendency for internal oxidation and nitridation of aluminum at higher temperatures could be an issue during alloy processing.

The properties were measured primarily on small laboratory-scale arcmelted castings using high-purity elemental feedstock: it remains to be demonstrated that similar properties can be obtained using commercial-scale processes and feedstock. The relatively small amount of material available necessitated the use of subsized specimens for screening of tensile and creep behavior trends as a function of alloy composition. Longer-term testing with standardized mechanical property specimen configurations are needed to confirm that oxidation and creep resistance can be maintained for the lifetimes (~ 10,000-100,000 h) required for many applications. Future work will pursue AFA alloy exposure in a range of sulfur- and carbon-containing high-temperature industrial environments to determine if the expected advantages of Al^sub 2^O^sub 3^ scale formation are conferred for these conditions. ACKNOWLEDGEMENTS

The authors thank J.H. Sehne ibel, R. Klueh, and I.G. Wright for helpful comments on this manuscript. This work was funded by the Fossil Energy Advanced Research Materials program. Additional funding and collaboration with the SHaRE User Facility at Oak Ridge National Laboratory is also acknowledged. Oak Ridge National Laboratory is managed by UT-Batteile, LLC for the U.S. Department of Energy (DOE) under contract DE-AC05000R22725. Notice: This submission was sponsored by a contractor of the U.S. government under contract DEAC05-000R22725 with the U.S. DOE. The U.S. government retains, and the publisher, by accepting this submission for publication, acknowledges that the U.S. government retains, a nonexclusive, paid-up, irrevocable, worldwide license to publish or reproduce the published form of this submission, or allow others to do so, for U.S. government purposes.

Author’s Note: Part of this research summary is based on a recent review paper (see Reference 1) and findings first reported in References 2-7.

How would you…

…describe the overall significance of this paper?

Alumina-forming austenitic stainless steels hold the potential to permit significantly increased operating temperatures in aggressive high-temperature oxidizing environments encountered in energy conversion and chemical process industry applications.

…describe this work to a materials science and engineering professional with no experience in your technical specialty?

We have discovered that creep-resistant, austenitic stainless steels capable of protective aluminum-oxide scale formation, rather than conventionally used chromium-oxide scales, are feasible. This results in significantly improved oxidation resistance and higher- temperature capability in many industrially relevant environments.

…describe this work to a layperson?

This paper describes the development of a new type of stainless sieel with improved corrosion resistance for high-temperature use in power generation. The corrosion protection is derived from aluminum additions to the alloy instead of typically used chromium additions.

References

1. M.R Brady et al., “Alumina-Forming Austenitics: A New Class of Heat-Resistant Stainless Steels,” Stainless Steel World Magazine (March 2008), pp. 23-29.

2.YYamamoto et al., Science, 316 (5823) (2007), pp. 433-436.

3. Y. Yimamoto et al., Met Mater. Trans. A, 38A (11) (2007). pp. 2737-2746.

4. MP Brady et al., Scripts Maler., 57 (12) (2007), pp. 1117- 1120.

5. Y. Yamamoto et al., Intermetallics, 16 (3) (2008), pp. 453- 462.

6. M.R Brady et al., Materials Science Forum 2008 (in press).

7. Y. Yamamoto et al., to be submitted to Acta Materialia.

8. G.Y. Lai, High Temperature Corrosion of Engineering Alloys (Materials Park, OH; ASM International, 1990).

9. B. Gleeson, Corrosion and Environmental Degradation, Volume II, ed. M. Schutze, Materials Science and Technology Series (Weinheim, Germany: Wiley-VCH, 2000), Chapter 5, pp. 173-228.

10. R Kofstad, editor, High Temperature Corrosion (London: Elsevier, 1988).

11. M.R Brady et al., Corrosion and Environmental Degradation, Volume II, ed. M. Schulze, Materials Science and Technology Series (Weinheim, Germany; Wiley-VCH, 2000). chapter 6, pp. 229-325.

12. J. Doychak, Intermetallic Compounds: Principles and Practice Vol. 1, eu. J.H. Westbrook and Rl. Fleischer (New York: John Wiley & Sons, 1994), pp. 977-1016.

13. G.H. Meier, Materials and Corrosion. 47 (11) (1996), pp. 595- 618.

14. G. Welsch et al., Oxidation and Corrosion of Intermetallic Alloys, ed. G. Welsch and RD. Desai (West Lafayette, IN: Purdue Research Foundation, 1996), pp. 121-266.

15. G.J. Yurek, Corrosion Mechanisms, ed. F. Mansfeld (New York: Marcel Dekker, Inc., 1987), pp. 398-446.

16. EJ. Opila. Mar. Sci. Forum, 461-464 (2004). pp. 765-773.

17. B.A. Pint, R. Peraldi. and RJ. Maziasz, Mat. Sci. Forum, 461- 464 (2004), p. 815

18. P.J. Maziasz et al., International Journal of Hydrogen Energy, 32 (16) (2007), pp. 3622-3630.

19. RG. Wilson, B.R. Knott, and C.D. Desforges, Mef. Mater. Trans. A. 9 (2) (1978), pp. 275-282.

20. T. Fujioka et al., U.S. patent 3,989,514 (1976).

21. J.A. McGurty, “Austenitic Iron Alloys,” U.S. patent 4,086.085 (25 April 1978).

22. J.C. Pivin et al., Con. Sa, 20 (1980), pp. 351-373.

23. V. Ramakrishnan. J. A. McGurty. and N. Jayaraman, Oxid. Met., 60 (1988), pp. 185-200.

24. FH. Stott, G.G. Wood, and J. Stringer, Oxid. Met., 44 (1-2) (1995), pp. 113-145.

25. C. Wagner, COTOS. Sci., 5 (1965), pp. 751-764.

26. P.J. Maziasz, JOM, 41 (7) (1989), pp. 14-20.

27. R.W. Swindeman et al., “Evolution of Advanced Austenitic Alloys Relative to Alloy Design Criteria for Steam Service: Part 1- Lean Stainless Steels,” Oak Ridge Natl. Lab. Rep. ORNL-6629/P1 (Oak Ridge, TN, 1990).

28. J.P. Shingledecker et al” Proc. ECCC Conference on Creep and Fracture in High Temperature Components-Design and Life Assessment Issues (Lancaster, PA: DEStech, 2005). pp. 99-109.

29. Allegheny Ludlum, TECHNICAL DATA BLUE SHEET, Stainless Steels, types 321, 347 and 348 (Pittsburgh, RA: ATI Allegheny Ludlum Corp., 2003), www.alleghenyludlum.com.

30. E. Essuman et al., Oxid. Met., 69 (3-4) (2008), pp. 143-162.

31. I. Kvernes, M. Oliveira, and R Kofstad, Corr. Sci., 17 (1997), pp. 237-252.

M.P. Brady, senior R&D staff member, M.L. Santelta, R&D staff member, P.J. Maziasz, distinguished research scientist, B.A. Pint, group leader and senior research staff, and H. Bei, R&D staff member, are with Oak Ridge National Laboratory, 1 Bethel Valley Road, Oak Ridge, TN 37831; Y. Yamamoto, research assistant professor, and C.T. Liu, distinguished research professor, are with the University of Tennessee; and Z.P. Lu, professor, Is with the University of Science and Technology in Beljing, China. Dr. Brady can be reached at bradymp@ ornl.gov.

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