International Journal of Powder Metallurgy Volume 46 Issue 6

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NOV.DEC.2010.IJPM cover_July_August IJPM cover 11/29/2010 10:28 AM Page 1

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November/December 2010

46/6

Newsmaker: Herbert Danninger, FAPMI Excellence in Metallography Award Gas-Atomized Iron-Base ODS Alloys Potential Effects of Retained Austenite on End-Quench Cooling Rates 440C Stainless Steels with Improved Hardness and Corrosion Resistance Vacuum Carburizing of Iron–Silicon Alloys

FRONT MATTER_ FRONT MATTER 11/29/2010 10:31 AM Page 15

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FRONT MATTER_ FRONT MATTER 11/29/2010 10:31 AM Page 1

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EDITORIAL REVIEW COMMITTEE P.W. Taubenblat, FAPMI, Chairman I.E. Anderson, FAPMI T. Ando S.G. Caldwell S.C. Deevi D. Dombrowski J.J. Dunkley Z. Fang B.L. Ferguson W. Frazier K. Kulkarni, FAPMI K.S. Kumar T.F. Murphy, FAPMI J.W. Newkirk P.D. Nurthen J.H. Perepezko P.K. Samal D.W. Smith, FAPMI R. Tandon T.A. Tomlin D.T. Whychell, Sr., FAPMI M. Wright, PMT A. Zavaliangos INTERNATIONAL LIAISON COMMITTEE D. Whittaker (UK) Chairman V. Arnhold (Germany) E.C. Barba (Mexico) P. Beiss, FAPMI (Germany) C. Blais (Canada) G.F. Bocchini (Italy) F. Chagnon (Canada) C-L Chu (Taiwan) O. Coube (Europe) H. Danninger, FAPMI (Austria) U. Engström (Sweden) O. Grinder (Sweden) S. Guo (China) F-L Han (China) K.S. Hwang (Taiwan) Y.D. Kim (Korea) G. L’Espérance, FAPMI (Canada) H. Miura (Japan) C.B. Molins (Spain) R.L. Orban (Romania) T.L. Pecanha (Brazil) F. Petzoldt (Germany) G.B. Schaffer (Australia) L. Sigl (Austria) Y. Takeda (Japan) G.S. Upadhyaya (India) Publisher C. James Trombino, CAE [email protected] Editor-in-Chief Alan Lawley, FAPMI [email protected] Managing Editor James P. Adams [email protected] Contributing Editor Peter K. Johnson [email protected] Advertising Manager Jessica S. Tamasi [email protected] Copy Editor Donni Magid [email protected] Production Assistant Dora Schember [email protected] Graphics Debby Stab [email protected] President of APMI International Dean Howard, PMT [email protected] Executive Director/CEO, APMI International C. James Trombino, CAE [email protected]

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46/6 November/December 2010

Editor’s Note Newsmaker ...Herbert Danninger, FAPMI PMT Spotlight On …Silvio Bartoletti Excellence in Metallography Award Consultants’ Corner John A. Shields, Jr.

RESEARCH & DEVELOPMENT 17 Microstructure Evolution of Gas-Atomized Iron-Base ODS Alloys J.R. Rieken, I.E. Anderson and M.J. Kramer

33 Potential Effects of Retained Austenite on End-Quench Cooling Rates in PM Steels F.J. Semel and D.A. Lados

ENGINEERING & TECHNOLOGY 43 As-Sintered AISI 440C Stainless Steels with Improved Hardness and Corrosion Resistance H. Ovri, C.J. Ohaukwu, K. Bahadirov, M. Larson and P. Kjeldsteen

51 Effect of Silicon on Vacuum-Carburizing Depth of Iron Compacts K. Widanka

56 59 60 61 62 64

DEPARTMENTS PM Industry News in Review Meetings and Conferences APMI Membership Application PM Bookshelf Table of Contents: Volume 46, Numbers 1–6, 2010 Advertisers’ Index Cover: As-HIPed microstructure of CR-144Hf-Y alloy showing residual porosity (dark spots). Photo courtesy Joel Rieken, Iowa State University.

The International Journal of Powder Metallurgy (ISSN No. 0888-7462) is a professional publication serving the scientific and technological needs and interests of the powder metallurgist and the metal powder producing and consuming industries. Advertising carried in the Journal is selected so as to meet these needs and interests. Unrelated advertising cannot be accepted. Published bimonthly by APMI International, 105 College Road East, Princeton, N.J. 08540-6692 USA. Telephone (609) 4527700. Periodical postage paid at Princeton, New Jersey, and at additional mailing offices. Copyright © 2010 by APMI International. Subscription rates to non-members; USA, Canada and Mexico: $100.00 individuals, $230.00 institutions; overseas: additional $40.00 postage; single issues $55.00. Printed in USA. Postmaster send address changes to the International Journal of Powder Metallurgy, 105 College Road East, Princeton, New Jersey 08540 USA USPS#267-120 ADVERTISING INFORMATION Jessica Tamasi, APMI International 105 College Road East, Princeton, New Jersey 08540-6692 USA Tel: (609) 452-7700 • Fax: (609) 987-8523 • E-mail: [email protected]

FRONT MATTER_ FRONT MATTER 11/29/2010 10:31 AM Page 2

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EDITOR’S NOTE

D

oes GEICO save you 15% on car insurance? Is metallography beneficial in identifying the microstructure and chemistry of PM materials? Take a moment to look at the 2010 Grand Prize entry in the Excellence in Metallography competition. Utilizing scanning electron microscopy, energy dispersive spectrography, auger electron spectrography, transmission electron microscopy, and X-ray diffraction, Rieken, Anderson, and Kramer probe the internal structure of gas-atomized oxide-dispersion-strengthened ferritic stainless steels. In combination, these metallographic techniques resulted in a complete and quantitative microstructural characterization of the PM steels down to the nanoscale level. In the “Newsmaker,” Peter Johnson profiles Herbert Danninger, a recipient of the 2010 APMI Fellow Award. A distinguished academic, Herbert has forged strong ties and interactions between the academic community and the PM industry in Europe. Returning to the “Consultants’ Corner,” John Shields provides detailed responses to two questions related to ferrous PM. The first elaborates on what is meant by “2 w/o max. other” in most MPIF materials, and the second discusses and compares intergranular embrittlement in wrought and PM steels of similar composition and hardness. In the “Research & Development” section, Rieken et al. detail the production of iron-base ODS alloys by means of gas-atomization reaction synthesis. After consolidation, a uniform distribution of nanoscale yttriumenriched oxide dispersoids is achieved throughout the microstructure. Also in this section, Semel and Lados employ process modeling simulations to predict the likely effect of porosity on retained austenite and end-quench cooling rates in PM steels. In the first of two contributions to the “Engineering & Technology” section, Ovri et al. describe the results of a study to develop as-sintered 440C stainless steels with enhanced hardness and corrosion resistance. The compositions and sintering temperatures are identified to give optimum combinations of the as-sintered properties. This section is completed with a paper by Widanka that quantifies the role of silicon on the vacuum carburizing of iron. Silicon additions up to 1 w/o increase the carburized depth by ~35% compared with iron in the absence of silicon.

Alan Lawley Editor-in-Chief

Subterfuge! That’s my take on the 2010 Nobel Prizes in Physics and Chemistry. While in no way diminishing the brilliance of the work resulting in each award, both clearly reflect achievements in Materials Science & Engineering utilizing basic tenets of Physics and Chemistry. The Physics Prize acknowledges two Russian-born scientists for creating graphene, a transparent two-dimensional form of carbon only one atom thick but more than 100 times stronger than steel, with unique properties: electrical conductivity similar to that of copper, and thermal conductivity better than that of other materials. Envisaged applications include computer chips, pollution monitoring, and improved flat-screen televisions. An American and two Japanese researchers received the Nobel Prize in Chemistry for their work in developing new and more efficient ways to link carbon atoms together in order to build complex molecules. This approach is a vital step in developing novel medicines, materials such as polymers and electronics. Both advances illustrate innovation manipulation of the relations between processing, internal architecture (microstructure), and properties, a cardinal tenet of Materials Science & Engineering.

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Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

FRONT MATTER_ FRONT MATTER 11/29/2010 10:31 AM Page 3

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Stain Free Solutions for PM Components The leading producer of metal powders, North American Höganäs, has introduced a series of products that address an increasingly problematic issue in component manufacturing: Stains. By the way, these products also facilitiate improved lubrication, enhanced machinability, increased productivity and scrap reduction.

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NEWSMAKER_ NEWSMAKER 11/29/2010 10:33 AM Page 4

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NEWSMAKER

HERBERT DANNINGER, FAPMI

By Peter K. Johnson*

Herbert Danninger, professor and managing director of the Institute of Chemical Technologies and Analytics, Vienna University of Technology (TUW), is a vocal advocate for stronger ties between the academic community and industry. In his teaching and research he has pursued the practical side of science to solve industrial problems. A recipient of the APMI International Fellow Award in 2010, he shares his views on PM technology issues. “My research work was and is carried out in part in close cooperation with industrial enterprises,” he says. These represent many aspects of PM processing, from leading metal powder makers and tool-steel producers to hardmetal suppliers and PM parts fabricators. He sees the need to convince PM parts customers that better properties cost more. “You cannot expect better performance with cheaper materials,” he stresses. Since receiving his Doctor of Technical Sciences (Dr. techn.) degree in 1980 from the TUW, Danninger has earned international recognition for his many accomplishments. His honors include the TUW Dr. Ernst Fehrer Prize for the development of liquid-phase-activated gravity sintering of stainless steels, the Skaupy Lecture award from the German Gemeinschaftsausschuss Pulvermetallurgie, and Doctor Honoris Causa of the Universitatea Tehnica din Cluj-Napoca in Romania. He has published more than 320 scientific papers in journals and conference proceedings and has given numerous oral presentations at international conferences, foreign universities, and institutes. Danninger’s interest in PM was cultivated while studying under Professor Gerhard Jangg, a prominent European powder metallurgist who had close links to PM companies. “Professor Jangg always *Contributing editor

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told us that you have to think using the brains of industry and anticipate the needs of companies before they recognize them,” he recalls. Danninger says he was also influenced by a book on refractory metals authored by Jangg, Richard Kieffer, and Peter Ettmayer. “Their style of writing was fascinating,” he says. “They described how temperamental these metals are.” He went on to write his doctoral thesis on “The Influence of Manufacturing Parameters on the Properties of Tungsten Heavy Alloys.” His teaching responsibilities encompass a basic class in inorganic chemical materials numbering 120–130 diploma engineering students, materials technology courses for Master’s degree students, and supervising Master’s and PhD thesis projects. About 40 percent of his time is devoted to teaching. A strong proponent of high-temperature sintering (>1,200ºC) since the early 1980s, Danninger says most European PM companies are devoting more development time to the process. His work points to the crucial importance of chemical reactions in the early stages of sintering by switching to oxygensensitive elements such as chromium, manganese, silicon, and vanadium. “These reactions remove surface oxides and, for standard constituents such as iron, molybdenum, and nickel, are ‘free of charge’ by occurring at temperatures below the standard sintering temperature. For chromium or manganese, the reduction temperatures are in the same order as belt-furnace sintering temperatures; therefore, high-temperature sintering is decidedly beneficial here,” he stresses. He refutes the adage that sintering destroys precision and cites the shrinkage controls used in metal injection molding (MIM). “You must choose sintering routes that minimize distortion,” he adds. ijpm

Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

NEWSMAKER_ NEWSMAKER 11/29/2010 10:33 AM Page 5

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NEWSMAKER: HERBERT DANNINGER

His sintering research has included liquid-phase sintering and related homogenization and pore-formation processes. Chemical reactions during sintering and heat-treatment processes are other areas of study. He has also devoted serious time to investigating the fatigue behavior of PM steels and tooling materials. His research with an industrial partner underscores the importance of fatigue testing with high loading cycles ranging from more than 100 × 106 cycles up to 109 cycles, he reports. “The old concepts of selling PM’s shape and precision are not

Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

sufficient,” he stresses. “You must offer increased performance, especially fatigue properties.” New alloys and the combination of high-temperature sintering and warm compaction can also offer improved properties. He believes that new material systems are needed because, in his opinion, the traditional iron–copper–carbon system has reached a plateau in performance. Danninger is confident about PM’s future and believes there is still much work to be done. “There is not a lack of ideas,” he says, “but rather a lack of time and funds.” ijpm

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NEWSMAKER_ NEWSMAKER 11/29/2010 10:33 AM Page 6

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SPOTLIGHT ON_ SPOTLIGHT ON 11/29/2010 10:34 AM Page 7

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SPOTLIGHT ON ...

SILVIO BARTOLETTI, PMT Education: Degrees in Industrial Chemistry, Colégio Oswaldo Cruz, São Paulo, 1965; Mechanical Engineering, Universidade Braz Cubas, São Paulo, 1974; Production Administration, Fundação Getulio Vargas, São Paulo, 1998 Why did you study powder metallurgy/particulate materials? My first contact with the powder metallurgy (PM) was in mid-1970 when I worked for a large householdappliance business; until then I knew nothing about PM. What caught my attention were self-lubricating bushings that were, and are still, used to replace bearings requiring lubrication. When I started my activities at Metalpó, I began to understand how a self-lubricating bushing works and why it is self-lubricating. One of the main reasons that led me to study PM technology more deeply was the diversity of applications of products made from metal powders, and the wide spectrum of materials available in the marketplace targeted at various applications, depending on the mechanical and metallurgical requirements. What was your first job in PM? What did you do? The Metalpó/Combustol group is made up of companies that focus on equipment for steel mills, petrochemicals, and organic-waste processing, among others. The group also has a refractories department, heat treatment, and PM. I started my PM career in 1974 working primarily in setting up the quality control department. After a brief stay I moved to the plant that made electrolytic copper powder. Here I was creating copper powders with densities between 1.0 and 2.7 g/cm3. At that time, all the bronze parts manufactured by Metalpó used electrolytic copper powder as the raw material. Besides copper powder, I developed red copper oxide (Cu2O), used as a raw material in manufacturing paints for seagoing vessels, and black copper oxide (CuO) for agricultural and other applications.

Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

Manufacturing tin powder by air atomization was also one of the improved techniques for obtaining density and size distributions within more accurate specifications. Controlling metal powder manufacturing process parameters, such as hydrogen loss (oxygen content), particle size, specific surface area, and average particle diameter, were also a part of my early work in PM. Describe your career path and companies worked for, and responsibilities. I started my career in industry in 1965 after I received my degree in industrial chemistry, working for a company in the Pirelli Group for four years. My first function was as a production scheduler, after which I was promoted to work in a testing laboratory for electric cables. In 1969 I transferred to ARNO S/A, a large company that is now part of the French group SEB. The company dealt in automotive products for the Delco Remy line, home appliances, and industrial motors. I served as an appliance assembly-line coordinator, and it was here that I had my first encounter with PM parts used in motors for household fans, food blenders, mixers, and other small appliances. In 1974 I began my career at Metalpó. My first job was setting up the quality department. In 1978 I took part in the building of a facility for water atomization of nonferrous powders, a technology acquired from a French company. I managed the factory until 2004. In 1986 I took over the management of Metalpó’s industrial area which encompassed the production of parts, Manager Metalpó Industria e Comercio Ltda. Rua Cel. Jose Rufino Freire 453-A Pirituba São Paulo, SP CEP 05159-900 Brazil Phone: 55-11-39047690 Fax: 55-11-39063173 E-mail: [email protected]

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SPOTLIGHT ON ...SILVIO BARTOLETTI, PMT

metal powders, tooling, and maintenance. I carried out this function until 2008. Today I am an engineering manager with primary responsibilities in new product development, the application of new technologies, ISO-TS specifications, tooling design, and tool fabrication. What gives you the most satisfaction in your career? Without doubt, it is the variety of tasks and challenges throughout my career and constant technological improvement due to knowledge of new materials that keep me motivated in the performance of my job. It is difficult to highlight any one reason for my satisfaction, but the details of making powder via electrolysis and water or air atomization have been a highlight, since they occurred at the start of my technical training in PM. What major changes/trend(s) in the PM industry have you seen? PM has developed significantly and is increasingly expanding its markets, especially in the automotive industry. I have witnessed the development of new materials and processes to achieve high densities and enhanced mechanical properties with increasing dimensional stability. Some of the materials and processes are high-compressibility and high-hardenability powders, special lubricants, warm die compaction, and sinter hardening. Another important segment of PM is that of soft magnetic composite materials for electromagnetic applications. They provide important benefits in reducing engine weight and cost, while maintaining strength due to new design concepts. These are therefore sustainable solutions from an ecological point of view, which could enable a change from the internal combustion engine to electric or hybrid vehicles. My team and I have developed a product with this material that is being used in fuel injectors. Metal injection molding is another trend, com-

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bining features of PM with plastic injection molding, capable of producing parts of complex geometry with the benefits of PM, namely, precision, near-net shaping, high productivity, and low cost. Why did you choose to pursue PMT certification? It is difficult for anyone to do a self-evaluation of our knowledge of PM. I believe the PMT examination embraces questions that reflect the knowledge that professionals must have. I find the diversity of technologies covered in the survey interesting, dealing with raw materials, applications, technologies, designs, and so on. I was honored and pleased when I received the news of my certification, which motivated me to continue to keep abreast of new materials and applications. How have you benefited from PMT certification in your career? Recognition by Metalpó senior management of my accomplishment was immediate. I received compliments, which keep me motivated to continue to improve technically. I hope to continue my career, and whenever possible, attend national and international events related to PM. What are your current interests, hobbies, and activities outside of work? I like to entertain relatives and friends at my house on weekends. One of my pastimes is to be in the kitchen preparing my favorite Italian dishes. I enjoy the beach, and I go whenever I have the opportunity. Once in a while I visit a home for the elderly or for children and provide support. As we are in a country where football (soccer) is a major sport, and as I am no different from most Brazilians, I like a good football game, occasionally live at a stadium but most often at home on TV. ijpm

Would you like to be featured here? Have you been PMT Certified for more than 2 years? Contact Dora Schember ([email protected]) for more information.

Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

Metallography Award_Zheng et al 11/29/2010 10:36 AM Page 9

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2010 EXCELLENCE IN METALLOGRAPHY AWARD

GRAND PRIZE

GAS-ATOMIZED CHEMICAL RESERVOIR ODS FERRITIC STAINLESS STEELS Joel R. Rieken*, Iver E. Anderson, FAPMI**, and Matthew J. Kramer***

ABSTRACT Gas-atomization reaction synthesis was used to surface oxidize ferritic stainless steel powders resulting in an ultrathin (<100 nm) metastable chromium-enriched oxide shell. Subsequently, the shell was dissociated and served as an oxygen reservoir for the formation of uniformly dispersed nanoscale Y-(Ti, Hf) oxide precipitates during heat treatment of the as-consolidated powders. Powders in the as-atomized and heat-treated conditions were characterized and compared using several electron microscopy techniques. Scanning (SEM), auger (AES), and transmission (TEM) electron microscopy were used in conjunction with energy dispersive spectroscopy (EDS) and X-ray powder diffraction (XRD) to effectively characterize the oxide layers and the size, shape, location, and chemical composition of the precipitates formed during heat treatment. The goal of these analyses was to determine the composition of the oxide shell and to confirm the exchange of oxygen from the shell to a distribution of finely spaced and highly stable nanoscale oxide particles. Selected examples of the use of these techniques are shown.

In its second year, this APMI award recognizes the individual(s) responsible for metallography used to support and provide evidence for the ideas set forth in a conference technical paper. A panel of judges evaluated all eligible manuscripts from PowderMet2010. An award presentation will be made at PowderMet2011 in San Francisco, California, May 18–21, 2011. The paper is published here in an abridged form, as edited by Thomas F. Murphy, FAPMI, Hoeganaes Corporation, the Metallography Competition Chairman. The unabridged version of the paper is published in the conference proceedings, Advances in Powder Metallurgy & Particulate Materials— 2010, available from the publications department of the Metal Powder Industries Federation, www.mpif.org.

Metallographic Analysis Chemical compositions of the bulk powders are shown in Table I. The three alloys (top to bottom in Table 1), will be referred to as A, B, and C. The relative sizes and spheroidal shape of the as-atomized particles are seen in Figure 1 along with the chemical compositions of the surface shells measured at increasing depth from the particle surface using AES. TABLE I. CHEMICAL COMPOSITION OF AS-ATOMIZED ALLOYS Alloy

Fe (a/o)

Cr (a/o)

W (a/o)

Ti (a/o)

Hf (a/o)

Y (a/o)

O (a/o)

CR-118Ti-Y CR-144Hf-Y CR-156Y-Hf

Bal. Bal. Bal.

15.84 16.16 15.84

0.94 -

0.50 -

0.27 0.11

0.20 0.08 0.18

1.67 0.23 0.38

*Student, Department of Materials Science and Engineering, Iowa State University, Ames, Iowa 50011; E-mail: [email protected], **Senior Metallurgist, ***Senior Scientist, Division of Materials Sciences and Engineering, Ames Laboratory, USDOE, Ames, Iowa 50011 Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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Metallography Award_Zheng et al 11/29/2010 10:36 AM Page 10

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2010 EXCELLENCE IN METALLOGRAPHY AWARD

SEM images of cross sections of the three alloy powders mixed with 75 v/o Cu powder and cold isostatically pressed (CIPed) are shown in Figure 2. The alloy particles are the round dark-gray particles in the matrix of lighter-gray copper particles. Initial evidence of phase precipitation is visible as a light-gray network in a few of the alloy particle cross sections.

Additionally, the three alloy powders were consolidated by hot isostatic pressing (HIP) and heat treated for formation of the stable Y-(Ti,Hf) oxides. Examples of the as-HIPed microstructures of each alloy are shown in Figures 3 and 4. In Figure 3, the as-HIPed microstructures of the three alloys are compared using SEM imaging. Grain boundaries, alloy segregation, and residual porosity are appar-

Figure 1. AES compositional-depth profiles and SEM images of as-atomized particles. A1 & A2 = CR118, B1 & B2 = CR144, C1 & C2 = CR156

Figure 2. As-atomized and CIPed cross sections of alloy powders. SEM

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Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

Metallography Award_Zheng et al 11/29/2010 10:37 AM Page 11

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2010 EXCELLENCE IN METALLOGRAPHY AWARD

ent in the micrographs as variations in gray scale. The TEM images in Figure 4 show the locations and distribution of the nanoscale oxides throughout the same set of HIPed alloys shown in Figure 3. The number and location of the small dark fea-

tures appear to change from sparce and uniform to more frequent and located along or near grain boundaries. Further examples of the effectiveness of metallographic analysis in characterizing the alloys are

Figure 3. SEM images of the three alloys after HIPing (700°C/300 MPa/4 h)

Figure 4. TEM brightfield images of the three alloys after HIPing (700°C/300 MPa/4 h) Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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Metallography Award_Zheng et al 11/29/2010 10:37 AM Page 12

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2010 EXCELLENCE IN METALLOGRAPHY AWARD

displayed in Figures 5 and 6. SEM images of the heat-treated alloys (1,300°C for 1 h) are compared in Figure 5. Residual precipitates are seen as either dark (titanium-enriched oxides), almost pinpoint features in the A alloy, to whitish, highaspect-ratio features (iron–hafnium intermetallic particles) in alloys B and C. A heat-treated sample (1,200°C for 2.5 h) of alloy B, CR-144Hf-Y, was used to demonstrate chemical analysis mapping using energy-filtered transmis-

sion electron microscopy (EFTEM), Figure 6. SUMMARY Examples of the metallographic techniques used to characterize these alloys have been presented. They demonstrate the effectiveness of metallography as a means of providing the information necessary to make sound and informed judgments about the alloys and their viability to perform in the intended applications. ijpm

Figure 5. SEM images of the three alloys after heat treatment at 1,300°C for 1 h

Figure 6. EFTEM images of alloy B heat treated at 1,200°C for 2.5 h

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Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

CONSULTANTS' CORNER_ CONSULTANTS' CORNER 11/29/2010 10:38 AM Page 13

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CONSULTANTS’ CORNER

JOHN A. SHIELDS, JR.* Q

Most MPIF materials allow “2 w/o max. others.” Why is there no guide as to what elements should be checked? It seems that laboratories will randomly choose elements for “w/o others.” No two laboratories give the same level for “w/o others” on a given sample. The 2% maximum of other elements, including “other minor alloying elements added for specific purposes” in the majority of material specifications in MPIF Standard 35, is a long-standing practice, one also implemented in ASTM B783, ISO 5755, JIS Z 2550, and others. Material standards typical of those for traditional ingot materials use the “other elements” category to establish limits on impurities. However, Standard 35 recognizes that powder metallurgy (PM) parts makers do not use identical design and manufacturing practices, raw materials, or equipment. Often, by exploiting unique designs, processes, and proprietary compositional modifications, parts makers carve their niche. Recognizing this aspect of PM parts manufacture in the compositional specification allows parts makers to meet specification requirements for dimensional, mechanical, and other properties by employing minor compositional adjustments to the basic alloy composition. Properly exploited, this approach provides value for the user. It also means the part designer has a responsibility to make appropriate material choices, including modifications to the specification requirements if necessary. Every specification in Standard 35 contains the following note at the bottom of the first page, immediately before the table of minimum and typical property values: “To select a material optimum in both properties and cost-effectiveness, it is essential that the part application be discussed with the PM parts manufacturer… Both the purchaser and the manufacturer should, in order to avoid possible misconceptions or misunderstandings, agree on the following conditions prior to the manufacturer of a PM part: minimum strength value, grade selection, chemical composition, proof testing, typ-

A

ical property values and processes that may affect the part application.” The designer and parts maker must work together during the design and quoting phases of a project to define which elements to measure during production. Once the design is fixed and production begins, the problem of monitoring chemistry falls to the analytical laboratory, whether it is captive or a contract laboratory. Here again, all parties must understand and agree upon the elements to be analyzed and the methods used for analysis. Often a laboratory offers a standard analysis for a particular class of materials, for example, alloy steels. In such an analysis, the laboratory analyzes a particular suite of alloy and impurity elements and reports those compositions. Oftentimes, this analysis is based on the needs of traditional ingot metallurgy materials, not PM materials. There is also no guarantee that all laboratories will use the same list of elements in a “standard” analysis, which can lead to misunderstanding and disagreement between different laboratories. The parts maker and parts designer must agree on the elements to be analyzed so these can be communicated to the analytical laboratory. If one wants to look at a broad spectrum of potential impurities, I have found it useful and cost effective to request what some laboratories call a “qual-quant” analysis. Advances in analytical instruments allow for rapid determination of optical and elemental spectra, permitting this approach. In a “qual-quant” analysis, the laboratory performs a preliminary qualitative analysis to determine which elements are present in the sample at levels above the instrument’s detection limit. These qualitative analyses can be obtained inexpensively for as many as 68 different potential impurities. The sample is then re-analyzed to quantify the specific elements present at levels above the detection limit. Typically, laboratories charge for this analysis on a per-element basis, up to a maximum cost.

*Principal, PentaMet Associates LLC, 4457 Brooks Road, Cleveland, OH 44105–6053, USA; E-mail: [email protected]

Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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CONSULTANTS' CORNER_ CONSULTANTS' CORNER 11/29/2010 10:38 AM Page 14

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CONSULTANTS’ CORNER

Instruments capable of “qual-quant” analysis include arc/spark optical emission, induction coupled plasma optical emission (ICP-OE), ICP mass spectroscopy (ICPMS), and atomic absorption graphite furnace (AAGF). This approach is useful when one wants to know the full array of intentional and unintentional alloying elements present in the material. This information can be used to choose the elements analyzed on a production basis, or it can provide a periodic check on production materials or parts to assure that unwanted contaminants have not crept in. Between-laboratory disagreements about total impurity content can arise from several sources. Differences in the suite of elements analyzed, and analytical techniques used to measure impurities, can contribute to disagreements. Even if two laboratories use the same analytical technique, variability caused by different material samples, instruments, operators, preparation techniques, etc., will all cause discrepancies, even when measuring the same impurity. If the laboratories employ different analytical techniques, an additional level of uncertainty is present. When comparing results between laboratories, it is important to recognize these statistical variances and allow for them. As always, the devil is in the details. Significant effort may be needed to understand such discrepancies, so it is important to focus only on those outside the normal statistical variability expected of the material and analytical technique. When producer and user disagree on analytical results, it is customary to engage a referee laboratory to arbitrate. These problems are not inevitable, but avoiding them requires forethought, good communication, and cooperation among users, parts makers, and analytical laboratories. The compositional requirements in MPIF Standard 35 are only the starting point. They should be augmented for each part as appropriate by defining limits for specific elements and incorporating these requirements on part drawings. The analytical laboratory (or laboratories) should also be part of the discussion, to be sure that the proper elements are quantified with the desired techniques. Laying the groundwork before launching PM parts production will minimize problems.

Q

Is there any reason why low-alloy ferrous PM materials hardened by conventional quenching/tempering or sinter hardening would be more susceptible to intergranular embrittlement than wrought steels of similar composition/hardness? Intergranular embrittlement in alloy steels takes two forms. 1 One is temper embrittlement (TE), which affects tempered alloy steels cooled slowly on

A 14

tempering through, or operating in, the 300°C–600°C (570°F–1,110°F) temperature range. TE can occur in heavy sections that experience slow cooling at their core during heat treatment. It has caused failure in power plant and chemical processing components during service in this range. TE does not occur in normalized or annealed alloy steels. Temper embrittlement raises the alloy’s ductile–brittle transition temperature (DBTT) and causes a shift in fracture mode from cleavage to intergranular (IG) fracture. Minor impurity elements from Groups IVA (Sn) and VA (Sb, As, P) cause TE by segregating to prior austenite grain boundaries in the quenched and tempered microstructure during the time the steel is in the susceptible temperature range. This segregation weakens grain boundaries, creating an intergranular path for crack propagation. Other elements affect an alloy’s susceptibility to TE. Manganese and silicon promote TE, while molybdenum in solid solution reduces it. Chromium and nickel, when present together, render an alloy more susceptible to TE than it is when either element is present alone. Susceptibility increases with increasing strength of the quenched and tempered alloy, driven by the increased discrepancy between the toughness of grain boundaries and that of the tempered martensitic microstructure. A second form of embrittlement, tempered martensite embrittlement, or alternatively 350°C (500°F) embrittlement (TME), occurs after tempering between 200°C and 400°C (390°F and 750°F)2. TME is less well understood, but is related to cementite (Fe3C) formation at prior austenite grain boundaries during tempering, with associated transformation of retained austenite to untempered martensite. These microstructural changes create low-toughness fracture paths at the prior austenite grain boundaries. Steels exposed to temperatures in the TME range can display intergranular, cleavage, or ductile fracture surfaces. Intergranular fracture seems to be associated with the presence of phosphorus.2 Both of these embrittlement phenomena can produce intergranular fracture in quenched-and-tempered steels. In both cases the appearance of the problem is related to minor impurity elements in the alloy. The question refers to steels of similar composition and hardness. This wording leaves a great deal of wiggle room, and invites the classic metallurgist’s response: “It all depends…” For a similar composition, accounting for not only major alloying elements but also minor elements that affect either of the embrittling phenomena, I would expect the two types of steels to behave in similar fashion. The PM microstructure seems not to complicate the situation to any degree. The porosity in PM materials Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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CONSULTANTS’ CORNER

tends to force crack propagation through the neck regions, which fail in a ductile manner. When density approaches the pore-free level, some intergranular fracture can be observed in low-alloy steels heat treated in the TME range. 3 In the study cited, detailed impurity analyses were not reported, so we cannot say unequivocally that the IG fracture was related to phosphorous contamination. Low-alloy steels with admixtures of nickel have not been reported to show IG fracture, though different observations have been reported about whether fatigue cracks tend to propagate through or deviate around nickel-rich phases.4 It appears that quenched-and-tempered PM lowalloy steels are no more sensitive to IG failure than conventional wrought quenched-and-tempered lowalloy steels, for the same levels of embrittling elements, strength, and thermal exposure to temperatures in the TE or TME range. Because of the potential for embrittlement, proper specification of the impurity content is critical. Fortunately, these are well-known and -understood phenomena, and there are established guidelines for maximum allowable levels of some elements, for example, arsenic, antimony, and phosphorous, to avoid them. As noted in the previous answer, it is important for material specifiers, parts makers, and

Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

analytical laboratories to cooperate closely, both to avoid manufacturing problems and to assure the materials do not contain impurities that promote intergranular fracture. 1. Metals Handbook Desk Edition, Second Edition, 1998, ASM International, Materials Park, OH, in ASM Desk Editions on the Web for Members Only (www.asminternational.org), ASM International (2002). 2. G. Krauss, Steels: Processing, Structure, and Performance, 2005, ASM International, Materials Park, OH, pp. 396–409. 3. I.W. Donaldson and M.L. Marucci, “Heat-Treat Properties of High-Density FLN2-4405,” Advances in Powder Metallurgy & Particulate Materials—2004, compiled by W.B. James and R.A. Chernenkoff, Metal Powder Industries Federation, Princeton, NJ, 2004, pp. 111–121. 4. D.A. Lados and D. Apelian, “Effects of Porosity and Microstructure on the Fatigue Crack Growth Behavior of Pre-alloyed and Admixed Fe-Ni-Mo PM Alloys,” Advances in Powder Metallurgy and Particulate Materials—2006, compiled by W.R. Gasbarre and J.W. von Arx, Metal Powder Industries Federation, Princeton, NJ, 2006, pp. 50–64. ijpm

Readers are invited to send in questions for future issues. Submit your questions to: Consultants’ Corner, APMI International, 105 College Road East, Princeton, NJ 08540-6692; Fax (609) 987-8523; E-mail: [email protected]

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AMETEK

AMETEK

AMETEK’s

AMETEK



















WWWAMETEKMETALSCOM

www.ametekmetals.com

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RESEARCH & DEVELOPMENT

MICROSTRUCTURE EVOLUTION OF GAS-ATOMIZED IRON-BASE ODS ALLOYS

Joel R. Rieken*, Iver E. Anderson, FAPMI**, and Matthew J. Kramer**

INTRODUCTION Oxide dispersion strengthening (ODS) has been used to enhance the elevated-temperature strength and creep resistance of ferritic stainless steels.1–3 Based on the extensive development of theses ODS ferritic stainless steels, it was concluded that the iron–chromium base system required the addition of yttrium to form dispersoids with high-temperature stability and a titanium addition to further refine the resulting oxide-dispersoid size.4–12 A tungsten addition was usually included as a solid-solution strengthener for the iron–chromium matrix phase.5 These iron-base alloys are being considered as structural materials for use in future-generation power reactors.10,13–15 However, a major drawback to the widespread use of iron-base ODS steels stems from the high cost associated with finished forms of these alloys (~US$340/kg for PM2000TM).16 Traditionally, ODS alloys are produced using an inefficient, time-intensive, high-energy mixing process known as mechanical alloying (MA).17–20 The goal of the current study is to demonstrate the feasibility of a new, simplified, and more-efficient powder-processing method for the direct production of precursor ferritic stainless steel powders with similar alloy compositions that, when consolidated, undergo microstructural transformations leading to the formation of an improved isotropic ODS microstructure. This new processing concept would allow for full consolidation of precursor oxide-dispersion-forming powders from the as-atomized state, thus eliminating the mechanicalalloying step from the ODS processing sequence. BACKGROUND This new simplified process utilizes a unique alloying method termed GARS, a rapid-solidification process for the production of precursor ferritic stainless steel powder particles.21,22 The powder particles are surface oxidized in situ during GARS and solidify with an ultrathin (<150 nm) surface oxide layer. The surface oxide layer forms as a metastable oxide phase which can then be dissociated during high-temperature consolidation of the powder. During consolidation, the metastable surface oxide layer is contained in the consolidated microstructure and

In a simplified process to produce precursor powders for oxide dispersion-strengthened (ODS) alloys, gas-atomization reaction synthesis (GARS) was used to induce a surface oxide layer on molten droplets of three differing erritic stainless steel alloys during break-up and rapid solidification. The chemistry of the surface oxide was identified using auger electron spectroscopy (AES) and scanning electron microscopy (SEM) with energy dispersive spectroscopy (EDS). The precursor iron-base powders were consolidated at 850°C and 1,300°C using hot isostatic pressing (HIPing). Consolidation at the lower temperature resulted in a fully dense microstructure, while preventing substantial prior-particle-boundary-oxide dissociation. Microstructural analysis of the alloys consolidated at the higher temperature confirmed a significant reduction in prior-particleboundary-oxide volume fraction, in comparison with the lower-temperature-consolidated sample. This provided evidence that a high-temperature internal oxygenexchange reaction occurred between the metastable prior-particle-boundary-oxide phase (chromium oxide) and the yttrium contained within each prior particle. This internal oxygen-exchange reaction is shown to result in the formation of yttrium-enriched oxide dispersoids throughout the alloy microstructure. The evolving microstructure was characterized using transmission electron microscopy (TEM) and high-energy X-ray diffraction (HE-XRD).

*Student, Department of Materials Science and Engineering, Iowa State University, Ames, Iowa 50011; E-mail: [email protected], **Senior Metallurgist, Senior Scientist, Division of Materials Sciences and Engineering, Ames Laboratory, USDOE, Ames, Iowa 50011

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resides along prior particle boundaries (PPBs). As the high-temperature-consolidation process continues, the PPB oxide phase dissociate, releasing oxygen for the formation of stable yttrium-containing nanosize oxide dispersoids. EXPERIMENTAL METHODS GARS The reactive-atomization gas contained a mixture of argon with small concentrations of oxygen for production of this experimental powder.23 The reactive-gas content and composition for each alloy are displayed in Table I. These chemical reservoir (CR) alloys were designed to exhibit a ferritic iron (α-Fe) matrix. The atomization charge was super heated to 1,700°C within an yttria-washed-zirconia crucible. On reaching the pouring temperature, the melt was tapped and pouring commenced through an yttria-stabilized-zirconia pour tube. Upon exiting the pour tube, the melt was immediately impinged upon by the reactive-atomization gas (Ar-(0.25 or 0.5)O2 v/o). The gas pressure was 6.9 MPa at the atomization nozzle manifold. The bulk oxygen content in the as-atomized powder particles was measured using an inert-gas-fusion (LECO) analyzer, and the composition of each alloy was verified using inductively coupled plasma/atomic emission spectroscopy (ICP/AES). Consolidation Cold isostatic pressing (CIP) was used to consolidate as-atomized powder samples for metallographic preparation prior to hot consolidation. Asatomized particles (20–53 µm) were blended with 70 v/o copper powders (–20 µm) and sealed in latex CIP bags. The blended powders were subsequently CIPed at 60 ksi for ~60 s. Following the CIP process each consolidated powder sample was impregnated with epoxy and cross sectioned for microstructural analysis. Powder samples screen classified to 20–53 µm dia. were consolidated to full density using hot isoTABLE I. AS-ATOMIZED POWDER COMPOSITION AND REACTIVEATOMIZATION GAS COMPOSITION FOR ALLOYS Alloy

Fe (a/o)

Cr (a/o)

Y (a/o)

Ti (a/o)

W (a/o)

O (a/o)

Gas (v/o)

CR-112 CR-118 CR-126

Bal. Bal. Bal.

15.53 15.70 15.13

0.09 0.20 0.09

0.49 0.56

0.9

1.10 1.12 0.56

Ar-0.5O2 Ar-0.5O2 Ar-0.25O2

18

static pressing (HIP). Prior to consolidation, each powder sample was inserted into a 316L stainless steel HIP can, which measured 25.4 mm dia. × 127 mm in length. Each HIP can was slowly evacuated using a diffusion-pumped vacuum system to a pressure of ~7 × 10-7 mbar and transported under vacuum to a laser welding apparatus, where the can was evacuated further using a turbo-molecular pump to a pressure of ~7 × 10-8 mbar and welded shut. Each sealed can was HIPed at 850°C or 1,300°C at a pressure of 303 MPa for a duration of 4.0 h at the selected temperature and peak pressure. Electron Microscopy Surface analysis of the as-atomized powder particles was conducted using a JEOL JAMP 7830F SEM with AES. The microstructures of the powders were examined using a JEOL 5910LV SEM with EDS. The composition of the PPB oxide phase was evaluated using a JEOL JXA-8200 WD/ED microanalyzer, and the HIP-consolidated alloy microstructure was analyzed using a Hitachi S2460N SEM with EDS. The nanoscale features present in the HIP-consolidated samples were characterized using a Tecnai G2 F20 TEM under bright field imaging (BFI) conditions at 200 keV. The TEM samples were ground flat using 400 and 600 SiC grit paper, and polished using 6.0 μm and 1.0 μm diamond polishing compound to a thickness of ~50 μm. The samples were then mechanically dimpled to a thickness of ~20 μm, and duel-jet polished using an electrolytic solution for stainless steels (700 ml methanol, 100 ml 2-butoxyethanol, and 35 ml perchloric acid) at –21°C. Synchrotron Phase Analysis HE-XRD phase analysis was conducted on the as-atomized powders and HIP-consolidated specimens at the Advanced Photon Source (APS), beamline 11-BM, located at Argonne National Laboratory–USDOE. The as-atomized powder samples contained particles sieved to 20–53 µm dia. The HIP-consolidated APS specimens were electrical-discharge machined into 0.8 mm dia. × 15.0 mm miniature rods. Diffraction data was collected using a continuous scan at room temperature covering 34 degrees with a scan speed of 0.01 degree/s and a wavelength of 0.045869 nm. A detailed description of the beamline 11-BM instrument can be found in the literature.24 Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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EXPERIMENTAL RESULTS CR Alloy Compositions The yttrium and oxygen concentrations are considered to be most integral in the CR alloys. CR112 contained 0.09 a/o Y and 1.10 a/o O. Yttrium concentration levels were approximately doubled in alloy CR-118, while maintaining a similar oxygen concentration compared with CR-112 (Table I). This similar oxygen concentration was achieved by using the same reactive atomization gas mixture (Ar-0.5 v/o O2). The yttrium concentration in CR126 was again lowered to 0.09 a/o, while the oxygen concentration was reduced to 0.56 a/o by halving the oxygen content in the reactive atomizing gas (Table I). Microstructure: Alloy CR-112 AES depth profiling of as-atomized CR-112 powders revealed significant enrichment in oxygen and chromium with a slight enrichment in yttrium found at the particle surfaces (Figure 1(a)). The average thickness of the resulting chromiumenriched oxide layer was 110±10 nm (approximated from the etching rate of SiO2). SEM/EDS was used to analyze the interface of the as-CIPed CR-112 powders in polished cross sections. The results verified increased concentrations of chromium and oxygen and decreased concentrations of iron at particle/particle interfaces (Figure 1(b)), but no enrichment in yttrium was detected, due to its lower concentration. Electron probe microanalysis (EPMA) with quantitative wavelength dispersive spectroscopy (WDS) was used to evaluate the chemistry of the PPB

oxides found within the HIP consolidated CR-112 microstructure. These results show that the lowtemperature HIP-consolidated microstructure contains a non-stoichiometric chromium-enriched oxide phase along the PPBs. The composition of this PPB oxide phase shifted towards stoichiometric Cr2O3 after high-temperature HIP consolidation. The level of the chromium-enriched PPB oxide phase was reduced from 6.0±0.1 v/o in the 850°C as-HIPed microstructure to 4.1±0.2 v/o in the 1,300°C as-HIPed microstructure, reflecting progress in the PPB oxide dissociation. Volume percent measurements were completed using quantitative image analysis. Examples of images in cross sections are shown in Figures 2(a) and 2(c). TEM was utilized to evaluate the nanoscale phases found in the low- and high-temperature HIP-consolidated samples. TEM micrographs revealed the formation of nanoscale oxide dispersoids throughout the microstructure in both the low- and high-temperature HIP-consolidated samples (Figures 3(a), 3(b), 3(c), and 3(d)). The size of the oxide dispersoids ranged from 3 to 30 nm in the high-temperature-consolidated sample with a number density on the order of 1021m-3 (assuming a TEM foil thickness of 100 nm). The oxide particles were not able to be identified by electron diffraction due to their small size, but subsequent HE-XRD results were successful. Microstructure: Alloy CR-118 AES depth profiling of the as-atomized CR-118 powders revealed an enrichment of oxygen, chromium, and titanium (very slight) at the sur-

Figure 1. (a) AES depth profile and (b) SEM cross-sectional analysis with EDS chemical line scans for as-atomized and as-CIPed CR-112 powders, respectively Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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Figure 2. Microstructural comparison of low-temperature as-HIPed CR-112 (outlined in green: SEM images (a) and (b)), and high-temperature as-HIPed CR-112 (outlined in blue: SEM images (c) and (d))

face of the particles (Figure 4(a)). This measurement revealed the average thickness of the oxide scale to be ~125±10 nm (approximated from the etching rate of SiO2). SEM with EDS analysis of asCIPed CR-118 powders (polished cross section) showed increased concentrations of titanium and oxygen and a decreased concentration of iron at particle/particle interfaces (Figure 4(b)), without any detectable increase in chromium. EPMA results showed the PPB oxide phase in the low-temperature as-HIPed microstructure to be titanium and chromium enriched, although no specific stoichiometric phase was identified. The composition of the PPB oxides in the high-temperature as-HIPed microstructure was shown to contain varying amounts of titanium with oxygen. The volume percent of the PPB oxide phase was decreased, as expected, from 5.9±0.2 v/o in the low-temperature as-HIPed microstructure to 2.5±0.2 v/o in the high-temperature as-HIPed microstructure (Figures 5(a) and Figure 5(c)).

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TEM analysis of the low-temperature as-HIPed sample showed very few yttrium-containing oxide dispersoids (Figures 6(a) and 6(b)). The dispersoids found in this sample were <10 nm in dia. and exhibited a globular morphology. TEM analysis of the high-temperature as-HIPed sample revealed a significant increase in the number of nanoscale yttrium-enriched oxide dispersoids throughout the alloy microstructure (Figures 6(c) and 6(d)). The size of the dispersoids ranged from 3 to 50 nm with a volume density on the order of 1021m-3 (assuming a TEM foil thickness of 100 nm). The morphology of dispersoids >10 nm was cuboidal, while smaller dispersoids maintained a globular morphology. A majority of the dispersoids found in this sample reflected a cuboidal morphology and were later identified using HE-XRD. Microstructure: Alloy CR-126 Surface analysis of the as-atomized CR-126 powders also revealed high concentrations of oxyVolume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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Figure 3. Microstructural comparison of low-temperature as-HIPed CR-112 (outlined in green: TEM BFI images (a) and (b)), and high-temperature as-HIPed CR-112 (outlined in blue: TEM BFI images (c) and (d))

gen, chromium, and titanium (Figure 7(a)), similar to CR-118. The average thickness of the resulting oxide layer was determined to be ~85±10 nm based on AES depth-profiling measurements (approximated from the etching rate of SiO2). SEM with EDS analysis of as-CIPed CR-126 powders verified increased concentrations of titanium and oxygen and decreased concentrations of iron at particle/particle interfaces (Figure 7(b)), without a detectable chromium increase. EPMA results confirmed titanium enrichment of the PPB oxide phase during HIP consolidation of Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

the powder particles. Consistent with the dissociation mechanism, the level of the PPB oxide phase also was reduced from 5.0±0.2 v/o in the low-temperature-consolidated as-HIPed microstructure to 2.3±0.1 v/o in the high-temperature as-HIPed microstructure (Figures 8(a) and 8(c)). TEM analysis revealed the formation of Fe17Y2 intermetallic precipitates in the low-temperature as-HIPed microstructure, which appeared to decorate prior as-solidified cell boundaries, and the formation of a low-number density of nanoscale yttrium-enriched oxide dispersoids (Figures 9(a) and

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9(b)) in the same area. TEM analysis of the hightemperature as-HIPed microstructure showed a significant increase in the number of yttrium-

enriched dispersoids throughout the alloy microstructure (Figures 9(c) and 9(d)). The size of the dispersoids ranged from 5 to 60 nm with a number

Figure 4. (a) AES depth profile and (b) SEM cross-sectional analysis with EDS line scans for as-atomized and as-CIPed CR-118 powders, respectively

Figure 5. Microstructural comparison of low-temperature as-HIPed CR-118 (outlined in green: SEM images (a) and (b)), and high-temperature as-HIPed CR-118 (outlined in blue: SEM images (c) and (d))

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Figure 6. Microstructural comparison of low-temperature as-HIPed CR-118 (outlined in green: TEM BFI images (a) and (b)), and high-temperature as-HIPed CR-118 (outlined in blue: TEM BFI images (c) and (d))

density on the order of 1020m-3 (assuming a TEM foil thickness of 100 nm). The morphology of the dispersoids appeared analogous to those found in alloy CR-118 where most exhibited a cuboidal morphology, but required HE-XRD for identification. Synchrotron Phase Analysis Phase analysis of the chemical-reservoir alloys was completed using through-penetrating HEXRD. The alloys were examined in the as-atomized (red), low-temperature as-HIPed (green), and highVolume 46, Issue 6, 2010 International Journal of Powder Metallurgy

temperature as-HIPed (blue) conditions (Figure 10). The resulting diffraction peaks were compared by means of the reciprocal lattice vector Q (Å-1), which is commonly used to compare X-ray diffraction data independent of wavelength. Rietveld refinement of the diffraction data was completed using a generalized structural analysis system (GSAS).25,26 The corresponding refinement figures of merit (goodness of fit (χ2), profile residual (Rp), and weighted profile residual (wRp)) used to calculated the α-iron lattice parameter and resulting

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yttrium-enriched dispersoid volume percent are listed in Table II. Refinement of the synchrotron diffraction data taken from as-atomized CR-112 powder (Figure

10(a)) showed that the particles contain a singlephase α-iron matrix. The ultrathin surface-oxide phase was not identified using HE-XRD. Phase analysis of the low-temperature as-consolidated

Figure 7. (a) AES depth profile and (b) SEM cross-sectional analysis with EDS line scans for as-atomized and as-CIPed CR-126 powders, respectively

Figure 8. Microstructural comparison of low-temperature as-HIPed CR-126 (outlined in green: SEM images (a) and (b)), and high-temperature asHIPed CR-126 (outlined in blue: SEM images (c) and (d))

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Figure 9. Microstructural comparison of low-temperature as-HIPed CR-126 (outlined in green: TEM BFI images (a) and (b)), and high-temperature as-HIPed CR-126 (outlined in blue: TEM BFI images (c) and (d))

HIPed microstructure recognized the presence of a minor amount of crystalline Cr2O3 and Fe17Y2 within the α-iron matrix. Examination of the hightemperature-consolidated HIPed microstructure revealed the presence of Cr2O3, (Y,Cr)2O3, and Cr2N within the matrix. The levels of (Y,Cr)2O3 and Cr2O3 in this sample were determined to be 0.13±0.013 v/o and 1.18±0.029 v/o, respectively, using GSAS refinement of the diffraction data. It should also be noted that, after high-temperature consolidation, no diffraction peaks corresponding to Fe17Y2 were present. Diffraction data comparing microstructurVolume 46, Issue 6, 2010 International Journal of Powder Metallurgy

al phase changes that occur during consolidation of precursor oxide-dispersion-forming CR-112 powders are summarized in Figure 10(a). Synchrotron diffraction data for the as-atomized CR-118 powders (Figure 10(b)) revealed a small amount of intermetallic Fe17Y2 present in the α-iron matrix. It should also be noted that no identifiable oxide phases could be resolved in the asatomized diffraction data. The product phases of the alloy seemed mostly unchanged during lowtemperature consolidation. During this consolidation step, the volume fraction of Fe17Y2 increased

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and several unidentified minor peaks began to emerge above background. The formation of Y2Ti2O7 and full dissolution of Fe17Y2 resulted TABLE II. RIETVELD REFINEMENT FIGURES OF MERIT, α-Fe MATRIX LATTICE PARAMETER AND DISPERSOID VOLUME PERCENT Alloy (Condition) CR-112 (powder) CR-118 (powder) CR-126 (powder) CR-112 (850°C HIP) CR-118 (850°C HIP) CR-126 (850°C HIP) CR-112 (1,300°C HIP) CR-118 (1,300°C HIP) CR-126 (1,300°C HIP)

Reduced χ2

Rp

wRp

4.983 3.993 4.488 3.336 3.792 3.679 2.800 4.143 1.721

0.1193 0.1467 0.1290 0.1367 0.1876 0.1562 0.1507 0.1118 0.1192

0.1548 0.1790 0.1681 0.1586 0.2106 0.1686 0.1698 0.1497 0.1481

α-Fe Lattice Y-enriched Parameter Dispersoid (Å) (v/o) 2.87282 2.87413 2.87717 2.87236 2.87258 2.87605 2.87221 2.87277 2.87624

0.13 0.94 0.58

from high temperature consolidation of the alloy powders. The level of Y2Ti2O7 was determined to be 0.94±0.019 v/o, using GSAS refinement (Table II). Diffraction data highlighting these microstructural phase changes are presented in Figure 10(b). The as-atomized CR-126 powders (Figure 10(c)) also were found to contain a small level of intermetallic Fe17Y2 within the rapidly solidified α-iron microstructure. The alloy microstructure seemingly was unchanged after low-temperature consolidation, analogous to the CR-118 microstructure. Hightemperature consolidation of this alloy also resulted in the formation of Y2Ti2O7 and TiN, in conjunction with full dissociation of Fe17Y2. The volume fraction of Y2Ti2O7 was determined to be 0.58±0.022 v/o, using GSAS refinement (Table II). Diffraction data comparing phase changes in the CR-126 microstructure are compiled in Figure 10(c).

Figure 10. Synchrotron HE-XRD phase analysis of as-atomized powders (Red), low-temperature HIPed microstructure (Green), and high-temperature HIPed microstructure (Blue): (a) CR-112, (b) CR-118, and (c) CR-126

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DISCUSSION The CR alloys contained approximately ~15.5 a/o Cr in order to stabilize the α-iron matrix over the consolidation-temperature range. These alloys also contained small additions of yttrium, which acted as a dispersoid-forming agent during hightemperature consolidation of the powders. Yttrium is basically insoluble in α-iron, but at small concentrations it is solute trapped in the α-iron matrix during rapid-solidification processing via closecoupled argon-gas atomization. Titanium additions (~0.5 a/o) were added to CR-118 and CR-126 in an effort to promote the nanoscale oxide-dispersoid phase. The addition of titanium led to the formation of Y2Ti2O7 dispersoids, as opposed to the (Y,Cr)2O3 dispersoids found in CR-112. These results are in close agreement with previous studies of MA ODS steels of similar composition.27,28 Tungsten was added to CR-126 as a solid-solution-strengthening element.5 The addition of tungsten seemed to limit the amount of yttrium that could be solute trapped during rapid solidification, but did not seem to influence the chemistry of the resulting nanoscale oxide-dispersoid phase upon hot consolidation. As-Atomized Surface Oxide The surfaces of the powder particles were oxidized in situ during GARS. This oxidizing reaction occurs during the primary break-up of moltenmetal ligaments into spheres. Oxidizing the particles in the superheated, fully molten state enhances the kinetics of the surface reaction and allows for oxide-layer formation prior to complete droplet solidification; this results in compressive forces in the surface oxide layers, which are believed to minimize oxide spallation during powder handling prior to consolidation. This high-temperature rapid reaction resulted in powder particles with an ultrathin surface oxide scale. Oxygen appeared to be consumed at the surface of the particles in the form of an amorphous oxide phase. AES surface analysis of the powder particles revealed an apparent oxide layer, but no crystal structure reflections could be detected using HE-XRD with a resolution of 2 × 10-4Q (Å-1). This result provides evidence that the surface oxide phase probably solidifies in an amorphous state, promoted by rapid solidification of the ultrathin (<150 nm) oxide layer. The composition of the surface oxide layer is also thought to be controlled by the rapid kinetics Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

of the oxidizing reaction. This rapid reaction seems to promote the formation of the kinetically favored oxide phase, rather than the most thermodynamically stable phase. This concept highlights the importance of limiting the concentration of yttrium in the melt prior to atomization, which in return reduces the available surface activity of yttrium relative to other reactive alloying additions (such as chromium or titanium) in the atomized droplets. This appeared to help prevent the powder particles from forming yttrium-enriched surface oxide layers during atomization and permitted formation of an evenly distributed ODS microstructure upon high-temperature consolidation by a delayed oxygen-exchange reaction. Surface analysis of CR-112 powders of the simple Fe-Cr-Y alloy showed an enriched layer of oxygen and chromium, with a small enrichment of yttrium. The composition of the surface layer is indicative of a chromium-enriched oxide but, as previously noted, no structural information could be confirmed using synchrotron diffraction analysis. The surface layer is thought to be a mixed amorphous oxide phase containing primarily chromium and oxygen, with small concentrations of yttrium. CR-118 powders were atomized using the same reactive gas as CR-112, but surface analysis showed an enriched layer of mostly oxygen and chromium, with small concentrations of titanium. Apparently, the extreme reactivity of titanium prevented the consumption of any yttrium during atomization. Analysis of CR-126 powders revealed a surface composition similar to that of CR-118, with an enriched layer of oxygen, chromium, and titanium. Both CR-118 and CR-126 likely contain an amorphous chromium-enriched surface oxide layer doped with small amounts of titanium. The difference in the average oxide layer thickness on the CR-112 and CR-118 powders is within statistical error of the AES measurement, but the surface oxide layer found on CR-126 powders is ~35 nm thinner than the oxide layer on the CR112 and CR-118 alloys. The thickness of the surface oxide layer is thought to be influenced directly by the concentration of oxygen in the reactiveatomizing gas. Reducing the concentration of oxygen in the reactive gas certainly reduces the activity of oxygen within the gas impacting the surface of the droplets during atomization, promoting a reduction in oxidation kinetics and the formation of a thinner surface oxide layer, given similar alloy droplets.

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Evolution of CR-Alloy Microstructures The microstructure of the CR alloys was manipulated during elevated-temperature consolidation of the powders. The level of the PPB oxide phase was reduced in the three CR-alloys during hightemperature consolidation. This result shows that the metastable PPB oxide could be dissociated, which in return releases oxygen, allowing for the formation of nanoscale yttrium-enriched oxide dispersoids throughout the microstructure. This delayed oxygen-exchange reaction is believed to be driven by the relative thermodynamic stability of oxides. It is theorized that the PPB oxide phase begins to dissociate during elevated-temperature consolidation, in an effort to reach equilibrium conditions with the α-iron matrix. During this process, the initial chromium-enriched PPB oxide begins to dissociate and releases oxygen into solid solution in the α-iron matrix. Oxygen diffuses away from the PPBs, driven by the concentration gradient or chemical potential within the consolidated microstructure. The solid-solubility limit of oxygen in α-iron is inherently low (~0.104 a/o), and the dissociation of the PPBs should, in theory, discontinue once the solubility limit is reached and an equilibrium concentration of oxygen is achieved; however, consumption of oxygen by yttrium allows the reaction to continue until all the yttrium has been depleted.29 This illustrates the importance of achieving a critical balance between the initial oxygen and yttrium concentrations in

Figure 11. Comparison of thermodynamic stability of oxides that could form within the CR-alloy microstructures during powder consolidation30

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the CR-alloy powders. The ideal balance of oxygen to yttrium should be defined by the stoichiometric oxygen:yttrium ratio linked to an explicit dispersoid-product phase. Optimization of this ratio will allow for complete dissociation of the PPB oxide phase, coupled with total conversion of yttrium into nanoscale oxides. This optimized reaction would eliminate non-ideal phases from the alloy microstructure (PPB oxides and Fe17Y2 precipitates), which can act as void nucleation sites during mechanical evaluation of these alloys.22 This internal oxygen-exchange reaction is possible since yttrium-containing oxides are thermodynamically stable, in comparison with chromium or titanium oxides.30 Figure 11 illustrates the free-energy-of-formation comparison for chromium, titanium, and yttrium oxide phases over the relevant consolidation-temperature range. FUTURE PLANS AND TECHNOLOGY OUTLOOK Modifications to the reactive-gas-atomization process will be tested on a series of future CRalloys. These tests will be used to formulate processing parameters that can establish an optimum ratio of oxygen to yttrium for a select powder-particle-size range. Careful selection of an ideal particle-size range should preface modifications to the processing parameters, knowing that oxygen content is directly linked to the specific surface area of the powders (Figure 12). Microsegregation of yttrium and titanium at assolidified cell boundaries can occur during rapid solidification of the CR-alloy powders (Figure 13). Refinement of this microsegregation could be used to increase the spatial and number density of nanoscale oxide dispersoids that form at these former cell boundaries during the elevated-temperature oxygen-exchange reaction. An increased density of oxide dispersoids should successfully impede dislocation movement along these boundaries, allowing for the development of a fine (<1 µm) and highly stable dislocation substructure. Nanoscale oxide dispersoids are known to stabilize dislocation substructures, hindering thermally induced recrystallization during elevated-temperature mechanical testing, promoting maximum strength gains in ODS alloys.31–33 A carefully selected powder-particle-size range will be used during consolidation of future CRalloys. This critical size range will be defined by the oxygen and yttrium contents and the as-solidified cell size in the powders. The goal of this selection Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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process will be to consolidate powders with an optimum oxygen to yttrium ratio, in conjunction with a sufficiently fine cell size (<1 µm). These powders will be consolidated and heat treated at elevated temperatures (≥850°C), in order to promote the oxygen-exchange reaction and yttriumenriched nanoscale-oxide-dispersoid formation. The ODS microstructure will then be subjected to thermal mechanical treatments in an effort to develop a fine (<1 µm) dislocation substructure for ultimate strengthening of the ODS microstructure. Elevated-temperature mechanical testing will be used to compare these new CR-alloys with more

Figure 12. Resulting oxygen content relative to as-atomized powder-particle size for different reactive-gas concentrations

conventional ODS ferritic alloys (MA956, MA957, and PM2000) fabricated using MA. This new reactive-gas-atomization process offers the potential for a considerable increase in processing capacity for the production of precursor oxide-dispersion-forming particulates compared with traditional MA methods (Figure 14). This more-efficient processing technique could significantly reduce the high cost of manufacturing ODS ferritic stainless steel alloys, thus making them a more viable and economical choice as structural materials in future-generation power reactors (advanced ultra supercritical steam and Generation IV fission). SUMMARY A new simplified processing technique involving gas atomization and in situ oxidation has been developed to produce precursor ferritic stainless steel powders that can be consolidated into an ODS isotropic microstructure. Microstructural results showed the ability to manipulate oxide and intermetallic phases within the alloy using hightemperature consolidation. HE-XRD phase analysis coupled with TEM microstructure analysis confirms the operation of an oxygen-exchange reaction between a less stable prior-particle-boundary oxide and yttrium. This oxygen-exchange reaction results in the formation of nanoscale yttrium-containing oxide dispersoids throughout the ferritic microstructure. The atomization processing parameters will need to be further optimized, in order to generate an ideal oxygen:yttrium ratio in improved CR-alloy powders. This ideal ratio

Figure 13. EPMA elemental composition map of as-atomized CR-118 powders: (a) microstructure of cross sections of as-atomized and CIPed powders. The lighter/ smaller particles are copper used during CIPing; (b) titanium microsegregation, and (c) yttrium microsegregation

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5.

6.

7.

8.

9.

10.

Figure 14. Comparison of approximate MA and gas-atomization processing rates20, 34

should permit full dissociation of both the PPB oxides and any previously formed Fe17Y2 precipitates, resulting in a stronger and more homogenous alloy microstructure. ACKNOWLEDGEMENT Support from the Department of Energy, Office of Fossil Energy (ARM program) through the Ames Laboratory (contract no. DE-AC02-07CH11358) is gratefully acknowledged. The high-energy X-ray work at beamline 11-BM of the APS was supported by the Department of Energy, Office of Science, Basic Energy Sciences (contract no. DE-AC0206CH11357). The authors also thank Danny Shechtman, James Anderegg, David Byrd, and Hal Sailsbury for their individual contributions to this paper. REFERENCES 1. D.T. Hoelzer, J. Bentley, M.A. Sokolov, M.K. Miller, G.R. Odette and M.J. Alinger, “Influence of Particle Dispersions on the High-Temperature Strength of Ferritic Alloys”, J. of Nuc. Mat., 2007, vol. 367–370, pp. 166–172. 2. M.A. Sokolov, D.T. Hoelzer, R.E. Stoller and D.A. McClintock, “Fracture Toughness and Tensile Properties of Nano-Structured Ferritic Steel 12YWT”, J. of Nuc. Mat., 2007, vol. 367–370, pp. 213–216. 3. R.L. Klueh, J.P. Shingledecker, R.W. Swinderman and D.T. Hoelzer, “Oxide Dispersion-Strengthened Steels: A Comparison of Some Commercial and Experimental Alloys”, J. of Nuc. Mat., 2005, vol. 341, pp. 103–114. 4. D.J. Larson, P.J. Maziasz, I.S. Kim and K. Miyahara, “Three-Dimensional Atom Probe Observation of Nanoscale Titanium-Oxygen Clustering in an Oxide-Dispersion-

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Stengthened Fe-12Cr-3W-0.4Ti + Y 2O 3 Ferritic Alloy”, Scripta Mater., 2001, vol. 44, pp. 359–364. I.S. Kim, B.Y. Choi, C.Y. Kang, T. Okuda, P.J. Maziasz and K. Miyahara, “Effect of Ti and W on the Mechancial Properties and Microstructure of 12% Cr base Mechanical-Alloyed Nano-sized ODS Ferritic Alloys”, ISIJ International, 2003, vol. 43, pp. 1,640–1,646. M.K. Miller, E.A. Kenik, K.F. Russell, L. Heatherly, D.T. Hoelzer and P.J. Maziasz, “Atom Probe Tomography of Nanoscale Particles in ODS Ferritic Alloys”, Mat. Sci. and Eng., 2003, vol. A353, pp. 140–145. M.K. Miller, D.T. Hoelzer, E.A. Kenik and K.F. Russell, “Nanometer Scale Precipitation in Ferritic MA/ODS Alloy MA957”, J. of Nuc. Mat., 2004, vol. 329–333, pp. 338–341. M.K. Miller, D.T. Hoelzer, E.A. Kenik and K.F. Russell, “Stability of Ferritic MA/ODS Alloys at High Temperature”, Intermet., 2005, vol. 13, pp. 387–392. M.K. Miller, K.F. Russell and D.T. Hoelzer, “Characterization of Precipitates in MA/ODS Ferritic Alloys”, J. of Nuc. Mat., 2006, vol. 351, pp. 261–268. G.R. Odette, M.J. Alinger and B.D. Wirth, “Recent Developments in Irradiation-Resistant Steels”, Annu. Rev. Mater. Res., 2008, vol. 38, pp. 471–503. M.J. Alinger, G.R. Odette and D.T. Hoelzer, “On The Role Of Alloy Composition and Processing Parameters In Nanocluster Formation and Dispersion Strengthening In Nanostructured Ferritic Alloys “, Fusion Materials Semiannual Progress Report, 2008, DOE-ER-0313/43, Oak Ridge, TN. M.J. Alinger, G.R. Odette and D.T. Hoelzer, “The Development and Stability of Y–Ti–O Nanoclusters in Mechanically Alloyed Fe–Cr based Ferritic Alloys”, J. of Nuc. Mat., 2004, vol. 329–333, pp. 382–386. S.J. Zinkle and N.M. Ghoniem, “Operating Temperature Windows for Fusion Reactor Structural Materials”, Fus. Eng. Des., 2000, vol. 51, pp. 55–71. E.A. Little, “Development of Radiation Resistant Materials for Advanced Nuclear Power Plants”, Mater. Sci. Tech., 2006, vol. 22, pp. 491–518. “A Technology Roadmap for Generation IV Nuclear Energy Systems”, U.S. DOE Nuclear Energy Research Advisory Committee and Generation IV International Forum (GIF), 2002, GIF-002-00, Washington, DC. J.T. Busby, “Economic Benefits of Advanced Materials in Nuclear Power Systems”, J. Nuc. Mater., 2009, vol. 392, pp. 301–306. J.S. Benjamin, and J.M. Larson, “Powder Metallurgy Techniques Applied to Superalloys”, J. Aircraft, 1977, vol. 14, pp. 613–623. J.S. Benjamin, “Mechanical Alloying—A Perspective”, Met. Powd. Report, 1990, vol. 45, pp. 122–127. P.S. Gilman and J.S. Benjamin, “Mechanical Alloying”, Annu. Rev. Mater. Sci., 1983, vol. 13, pp. 279–300. C. Suryanarayana, “Mechanical Alloying and Milling”, Prog. Mat. Sci., 2001, vol. 46, pp. 1–184. I.E. Anderson, B.K. Lograsso and T.W. Ellis, “Gas Atomization Synthesis of Refractory or Intermetallic Compounds and Supersaturated Solid Solutions”, U.S. Patent No. 5,368,657, November 29, 1994. J.R. Rieken, I.E. Anderson, M.J. Kramer, Y.Q. Wu, J.W. Anderegg, A. Kracher and M.F. Besser, “Atomized Precursor Alloy Powder for Oxide Dispersion-Strengthened

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Ferritic Stainless Steel”, Advances in Powder Metallurgy and Particulate Material—2008, compiled by R. Lawcock, A. Lawley and P.J. McGeehan, Metal Powder Industries Federation, Princeton, NJ, 2008, part 10, pp. 324–342. Materials Preparation Center, Ames Laboratory, US DOE Basic Energy Sciences, Ames, IA, USA, available from: www.mpc.ameslab.gov. J. Wang, B.H. Toby, P.L. Lee, L. Ribaud, S.M. Antao, C. Kurtz, M. Ramanathan, R.B. Von Dreele and M.A. Beno, “A Dedicated Powder Diffraction Beamline at the Advanced Photon Source: Commissioning and Early Operational Results”, Rev. Sci. Instrum., 2008, vol. 79, pp. 85–105. B.H. Toby, “EXPGUI, a Graphical Interface for GSAS”, J. Appl. Cryst., 2001, vol. 34, pp. 210–221. A.C. Larson and R.B. Von Dreele, “General Structure Analysis System (GSAS)”, 2004, LAUR 86-748, Los Alamos National Laboratory, Los Alamos, NM. S. Nomura, T. Okuda, S. Shikakura, M. Fujiwara and K. Asabe, “Effect of Alloying Elements on High-Temperature Strength of Oxide Dispersion Strengthened Chromium (12Cr) Ferritic/Martensitic Steels”, Solid State Powder Processing, edited by A.H. Clauer and J.J. Debarbadillo, TMS, Warrendale, PA, 1990, vol. 1, pp. 203–211. Y. Uchida, S. Ohnuki, N. Hasimoto, T. Suda, T. Nagai, T. Shibayama, K. Hamada, N. Akasaka, S. Yamashita, S.

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Ohtsuka and T. Yoshitake, “Effect of Minor Alloying Element on Dispersing Nano-Particles in ODS Steel”, Structural and Refractory Materials for Fusion and Fission Technologies, edited by J. Aktaa, M. Samaras, M. Serrano de Caro, M. Victoria and B. Wirth, MRS, Warrendale, PA, 2007, vol. 981E, pp. 09–14. A.U. Seybolt and R.L. Fullman, “A Rationalization of the Oxygen Solid Solubility in Some Transition Metals”, J. Met., 1954, vol. 6, pp. 548–549. F. Sauert, E. Schultze-Rhonhof and W.S. Sheng, Thermochemical Data of Pure Substances, Second Edition, 1992, VCH Verlagsgesellschaft mbH, New York, NY. A. Kelly and R.B. Nicholson, Strengthening Methods in Crystals, 1971, John Wiley & Sons, Inc., New York, NY. B.A. Wilcox and R.I. Jaffee, “Direct and Indirect Strengthening Effects of Thorium Oxide (ThO2) Particles in Dispersion-Hardened Nickel”, Int. Conf. on Strength of Metals & Alloys, Trans. Japan Inst. Met., Sendai, Japan, 1967, vol. 9, pp. 575–579. J.E. White and R.D. Carnahan, “A Microplasticity Study of Dispersion Strengthening in TD-Nickel”, Trans. Met. Soc. AIME, 1964, vol. 230, pp. 1,298–1,306. R.M. German, Powder Metallurgy & Particulate Materials Processing, 2005, Metal Powder Industries Federation, Princeton, NJ. ijpm

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RESEARCH & DEVELOPMENT

POTENTIAL EFFECTS OF RETAINED AUSTENITE ON END-QUENCH COOLING RATES IN PM STEELS Fred J. Semel* and Diana A. Lados**

INTRODUCTION In an earlier paper,1 process modeling was used to show that the thermal-conductivity increases that typically accompany the martensite transformation in steels affect the cooling rate in the Jominy end-quench test. A one-dimensional model that included the effects of material property variations with temperature was presented, which predicted a small increase in cooling rates with increases in the Ms temperature for fully dense steels, and significantly increased rates for PM steels. The model was based on earlier studies of the end-quench test by Saritaş et al.2 that showed increased cooling rates in PM steels compared with fully dense ones. Later studies by Lawley et al.3 went on to show water penetration of the pores as a causative mechanism. By incorporating a simple theory of this mechanism with speculated increases in the Ms temperature due to porosity in accordance with the findings of Warke et al.,4 it was possible to obtain cooling curves that displayed a marked resemblance to those observed experimentally by Saritaş. Basically, the current paper adds the results of two improvements to the earlier treatment of the subject. One involves a minor but potentially important change in the simulation model to take account of the effects on the subject phenomena of variations in the retained austenite content of the steels of interest. These effects were not included in the earlier study because the data needed to estimate the associated retained austenite contents were not then available. The other is in the form of a subsidiary method to estimate the separate contributions to the end-quench cooling rate of the several variables that are indicated to be important. The first iteration of the study was prompted by gas-cooling trials that the present authors conducted in otherwise unrelated work. The results of the trials indicated increased hardenability with decreasing density. Significantly, a subsequent limited survey of the open literature revealed that Bocchini et al.5 had also observed a similar effect of density on hardenability using gas as a coolant. The apparent implica-

This paper updates an earlier one in which process modeling was used to make a case for the idea that increases in the martensite start temperature (Ms) due to porosity were partially responsible for the increased end-quench cooling rates that were previously observed in various PM steels. In this paper, it is shown that the likely effects of porosity in reducing the retained austenite content in comparison with that of the porefree condition should also be expected to make a modest contribution to increasing the relative end-quench cooling rates of PM steels. In addition, a subsidiary method is used to estimate the separate contributions of the aforementioned effects of pores and those of the “water penetration of the pores” mechanism originally proposed to explain cooling-rate increases. Estimates are also presented of the effects of likely increases in the thermal conductivity of the PM steels due to various causes including alloy inhomogenities and/or decarburization during processing.

*Retired, formerly Hoeganaes Corporation, Cinnaminson, New Jersey 08077, USA; E-mail: [email protected], **Assistant Professor of Mechanical Engineering, Worcester Polytechnic Institute, Director, Integrative Materials Design Center, Worcester, Massachusetts 01609, USA; E-mail: [email protected]

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tion of these findings was that water cooling was unnecessary to effect increased cooling rates with increasing porosity; thus, either the water penetration explanation of the Saritaş findings was flawed, or there was something in addition to it, common to both water and gas cooling, that was contributing. A critical review of the studies that led to the concept of water penetration, including a follow-up study by Zavaliangos et al.,6 favored the conclusion that the explanation was almost certainly correct but incomplete. This study was thought to be especially relevant in that it not only supported this conclusion but suggested process modeling as the most direct means to study the issue further, and provided a simple basis for, if not an actual explanation of, the gas-cooling results. The study consisted mainly of comparative simulations based on finite difference methods and employed Neumann-type boundary conditions at the quench ends of the models to reflect the distinctly convective nature of the Jominy cooling process. Two cases were examined. In one, the presence of porosity was compared with the pore-free condition without regard for the water-penetration mechanism. The findings unexpectedly indicated that the reduction in thermal conductivity occasioned by porosity would result in modest increases in cooling rates, at least, near the quench end of the specimen. Further into the specimen, and more in line with expectation, the indications were of increasingly slower cooling in the balance of the material. The significant implication with regard to gas cooling was that neither water penetration of the pores nor even the use of water as a coolant are apparently necessary to effect faster cooling in a porous specimen. Evidently, the presence of porosity itself is sufficient to do this, albeit to a limited extent. In the second of the two cases examined, a first-order attempt was made to model the water-penetration mechanism. This was successful to the extent that the findings indicated significant cooling-rate increases, even in the face of substantially reduced thermal conductivities, but was otherwise unsuccessful in that the resultant cooling curves did not, even remotely, resemble the earlier experimental findings. This last result was considered to be especially relevant because other than the cooling rate increases, perhaps the most striking feature of the findings by Saritaş was the shape of the curves. As reference to the paper confirms, the greatest and

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most consistent indications in this regard were for cooling at the 5 mm depth below the quench ends of the specimens. All three of the PM compositions studied (FL-4405, FLC2-4405, and FLN2-4405) undercooled the wrought standard of comparison at this depth, but did so mainly in the middle temperature range from about 550°C to 200°C. Otherwise, at both higher and lower temperatures, the rates and associated differences in the temperature values of the PM specimens vs. the standard were smaller and appeared to approach those of the latter asymptotically. Otherwise, the findings were not as consistent as to the effects of density and alloy content. Nevertheless, the general trend in the data was for the degree of undercooling, and/or the magnitude of the associated cooling rate increases, to increase, with decreasing density and/or increasing alloy content. What was especially intriguing about these findings was that, while they were at least consistent with water penetration as a possible extrinsic cause of the rate increases, the fact that the cooling curves showed that the greatest undercooling was in a middle temperature range suggested the possibility of one or more additional intrinsic causes. Likely candidates included a thermal conductivity changes and heat of transformation effects that accompany a phase change. Saritaş presented evidence to show that all of the specimens had transformed to predominantly martensitic microstructures at the 5 mm depth. Thus, the phase change in question was the austenite-tomartensite transformation. Interestingly, this transformation starts at, or just below, the middle of the subject temperature range in all of the steels under consideration. In addition, unlike the waterpenetration mechanism, the martensite transformation, along with the various phenomena that accompany it, is common to both the water- and gas-quenching processes. SIMULATION PROCEDURE AND ESTIMATES OF MATERIAL PROPERTIES Preliminary studies with the aim of selecting a suitable simulation procedure indicated that a one-dimensional model would be adequate. The governing mathematical equations were derived in the first iteration of the study and need not be repeated here. They included terms to account for the enthalpy effects of the austenite-to-martensite transformation and employed a Neumann-type boundary condition at the quench end of the Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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POTENTIAL EFFECTS OF RETAINED AUSTENITE ON END-QUENCH COOLING RATES IN PM STEEL

model. The heat-transfer coefficient was back calculated from the findings of Saritaş. The simulations were based on finite difference approximations of the equations. In the general case, the respective time and length increments of the approximations were 1 × 10-4 s and 2.5 × 10-4 m. To avoid excessive computer run times, the number of iterations of the recurrence equation was limited to 106. Comparison of the model with known closed-form solutions of the heat equation suggested that these values were of the correct order to produce results with reasonable precision. Application of the equations requires a knowledge of the temperature variations of the thermal conductivity, enthalpy of transformation, specific heat, and density of each of the steels of interest. Virtually none of these data were known for the standard PM steel grades. However, databases were available for many of the standard wrought grades, and the general effects of porosity on each of the subject properties were known. Consequently, it was decided to estimate them. Since the various assumptions and methods that were used to do this were set forth in considerable detail previously,1 only those particulars that are essential to an understanding of the present study have been repeated here. The standard for comparison in the Saritaş study and, hence by default, in both this and the first study, was wrought SAE 4150. The principal PM steel of interest was FLN2-4405. These steels have the same carbon content and essentially the same total alloy content, but differ in terms of the specific alloying elements that make up the total. SAE 4150 is silicon killed and nominally contains 0.85 w/o Mn, 1.0 w/o Cr, and 0.20 w/o Mo.7 The alloy base of FLN2-4405 is water atomized and annealed to remove residual oxygen. As finally constituted, the steel nominally contains 0.15 w/o Mn and 0.80 w/o Mo, which are prealloyed, and 2 w/o Ni which is admixed. A review of a materials-properties compendium8 which listed data at both ambient and elevated temperatures indicated that steels of moderately different carbon and total alloy contents did not differ significantly in density and specific heat at either a particular temperature or with increasing temperature. The important implication was that similar steels could be expected to exhibit essentially similar temperature variations of density and specific heat. Since enthalpy is the temperature integral of specific heat, this would also be true. Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

However, the indications were slightly different in the case of thermal conductivity. Although the data indicated similar trends with increasing temperature from steel to steel, they also showed a greater sensitivity to composition at any given temperature. In general, thermal conductivity decreased with increasing alloy content, especially the carbon content. A linear regression analysis of the data indicated that the rate of decrease was ~5% per 0.1 w/o C. As to actually estimating the various properties needed for the simulations, a literature survey led to a paper based on a computer-modeled database that contained the temperature variations of the properties of interest, albeit for SAE 4140 rather than 4150.9 However, as suggested by their designations, reference to the relevant specifications7 confirmed that the only nominal difference in compositions of the two steels was a 0.1 w/o C. Thus, after applying the regression analysis result to correct the thermal conductivity data for this difference, the 4140 data were used to generate the property estimates that form the basis of the balance of the study. Based on a limited survey of the literature,13 the retained austenite content of 4150 was taken to be 6 v/o. Although information for 4140 was explicit in terms of the behaviors of the individual phases,9 it was non-specific as to the functional relationships needed to calculate the properties of the phase mixtures that arise in a phase transformation. A further review of the open literature, however, provided this information for the austenite-to-martensite transformation in the form of the following equation:10 Volume Fraction of Un-Transformed Austenite = exp [ χ · (Ms– T )] for Mf ≤ T ≤ Ms. (1(a)) Here, T is the temperature (°C) corresponding to the volume fraction of un-transformed austenite; Ms and Mf are, respectively, the martensite start and finish temperatures; and χ is a constant related to the retained-austenite content. If γr is the volume fraction of retained austenite, then: ln(γr) χ = ––––––– MsMf

(1(b))

In the first iteration of the study, all of the steels were assumed to have the same value of χ. This was based on the following considerations: (1) the

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conjecture that the associated transformations went to virtual completion with the result that the retained austenite content was, in all cases, ~0.5 v/o (i.e., γr = 0.005); (2) an indication in the open literature to the effect that the numerical difference between Ms and Mf of plain carbon and lowalloy steels is approximately constant at ~215°C.11 Thus, the value of χ that was universally used in the earlier study was ln (0.005)/215 = - 0.0246. In the present study, the value of χ was again determined in accordance with Equation (1(b)) but for variable values of the Ms temperature and the retained austenite content. The Ms temperatures of the compositions of interest were calculated from the relation:12 Ms (°C) = 570-474 w/o C-33 w/o Mn-21 w/o Mo-18 w/o Cr-17 w/o Ni. (2) In addition, in contrast with the first study, since the retained austenite content is typically measured at room temperature, the Mf temperature in this study was taken to be equal to 20°C in all cases. As an example, Figure 1 shows the reported variation in thermal conductivity of 4140 compared with that calculated for 4150 at two levels of retained austenite. The 4150 curve (unconnected points) shows the variation at the 0.5 v/o level while the solid curve shows the variation at the 6 v/o level. The uniformly higher values for 4140 in Figure 1 compared with those for 4150 at the 6 v/o level are a consequence of three effects: the previously noted correction to the 4140 data to account for the higher carbon content of 4150; the lower Ms of 4150; and its higher retained-austenite content. All three effects are important. However, of the three, the latter two are the more important in terms of the objectives of the study. Accordingly, variations in thermal conductivity can be expected to affect the cooling rate and, as clearly indicated by the present findings, variations in either or both the Ms temperature and the retained-austenite content affect the thermal conductivity. Thus, it follows that the values of each of these parameters, as well as anything which affects them, should likewise have an effect on the cooling rate. POROSITY EFFECTS Porosity affects all of the material properties of interest. Its effects on density and heat capacity

36

are well known, and easily estimated. Its direct effects on thermal conductivity are not as well known but have been studied, and are likewise relatively easy to estimate.14,15 The specific relationships that were used for the PM steels were set out previously.1 As will be discussed, porosity also has indirect effects on thermal conductivity as a result of its effects on the Ms temperature and the levels of retained austenite. In general, both the Ms temperature and the level of retained austenite of wrought steels are known to be affected by several factors. Other than composition, these include the austenitic grain size and the presence of superimposed magnetic fields, extrinsically or intrinsically imposed stress fields, and plastic deformation.16 Significantly, recent findings have indicated that both properties are also subject to porosity effects. Warke utilized dilatometry to examine the martensite transformation and X-ray diffraction to determine the attendant retained-austenite contents of both powderforged and porous specimens of FL-4605 (0.5 w/o C steel based on prealloyed powder containing nominally 0.5 w/o Mo and 1.8 w/o Ni). They observed significant increases in the Ms temperature and decreases in the level of retained austenite with increasing porosity. The Ms determinations were conducted at three different cooling rates in the range from 40°C to 180°C/s. These authors reported the resulting average values, and noted that the magnitude of the averages exceeded the scatter in the data at each of the cooling rates by a factor of 3 or more. Both the observed Ms and retained-austenite levels were approximate linear functions of the pore content. The least squares slope of the Ms variation was +2.3°C per 1 v/o porosity. The slope of the retained-austenite varia-

Figure 1. Thermal conductivity of 4140 (experimental) and 4150 (calculated) as a function of temperature. Unconnected points = 0.5 v/o retained austenite. Solid line = 6 v/o retained austenite Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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tion was –0.58% per 1 v/o porosity. These findings have yet to be confirmed. However, they are regarded as theoretically reasonable because the presence of porosity is expected to reduce the magnitude of the internal stresses that act against the transformation, and hence modify the way the transformation progresses. In fact, preliminary studies of the effects on cooling rate of similar Ms temperature increases to the ones observed by Warke. had anticipated the likelihood of just such porosity effects well in advance of their actually being reported. Thus, it will be appreciated that the result of Warke on Ms were immediately recognized as being especially relevant to the presence study. Ostensibly, this would have also been true of the retained austenite results as well. However, as noted at the outset, the first iteration of the study did not include them. This was because both findings were originally communicated privately and, as it happened, the Ms results predated the retained-austenite data by several months. By the time the austenite results became available, the report of the aforementioned first iteration of the study had already been prepared and submitted for publication. In any event, the Ms results of Warke were incorporated in the first iteration of the study by simply adding the least squares result to Equation (2). Precisely the same procedure was used in the present iteration. A different procedure was used in the case of the retained-austenite results. The objection here to using a simple straight-line prediction was that it led to untenable (negative) retained-austenite values at relatively modest porosity levels. Instead, a curvilinear relationship that precluded this possibility was derived by setting the porosity rate of change of the retained austenite proportional to the instantaneous value and integrating. If γr is the volume (or weight) fraction of retained austenite, and Vf is the volume fraction of pores, then: γr = γr,o exp(α·Vf ),

(3(a))

where γr,o is the volume (or weight) fraction of retained austenite of the pore-free composition, and α is a constant, characteristic of the variation. Data on the retained-austenite content of FLN24405 were not available to evaluate these two constants. However, given the compositional similarity to FL-4605, both were approximated on the basis of the Warke results. Regression analysis of the latter provided an estimate of α, and since the Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

nominal alloy content of FLN2-4405 is about 15% higher than that of FL-4605, its pore-free retainedaustenite content was correspondingly estimated to be ~15% higher. The resulting relationship is: γr = 0.10 exp(–10.99·Vf ).

(3(b))

As an example, at 5 v/o porosity (Vf = 0.05), the retained-austenite content of FLN2-4405 would, according to this result, be 0.058 (5.8 v/o). HEAT-TRANSFER COEFFICIENT USED IN THE SIMULATIONS The value of the heat-transfer coefficient required in the quench-end boundary condition has a pronounced effect on the output of the associated equations. In order to simulate actual experimental findings, it is essential to use the same heat-transfer coefficient as that which characterized the trials that led to the findings. In the case of the Saritaş study, as in most studies of the sort, the coefficient was not measured and was consequently unknown. However, it was inherent in the findings and hence could be estimated from them. The procedure used to do this involved using the model to back-calculate the value needed to produce a cooling curve that was a reasonable approximation of the one reported by Saritaş for their SAE-4150 standard of comparison. In the first iteration, the retained austenite in the 4150 was assumed to be 0.5 v/o. Increasing the level to 6 v/o, as in the present study, necessitated recalculating the heat-transfer coefficient, and led to an overall increase in its value of ~7%. The results were that the coefficient increases linearly with temperature from 10 to 11,250 J·s-1·m-2·K-1 in the range from 820°C to 550°C and then remains constant at 11,250 J·s1·m-2·K-1 in the range from 550°C to 500°C. Thereafter it decreases to 10,050 J·s-1·m-2·K-1 at 20°C according to the polynomial relation: h = –0.006944T2 + 6.1111T + 9930.6.

(4)

Compared with the average of three of the reported cooling curves for SAE-4150, the values produced a simulated curve for which the absolute difference from selected points of the latter averaged 2.1°C. The largest single difference was 4.5°C. It is also relevant to note that these values are apparent values; in addition to convection they necessarily reflect the radiant-heat losses that occurred during cooling, which are not otherwise accounted for by the model.

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DETAILS OF SIMULATION STUDIES UTILIZING THE MODIFIED MODEL Four subsidiary studies were conducted with the modified model. However, only the results of two need be reported here. The other two were similar to studies that were reported in connection with the original model and produced marginally similar results. In terms of the effects of variations in the retained-austenite content, in both cases the findings were in reasonable accord with expectation, in that reductions in the retained-austenite content produced increases in the cooling rate. Study 1 This was the analog of an earlier study and produced somewhat similar results. However, its importance here is twofold: in essence, this is the primary study of the inquiry as a whole; and, it provided the data that form the basis of the second study which, as will be seen, offers insight into how the cooling-rate increases break down in terms of the variables that contributed to them. The study per se consisted of three simulations: one of SAE-4150, and one each at 10 and 15 v/o porosity levels in FLN2-4405. The aim was to determine if the model would produce similar cooling-rate differences to the ones reported by Saritaş. Apart from variations in material properties, and water-penetration effects that were incorporated in the model, the simulations required Ms temperature and retained-austenite inputs and offered the possibility of varying the base thermal conductivity of 4405. Based on Equation (2), the Ms temperature of SAE-4150 was calculated to be 280°C. In comparison, Table I lists the estimated values of the Ms temperatures (also based on Equation (2) but rounded to the nearest multiple of 5) of FLN2-4405 for the porosity levels of interest and for different degrees of homogenization of the steels in relation to the admixed nickel. In general, Table I shows that Ms increases with increases in the level of porosity and decreases with increases TABLE I. MS TEMPERATURES OF FLN2-4405 VS. POROSITY AND DISSOLUTION OF ADMIXED NICKEL Porosity (v/o) 0 10 15

38

Nickel Content in Solid Solution 0.5 w/o 1.0 w/o 2.0 w/o Ms Temperature (°C) 305 325 340

295 320 330

280 300 315

in the nickel content in solid solution. Saritaş used normal sintering conditions (1,120°C for 0.5 h at temperature) to prepare their Jominy specimens. However, as is well known, such conditions are not adequate to homogenize admixed 2 w/o Ni in FLN2-4405. Although later metallographic studies by Lawley suggested that most, if not all, the nickel had dissolved (at least locally) under these conditions, it was also clear from the comparative hardness profiles in the Saritaş study that only a fraction of the nickel was actually effective as a completely homogenized alloy addition. Consequently, it was decided to use the Ms values cited here for the lowest dissolvednickel content to characterize the behavior of 4405 in the simulations. In addition, it was also considered that the lower effective-total-alloy content associated with this level of nickel was likely to result in a modestly higher thermal-conductivity value as well. Thus, the base thermal conductivity of 4405 was increased by 5% relative to that of 4150 for the simulations. In accordance with Equation (3(b)), the retained austenite contents of 4405 at 10 and 15 v/o porosity levels were taken to be 3.3 and 1.9 v/o, respectively. Study 2 In the interim between the first iteration of the study and the present one, a method was devised to use the simulation program to estimate the contributions of the variables to the cooling-rate changes predicted by the model. The objective of the present study was to apply the method to the results of the earlier study. The method was based on the use of the maximum undercooling between the cooling curves of the PM variants and the standard of comparison to characterize the cooling-rate increases, and the fact that this value is evidently a single-valued function of the variables. Thus, if ΔTM is the maximum undercooling, and ΔTM = ΔTM(Vf , Ms, γc , f ) where Vf , Ms, and γc have been defined previously, and f is the fractional increase in the thermal conductivity, then by starting with the total derivative of ΔTM:

o∫dΔTM = ΔTM =

dΔT

∂ΔTM ΔVf + ∫ –––––– dMs + ∂Ms

∂ΔT

∂ΔTM dγc + ∫ –––––– df . ∂f

M ∫ –––––– ∂Vf M ∫ –––––– ∂γ c

(5)

The line integral on the left is over the path defined Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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by the several variables, and the integrals on the right are all Riemann integrals, dependent only on their respective limits. According to this result, the individual contributions of the several variables to a particular value of the maximum undercooling can be determined by separately evaluating each of the various integrals. In each case, the latter can be accomplished by using the simulation program to determine the value of ΔTM corresponding to any selected value of the associated variable, while holding the other variables constant. In the actual study, ΔTM of each variable was determined in this manner over a range of the variable’s values and submitted to regression analysis to generate an analytical expression that was, in effect, a functional estimate of the associated integral. These estimates were then used to analyze the predicted ΔTM values for the PM variants vs. the standard of comparison in Study 1. RESULTS AND DISCUSSION The cooling curves generated in Study 1 are shown in Figure 2. The blue curve shows the simulated cooling behavior of the SAE-4150 standard of comparison, and the red and green curves show the simulated cooling behaviors for FLN2-4405. A cursory review of Figure 2 confirms that each of the hypothetical PM steels undercooled 4150 in both the intermediate and low temperature ranges. At 10 v/o porosity (red curve), the maximum undercooling was a little over 37°C and occurred below the estimated Ms temperature of the steel at ~314°C. Undercooling in the low temperature range was just under 10°C. At the 15 v/o porosity level (green curve), the maximum undercooling rounded up to 47°C, and also occurred below the Ms temperature of the steel at ~316°C. In the low temperature range, the undercooling was a little over 12°C. As reference to Saritaş confirms, the overall similarity of these findings to those in their Figure 5 is apparent. The functional estimates of the integrals in Equation (5) needed in Study 2 and the graphical results of the simulations that led to them are presented in Figure 3. The four

Figure 2. Simulated cooling curves at a depth of 5 mm of SAE 4150 and hypothetical FLN2-4405 at porosity levels of 10 v/o and 15 v/o Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

graphs detail the individual effects of each of the four variables of interest on the maximum undercooling to be expected in a hypothetical variant of FLN2-4405 and the SAE 4150 standard of comparison. The MS Excel regression results accompanying each of the variations comprise the functional estimates of the integrals. Otherwise, the graphs indicating the effects of Ms and the retained-austenite content are self explanatory. The graph labeled “Water Penetration Effects on Cooling” in Figure 3, indicates the porosity effects due to the water-penetration mechanism and does so independent of the relation of porosity to Ms and the retained-austenite contents and their effects on cooling. The fractional change in thermal conductivity in the other graphs is indicated by the ratio of the thermal conductivity of the hypothetical variant, k, to that of the standard of comparison, ko (i.e., by k/ko). Note also that the “y” intercept values of the four variations correspond to the base properties of the SAE 4150 standard of comparison. This means that “Undercooling” = 0 at the following values of the variables: porosity = 0; Ms = 280°C; v/o austenite = 6; and, k/ko = 1. The functional estimates of the integrals (regression results) were substituted into Equation (5) and used to estimate the maximum undercooling to be expected for the particular combinations of variables that characterized the PM variants of Study 1. The results, along with the percentage contributions of the individual variables to the estimates, are shown in Table II. The maximum undercooling estimates, given by the totals in the last row of Table II, differed from the values previously reported in Study 1 for the same values of the variables. However, the associated errors were small, namely, 0.6% and 1.4%, respectively, and were not unreasonable, given the fact that both the present and earlier values were basically the result of finite difference approximations. Otherwise, as will be evident from a review of the balance of the data in Table II, the general outcome of the analysis was that the percentage contributions of the individual variables to these values were similar. Accordingly, the indications were that ~65% of the undercooling was due to porosity and attendant water penetration, a little less than 30% to the combined effects of Ms and the retained austenite content, and the balance to the assumed fractional increase in thermal conductivity. Given that the porosity was also respon-

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sible for about half of the Ms, and all of the retained-austenite differences that were considered to exist between the PM variants and the standard of comparison, the overall implication of the data was that ~85% of the undercooling was ultimately due to the combined direct and indirect effects of porosity. In addition, given that the increased thermal conductivity of the PM variants and the rest of the Ms differences were the result of hypothesized nickel heterogeneities in the variants, a secondary outcome of the data was that a small but significant part of the undercooling (~15%) was essentially the result of extraneous effects of the admixed nickel in FLN2-4405. SUMMARY AND CONCLUSIONS The present work reflects the second of two studies with the overall objective of using process modeling to examine the possibility that in addition to water penetration of pores, there are factors that

are responsible for earlier observed increases in end-quench-cooling rates of PM steels compared with a wrought standard of comparison. The first study was successful in showing that increases in Ms and decreases in thermal conductivity of PM steels due to porosity effects and/or nickel heterogeneities were two such possibilities. However, because of a paucity of applicable data at the time, the first study neglected to examine the potential effects of porosity on the retained-austenite contents of the steels as an additional possibility. A major objective of the present study was to correct this oversight. A secondary objective was to attempt to assess the separate contributions of the several factors that the study indicated were important, including the water-penetration mechanism. Apart from the steps needed to modify the process model to account for variations in the retained austenite content, the study consisted essentially of two subsidiary studies. The first was

Figure 3. Graphs showing the variation of maximum undercooling of each of the variables and associated functional relationships

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TABLE II. ESTIMATED CONTRIBUTIONS OF VARIABLES TO ΔTM OF PM VARIANTS (STUDY 2) PM Variant 1 Variable v/o Porosity Ms (°C) v/o Austenite k/ko Total

Variable ΔT Value (°C) 10.0 325 3.3 1.05 -

24.4 8.4 2.1 2.6 37.5

PM Variant 2

% of Total

Variable Value

ΔT (°C)

% of Total

65.1 22.4 5.5 6.9 100.0

15.0 340 1.9 1.05 -

29.9 10.0 3.3 2.6 45.9

65.2 21.9 7.2 5.7 100.0

directed to reproducing the findings of Saritaş. As expected, the simulations produced results that were reasonable facsimiles of the experimental findings. In addition to the water-penetration mechanism, the general implications were that increases in Ms, decreases in the retained-austenite content, and increases in thermal conductivity had all contributed to the observed cooling-rate increases. In the cases of the PM variants of the studies, the Ms increases were speculated to be due, in part, to an effect of porosity and, in part, to nickel heterogeneities in the steel. Otherwise, the decreases in retained austenite were considered to be due solely to an effect of porosity, and the increases in thermal conductivity to a separate effect of the nickel heterogeneities. The second study sought to quantify the contributions of the several factors to the cooling-rate increases as manifest by the maximum undercooling values. As applied to the PM variants of the latter findings, the results of the study suggested that ~65% of the increased cooling was due to water penetration of the pores, ~20% to the combined effects of increases in Ms and decreases in the retained-austenite content, due also to porosity, and the balance (~15%) to increases in both Ms and thermal conductivity due to the presumed nickel heterogeneities in the steel. As a final matter, in view of the fact that virtually all of the above is the product of speculation, it now remains to confirm or refute the findings of the simulations experimentally. ACKNOWLEDGEMENTS The authors wish to thank Professors A. Zavaliangos, R.D. Doherty, and A. Lawley, Drexel University, for input, and Hoeganaes Corporation for access to its technical library and services, which helped both in the study and the preparation of the manuscript. Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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REFERENCES 1. F.J. Semel and D.A. Lados, “Simulated Effects of Martensite Start Temperature, Thermal Conductivity, and Pore Content on End Quench Cooling Rate”, Powder Metallurgy, 2009, vol. 52, no. 4, pp. 282–290. 2. S. Saritaş, R.D. Doherty and A. Lawley, “Effect of Porosity on The Hardenability of P/M Steels”, Int J Powder Metall., 2002, vol. 38, no. 1, pp. 31–46. 3. A. Lawley, D. Stiles, M.L. Marucci and R.J. Causton, “Effects of Porosity on the Thermal Response, Hardness, Hardenability and Microstructure of P/M Steels”, Advances in Powder Metallurgy & Particulate Materials—2003, compiled by R. Lawcock and M. Wright, Metal Powder Industries Federation, Princeton, NJ, 2003, part 7, pp. 144–153. 4. V.S. Warke, R.S. Sisson and M.M. Makhlouf, “A Model For Converting Dilatometric Strain Measurements to Fraction Phase Formed During the Transformation of Austenite to Martensite in Powder Metallurgy Steels”, Metall. Trans. A, 2009, Vol. 40A, pp. 569–572. 5. G.F. Bocchini, A. Baggioli, R. Gerosa, B. Rivolta and G. Silva, “Cooling Rates of P/M Steels”, Int J Powder Metall., 2004, vol. 40, no. 5, pp. 57–65. 6. A. Zavaliangos, R.D. Doherty and A. Lawley, “Effect of Water Penetration on the Cooling of P/M Jominy Bars”, Process Modeling in Powder Metallurgy & Particulate Materials, compiled by A. Lawley, J.E. Smugeresky and L. Smith, Metal Powder Industries Federation, Princeton, NJ, 2002, pp. 92–98. 7. M. Hunt, “Guide to Engineered Materials”, Advanced Materials & Processes, 2000, vol. 158, no. 6, p. 59. 8. C.J. Smithells, E.A. Brandes and G.B. Brooks, Metals Reference Book, 7th edition, 1992, Elsevier Science & Technical Books., London, UK, pp. 14-08–14-29. 9. Z. Guo, J.Ph. Schillé, N. Saunders and A.P. Miodownik, “Modeling Material Properties Leading to Distortion Prediction During Heat Treatment of Steels for Automotive Applications” The 2nd International Conference on Heat Treatment and Surface Engineering in Automotive Applications, 2005, http://www.sentesoftware.co.uk/ downloads/articles-and-papers.aspx. 10. D.P. Koistinen and R.E. Marburger, “Austenite to Martensite Transformation”, Acta Metallurgica, 1959, vol. 7, p. 59. 11. K.J. Irvine, F.B. Pickering and J. Gladstone, “The Effect of Composition on the Structure and Properties of Martensite”, JISI, 1960, vol. 193, pp. 66–81. 12. W. Steven and A.G. Haynes, “Martensite Start Temperatures of Alloy Steels”, JISI, 1956, vol. 183, pp. 349–359. 13. M.A. Zaccone and G. Krauss, “Elastic Limit and Microplastic Response of Hardened Steels”, Metall. Trans. A, 1993, vol. 24A, pp. 2,263–2,277. 14. P. Grootenhuis, R.W. Powell and R.P. Tye, “Thermal and Electrical Conductivity of Porous Metals Made by Powder Metallurgy Methods”, Proc. Physical. Soc., vol. 65B, 1952, pp. 505–511. 15. J.C.Y. Koh and A. Fortini, “Prediction of Thermal Conductivity and Electrical Resistivity of Porous Metallic Materials”, Int. J. Heat and Mass Transfer, 1973, vol. 16, pp. 2,013–2,021. 16. V. Raghavan, “Kinetics of Martensitic Transformations,” Martensite, First Edition, edited by G.B. Olson and W.S. Owen, ASM International, Materials Park, OH, 1992, p. 200. ijpm

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ENGINEERING & TECHNOLOGY

AS-SINTERED AISI 440C STAINLESS STEELS WITH IMPROVED HARDNESS AND CORROSION RESISTANCE Henry Ovri*, Chukwuemeka J. Ohaukwu**, Kudrathon Bahadirov***, Mikael Larson**** and Peter Kjeldsteen*****

INTRODUCTION AISI 440C (0.95–1.20 w/o C-17 w/o Cr-0.15 w/o Mn-0.0–0.75 w/o Mo-0.7–1.0 w/o Si), is a high-carbon, high-hardenability martensitic stainless steel characterized by good corrosion resistance in mild domestic and industrial environments, including fresh water, organic materials, mild acids, and various petroleum products. Other attractive properties include high strength, hardness, and wear resistance in the hardened-and-tempered condition. Typical applications include ball bearings and races, bushings, cutlery, chisels, knife blades, pump parts, surgical instruments, and valve seats. The limited formability of this group of wrought stainless steels, however, significantly restricts their utility. This drawback can be potentially overcome through PM processing and metal injection molding (MIM) because of their intrinsic near-net-shape capability.1,2 Full densification is necessary for obtaining optimum properties in structural parts. To this end, liquid-phase sintering additions to the powder prior to compaction and sintering are known to be an effective way to achieve full densification.3,4 Preliminary work on a PM AISI 440C steel with 1.2 w/o C revealed the potential for achieving the required hardness in the as-sintered condition. Two groups of samples were compacted at a pressure of 550 MPa and sintered at 1,170°C and 1,220°C, respectively. About 500 ppm of boron in the form of FeB was added as a liquid-phase-sintering agent. The samples sintered at 1,220°C exhibited excessive liquid phase and were warped. Sintered densities of up to 99% of the pore-free density were achieved in some of the samples. While the as-sintered hardness (576 HV) was acceptable and comparable with that of the heat-treated alloy,3,4 the corrosion resistance was inferior and unacceptable; the alloy was completely covered with rust within 24 h in a salt-spray accelerat-

Preliminary work on powder metallurgy (PM) AISI 440C steel with 1.2 w/o C indicated that it is possible to achieve acceptable hardness in the as-sintered condition, but with inferior corrosion resistance. To address this limitation, a series of simulations were carried out using Thermocalc software to predict alloy compositions and sintering temperatures for achieving optimum properties. The simulations served as the basis for the design of a Taguchi trial run with an L9 orthogonal array. Test samples were characterized utilizing microindentation hardness and salt-spray-corrosion testing. Microstructural analysis was also used to characterize and interpret the observed behavior. The results show that AISI 440C with 0.95 w/o C and 200 ppm B sintered at 1,220°C gives an optimum combination of the assintered properties.

*Student, Institute of Materials Research (WME), Helmholtz-Zentrum Geesthacht, Max-Planck-Str. 1, D-21502 Geesthacht, Germany; E-mail: [email protected], **Student, Institut für Werkstoffphysik und -technologie Technische Universität Hamburg-Harburg, Eißendorfer Straße 42, 21073 Hamburg, Germany, ***Student, Department of Ceramic and Glass Engineering, Campus Universitário de Santiago, 3810-193 Aveiro, Portugal, ****Associate Professor, Department of Mechanics and Manufacturing Engineering, Aalborg University, Pontoppidansträde 101, 9220 Aalborg, Denmark, *****R&D Manager, Sintex A/S, Jyllandsvej 14, 9500 Hobro, Denmark

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ed corrosion test. The primary focus of this work was therefore to develop a PM 440C martensitic stainless steel with high hardness and improved corrosion resistance in the as-sintered state. Optical micrographs of the preliminary test samples sintered at 1,170°C confirmed the presence of large amounts of carbides around the prior austenite grain boundaries. These carbides, usually in the form of M23C6 and M7C3, lead to a depletion of chromium at the grain boundaries.1,5,6 It has been argued that high sintering temperatures in the range 1,260°C–1,350°C generally lead to superior corrosion resistance and enhanced mechanical properties.1,2,4,7 These studies suggest that high sintering temperatures lead to a more complete reaction between residual oxygen and carbon, and hence result in higher property levels. In the present study, a number of simulations were carried out with the aid of Thermocalc software in order to predict alloy compositions (in terms of carbon and boron content) and sintering temperatures that will result in optimum selected properties. Thermocalc software is a comprehensive combination of thermodynamic models that makes it possible to predict material compositions, microstructures, and properties resulting from various processing conditions.8 The results of the simulations formed the basis for the design of a Taguchi trial run using an L9 orthogonal array. An orthogonal array is a systematic experimental setup designed by Taguchi to facilitate the identification of key parameters that have the most effect on a given property, without a prohibitively high level of experimentation. It also facilitates the determination of optimum working conditions/ parameters.9,10 METHODS Simulations with Thermocalc The influence of boron and carbon on the amount of liquid phase formed as a function of the sintering temperature was simulated. Four levels of boron (200, 300, 400, and 500 ppm) and three carbon levels (0.95, 1.05, and 1.20 w/o) were used in the simulation. These variations in the boron and carbon levels were selected after careful analysis of the preliminary test results. The simulations showed that the amount of liquid phase increased with an increase in the boron content, the effect being most significant at 1,260°C, Figure 1. Although the influence of boron is much smaller between 1,170°C and 1,220°C, it is worth noting

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that a lower amount of liquid phase is formed at a boron content of 200 ppm for the different carbon levels. By decreasing the boron content to 200 ppm while keeping the carbon content constant, sintering temperatures of up to 1,220°C can be used without excessive liquid phase. This is based on the observation that a liquid-phase content of 3–4 w/o was sufficient to achieve liquid-phase sintering in the preliminary tests. Less than 2.6 w/o of liquid phase is formed at 1,220°C for the three carbon levels simulated. The amount of liquid phase formed at 1,260°C was >12 w/o in all cases, so sintering at this temperature may not be possible. It was also observed that an increase in the carbon content results in a gradual increase in the amount of liquid phase formed. This effect, however, is not as pronounced as that attributed to boron. Boron is obviously a more effective liquidphase former than carbon. Corrosion resistance is sensitive to the level of chromium dissolved in the face-centered cubic (FCC) matrix before cooling; a minimum of 12 w/o Cr in the matrix is required for corrosion protection.5,11,12 The precipitation of chromium in the form of carbides and borides is primarily responsible for the loss of corrosion resistance in stainless steels. The simulations predicted the presence of M23C6, M7C3, Cr5B3, and CrB at temperatures within the sintering range. The chromium content in the FCC phase was found to gradually increase with a decrease in the carbon level but to increase sharply with temperature, Figure 2(a). Based on these simulations, we presumed that by using an alloy powder with reduced carbon and boron levels, and by sintering at temperatures ~1,220° followed by rapid cooling, improvement in corrosion resistance may be expected as more chromium may be retained in the FCC matrix. According to Sourmail and Bhadeshia,5 M23C6 (the main carbide phase at temperatures <1,000°C), requires long-range diffusion in order to precipitate; hence, this phenomenon can be avoided by rapid cooling from high temperatures. Since reducing the carbon content may lead to a decrease in hardness, we also determined the amount of carbon in the FCC phase. Only a small increase in the carbon content of the FCC phase was observed when the carbon content of the alloy was increased, Figure 2(b). Consequently, since most of the carbon ends up in the form of carbides, it may be beneficial to decrease the carbon level as Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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AS-SINTERED AISI 440C STAINLESS STEELS WITH IMPROVED HARDNESS AND CORROSION RESISTANCE

Figure 1. Liquid-phase content in (a) 0.95 w/o C and (b) 1.2 w/o C AISI 440C

Figure 2. (a) chromium content and (b) carbon content in FCC phase vs. temperature for different alloy compositions (Thermocalc simulations)

this will result in a decrease in the amount of carbides formed. Taguchi Trials The simulations predict that sintering temperature, and the carbon and boron levels will have a significant effect on the chemical composition of the FCC phase and, by extension, the corrosion resistance and mechanical properties of the test steel. A Taguchi trial run for an L9 orthogonal array with three (3) control factors at three (3) levels was designed to determine the combinations of these parameters that result in optimum values of the desired properties. The control factors and the corresponding levels are shown in Table I, while details of the orthogonal array are given in Table II. Test samples were prepared using the experimental setup for the L9 orthogonal array (an array with nine experiments) as proposed by Taguchi.13 The alloy combinations cited in Table II were Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

achieved by mixing the required amounts of prealloyed 434L powder (16.8 w/o Cr-0.024 w/o C-0.94 w/o Mo-0.82 w/o Si-0.13 w/o Mn), FeB, pure graphite, and Acrawax (lubricant) in an automated mixer for 15 min. Dog-bone tensile specimens were then compacted at a pressure of 650 MPa. Sintering of the samples at the specified temperatures was carried out in a belt furnace in hydrogen at a dew point of –65°C at the inlet. The dew point was later increased to –35°C when the samples were in the TABLE I: CONTROL FACTORS AND CORRESPONDING LEVELS Control Factor Temperature (°C) Carbon Content (w/o) Boron Content (ppm)

1

Level 2

3

1,170 0.95 200

1,195 1.05 300

1,220 1.20 400

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TABLE II. TAGUCHI TRIAL RUN FOR L9 ORTHOGONAL ARRAY, SINTERED PROPERTIES AND CALCULATED S:N RATIO Mean Sintering Carbon Boron Sintered Open Closed Total Mean Mean Temperature Content Content Density Porosity Porosity Porosity Hardness Corrosion* S:N Ratio S:N Ratio (°C) (w/o) (ppm) (g/cm3) (v/o) (v/o) (v/o) (HV) (h) (Hardness) (Corrosion) 1 2 3 4 5 6 7 8 9

1,170 1,170 1,170 1,195 1,195 1,195 1,220 1,220 1,220

0.95 1.05 1.2 0.95 1.05 1.2 0.95 1.05 1.2

200 300 400 400 200 300 300 400 200

7.60 7.62 7.63 7.43 6.90 6.96 7.61 7.50 6.97

0.1 0.0 0.1 0.2 6.7 2.5 0.1 0.2 3.2

1.2 1.0 0.9 3.3 6.3 7.0 1.1 2.4 6.3

1.3 1.0 1.0 3.5 13.0 9.5 1.2 2.6 9.5

478 428 419 630 454 577 671 636 462

26 18 18 18 114 34.5 114 43 34.5

53.59 52.62 52.44 55.96 53.15 55.21 56.53 56.07 53.28

31.14 27.87 27.87 27.87 45.67 8.89 38.05 8.89 8.89

* Time for 50% of surface to be covered with rust muffle. Metallographic specimens were prepared utilizing standard procedures and examined by means of optical microscopy. Etched specimens were also used for the determination of microindentation hardness. Density measurements were performed on each batch of specimens. Salt-spray corrosion testing (ISO 9227) was carried out and the percent area of the bars covered by red rust was recorded as a function of time. The mean corrosion time required for 50% of the sur-

face of the specimen to be covered by rust is given in Table II. RESULTS AND DISCUSSION A summary of the sintered properties is given in Table II. These data were used in the Taguchi analysis. Representative optical micrographs are shown in Figure 3. It can be seen from Table II that the specimens sintered at 1,170°C exhibited the least corrosion resistance. After sintering at 1,170°C, 50% of the sample surfaces were rusted

Figure 3. Optical micrographs of sintered samples, showing carbides in martensitic matrix. Etched for 10 s in Kalling’s I reagent

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after less than 26 h compared with up to 114 h in some of the specimens sintered at 1,195°C and 1,220°C. A similar trend was observed in hardness. The samples sintered at 1,170°C had the lowest microindentation hardness (HV) levels. Optimization entails finding the best combination of control factors (factors within our control, namely, temperature, carbon, and boron content) which is least sensitive to noise (factors which have an effect on performance but are beyond our control). Various performance statistics are available,9,13 but the one adopted in this analysis was “larger is better,” since higher values of mechanical properties and corrosion resistance are more desirable. This performance statistic is calculated from the relation: S:N = –10log (

Σ /) y-2

n

(1)

where S:N = signal/noise ratio, y = primary response (experimental trial value), and n = no of repetitions of each experiment. The calculated S:N ratios for the parameters test-

ed are also shown in Table II. The mean value of the S:N ratios for each level of the control factors was then calculated by averaging the S:N ratios at that particular level. For example, the average S:N ratio for 1,170°C is the average of the S:N ratios (Table II) from all the tests at this temperature. The information is plotted in Figure 4 and Figure 5. It can be seen from Figures 4 and 5 that the maximum hardness is achieved at a sintering temperature of 1,220°C, a carbon content of 0.95 w/o, and boron content of 400 ppm. A maximum in corrosion resistance is achieved at a sintering temperature of 1,220°C, a carbon content of 1.05 w/o, and a boron content of 200 ppm. From Figures 4 and 5, it can also be deduced that sintering temperature had the largest effect on hardness and corrosion. This is because the sintering temperature has the highest range of the average S:N ratio. Improvements in hardness will, therefore, depend primarily on the sintering temperature and to a lesser degree on the carbon and boron levels. Also, low carbon (0.95 w/o) and boron (200 ppm) levels were found to be beneficial. Thus, it is expected that the optimum combination of hardness and

Figure 4. Average S:N ratios showing effects of temperature, carbon content and boron content on microindentation hardness

Figure 5. Average S:N ratios showing effects of temperature, carbon, and boron content on corrosion resistance

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corrosion resistance will be achieved with 0.95 w/o C, 200 ppm B and a sintering temperature of 1,220°C. This conclusion agrees well with the results obtained from the Thermocalc simulations. The increase in hardness and corrosion resistance with increasing sintering temperature can be attributed to the increase in solubility of carbon and chromium in the FCC matrix, Figure 2. Furthermore, it can be seen from Figure 6 that there are more carbides precipitated at the low sintering temperature (1,170°C) than at higher temperatures; more carbides were also precipitated by the alloy with higher carbon content. Examination of the optical micrographs following sintering at 1,170°C reveals the presence of large carbides around the martensite grain boundaries, Figure 3. This is in contrast to the relatively small carbide particles and small grain size seen at higher temperatures. These carbides are most likely type M7C3 carbides since the M23C6 carbides are completely dissolved at the sintering temperatures used. Although M23C6 carbides are re-precipitated during cooling, the relatively fast cooling (1°C/min) will likely prevent this from occurring to any significant degree. These results agree well with published data.1,5,6 Analysis of the micrographs in Figure 3 also shows that both the grain size and carbide size decrease with increase in the sintering temperature and a decrease in boron content; the reason for this behavior is, however, not clear. Thus, the decrease in the weight fraction of carbides precipitated/retained in the FCC matrix with increasing sintering temperature may also provide an explanation for the improvement in properties. Improvements in corrosion resistance have also

been associated with high temperatures; it is argued that a more complete reaction between residual oxygen and carbon is achieved at sintering temperature >1,205°C.1,2,4 These experimental results show that the average percentage loss in carbon at sintering temperatures of 1,170°C, 1,195°C, and 1,220°C are 16%, 23%, and 17%, respectively. It is not totally clear why sintering at 1,195°C resulted in the highest carbon depletion; however, only a small dependence of carbon depletion on sintering temperature is observed. The results also show that the sintered density has an effect on corrosion resistance. All the samples with a low sintered density and a high level of open porosity exhibited a relatively higher corrosion resistance, Table II. Although there have been disagreements over the effects of sintered density on corrosion resistance, based on long-time saltimmersion testing, it has been established that increasing density in the presence of open porosity is not beneficial to corrosion resistance.2 Improvement in corrosion resistance, in particular crevice corrosion, is usually associated with the closure of surface pores. A decrease in crevice corrosion with increasing density, however, has also been reported and was attributed to increasing oxygen depletion in the pores and failure to maintain a passive layer.4 The results obtained in the present study are in agreement with these observations. In contrast, hardness was found to decrease with an increase in the carbon content at the different sintering temperatures, Table II. This deviation from the expected trend may be related to an increase in retained austenite as a result of the increase in the carbon level. It has been suggested

Figure 6. Amount of carbide precipitation in FCC matrix (a) 0.95 w/o C and (b) 1.2 w/o C. Simulated (Thermocalc) plots for 440C

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that solid-solution strengthening due to carbon and chromium in the matrix leads to a lowering of the Ms temperature. A decrease in the Ms temperature, however, increases the tendency for retained austenite.14,15 Since austenite is a relatively soft phase, an increase in its volume fraction is expected to result in a decrease in hardness. CONCLUSIONS This study was aimed at developing press-and-sinter 440C alloys with optimum hardness and corrosion resistance without heat treatment. On the basis of Thermocalc simulations, Taguchi analysis, and mechanical-property and corrosion-resistance measurements, the following conclusions can be drawn: (1) Liquid-phase sintering of AISI 440C with boron as a liquid-phase-sintering agent is beneficial in obtaining a high sintered density; however, alloy combinations with high carbon (1.20 w/o) and boron (500 ppm) levels result in excessive liquid phase. The results show that, by decreasing the boron content to ≤400 ppm, it is possible to utilize high sintering temperatures of ~1,220°C. (2) Sintering temperature has a critical influence on hardness and corrosion resistance. Chromium and carbon, responsible for attendant corrosion resistance and hardness, increase in the matrix by increasing the sintering temperature. A high sintered density in the presence of open porosity was also found to be detrimental to corrosion resistance. (3) Since different combinations of composition and sintering temperature are required for maximum values of hardness and corrosion resistance, there is a need for tradeoffs. Analysis of the results shows that a control-factor combination of 0.95 w/o C, 200 ppm B, and a sintering temperature of 1,220°C may give optimum combination of the two properties. (4) The as-sintered hardness is comparable with that of heat-treated steel and some improvement in corrosion resistance is also observed at some combinations of sintering temperature and alloy composition. The results of the simulations and Taguchi analysis suggest that a significant improvement in these properties can be achieved utilizing cooling rates higher than those used in our experiments. Higher cooling rates will significantly reduce the level of carbides precipitated. REFERENCES 1. P.K. Samal, J. Valko and J. Pannell, “Processing and Properties of PM 440C Stainless Steel”, Advances in Powder Metallurgy & Particulate Materials—2009, compiled by T.J. Jesberger and S. Mashl, Metal Powder Industries Federation, Princeton, NJ, 2009, vol.

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1, part 7, pp. 112–121. 2. ASM Specialty Handbook: Stainless Steels, edited by J.R. Davies, 2005, ASM International, Materials Park, OH. 3. ASM Handbook, Vol. 4: Heat Treating, 2007, ASM International, Materials Park, OH. 4. ASM Handbook, Vol. 7: Powder Metal Technologies and Applications, 1998, ASM International, Materials Park, OH. 5. T. Sourmail and H.K.D.H. Bhadeshia, “Stainless Steels,” University of Cambridge, http://www.msm.cam.ac.uk/ phase-trans/2005/Stainless_steels/stainless.html. 6. “Martensitic Stainless Steels for Knives Applications”, Computational Thermodynamics Inc., http://www. calphad.com/martensitic_stainless_steel_for_knives_part_ 1.html. 7. E. Maahn, S.K. Jensen, R.M. Larsen and T. Mathiesen, “Factors Affecting the Corrosion Resistance of Sintered Stainless Steel”, Advances in Powder Metallurgy and Particulate Materials—1994, compiled by C. Lall and A.J. Neupaver, Metal Powder Industries Federation, Princeton, NJ, 1994, vol. 7, pp. 253–271. 8. “Thermocalc Software”, http://www.thermocalc.com/.

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9. “Introduction to Taguchi Method”, http://www.ee.iitb.ac. in/~apte/CV_PRA_TAGUCHI_INTRO.htm. 10. “Taguchi Method”, http://en.wikipedia.org/wiki/ Taguchi_methods. 11. M.F. Ashby and D.R.H. Jones, Engineering Materials 2, 1992, Pergamon Press, Oxford, UK. 12. F.C. Aria and C.E. Pinedo, “Influence of Heat Treatment on the Corrosion Resistance of Martensitic Stainless Steel type AISI 420”, J. Matls Sci. Letters, 2003, vol. 22, pp. 1,151–1,153. 13. S. Fraley, M. Oom, B. Terrien and J. Zalewski, “Design of Experiments via Taguchi Methods”, http://controls.engin. umich.edu/wiki/index.php/Design_of_experiments _via_taguchi_methods:_orthogonal_arrays. 14. M.C. Tsai, C.S. Chiou, J.S. Du and J.R. Yang, “Phase Transformations in AISI 410 Stainless Steel”, Matls Sci and Engr., 2002, vol. A332, pp. 1–10. 15. A. Litwinchuk, F.X. Kayser, H.H. Baker and A. Henkin, “The Rockwell C Hardness of Quenched High Purity IronCarbon Alloys Containing 0.09–1.91% C”, J. Matls Sci., 1976, vol. 11, pp. 1,200–1,206. ijpm

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ENGINEERING & TECHNOLOGY

EFFECT OF SILICON ON VACUUM-CARBURIZING DEPTH OF IRON COMPACTS Krzysztof Widanka*

INTRODUCTION In iron-base alloys, silicon is a ferrite-forming element. It exhibits a relatively high solid solubility in α-iron (18.5 w/o over the temperature range 1,020°C to 1,200°C), decreasing to 15 w/o at ambient temperature. In contrast, the solubility of silicon in γ-iron is ~2 w/o.1 At silicon contents >2.15 w/o, γ-iron is no longer present in iron–silicon alloys. In the Fe-Fe3C phase diagram, silicon restricts the austenite phase field, shifting the eutectoid composition to carbon levels <0.77 w/o and to temperatures >727°C. In wrought structural steels, silicon is an important alloying element that increases heat resistance and high-temperature creep resistance, especially with other alloying additives such as chromium and molybdenum. It is the basic additive in valve steels containing up to ~3 w/o Si which are often exposed to oxidizing hot-exhaust gases. Silicon is also frequently used in acid-resistant steels and alloys. Iron alloys containing 12 w/o to 18 w/o Si are resistant to sulphuric and nitric acids. Silicon steels are also widely used as soft magnetic steels and electrical steels. Sintered steels with up to 1 w/o Si are used primarily in stainless steels with chromium and in acid-resistant steels with chromium, nickel, and molybdenum.2 Iron-base soft magnetic alloys are another important group of materials manufactured by powder metallurgy (PM) with silicon as the main alloying addition. In these alloys, the average silicon content is in the range 1 w/o to 4 w/o (iron–silicon alloys) and 1 w/o to 3 w/o (iron–silicon–phosphorous alloys).2 Sintered parts can be vacuum carburized to enhance the control of carburizing depth in comparison with conventional gas carburizing since it allows the thickness of the carburized layer to be produced in a more predictable way. This method ensures a faster carburizing process, due primarily to the higher temperature and low hydrocarbongas pressure during the process.3–7 The carburized depth is related to the carbon diffusion rate into the material, which depends not only on the process parameters but also to a large extent on the characteristics of the material. In the case of products made by PM, the carburized depth depends, in addition to chemical composition, on sintered densi-

This study focused on the effect of silicon on vacuum carburizing of iron compacts with a density >7.2 g/cm3. An attempt was made to determine the effectiveness of silicon on increasing the carbon diffusion rate into the compacts. To reduce the influence of porosity, the level of interconnected porosity was minimized. Vacuum carburizing of the compacts with silicon additions was carried out at 1,050°C in a laboratory vacuum furnace. The effect of silicon over the range 0.5 w/o to 2.0 w/o on the vacuum-carburizing depth was analyzed. It was found that silicon additions up to 1 w/o increased the carburized depth by ~35% compared with iron in the absence of silicon.

*Professor, Wroclaw University of Technology, Institute of Materials Science and Applied Mechanics, Materials Science Division, Wroclaw, Poland; E-mail: [email protected]

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ty and the level of interconnected porosity. In this study, an attempt was made to evaluate the effectiveness of silicon on the vacuum carburizing of iron compacts with a density >7.2 g/cm3. The total level of porosity in the compacts was ~8 v/o with minimal interconnected porosity. The main purpose was to analyze the possible effect of silicon additions on the carbon-diffusion rate into the iron by excluding the effect of interconnected porosity. Based on the open literature, silicon increases the activity of carbon in iron–carbon–silicon alloys. Therefore, it was assumed that increased carbon activity could accelerate the rate of diffusion into the iron, resulting in a faster and deeper carburized layer. MATERIALS AND EXPERIMENT Cylindrical test specimens 10 mm dia. × 10 mm were fabricated from a blend of commercial iron powder ASC100.29 (Höganäs AB) with an average particle size of 90 µm, and silicon powder Si AX0.5 (H.C. Starck) with an average particle size of 3.5 μm. Two-sided compaction was utilized at pressures in the range 750 to 850 MPa to obtain compacts with a green density >7.2 g/cm3. 0.7 w/o kenolube P11 was used as a lubricant. Chemical compositions and densities of the compacts are given in Table I. For each composition three specimens were prepared. Vacuum carburizing combined with simultaneous sintering was carried out in a laboratory vacuum furnace (Seco/Warwick). The processing parameters were selected on the basis of the results of our previous research in order to produce a carbon level at the surface in the range 0.7

w/o to 0.8 w/o. The carburizing parameters are listed in Table II. The carburizing atmosphere consisted of propane diluted with nitrogen. The working pressure in the furnace chamber was 2,000 Pa. A stable working pressure in the chamber during carburizing was maintained by cyclical dosing of the gas at a constant flow rate of 110 dm3h-1. After carburizing, the specimens were cooled in nitrogen at a rate ~7°C/s. A schematic of the vacuum carburizing process is shown in Figure 1. The carburized depth was determined by means of optical microscopy. Specifically the distance from the surface to the interior at which the carbon content was ~0.4 w/o was measured. This distance was identified on transverse cross sections of the specimens along a diameter by means of computer images (Multiscan, Computer Scanning Systems) corresponded to 50% pearlite (based on area) in the microstructure. For each measurement of the pearlite fraction, a minimum of five images of the microstructure were analyzed. RESULTS AND DISCUSSION Results for the silicon-containing compacts are given in Table III. The case-depth results for the

TABLE I. CHEMICAL COMPOSITION AND DENSITY OF COMPACTS Chemical Composition Si (w/o) Fe (w/o)

Alloy Fe-0.5 Si Fe-1.0 Si Fe-1.5 Si Fe-2.0 Si

0.5 1.0 1.5 2.0

Green Density (g/cm3)

remainder remainder remainder remainder

7.29 7.24 7.22 7.20

TABLE II. VACUUM-CARBURIZING VARIANTS Variant

Carburizing Temperature (°C)

Boost Time (min)

Diffusion Time (min)

Total Carburizing Time (min)

I II III

1,050 1,050 1,050

20 40 60

100 200 300

120 240 360

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Figure 1. Schematic of vacuum carburizing process: (a) variation of temperature, (b) variation of pressure Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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EFFECT OF SILICON ON VACUUM-CARBURIZING DEPTH OF IRON COMPACTS

pure iron compacts at a density of 7.25 g/cm3 are presented in Table IV. The largest case depth was found in Fe-1.0 w/o Si for all the carburizing variants. The thickness TABLE III. CASE DEPTH OF SILICON-CONTAINING COMPACTS Case Depth (mm) Carburization Variant Fe-0.5 w/o Si Fe-1.0 w/o Si Fe-1.5 w/o Si Fe-2.0 w/o Si I II III

2.5±0.1 3.4±0.2 4.0±0.1

2.7±0.2 3.6±0.1 4.3±0.2

2.6±0.1 3.4±0.2 4.0±0.2

TABLE IV. CASE DEPTH IN PURE-IRON COMPACTS Carburization Variant

Case Depth (mm)

I II III

2.0±0.1 2.7±0.2 3.2±0.1

2.4±0.1 3.2±0.2 3.8±0.1

values are ~35% larger than those obtained in the pure-iron compacts. Figures 2(a) and 2(b) show representative macro- and microstructures of the carburized layer in the Fe-0.5 w/o Si alloy after vacuum carburizing utilizing variant I. Figure 3 shows representative optical microstructures of the carburized layer for the Fe-1.5 w/o Si alloy after vacuum carburizing utilizing variant I. The microstructure of the carburized layer on Fe-0.5 w/o Si and Fe-1.0 w/o Si specimens consists of pearlite and a small amount of cementite precipitated during cooling on prior austenite grain boundaries (Figure 2(b)). The higher magnification micrograph of the carburized layer on Fe-1.5 w/o Si in Figure 3 and in Fe-2.0 w/o Si consists of pearlite and ferrite, and the level of ferrite increased with increasing silicon content. The presence of ferrite in the microstruc-

Figure 2. (a) macrostructure and (b) microstructure of Fe-0.5 w/o Si after vacuum carburizing at 1,050°C for 2 h

Figure 3. Microstructures of carburized layer on Fe-1.5 w/o Si after vacuum carburizing at 1,050°C for 2 h Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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ture was lower than that for an assumed carbon concentration of 0.8 w/o in the surface layer; in this case it can equal ~0.6%, based on the carbon–iron–silicon ternary-phase diagram.8 This probably results from the fact that surface layers in these specimens are less saturated with carbon during the boost time than are the surface layers, which contain 0.5 w/o Si and 1.0 w/o Si. This is because silicon increases the carbon activity and decreases its solubility in austenite, thus reducing the surface concentration in the carburized layer. However, silicon reduces the carbon stream for diffusion and increases the carbon-diffusion coefficient and accelerates carbon flow from the surface layer to the core. The ferrite seen in the carburized surface layer of Fe-1.5 w/o Si and Fe-2.0 w/o Si appears to confirm the influence of silicon on the carburization of the silicon alloys. From analysis of the relation between the carburized-layer thickness and silicon content, it can be seen that the carburized-layer thickness shows a maximum at 1 w/o Si. This can be explained by the smaller level of carbon required for diffusion. The smaller concentration of carbon in the carburized layer, the lower its rate of diffusion into the material. If the level of carbon is too small for diffusion, notwithstanding its higher diffusion coefficient in the presence of silicon, the carburized depth (measured as the distance from the surface to the depth at which the carbon concentration is ~0.4 w/o), is smaller in the compacts with a silicon content >1 w/o than in those with a silicon content <1 w/o. On the basis of these results, the factor of proportionality (k) was determined in equation (1) for the case depth of all the silicon-containing specimens and pure iron: D = k √–t

(1)

where D = case depth (mm), k = factor of proportionality, and t = total carburizing time (min). The values of k are given in Table V. The largest value of k, namely, 0.24, was obtained for Fe-1.0 w/o Si. Figure 4 shows a representative diagram of the relationship between carburized depth and total carburizing time (boost time + diffusion time) in Fe-0.5 w/o Si. CONCLUSIONS 1. The results indicate that 0.5 w/o to 2.0 w/o of

54

TABLE V. k FACTOR FOR IRON–SILICON ALLOYS AND PURE IRON Specimen

Silicon Content (w/o)

k

Fe-0.5 Si Fe-1.0 Si Fe-1.5 Si Fe-2.0 Si “Pure” Fe

0.5 1.0 1.5 2.0 -

0.214 0.236 0.221 0.206 0.174

Figure 4. Relationship between carburized depth and total carburizing time (boost time + diffusion time) for Fe-0.5 w/o Si

silicon additions increase the carburized depth of iron-base compacts. The largest increase in the carburized layer thickness was 35% in comparison with that in pure iron for a 1 w/o Si addition. 2. In comparison with pure iron, thicker carburized layers in silicon-containing specimens result from an increase in the carbon-diffusion coefficient in austenite due to a higher carbon activity in the presence of silicon. 3. Analysis of the relationship between the carburized-layer thickness and silicon content indicates that the layer thickness gradually increases up to 1 w/o Si and then declines slowly. This decrease is attributed to a lower surface-carbon concentration caused by the increased silicon content and is confirmed by the presence of ferrite in the carburized layer in compacts containing >1 w/o Si. REFERENCES 1. Binary Alloy Phase Diagrams, Second Edition, 1990, ASM International, Materials Park, OH. 2. P. Beiss, R. Ruthardt and H. Warlimont, Powder Metallurgy Data, Landolt-Börnstein Online, Springer Berlin Heidelberg, Germany, 2003, vol. 2A1.

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3. D.H. Herring and P.T. Hansen, “Heat Treating Ferrous P/M Parts”, Advanced Materials & Processes, 1998, vol. 153, no. 4, pp. 44CC–44GG. 4. D.H. Herring and J. St. Pierre, “Vacuum Carburizing of P/M Steels”, Progress in Powder Metallurgy 1987, compiled by C.L. Freeby and H. Hjorth, Metal Powder Industries Federation, Princeton, NJ., 1987, vol. 43, pp. 525–537. 5. R.G. Weber, “Vacuum Carburizing and Carbonitriding of Powder Metallurgy Ferrous Alloys”, Powder Metall. Intl.,

1983, vol. 15, no. 2, pp. 94–97. 6. D.H. Herring, “Pros and Cons of Atmosphere and Vacuum Carburizing”, Industrial Heating, 2002, vol. 69, no. 1, p. 45. 7. F. Preisser, R. Seemann and W.R. Zenker, “Update on Vacuum-Based Carburizing”, Advanced Materials & Processes, 1998, vol. 153, no. 6, pp. 84ll–84LL. 8. V. Raghavan, “The Carbon-Iron-Silicon System”, J. Alloy Phase Diagram, 1986, India, vol. 2 no. 2, pp. 97–107. ijpm

WANTED: USED HIGH TEMPERATURE FURNACE Used high temperature furnace wanted having the following specifications: 2400° - 2600° Fahrenheit or higher pusher furnace, with hydrogen and nitrogen protective atmospheres; ideally capable of running 12” x 12” x 1” ceramic pusher plate design; requires 4” - 6” of working height; will consider all sizes and manufacturers. Contact-Robert Henkle at (414) 298-8140 or at [email protected]

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PM INDUSTRY NEWS IN REVIEW The following items have appeared in PM Newsbytes since the previous issue of the Journal. To read a fuller treatment of any of these items, go to www.apmiinternational.org, login to the “Members Only” section, and click on “Expanded Stories from PM Newsbytes.”

Nonferrous Powder Plant Unaffected by Zinc Refinery Shutdown Horsehead Corporation’s metal powder plant in Palmerton, Pa., continues operating at a high level and is not impacted by the temporary shutdown of the company’s Monaca, Pa., zinc oxide and metal refinery. Horsehead makes airatomized zinc, brass, copper, phos-copper, and nickel silver powders and infiltrants in Palmerton.

Miba Gains, Acquires Friction Business Miba AG, Laarkirchen, Austria, reports a 36.8 percent rise in sales for the first half of its fiscal 2010–11 year to 203.1 million euros (about $265 million). All business segments scored robust incoming orders. Continuing its expansion path, Miba also reports the acquisition of the HOERBIGER Group’s off-road vehicle friction lining business in Schongau, Germany.

Coatings Company Sells 41 Percent Interest Abakan Inc., a Miami, Floridabased development company, will acquire a 41% interest in Powdermet, Inc., Euclid, Ohio. The stock purchase agreement covers an initial payment of $500,000 with a closing on or before September 30 for the remaining portion of the price.

Joint Venture for Electric Car Battery Materials Formed H.C. Starck GmbH, Goslar, Germany, will form a joint venture company, CS Energy Materials, with Chisso Corporation in Japan to develop and make lithium mixed oxides for high-performance lithium-ion batteries for electric cars. Starck will have a 49 percent interest in the new company and Chisso 51 percent.

Silicon Carbide and Silicon Nitride Powder Agreement H.C. Starck Ceramics, a subsidiary of H.C. Starck GmbH & Co. KG, Goslar, Germany, announces a new cooperation agreement with Krahn Chemie, a distributor in Hamburg, Germany, covering the European sales of alpha-silicon carbide and silicon nitride powders mixed with sintering additives and pressing agents. The products, used in high-performance ceramics and as pressready materials, are marketed under the trade names Starceram S (α-SiC) and Starceram N (Si3N4).

56

Ecka Granules Group Selects Buyer The insolvency administrators of the Ecka Granules Group, Fürth, Germany, have signed a definitive agreement selling substantially all of the company to Platinum Equity LLC, Los Angeles, Calif., owner of SCM Metal Products, industry sources report. Although details of the acquisition were not released, completion of the transaction is expected in the fourth quarter pending customary closing conditions.

Microwave Marketing Agreement Spheric Technologies, Inc., Phoenix, Ariz., and Victec Europe Limited, Befordshire, England, have signed an international cooperative marketing agreement to develop microwave technology, equipment, and applications. Each firm has exclusive territorial rights for the sale of microwave furnaces built by Synotherm Corp. Drexel Researchers Receive Major Grant for Battery Anode Materials The U.S. Department of Energy (DOE) has awarded a $1 million grant to Yury Gogotsi, Trustee Chair Professor, and Michael Barsoum, A.W. Grosvenor Professor, Department of Materials Science & Engineering, Drexel University, to develop novel battery anode materials. The grant is from the DOE Batteries for Advanced Transportation Technologies program aimed at increasing the performance and decreasing the cost of batteries for plug-in electric and hybrid electric vehicles. North American Overview Report at World Congress The North American powder metallurgy industry’s strong rebound continued into the summer of 2010, reported MPIF President Michael E. Lutheran in his comments at the Global Review session at the PM2010 World Congress in Florence, Italy. Iron

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powder shipments soared 62 percent to 240,000 short tons through August and copper and copper-base powder shipments rose 38 percent to almost 8,000 short tons for the first six months. DOE Awards $2.8 Million for Demonstration Industrial Fuel Cell Fuel Cell Energy, Inc., Danbury, Conn., with partner ACuPowder International, LLC, Union, N.J., have been awarded $2.8 million Department of Energy funding to demonstrate a combined power & heat fuel cell for industrial applications. The two firms will jointly develop and demonstrate the efficient use of a 300kw DFC300 fuel cell at the ACuPowder facility by utilization of clean electricity, hydrogen (co-produced by the unit), and usable highquality heat. PM Titanium Project Aimed at Increasing Dehumidifying Performance ADMA Products Inc., Hudson, Ohio, has received a $4.4 million award from the U.S. Department of Energy to develop a high-efficiency online membrane air dehumidifier for enabling sensible cooling in warm and humid climates. Serving as team leader for the three-year project, ADMA will work with the Pacific Northwest National Laboratory, Richland, Wash. Tungsten Mine Opens with Extended Life North American Tungsten Corporation Ltd. reports an increase of the probable mineral reserve at its Cantung Mine to 1.693 million tons, which extends the mine life to more than four years at its current Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

mill capacity. The mine is scheduled to begin operating again during October. Sandvik, Carpenter Establish Powder Technology Partnership Carpenter Technology and Sandvik Materials Technology have announced the formation of a strategic partnership that, according to the partners, will result in the creation of an integrated value chain from R&D and powder production to sales and application development for near-net-shape solutions. Under the agreement Sandvik has acquired a 40 percent share in Carpenter Powder Products AB, Torshälla, Sweden, while Carpenter has acquired a 40 percent share in Sandvik Powdermet AB, Surahammar, Sweden. Sodium-Doped Sputtering Targets The efficiency of CIS/CIGS (copper–indium–selenium/copper–in dium–gallium–selenide) solar cells can be significantly improved by sputtering a precisely controlled layer of sodium-doped molybdenum, according to a report by Plansee AG, Reutte, Austria. The company offers MoNa sputtering targets with full density, high purity, and a uniform fine-grain structure. Powder Maker Reports Record Third Quarter Swedish powder maker Höganäs AB reports its best third-quarter sales ever, MSEK 1,728 (about $258 million), a 45 percent increase over 2009. Net sales for the first nine months jumped 58 percent to MSEK 5,059 (about $755 million).

Carbide Joint Venture Established After signing a joint-venture agreement on October 1 with CB Carbide, Taiwan, Plansee’s CERATIZIT Group, Luxembourg, owns a 50 percent stake in one of Asia’s leading carbide manufacturers with eight production locations in Taiwan and China. The agreement will increase the Plansee division’s sales in Asia by 25 percent. GKN Automotive Unit Gains The Automotive unit of GKN plc, London, UK, which includes its Driveline and Powder Metallurgy businesses, posts 2010 thirdquarter sales of £800 million (about $1.28 billion), a 34 percent gain from 2009. The company reports PM’s trading margin at 7.9 percent. New Furnace Installed Surface Combustion, Inc., Maumee, Ohio, reports the installation of a new Allcase Batch Integral Quench furnace line at Bluewater Thermal Solutions’ plant in Greensburg, Ind. The furnace line includes two furnaces with atmosphere top-cool and oil-quench capabilities, two temper furnaces, an endothermic gas generator, a spray/dunk washer, and companion equipment. MIM Feedstock Maker Upgrades Web Site Advanced Metalworking Practices (AMP), LLC, Carmel, Ind., has redesigned its Web site adding new information on ADVAMET® feedstock for metal injection molding. The new site also includes AMP’s ISO 90012008 certification and feedstock material data sheets.

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Thorium-Free TIG Welding Electrodes PLANSEE, Reutte, Austria, reports that customers are switching to thorium-free tungsten welding electrodes to avoid health and handling risks associated with thorium, a weakly radioactive element. A thoriumfree electrode, such as the lanthanum-containing WL15 product, is a recommended alternative providing improved ignition properties and wear resistance, notes the company.

58

Hoeganaes Expands in China Hoeganaes Corporation, Cinnaminson, N.J., has opened a metal powder premixing and market-support facility near Shanghai in Danyang, Jiangsu Province, China. Serving Chinese customers with quick delivery of press-ready premixes, the new facility also includes a laboratory with technical support and qualityassurance services, and warehousing for the company’s full range of ferrous powder products.

New Workhorse Attritor Union Process, Inc., Akron, Ohio, has introduced the C-20 attritor for applications requiring continuous production of large amounts of materials such as ceramics, coatings, and metal oxides. Labeled a ‘workhorse’ machine, it uses 1/8-3/8 inch grinding media. ijpm

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MEETINGS AND CONFERENCES

2011 MIM2011 INTERNATIONAL CONFERENCE ON INJECTION MOLDING OF METALS, CERAMICS AND CARBIDES March 14–16 Lake Buena Vista (Orlando), FL MPIF* (NOTE: NEW DATES & LOCATION) PowderMet2011: MPIF/APMI INTERNATIONAL CONFERENCE ON POWDER METALLURGY & PARTICULATE MATERIALS May 18–21 San Francisco, CA MPIF* INTERNATIONAL CONFERENCE ON TUNGSTEN, REFRACTORY & HARDMATERIALS VIII Co-located with PowderMet2011 May 18–21 San Francisco, CA MPIF* ICM11 THE 11TH INTERNATIONAL CONFERENCE ON THE MECHANICAL BEHAVIOR OF MATERIALS June 5–8 Lake Como, Italy www.icm11.org TI – 2011 12TH WORLD CONFERENCE ON TITANIUM June 19–25 Beijing, China http://www.ti-2011.com/

BASIC PM SHORT COURSE July MPIF* 7TH INTERNATIONAL SYMPOSIUM ON RHENIUM AND TECHNETIUM July 4–8 Moscow, Russia [email protected] THERMEC 2011 7TH INTERNATIONAL CONFERENCE ON PROCESSING & MANUFACTURING OF ADVANCED MATERIALS August 1–5 Quebec City, Canada www.thermec2011.ca/ INTERNATIONAL CONFERENCE ON SINTERING 2011 August 28–September 1 Jeju Island, Korea www.sintering2011.org PM MACHINABILITY SEMINAR September MPIF* ILASS 2011 24TH EUROPEAN CONFERENCE ON LIQUID ATOMIZATION AND SPRAY SYSTEMS September 5–7 Estoril, Portugal www.ilass2011.org PM COMPACTING/TOOLING SEMINAR November MPIF*

2012 PowderMet2012: MPIF/APMI INTERNATIONAL CONFERENCE ON POWDER METALLURGY & PARTICULATE MATERIALS June 10–13 Nashville, TN MPIF* SUPERALLOYS 2012: TWELFTH INTERNATIONAL SYMPOSIUM ON SUPERALLOYS September 9–13 Champion, PA PM2012 YOKOHAMA POWDER METALLURGY WORLD CONGRESS & EXHIBITION October 14–18 Yokohama, Japan www.pm2012.jp/

2013 PowderMet2013: MPIF/APMI INTERNATIONAL CONFERENCE ON POWDER METALLURGY & PARTICULATE MATERIALS June 23–26 Chicago, IL MPIF*

2014 PM2014 WORLD CONGRESS May 18–22 Orlando, FL MPIF*

*Metal Powder Industries Federation 105 College Road East, Princeton, New Jersey 08540-6692 USA (609) 452-7700 Fax (609) 987-8523 Visit www.mpif.org for updates and registration. Dates and locations may change Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

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YEARLY CONTENTS INTERNATIONAL JOURNAL OF POWDER METALLURGY TABLE OF CONTENTS FOR VOLUME 46, NUMBERS 1–6, 2010

46/1 JANUARY/ FEBRUARY 2010 2 Editor’s Note 4 Consultants’ Corner Kenneth J.A. Brookes HEALTH & ENVIRONMENT 9 Update on REACH, the CLP Regulation, and Their Implementation in the European Union PM Industry O. Coube and P. Brewin RESEARCH & DEVELOPMENT 17 Optimization of Compressibility and Hardenability by Admixing and Prealloying N. Giguere and C. Blais ENGINEERING & TECHNOLOGY 31 Effect of Sintering Temperature on Static and Dynamic Properties of Sinter-Hardened PM Steels F. Chagnon OUTSTANDING TECHNICAL PAPER: PowderMet2009 43 Influence of Chemical Composition and Austenitizing Temperature on Hardenability of PM Steels P.K. Sokolowski and B.A. Lindsley DEPARTMENTS 55 PM Industry News in Review 56 Web Site Directory 64 Advertisers’ Index

46/2 MARCH/APRIL 2010 2 4 9 11

15 21

29 39

45 46 47 48

27 37

49

61

73 75 76 78 79 80

62

FOCUS: Microminiature Powder Injection Molding— Part II Full-Density Nanopowder Agglomerate Sintering of Injection Molded Iron–Nickel J.-S. Lee, B.-H. Cha and W.-K. You A Review of Computer Simulations in Powder Injection Molding S.J. Park, S. Ahn, T.G. Kang, S.-T. Chung, Y.-S. Kwon, S.H. Chung, S.-G. Kim, S. Kim, S.V. Atre, S. Lee and R.M. German Characterization and Simulation of Macroscale Mold-Filling Defects in Microminiature Powder Injection Molding S.G. Laddha, C. Wu, S.-J. Park, S. Lee, S. Ahn, R.M. German and S.V. Atre Sintering of Powder Injection Molded 316L Stainless Steel: Experimental Investigation and Simulation X. Kong, T. Barriere, J.C. Gelin and C. Quinard DEPARTMENTS PM Industry News in Review Meetings and Conferences Instructions for Authors APMI Membership Application PM Bookshelf Advertisers’ Index

FOCUS: Microminiature Powder Injection Molding—Part I Materials for Microminiature Powder Injection Molded Medical and Dental Devices R.M. German Metal and Ceramic Parts Fabricated by Microminiature Powder Injection Molding V. Piotter, T. Hanemann, R. Heldele, M. Mueller, T. Mueller, K. Plewa and A. Ruh High-Strength Powder Injection Molded 316L Stainless Steel L.-H. Cheng and K.-S. Hwang Nitriding Response of Microminiature Powder Injection Molded Titanium T. Osada and H. Miura DEPARTMENTS PM Industry News in Review Meetings and Conferences PM Bookshelf Advertisers’ Index

46/4 JULY/AUGUST 2010

46/3 MAY/JUNE 2010 2 Editor’s Note 5 Consultants’ Corner David Whittaker 9 PM’s New Growth Engine—Technology Development Peter K. Johnson 17 Exhibitor Showcase: PowderMet2010

Editor’s Note Newsmaker Animesh Bose PMT Spotlight On …Todd M. Jensen, PMTII Consultants’ Corner Joseph Tunick Strauss

2 4 6 7 9 11 13

Editor’s Note PMT Spotlight On …Jason R. Forster 2010 APMI Fellow Awards An Appreciation—Alan Lawley Consultants’ Corner Pierre W. Taubenblat 2010 Poster Awards 2010 PM Design Excellence Awards Competition Winners 20 PM World Congress in Florence 22 Axel Madsen/CPMT Scholar Reports ENGINEERING & TECHNOLOGY 25 State of the PM Industry in North America—2010 M.E. Lutheran 29 Industrial Sintering of Hybrid Low-Carbon 3Cr-0.5Mo-xMn Steels M. Selecká and A. Šalak RESEARCH & DEVELOPMENT 43 Solution Annealing and Aging of a MIM CoCrMo Alloy P.V. Muterlle, I. Lonardelli, M. Perina, M. Zendron, R. Bardini and A. Molinari DEPARTMENTS 53 PM Industry News in Review 55 Meetings and Conferences 56 Advertisers’ Index

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YEARLY CONTENTS 46/5 SEPTEMBER/OCTOBER 2010 2 Editor’s Note 4 Newsmaker ...Paul Beiss, FAPMI 7 Consultants’ Corner James G. Marsden, FAPMI FOCUS: PM Titanium 9 Powder Metallurgy Titanium—Challenges and Opportunities Z.Z. Fang 11 Status of Metal Powder Injection Molding of Titanium Randall M. German 19 Review of Titanium-Powder-Production Methods C.G. McCracken, C. Motchenbacher and D.P. Barbis 29 Cold Compaction and Sintering of Titanium and Its Alloys for Near-Net-Shape or Preform Fabrication M. Qian 45 A Critical Review of Mechanical Properties of Powder Metallurgy Titanium H. Wang, Z.Z. Fang and P. Sun 58 61 62 63 64

DEPARTMENTS PM Industry News in Review Meetings and Conferences APMI Membership Application PM Bookshelf Advertisers’ Index

Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

46/6 NOVEMBER/DECEMBER 2010 2 4 7 9 13

Editor’s Note Newsmaker ...Herbert Danninger, FAPMI PMT Spotlight On …Silvio Bartoletti Excellence in Metallography Award Consultants’ Corner John A. Shields, Jr.

RESEARCH & DEVELOPMENT 17 Microstructure Evolution of Gas-Atomized Iron-Base ODS Alloys J.R. Rieken, I.E. Anderson and M.J. Kramer 33 Potential Effects of Retained Austenite on End-Quench Cooling Rates in PM Steels F.J. Semel and D.A. Lados ENGINEERING & TECHNOLOGY 43 As-Sintered AISI 440C Stainless Steels with Improved Hardness and Corrosion Resistance H. Ovri, C.J. Ohaukwu, K. Bahadirov, M. Larson and P. Kjeldsteen 51 Effect of Silicon on Vacuum-Carburizing Depth of Iron Compacts K. Widanka 56 59 60 61 62 64

DEPARTMENTS PM Industry News in Review Meetings and Conferences APMI Membership Application PM Bookshelf Table of Contents: Volume 46, Numbers 1–6, 2010 Advertisers’ Index

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Ace Iron & Metal Co. Inc. ______________(269) 342-0185_____________________________________________________5 ACuPowder International, LLC __________(908) 851-4597 ________www.acupowder.com___________________________49 Ametek Specialty Metal Products _______(724) 225-6622 ________www.ametekmetals.com ________________________16 Asbury Carbons _____________________(908) 537-2908 ________www.asbury.com ______________________________42 Elnik Systems_______________________(973) 239-6066 ________www.elnik.com _______________________________41 Global Titanium _____________________(313) 366-5305 ________www.globaltitanium.com ________________________50 Hascor International Group ____________+10 210 225 6120 _____www.hascor.com _______________________________6 Hoeganaes Corporation _______________(856) 786-2574 ________www.hoeganaes.com______________Inside Front Cover Magnequench_______________________(65) 6415 0670 ________www.mqitechnology.com ________________________15 North American Höganäs Inc. __________(814) 479-2003 ________www.nah.com _________________________________3 Rio Tinto Metal Powders/ Quebec Metal Powders Limited ________(734) 953-0082 ________www.qmp-powders.com _________________Back Cover Robert Henkle ___________________________________________rhenkle@reinhartlaw.com _______________________55 SCM Metal Products, Inc.______________(919) 544-7996 ________www.scmmetals.com ______________Inside Back Cover Timcal _____________________________+41-91-873-2009 ______www.timcal.com_______________________________31

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To:___________________________________ Fax #: ______________________________________ Company: _________________________________________________________________________ Please send me more information on:_____________________________________________________ _________________________________________________________________________________ as advertised in the __________ issue of the International Journal of Powder Metallurgy. Please send information to: Name: Title: ________________________________________________________________________ Company: _________________________________________________________________________ Address:___________________________________________________________________________ City:____________________________ State:_______________ Postal Code: ___________________ Country:___________________________________________________________________________ Phone:___________________ Fax:___________________ E-Mail: ___________________________

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Volume 46, Issue 6, 2010 International Journal of Powder Metallurgy

58-64,C3,C4_MEETINGS_CONFERENCES 11/29/2010 10:51 AM Page 65

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