International Journal of Powder Metallurgy Volume 45 Issue 6

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45/6 Excellence in Metallography Awards Particle-Reinforced Precipitation-Hardening High-Speed Steels Diamond Cutting Tools with a Ni3Al+ Copper Matrix Surface-Treated Nd-Fe-B Magnets: Corrosion Behavior

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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) P. Blanchard (France) G.F. Bocchini (Italy) F. Chagnon (Canada) C-L Chu (Taiwan) O. Coube (Europe) H. Danninger (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 Nicholas T. Mares [email protected] Executive Director/CEO, APMI International C. James Trombino, CAE [email protected]

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45/6 November/December 2009

Editor’s Note PM Industry News in Review PMT Spotlight On …Leander F. Pease III, FAPMI Excellence in Metallography Awards Consultants’ Corner Myron I. (Mike) Jaffe

ENGINEERING & TECHNOLOGY 27 Particle-Reinforced Carbon-Free Precipitation-Hardening High-Speed Steels H. Danninger, C. Harold, F. Rouzbahani, H. Ponemayr, M. Daxelmüller, F. Simancik and K. Izdinsky

RESEARCH & DEVELOPMENT 37 Diamond Cutting Tools with a Ni3Al+ Copper Matrix K.S. Hwang and T.H. Yang

45 Corrosion Behavior of Sintered Surface-Treated Nd-Fe-B Magnets E.A. Martins, J.L. Rossi, M.C.L. de Oliveira, I. Costa and H.G. de Melo

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DEPARTMENTS Meetings and Conferences APMI Membership Application Table of Contents: Volume 45, Numbers 1–6, 2009 Advertisers’ Index Cover: Boundary region between a sinter-brazed joint and PM part. Photo courtesy Thomas F. Murphy, Hoeganaes Corporation.

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 © 2009 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]

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EDITOR’S NOTE

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n 2009, APMI International inaugurated its “Excellence in Metallography” award to recognize seminal metallographic input to a conference paper. The first recipient is Tom Murphy, Hoeganaes Corporation. It is my judgment that his detailed assessment of the use of metallography in identifying common problems in PM steels (PowderMet2009) has set the bar at a challenging level for subsequent years. Mike Jaffe is no stranger to the “Consultants’ Corner.” In his latest contribution, Mike provides pragmatic insight and guidance to readers’ questions in three diverse areas of PM processing: finish machining near-netshape PM parts to tight tolerances; avoiding failure when using shoulder dies (aka flange or step dies); and the effect of pores, oil, or resin impregnation on microindentation hardness determination. The “PMT Spotlight” is on Lanny Pease, a long-time consultant and sage to the PM industry. This is the first time that this column features an individual with PMT Certification at Level II. In the “Engineering & Technology” section, Danninger et al. detail the development of particle-reinforced carbon-free precipitation-hardening highspeed steels. Using fine alumina particles, the reinforced grades exhibit higher apparent hardness and enhanced cutting performance compared with monolithic alloy variants, with only a modest reduction in transverse rupture strength. In the first of two contributions to the “Research & Development” section, Hwang and Yang evaluate the performance of an intermetallic (Ni3Al) matrix alloyed with copper as an alternative to a conventional cobalt matrix in diamond cutting tools. In dry cutting, the performance of Ni3Al alloyed with copper is superior to that of unalloyed Ni3Al or a cobalt matrix. The second contribution, by Martins et al., focuses on the corrosion response of sintered neodymium–iron–boron magnets. Surface treatment in a Cr (III) solution improved corrosion resistance but was not as effective as phosphating. This difference is understood by means of microstructural characterization and the results of electrochemical tests.

Alan Lawley Editor-in-Chief

Again this year, R&D Magazine’s Annual R&D Awards illustrate the importance of technology in our society and the prominent role of materials, including particulates. These global awards recognize the most significant innovations and products from academe, government, and industry. A sampling of the list includes: Silicon carbide high-temperature (250°C) power modules (50 kW). Sandia National Laboratory (www.sandia.gov). Austenitic stainless steels with enhanced corrosion/oxidation resistance derived from a thin protective surface layer of alumina. Oak Ridge National Laboratory (www.ornl.gov). Super-hard steel wear plate for mining processing and industrial applications. Consisting of a nanostructured weld overlay, the steel exhibits extreme resistance to abrasion while maintaining toughness. The NanoSteel Company (www.nanosteel.com). NanoCoral dendritic platinum nanostructures for renewable energy applications. The extended crystalline metal structures include nanowire networks, holey (porous) nanosheets, and dendritic nanosheets, and are expected to impact the development of sensors, solar cells, electronics, and catalysis. Sandia National Laboratory (www.sandia.gov). The National Laboratories and small innovative technology companies continue to populate the R&D 100 Awards list!

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Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

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June 27–30 The Westin Diplomat Hollywood (Ft. Lauderdale), Florida

2010 International Conference on Powder Metallurgy & Particulate Materials For complete program and registration information contact: METAL POWDER INDUSTRIES FEDERATION ~ APMI INTERNATIONAL INTERNATIONAL 105 College Road East, Princeton, New Jersey 08540 USA Tel: 609-452-7700 ~ Fax: 609-987-8523 ~ www.mpif.org

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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.”

Investors Buy Makin Metal Powders Makin Metal Powders Ltd., Rochdale, England, has emerged from administrative control of KPMG, Manchester, according to industry and media sources. Several private investors in the U.K. and key Makin executives have purchased the nonferrous powder maker following talks with potential buyers from the U.S., Germany, and the Far East. Company Exits Copper Powder Production Citing dire economic conditions, United States Metal Powders Inc., Flemington, N.J., has ceased copper-based powder production. The company will concentrate on making atomized aluminum powder, shot, and flake under its subsidiaries AMPAL, Inc., in Palmerton, Pa., and Poudres Hermillon in France. PM Joint Venture Set for Dissolution Reflecting a marriage not made in heaven, Plansee Holding AG, Reutte, Austria, and Mitsubishi Materials Corp., Tokyo, Japan, have decided to dissolve Plansee Mitsubishi Materials Global Sinter Holding S.A. Each company owns a 50 percent stake in the joint venture PM parts operation, scheduled for dissolution on December 1, 2009.

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Pending Sale of Crucible Powder Operations Allegheny Technologies Inc. (ATI), Pittsburgh, Pa., has announced its intention to buy Crucible Compaction Metals and Crucible Research for $40.95 million in an auction held on September 21 as part of a U.S. Bankruptcy proceeding. The transaction is expected to close no later than October 31, 2009. Hoeganaes Expands in China To serve the growing PM market in China, Hoeganaes Corp., Cinnaminson, N.J., will open a premixing unit and warehouse in Danyang, Jiangsu Province, about 200 km from Shanghai. Construction will begin during the first quarter of 2010 and is scheduled for completion by mid-year. Nonferrous Powder Producer Refocuses on PM Business Horsehead Corporation is refocusing efforts to serve the PM industry, reports Paul Wagar, the new general manager of metal powder sales. The company is a major supplier of zinc and brass powders as well as copper, bronze, infiltrants, phos-copper, and nickel silver powders. MIM Operations Bought Parmatech-Proform Corporation, Petaluma, Calif., an ATW company, has purchased the metal injec-

tion molding business of Morgan Advanced Ceramics (MAC), New Bedford, Mass. Proform is a wholly owned subsidiary of ATW Companies, Warwick, R.I. Glass-Melting Electrodes Expertise H.C. Starck, Goslar, Germany, showed its expertise in supplying refractory metals for the glass industry during a recent conference at Ohio State University. Molybdenum, tungsten, and their alloys are used in melting, homogenizing, feeding, and shaping glass products. Despite 3Q Sales Dip Höganäs Sees Conditions Improving Net sales in the third quarter for powder maker Höganäs AB, Sweden, declined 24 percent to MSEK 1,193 (about $175 million). However, all markets continued to improve, according to CEO Alrik Danielsson, with September somewhat stronger than expected. Plansee USA Celebrates 70th Anniversary Plansee USA, Franklin, Mass., is celebrating 70 years in business. The manufacturer of refractory metal products was originally incorporated in 1939 under the name American Electro Metal Corp. by Paul Schwarzkopf, founder of Plansee AB, Reutte, Austria.

Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

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PM INDUSTRY NEWS IN REVIEW

Crucible Sales Finalized In conjunction with bankruptcy proceedings, Crucible Materials Corp., Syracuse, N.Y., has completed the sale of its nickelbased superalloy powder and research facilities in Pittsburgh, Pa., as well as a steel service center in Romeoville, Ill. Allegheny Technologies Inc., Pittsburgh, Pa., purchased the powder facilities, and Erasteel, a subsidiary of ERAMET, Paris, France, acquired the steel service center.

Net sales hit $1.4 billion, an increase over the first and second quarters of 2009.

Strong Numbers for FederalMogul Automotive supplier FederalMogul Corporation, Southfield, Mich., reports strong financial performance with positive thirdquarter earnings and cash flow.

New Powder-Screening Machine Minox/Elcan, Mamaroneck, N.Y., offers the MTS 2600 Tumbler 8.5-foot screening machine available in up to five decks. The new machine offers a

Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

High-Temperature Furnace Sale Centorr Vacuum Industries, Nashua, N.H., is building two high-temperature sintering furnaces for customers in the U.S. and South America. The models include a two-cubic-foot furnace and a 12-cubic-foot furnace for sintering high-temperature SiC and B4C materials.

three-dimensional motion that gives up to four times the output per square foot of screen area and is able to accommodate friable materials from 20 microns to 20 mm. Sales Agent Appointed Makin Metal Powders (UK) Ltd., Rochdale, United Kingdom, has signed up Jet Metals, Inc., St. Marys, Pa., as sales agent for its copper and copper-base powders. Founded more than 50 years ago, Makin recently emerged from administrative control. Award-Winning PIM Paper The Institute of Materials, Minerals and Mining has granted its second biannual bestpaper award in the journal

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PM INDUSTRY NEWS IN REVIEW

Powder Metallurgy to the paper “Optimisation of Process Conditions in Powder Injection Molding (PIM) of Microsystem Components Using a Robust Design Method.” Written by R. Urval, S. Lee, S.V. Atre, S.-J. Park, and R.M. German, the paper was published in the June 2008 issue. New Materials for Sputtering Targets Plansee Metall GmbH, Reutte, Austria, a manufacturer of sputtering targets for all types of metallic layers used in LCD displays, has added aluminum and copper sputtering targets to its product line. Made from molybdenum, copper, and aluminum, targets are available in different sizes for use in coating flat panel displays, from small screen displays for cell phones to large flat-screen TVs. PM2010 World Congress Abstract Deadline Approaches November 30, 2009, is the deadline for submitting technical paper and poster abstracts for the PM2010 World Congress to be held October 10–14, 2010, in Florence, Italy, according to the European Powder Metallurgy Association (EPMA), the event’s sponsor. ijpm

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Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

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SPOTLIGHT ON ...

LEANDER F. PEASE III, FAPMI PMTII Education, the study of powder metallurgy (PM) and particulate materials, and interest in engineering/science. I am proud to say that I graduated from Classical High School, Providence, Rhode Island, in 1955. In the early 1950s, we all read that nuclear power was the future and that it would be so cheap that they would not bother with electric meters in people’s homes. That thought encouraged me to pursue and receive an ScB in Physics at MIT in 1959. During summer vacations I worked at Metals and Controls Nuclear on fuel for nuclear reactors. We had regular visits from Admiral Rickover, for whom I once delivered a gift pen to an employee at the train station in Providence, which person was being remembered for having helped the admiral in his travels. In the fall of 1958, having taken note of the importance of metallurgy in making nuclear fuel, I enrolled in John Wulff’s introductory metallurgy course at MIT. John was a physicist-turned-metallurgist and asked if I wanted to go to graduate school in metallurgy. He and Jere Brophy became my advisors, leading to an MS in Metallurgy in 1962 and an ScD in Physical Metallurgy in 1963. John Wulff was well known in PM and taught the subject. However, thinking it was not an arm’s length transaction to take a course from one’s thesis advisor, I never took the PM course. During the spring of 1963, I was a Ford Post-Doctoral Fellow working on superconductivity, an activity which I am well pleased not to have continued. From the fall of 1963 to 1968, I was a parttime MBA student, some days and some nights, at the Sloan School of Management at MIT and Northeastern University. The MBA was awarded from Northeastern University in 1968. What was your first job in PM? What did you do? Jere Brophy was a consultant to Wakefield Bearing Corporation and introduced me to its owner, Nathaniel Clapp, who hired me in the fall of 1963. Brophy’s advice was to bone up on PM by reading W.D. Jones’ tome. I learned some things from Jones but I learned a lot more seeing PM parts being made and working on various Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

production problems and customer issues. My position was vice president, technology. Describe your career path and companies worked for, and responsibilities. In 1968 I moved to the Kennecott Copper central research laboratory in Lexington, Massachusetts. The position was program manager, with the task of finding large-scale uses for the newly developed QMP Atomet powders. Kennecott was a joint owner (with New Jersey Zinc) of the iron-powder-making operation in Sorel, Québec. We investigated powder forging of loose powder, sonic vibratory compaction, and spark sintering. In the fall of 1970, Raymond Collette, William Hooper, and I found financing through the W.H. Nichols Co. to start a conventional press-and-sinter business, Sinterbond Inc., in Gloucester, Massachusetts. The Nichols company had found me because I was listed in APMI’s Who’s Who in Powder Metallurgy! In 1973, we sold our interest to the company, which moved the operation to Portland, Maine, where it continues as a division of Parker Hannifin Corp. In March 1973, Raymond Collette and I started Powder-Tech Associates, to offer consulting and engineering services, including a PM pilot plant with presses, furnaces, and a metallurgical laboratory. With Ray’s passing in 1991, my wife Barbara joined the company as treasurer, in which position she handles all of the company’s financial matters. What gives you the most satisfaction in your career? I enjoy working on, and solving, technical problems related to PM. We answer a broad range of questions, for example, why did this part crack? was it green or sinLeander F. Pease III, FAPMI PMTII President Powder-Tech Associates, Inc. 31 Flagship Drive North Andover, Massachusetts 01845-6194 Phone: 978-685-6027 Fax: 978-683-5733 E-mail: [email protected]

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SPOTLIGHT ON ...LEANDER F. PEASE III, FAPMI PMTII

tered? why did this part not machine like it used to? and is the part likely to pass MPIF Standard 35? Our pilot plant offers simple shapes for use in machining prototype parts. I enjoy talking with personnel from the different PM houses or end users about the blanks that they need. Since 1963, I have had a chance to meet many people in our industry. These contacts have helped me to answer questions or problems by referring inquiries to the individuals who know more than I do. Bill West and I co-authored an introductory PM text in 2002, and I am pleased to continue to get positive feedback from readers, including the fact that it helped them with their preparation for PMT certification. List your MPIF/APMI activities. Starting in 1978, I became a consultant to, and exofficio member of, the MPIF Standards Committees for structural parts and, later, for bearings and metal injection molding (MIM). For the standardization work, we prepared test bars for sintering by others, tested them, and then wrote the reports which the Standards Committee members used in setting the values that are cited in MPIF Standard 35. We also prepared the materials for the Test Method Assurance Program. For many years I was a member of the Program Committee which organizes the annual MPIF conferences and World Congresses. My service on the MPIF Technical Board under Bill Jandeska led to representing APMI on the Technical Board and becoming an APMI Director. After achieving PMT II status, I enjoyed working with Professor Alan Lawley and others in preparing questions for the PMT I and PMT II examinations. Finally, I was nominated for, and was elected, a Fellow of APMI in 2001. What major changes/trend(s) in the PM industry have you seen? In 1963 using sponge iron powder, any density >6.8 g/cm 3 was double pressed-and-sintered. Today, with small amounts of special lubricants and heated tooling, for reasonably shaped parts, we achieve 7.4 g/cm3. In 1963, sintering atmospheres were endo gas and exo gas with their battles to control the carbon content. These atmospheres have been supplanted by dry (low– dew point) cryogenic nitrogen and hydrogen with attendant control of carbon. The blisters and alligators of 1963 are largely gone today, with rapid burn-off preheat zones, adding oxidants to the delubrication zone, and avoiding zinc stearate as a lubricant. Alloy sintering belts still result in most parts manufacturers sintering iron at

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1,110°C–1,121°C (2,030°F–2,050°F), but many companies now have ceramic belts or walkingbeam furnaces that run at 1,316°C (2,400°F). In 1971, Raymond Wiech and Raymond Millett began making mixtures of metal powders and binders at Cer-Plas, from which sprang Parmatech and the MIM industry. They were able to make the powders flow around corners and provide all the shapes that the press-and-sinter community could only dream about. Why did you choose to pursue PMT certification? For a consulting and engineering company, promotion of the company comes in many different ways: writing technical papers, giving lectures, and attending PM meetings. It was therefore logical that showing our clients that we had the two PMT certifications was another good way to promote our business. Also, preparation and review for the two PMT certification examinations is beneficial on a personal level. How have you benefited from PMT certification in your career? My career centers primarily on managing our small company. It is hard to say that a certain piece of business came from having PMT certification. It is one among several ways in which we try to show technical competence and integrity to our clients. What are your current interests, hobbies, and activities outside of work? As I write this in June 2009, the Pease’s peas are in blossom and gardening is my favorite hobby. We raise, eat, freeze, or store most of the fruits and vegetables that we consume. My major interest is running Powder-Tech Associates and, with my wife’s help, looking out for the welfare of three children and five grandchildren. Our 47-year-old daughter Jennifer suffers from a severe developmental disability, or, as we used to say more plainly, mental retardation. Jennifer is at home with us every other weekend. My wife Barbara and I spend time looking after her here and at the group home in which she lives. Barbara and I serve, or have served, on various human rights committees, area boards and regional boards, as well as reminding the legislature of the needs of people like Jennifer. 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 45, Issue 6, 2009 International Journal of Powder Metallurgy

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2009 EXCELLENCE IN METALLOGRAPHY AWARDS

GRAND PRIZE

THE USE OF METALLOGRAPHY TO IDENTIFY COMMON PROBLEMS IN PM STEELS Thomas F. Murphy, FAPMI*

ABSTRACT Metallography has proven to be one of the most accurate and effective means of evaluating materials and explaining their behavior. It is particularly valuable in the study of PM materials where it is the only set of test methods capable of generating information on the three elements responsible for determining PM properties, i.e., part density/pore structure, chemical composition as it relates to the alloying method, and the microstructure. Many of the procedures used in metallographic testing are well established and these procedures aid the metallographer in all aspects of the analysis, from sample selection, to preparation, examination, and, finally, to storage of images and documentation of findings. Of particular importance are the procedures describing correct sample preparation. The importance of proper sample preparation cannot be emphasized enough. Without the use of proper preparation techniques, test results are questionable. The product of improper grinding and polishing is shown in this paper, along with several examples of problem-solving and documentation techniques. Metallographic capabilities have also been advanced through improvements in computer and digital-imaging technologies. The almost universal use of digital-imaging systems, in conjunction with large-capacity computer storage of images, has produced several opportunities to advance metallographic testing. One of these opportunities, the production of large multi-image montages, will be discussed along with a solution to a problem: the estimation of grain size using a computer monitor for feature counting. The value of metallographic analysis is enhanced through these technological improvements, but the analyst must be aware of what is needed to perform the test correctly while recognizing the hazards of relying on computers and automation.

APMI is pleased to announce the new “Excellence in Metallography Award” to recognize the individual(s) responsible for metallography used to support and provide evidence for the ideas set forth in a conference technical paper. A specially selected Judging Committee evaluated all eligible manuscripts. An award presentation will be made at PowderMet2010 in Hollywood (Ft. Lauderdale), Florida, June 27–30, 2010. Awards contributed by: Buehler Ltd. Buehler Worldwide Headquarters 41 Waukegan Rd. P.O. Box 1 Lake Bluff, Illinois 60044-1699 USA Phone: 847-295-6500 Fax: 847-295-7979 www.buehler.com Precision Surfaces International, Inc. 922 Ashland Houston, Texas 77008 USA Phone: 713-426-2220 Fax: 713-426-2223 www.psidragon.com

* Scientist, Research & Development, Hoeganaes Corporation, 1001 Taylors Lane, Cinnaminson, New Jersey 08077-2017, USA; E-mail: [email protected].

Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

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2009 EXCELLENCE IN METALLOGRAPHY AWARDS

INTRODUCTION Metallography is probably the most effective diagnostic tool used in the evaluation of metal powders and the products made from them. It is particularly effective in failure analysis, helping explain material behavior, and as a monitor for control and improvement of products and processes. Samples to be analyzed can be chosen on a random basis or they may be selected for a specific reason, such as those containing a defect or having abnormal properties. The examination can be accomplished using multiple techniques with both light optical (LOM) and scanning electron (SEM) microscopes. It must be remembered that, regardless of the method chosen for examination, the information contained in the sample must be protected from damage and/or prepared correctly using the best preparation procedures. In many cases, one sample may be the only source of evidence available for examination. The process of metallographic analysis usually begins with sample preparation. When preparing for examination using a light microscope, a flat cross section is frequently removed from a bulk specimen using an abrasive saw. Mounting, grinding, and polishing of the planar section generally follow the sectioning step. Adherence to a well-thought-out preparation procedure usually results in a planar surface that is a true representation of the microstructure. Depending on the features of interest, examination can be in either the as-polished or chemically etched and/or stained condition. Additional analysis of these

Figure 1. Example of a poorly prepared cross section. The pores are filled with cut metal and grinding/polishing debris. LOM: unetched, brightfield illumination

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cross sections can be accomplished in an SEM, where both planar and non-planar surfaces, such as fracture or particle surfaces, can be used. Regardless of the choice of microscope utilized for the evaluation, improper sample handling and/or preparation can lead to incorrect assumptions and a compromised analysis. The microscope cannot compensate for poor preparation or careless handling. An example of a poorly prepared sample is shown in Figure 1, where the porosity is incorrectly presented. Cut metal and grinding/ polishing debris are trapped inside the pores and some pores remain unopened. The first half of this paper shows examples of metallographic techniques used in the analysis of problems and situations common to the PM industry. They are intended to demonstrate the versatility of the metallographic techniques showing a variety of microstructures and magnifications. The second part of the paper offers two opportunities aimed at improving metallographic analysis and documentation. They utilize the advancements in computing power, computer storage, and digital imaging now common in modern metallography laboratories. The techniques describe suggestions for performing grain-size measurement and the creation of multi-image montages from numerous single images. PRACTICAL EXAMPLES Figures 2 through 6 show several examples of samples evaluated using LOM both with and without chemical etching. As seen in these photomicrographs, the use of color imaging is sometimes required to see features or regions of interest accurately. In others, grayscale images are sufficient to capture the important characteristics. Several microscopy techniques and a variety of magnifications are used in these figures to illustrate the appearance of important characteristics within the microstructure. As a complement to the light microscopy images in these examples, Figure 7 illustrates the value of imaging fracture surfaces using an SEM. In all of the examples, the illumination technique and etching condition are shown along with a linear scale to allow estimation of feature sizes. Example 1—Sinter Quality Figure 2 illustrates the visual difference between a poorly sintered sample compared with

Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

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2009 EXCELLENCE IN METALLOGRAPHY AWARDS

Figure 2. Comparison of (a) poor and (b) well-sintered samples. LOM: unetched, brightfield illumination

one that is well sintered. The image in 2(a) contains unsintered particle boundaries, angular pore edges, and clusters or necklaces of small pores located at prior particle boundaries. In 2(b), the image shows the improvement in sintering with the disappearance of particle boundaries and small pores, in addition to a smoothing of the individual pores as the surface area of the pore network in reduced. Example 2—Low-Temperature Sinter Figure 3 is an image of an admixed iron–copper–carbon alloy after sintering at a temperature

Figure 3. Iron–copper–carbon admix after a low temperature sinter. LOM: unetched, brightfield illumination

Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

below the melting point of copper (1,083°C). The temperature was sufficient to produce a marginal sinter, but the copper particles (orange features) remain unmelted. Example 3—Cracking in a Sinter-brazed Joint The interface between a porous PM part (textured, upper portion of the image) and the sinterbraze compound (smooth region along bottom edge) is shown in Figure 4. The problem with this assembly was cracking in the area adjacent to the brazed joint. Several cracks can be seen in the textured portion of the image.

Figure 4. Boundary region between a sinter-brazed joint and a PM part. (2 v/o nital + 4 w/o picral etch). LOM: differential interference contrast

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2009 EXCELLENCE IN METALLOGRAPHY AWARDS

The section was prepared using standard procedures and etched with a combination of 2 v/o nital and 4 w/o picral to reveal the microstructure. After etching, the mount was quenched in a liquid-nitrogen bath to transform any retained austenite to martensite. Surface relief can be seen in areas containing newly transformed martensite that was formed upon cryogenic quenching. It was determined that the cracking near the joint was caused by a large amount of austenite remaining after cooling from the sinter-braze operation. A reduction in the carbon content of the PM part resulted in a decrease in the amount of retained austenite and elimination of the cracking problem. Example 4—Induction Hardened Surface The low-magnification image in Figure 5 shows

Figure 5. Low-magnification image showing martensitic surface and pearlitic core of induction-hardened gear. (2 v/o nital + 4 w/o picral etch). LOM: brightfield illumination

the surface and core of an induction-hardened part after etching. The martensitic surface and the pearlitic core are clearly recognizable by the color of the etched microstructure. The hardened (martensitic) layer appears as the light tan area, the pearlitic core is brown (pearlite) with small white regions (ferrite), and the mounting material appears black and borders the martensitic surface. The depth of the hardened surface can be estimated using this image because the difference in the appearance of the transformation products is substantial. The amount of detail and resolution in this image is sufficient to accomplish the task. Example 5—Multiple-Image Montage of an Internal Green Crack Figure 6 is a 3-image montage (3 images × 1 image) showing a crack running through the interior of a green part. In this example, the shape and size of the feature of interest, the crack, necessitated the creation of a high-aspect-ratio image rather than one with the more standard 1.3:1 aspect ratio, similar to what is seen in Figures 1 through 5. Creation of the montage required careful movement of the microscope stage, uniform illumination of the image, and careful placement of each image with its neighbor. Example 6—Appearance of PM Fracture Surfaces In this example, the SEM was used to document the characteristic differences contained in the fracture surfaces of three materials. Chemical composition, processing, microstructure, and, in one case, handling of the part prior to sintering, all may contribute to the appearance of the fracture surface.

Figure 6. A 3-image montage showing a crack in a green part. LOM: unetched, brightfield illumination

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Figure 7(a) shows the appearance of a green crack after sintering. The particle surfaces are smooth and show no particle contact along the separated 3D area. The surface shown in Figure 7(b) has the appearance of an intergranular brittle surface. The surface is a combination of interior pore surfaces and flat features that are grain faces. A mixture of pore surfaces, ductile microvoids, and transgranular cleavage is seen in Figure 7(c). These images provide a significant amount of information on the manner by which the materials failed under stress. Although metallographic analysis is usually performed using LOM, a significant amount of additional information can be gained using the SEM. Where flat or nearly flat surfaces are usually required for examination using the light microscope, the SEM can be used to examine irregular samples, such as particle and fracture surfaces, in addition to planar surfaces. An additional benefit of SEM analysis is the possibility of performing chemical analysis on small, localized areas. The combination of LOM and SEM can provide significant information about the behavior of a material, the distribution of alloying elements, and many other characteristics.

Figure 7. Fracture surfaces showing three modes of failure. (a) sintered green crack, (b) intergranular brittle failure, and (c) combination of ductile and transgranular brittle failure. SEM: secondary electron images

Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

EFFECT OF DIGITAL IMAGING AND COMPUTING IMPROVEMENTS Digital imaging, combined with the larger storage capacity for photomicrographs and improvements in software capabilities, has opened many doors for the metallographer. Also, with improvements in resolution and color reproduction, digital-image capture is almost universally accepted throughout most areas of microscopy. Time-consuming image manipulations, previously done using a darkroom, can now be accomplished using computer software, usually with a significant time saving and often with improvements in quality and consistency. However, these advancements come with a price. The operator must now have a better understanding of the test being performed and what part the computer is playing to aid in the testing. Simply using a digital camera to acquire an image with a computer program to make measurements does not ensure accuracy of the test. An example of a difficulty caused through these advances, along with an opportunity, is presented in this section of the paper. The difficulty is meas-

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uring grain size using a computer monitor for counting, and the opportunity is the creation of multi-image montages to represent large-area sample surfaces. Grain Size Determination The procedure chosen for the estimation of grain size closely follows ASTM Standard Test Method E 112,1 specifically the Abrams Method described in section 14.3. Briefly, the image of a prepared and etched sample is either photographed or projected onto a screen or computer monitor with an overlay of three open concentric circles superimposed on the image. The circles act as the probes for sampling the image. As the operator moves along the perimeter of a circle, a manual counter is incremented once each time the circle crosses a grain boundary. Counts are made for the total length of the three circles and these values are retained on an individual field basis. Specifics on the application of the overlay and the counting details are described in ASTM E 112. In cases where an automated image-analysis system is used and the grain boundaries or individual grains are accurately and completely detected, this procedure is of little or no use. In that situation, the virtual length of the concentric circle overlay is accommodated by system calibration. If the grain boundaries cannot be automatically detected, this manual counting procedure is appropriate. Grain-size test is a stereological measurement, which uses the number of points encountered over a known line length (PL) to provide an average distance.2,3 In this case, the line length is the total of the three circle perimeters, as seen in Figure 8, and the points are the locations where grain boundaries are intersected by the circles. A simple division of the line length by the number of grain boundaries intersected gives the average intercept distance. This distance is a measure of grain size and can be used to calculate the ASTM grain-size number. Therefore, to perform the test accurately, two conditions must be met. The magnification of the image being analyzed must be accurately known and the grains must be clearly defined by etching, staining, or decoration of the grain boundaries. Calibration In performing the test, the magnified image is

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photographed or projected onto a glass screen or computer monitor and the three-circle overlay placed on the image. A serious problem arises when a computer monitor is used for counting and the magnification of the image on the monitor is not known. The magnification determines the virtual length of the sampling probes as they are projected on the image of the metallographically prepared specimen. In optical systems fitted with glass viewing screens, the magnifications of the optics and projection distances are known; therefore, the final magnification is also known and the concentric-circle sample probes can be used without difficulty. The situation is more complicated where computer monitors are used for viewing and counting. Compared with enclosed optical systems where all contributors to the final magnification are known, the addition of the camera, the camera-tube length, and monitor, each contribute to the final magnification in computerized systems. Under normal imaging and documentation circumstances using a digital camera and computer system, the magnification effects of all the contributors are compensated for during system calibration. Calibration assigns a linear distanceper-pixel value at each magnification by determining the specific number of pixels used to measure a known linear distance. Using this pixel-equivalent value, a line of appropriate or specified length is generated and it, along with a scale value, is annotated onto the image as the scale marker.

Figure 8. Illustration of concentric circles with horizontal and vertical reference markers. (ASTM E 112)

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Thus, the size of all features in the magnified image can be related to the scale on the image. Unfortunately, this calibration value is not useful in performing the grain-size analysis on the screen-projected images because the magnification on the screen is not known. Without knowledge of this magnification, the virtual lengths of the concentric circles relative to the image size are also not known. To remedy this situation, the length of the circle perimeters must be put into the same scale as the sample image, i.e., the length of the concentric circles must be reduced in apparent length by the same amount that the sample image is increased in apparent size. This creates a virtual line length that is determined by the magnification on the monitor and makes unnecessary the use of a scale marker on a photomicrograph containing the concentric circles. The process of calculating the magnification on the monitor can be accomplished easily in the following way. A standard-length linear scale, a stage micrometer usually one or several millimeters in length, is used as the standard reference distance. The scale is imaged at the appropriate magnification and a hard-copy print is made at the actual size (1:1) of the image on the monitor. This can be done for each optical combination used in testing if needed. The printed image of the stage micrometer is measured and the magnified length divided by the actual length of the stage micrometer. The result is the total magnification on the screen. An example can be seen in Figure 9 where an image of a 1 mm stage micrometer is shown. The length

Figure 9. Example of a magnified 1 mm stage micrometer

Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

of the micrometer image on the print is measured at 100 mm and the calculation is 100 mm/1 mm with the result being a magnification of 100x on the monitor. As a result, the virtual length of the concentric circles in the overlay is 1/100 the actual size. Using multiple magnifications on the same microscope it was found that the final magnification on the monitor divided by the total magnification of the optical components yielded a system constant. This constant is the effect of all nonoptical contributions in the system. The overlay can be created using any software drawing program and usually can be inserted into the field of view using an image-analysis system or imaging program. The total length of the three circle perimeters in the standard is 500 mm, but this can be altered if the visible field can accommodate larger or smaller circles. Larger circles take advantage of the larger area available with many imaging systems, although the relationship of the circles to one another must be the same as in the standard. Specifically, the diameter of the smallest circle is approximately one-half the intermediate circle and one-third the outside circle. Additionally, the circles are concentric, and the distance between the circles is constant. The test is performed by first randomly choosing a field, overlaying the concentric circles, then incrementing a counter at each occurrence where a grain boundary is intersected by a circle. This is a PL count where PL is the number of points encountered divided by a unit line length, and the boundaries are the points and the circle perimeters are the lines. In situations where a line crosses a pore located between grains, the pore is counted as one boundary. Sampled pores located within grains are not counted as boundaries. A sufficient number of fields must be analyzed to provide a total of at least 400 counts or the region measured must cover a specific area of interest. In cases where the sample contains porosity, the volume fraction of porosity is subtracted from the total line length and the adjusted line length is used in the calculation. At the conclusion of counting, the average number of counts-per-field is calculated and this is entered in equation (1) which gives the value of l , the average intercept distance between boundaries: l(adjusted)mm _ l = _____________________ magnification × C × N

(1)

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where l(adjusted) is the total line length with the volume fraction of porosity removed, the magnification is the total magnification of the optics, C_ is the non-optical constant for the system, and N is the average number of counts/field. To convert the line length value from millimeters to micrometers, multiply the total length of the circle perimeters in the numerator by 1,000. Delineation of the Grains Once the optical/viewing system is calibrated, the test is the same for any material. The task now is to faithfully and completely define the grain boundaries or the individual grains in order to perform the counts accurately. This can be performed in several ways and is determined by the material being analyzed. Standard test method ASTM E 112 is applicable to ferritic and austenitic alloys, and, with special sample processing, the prior austenite-grain size. The most common technique with ferrous alloys is to etch the prepared sample using a combination of nitric acid and ethanol (nital), in a concentration from 1 to 5 v/o. Figure 10 is an example of a nital-etched pore-free ferritic material. The majority of the grain boundaries are well defined; however, a few are faint and may cause a problem in counting. Several areas are emphasized within the red circles. An alternative approach is to pre-etch the sample, then stain the surface to separate the grains by color. A powder-forged (PF) iron powder is seen in Figure 11 after pre-etching with 2 v/o nital, then staining with Beraha’s 3-10. Examination is made using polarized light with a sensitive tint filter. The variation in color is a result of the thickness of the deposited coating, which is caused by variations in grain orientation. The three-circle overlay is superimposed onto this image and the coincidence of the grain boundaries and circle perimeters can be readily seen. The measurement of prior austenite-grain size can be difficult with some materials. Several etchants are described in the literature;4,5 however, they are not effective with all materials. One technique used to avoid the difficulties in the choice of etchant is to process the samples to a hypereutectoid carbon content and to decorate the austenite grain boundaries with precipitated carbides upon cooling. The material used in making the sample shown

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in Figure 12 was FL-4900 with a 0.9 w/o graphite addition. It was sintered at 1,120°C (2,050°F) and cooled to room temperature. After cooling, it was austenitized at 950°C (1,750°F), the temperature reduced to 745°C (1,375°F), and oil quenched. The reduction in temperature from 950°C to 745°C placed the sample in the two-phase hypereutectoid region (austenite + carbide) and resulted in the precipitation of proeutectoid carbides on the austenite grain boundaries. After metallo-

Figure 10. Pore-free ferritic material. This is a rolled iron powder that was recrystallized after rolling. (2 v/o nital etch). LOM: brightfield illumination

Figure 11. PF iron powder pre-etched, then stained to reveal the grain structure. The three-circle overlay is shown (2 v/o nital etch, then Beraha’s 3-10), a solution of 3 g potassium metabisulphite (K2S2O2) and 10 g anhydrous sodium thiosulphate (Na2S2O3 in 100 ml H2O). LOM: polarized light and a sensitive tint filter

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Figure 12. Porous FL-4900 with 0.9 w/o graphite added. Grain boundaries are decorated with carbides (2 v/o nital + 4 w/o picral). LOM: brightfield illumination

graphic sample preparation and etching, the grain structure was well defined and the estimation of the grain size was accomplished without difficulty. Montage Creation As noted previously, the use of digital imaging with image manipulation and computer storage has enabled the user to create multi-image montages, which can encompass large areas with the capability of retaining the original resolution of each individual image. Previously, montages were useful for documenting specific features of inter-

est with several photomicrographs, such as the green crack shown in Figure 6. However, the final montage contained little more than the smallest number of images needed to cover a particular feature. Technology has now provided the metallographer with the ability to use the montage to analyze large areas with individual, large images. These can be useful in examining spatial arrangements, non-uniform distributions, neighborhood relationships, feature distributions, and many others. Louis and Gokhale6 described the use of montages in the characterization of nearest neighborhood relationships, radial-distribution functions (RDF), k-functions, and pair-correlation functions. In addition, montages can be constructed of images from both light and electron microscopes. One drawback to the use of these images is that they often occupy a large amount of computer storage. It is not unusual to have a single montage occupy 100 to several hundred Mbytes of computer memory, especially if the montage is a color image. A cross section through a tested tensile bar is displayed in Figure 13 with several measurements showing the reduction in area near the fracture surface in this PM material. The precise location of these measurements would be difficult, if not impossible, without the use of a montage. In addition, a single low-magnification image would not recreate the cross section with the same level of accuracy. The cross section used to create the montage in

Figure 13. Image of cross section through failed tensile bar. Measurements confirm a measurable reduction in area before failure. The section was electroless nickel plated. LOM: unetched, brightfield illumination, total images—10 fields horizontal × 8 fields vertical

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Figure 14. Montage showing the same section as in Figure 13, but at a different resolution. Red dots are markers locating short, disconnected internal cracks. Inset image (a) shows the actual resolution of the montage in the area surrounded by the red box. The section is protected with a nickel coating, making it possible, in addition to other tests, to quantitatively analyze the entire fracture surface using this image. LOM: unetched, brightfield illumination, total images— 9 fields horizontal × 15 fields vertical

Figure 15. In comparison with Figure 14, this sample was made using the same material; however, it has undergone a different heat treatment. Red dots again represent locations of internal cracks and show a different population and spatial distribution throughout the cross section, especially in the region near the fracture surface. As in Figure 14, the surface was protected with a nickel plate. Resolution is the same as Figure 14(a), and all aspects of the image can be analyzed by first enlarging the single montage to an appropriate size, then moving from area to area selecting important features. LOM: unetched, brightfield illumination, total images—9 fields horizontal × 15 fields vertical

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Figure 13 was also the source of Figure 14, but at a different resolution. Figure 13 embraced a total of 80 images and extended a large distance from the fractured end, whereas Figure 14 was composed of 135 images and covered a short distance from the fracture surface. A higher optical magnification was used in Figure 14; therefore, the resolution of the image in Figure 14 is higher than in Figure 13. The small inset (Figure 14(a)) shows the actual information contained in the larger montage. The detail of each single image remains in the final montage. Figure 14 was used in a comparison with Figure 15, which was a similar material that had

undergone a different heat treatment. The small red dots in both images delineate the locations of short internal cracks. A comparison of the neighborhoods containing cracks (red dots), the total populations, and the spatial distributions can be made using these images. These comparisons would be difficult to make without the use of the large montages. As previously noted, montage creation is not limited to images from a light microscope. Those generated from SEM images can be useful in describing surfaces and feature distributions at different scales. Figure 16 shows a section of a tensile-fracture surface from a PF material. The

Figure 16. Multi-image montage (top) of tensile-fracture surface from PF specimen. The more detailed image (a) is seen at approximately the magnification used for original acquisition. Details of the original images are maintained. SEM: secondary electron image, total images—6 fields horizontal × 3 fields vertical

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same situation exists as was seen with the LOM images in Figures 13 through 15. Details of the original images are maintained, as can be seen in the higher magnification image (Figure 16(a)) from the cropped area within the white box. CONCLUSIONS • Metallography is the most effective and efficient analytical tool for use in understanding material behavior. • It is the only set of techniques capable of providing information in the three areas responsible for PM properties, namely, density/ porosity, chemical composition/alloying method, and microstructure. • The usefulness of the analyses is enhanced by a greater understanding of the tests. • Improvements in digital-image capture and computer storage of images have led to more opportunities for the metallographer. • Grain-size testing can be simplified and improved through the use of digital-imaging systems. • Large multi-image montages provide access to examination of photomicrographs at multiple scales, similar to progressing from low to high magnification, without the need of the sample or a microscope. • Montages allow for evaluation of non-uniform and nonrandom distributions, in addition to spatial and neighborhood relationships. • Information contained in each individual image can be retained in the montage. • A drawback of the montages could be the size of the file. Images 100 Mbytes and larger in size are not uncommon. REFERENCES 1. Standard Test Method for Determining Average Grain Size; ASTM E 112 – 96 (2004), Annual Book of ASTM Standards, 2009, Volume 03.01, ASTM International, West Conshohocken, PA. 2. E.E. Underwood, Quantitative Stereology, 1970, AddisonWesley Publishing Company, Inc., Reading, MA, pp. 6–9. 3. F. Schuckher, “Grain Size”, Quantitative Microscopy, edited by R.T. DeHoff and F.N. Rhines, 1968, McGraw-Hill Book Company, New York, NY, pp. 201–265. 4. G.F. VanderVoort, Metallography Principles and Practice, 1984, McGraw-Hill Book Company, New York, NY, pp. 219–223 and 638–640. 5. S.J. Lawrence, “Delineating Prior Austenite Grain Boundaries in Steels”, Microscopy and Microanalysis, 2004, vol. 10 (Suppl 2), pp. 750–751.

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6. P. Louis & A.M. Gokhale, “Application of Image Analysis for Characterization of Spatial Arrangements of Features in Microstructure”, Metallurgical and Materials Transactions A, 1995, vol. 26A, pp. 1,449–1,456. ijpm

The International Journal of Powder Metallurgy also recognizes the “Awards of Merit” from PowderMet2009: Microstructure Evolution of Gas-Atomized Iron-Based ODS Alloys J.R. Rieken, I.E. Anderson, M.J. Kramer, J.W. Anderegg and D. Shechtman

Sinter-Hardening Response of Leaner Alloy Systems B. Lindsley Metallographer: G.J. Golin These papers were presented at PowderMet2009 and published in Advances in Powder Metallurgy & Particulate Materials—2009, Proceedings of the PowderMet2009 International Conference on Powder Metallurgy & Particulate Materials, which are available from the Publications Department of MPIF (www.mpif.org).

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NOW AVAILABLE!

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CONSULTANTS’ CORNER

M.I. “Mike” JAFFE*

AND SAMUEL M. “Sandy” JAFFE

Q

What kind of cutting tool and attendant parameters should be used to finish machine near-net-shape PM parts to tight tolerances? This is a good question with, unfortunately, not a simple answer. The main issue from our experience is the vast differences in the metal powder that you are trying to machine. Unlike wrought materials with consistent properties such as hardness, wear resistance, and microstructure, PM parts tend to exhibit vastly different properties. This is due to density variations, sintering conditions, and the actual state that the component may be in at the time of machining. That is not to say that a component will be different lot to lot, but the magnitude of the variables in the PM process make it difficult for cutting-tool manufacturers to develop a grade of material that will work with PM as well as it does in the wrought condition. The main rule of thumb that we use when selecting a cutting-material grade is that the paradigms that apply to cutting wrought materials rarely apply when machining PM materials. Cutting grades and tip configurations that are not supposed to work with a given material composition often work satisfactorily, and conversely the latest material with a high-tech shape often may fail almost immediately. Cubic boron nitride (CBN), a material not normally used for softer materials or when interrupted cutting conditions exist, has worked well for conventional as well as sintered materials, and for heat treated PM materials. Elaborate single-point cutting inserts with special chip-breaking configurations are often not needed with PM components, as the material rarely produces long stringy chips. Some of these shapes can reduce the cutting forces, which can be helpful in reducing distortion of the component during

A

cutting, allowing tighter tolerances to be held. For drilling, milling, and single-point cutting, a carbide with a wear coating or CBN will normally work well, as long as the tool is fed into the PM component smoothly. By using these hard materials it is often possible to machine PM components without the use of a liquid coolant, especially if a blast of compressed air is used to clear away the chips formed during machining. In the absence of water-based coolants corrosion issues may be eliminated or reduced. For tight-tolerance finish machining of holes, the use of honing sticks consisting of a diamondor borazon-coated split sleeve with a tapered ID, slid over a tapered mandrel, has worked satisfactorily in holding micron-level tolerances with exceptional tool life. Surface finish is determined by the grit size on the final honing-stick sleeve. Frequently, uncoated conventional steel cutting tools will not exhibit a long cutting life when machining sintered or heat-treated PM materials. They may, however, work well in the “green” machining of unsintered or presintered components. Since this question is concerned with finish machining, this may not apply, as it is often a rough machining operation. When machining PM components, cutting speeds and feeds also defy many published guidelines in the form of figures and charts. Typically, cutting speeds may be significantly faster than those used for cutting wrought materials of the same general chemistry. However, the cutting feed may need to be reduced due to the difficulty in holding the component in the chucking device or to distortion of the PM component.

*M. I. (Mike) Jaffe, Box 240, 144 Brewer Hill, Mill River, Massachusetts 01244-0240, USA; Phone: 413-229-3134; E-mail:[email protected]

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Q A

Can you provide guidelines in the use of shoulder dies (aka flange or step dies) to avoid failure? In general, shoulder dies in small presses (≤18.1 mt (20 st)) seldom fail by cracking or by shearing at the flange if there is no catastrophic event such as a double compact, a machine-function failure, or a serious foreign body in the powder. The expense of the carbide insert and the steel case or ring is not burdensome so the die assembly can be rugged (Figure 1). Compared with the part shape, the die insert and the ring are relatively massive. The shrinkage compressive stresses on the insert are reasonably uniform and the compressive forces from compaction will be sustainable. When unexplained die failures in larger presses became a serious problem, we commissioned a program to analyze the horizontal forces in a die and its ring caused by thermal shrinkage and the internal forces due to compaction. Assuming a round die, the program used basic hoop-stress analysis. We are sure that today a more thorough analysis could be performed utilizing finite element analysis. Our program identified the compressive forces set up in the die, the tensile forces in the ring, possible tension in the die caused by the compressive forces in compaction, insufficient shrink forces (insufficient shrink interface or too small a ring, or both), incremental movement of the die caused by shrinkage and by the compressive forces, and the difference in motion at the interface in a flanged die (Figure 1). The system recognized either steel or carbide dies. In general, we found the following: a. A thin-wall die case (shrink ring) shown in Figure 2 can allow the die to go into tension during compaction, yet the die is in compression when idle. It will be a cyclical stress, which may result in radial cracking of the insert. It assumed that the horizontal force

Figure 2. Thin-wall die case

Figure 3. Distortion caused by shrinkage of heavy support section

was about 1/3 of the vertical compacting force. b. A flanged die can develop high shear stresses at the interface of the flange, Figure 1 c. A ring with a heavy bottom step to support a flanged die can result in severe distortion of the ring when the die is “shrunk in.” This can reduce the shrinkage force near the top and increase the shrinkage force near the bottom, resulting in the distortion shown in Figure 3. This can go unnoticed since the die is usually ground at the top and bottom after assembly. This was proven by removing the heavy step in the ring on a damaged die. It was observed that the top surface was concave, and the ID of the die decreased somewhat at the top. Although this analysis was based on several assumptions and had limitations, such as not considering the vertical interfaces directly, it was useful in solving many practical problems in die failure and troubleshooting.

Q A Figure 1. Shoulder die configuration

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What is the effect of a pore, oil, or resin under a microindentation hardness site? Sintered PM parts consist of metallic phases, pores, and on occasion oil and/or resin. As the pore or oil do not resist deformation and the resin is much softer than the metal, the macro hardness will be a composite reading of the metal Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

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CONSULTANTS’ CORNER

and the nonmetallic constituents. Since a Rockwell or Brinell test makes an indent that is large compared with a microindentation, it will cover both metal and nonmetallic constituents, and will therefore give a lower reading than that of the equivalent pore-free (solid) metal. For example, in a medium-density hardened PM part, a Rockwell C (HRC) reading of 39 corresponds to an indent circle ~0.4 mm dia. (0.016 in.). In comparison, the Diamond Pyramid Hardness (DPH) or Vickers (DPV) microindentation hardness under a load of 100 g would result in a square with diagonals of magnitude 0.014 mm (0.00055 in.). Figure 4(a) shows a Rockwell C indent at 150 kg load. Knoop and DPV indents at 100 g load in a sintered nickel steel part of apparent hardness ~HRC39 are shown in Figure 4(b). The Knoop indent converts to HRC 61 and the DPV indent to HRC 68. It is clear that the HRC indent covers many particles as well as nonmetallic areas while the microindentation can be located on one particle. In microindentation hardness tests (Knoop and DPV) conducted at different laboratories, the repeatability (r) and reproducilibity (R) were significantly higher for hardened PM parts than for the sintered counterparts (MPIF Standard 51); for example, 5 points in R for the HRC-equivalent DPV and ~20 points in R for the HRC-equivalent Knoop. The light nickel-rich areas should be avoided. With solid materials a Rockwell hardness reading can be taken on most flat, reasonably smooth surfaces. For a DPH or DPV reading, the measurement depends upon an optical reading, hence a smooth, polished surface is needed. For microin-

dentation readings on a solid material, a well-polished and lightly etched surface that identifies the different phases is needed. However, for a PM part a well-polished and prepared surface is essential. It is necessary to remove any smearing of the metal that may cover the pores as a reading here would be misleading. PM parts must be well ground (sanded), polished, etched to remove smears, re-polished, and lightly etched to show the microstructure. This sequence may have to be repeated more than once. With a correctly prepared sample the small indenter can be placed on a metal particle, avoiding the smear. The indent appearance should be clean and sharp, be it DPH or Knoop. We prefer to use DPH since the measured indent length is shorter and can more easily fit in a small area. In Figure 4(a), the diagonal for DPH on a particle of HRC 60 with a 100 g load is ~0.016 mm, whereas the long length for Knoop under the same conditions would be ~0.044 mm. If it is distorted, and in the case of DPH not perfectly square, it should not be considered as it may have a pore under it or some other anomaly. In most cases, a series of readings should be taken, not too close together (conforming to ASTM B933, B384 and MPIF STD. 51). The lowest reading (largest indent) should be excluded and the other readings averaged to obtain the final result. 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]

Figure 4. (a) Rockwell C scale indent, (b) Knoop and DPV indents. Sintered nickel steel. Optical micrographs courtesy of Powder-Tech Associates, Inc.

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ENGINEERING & TECHNOLOGY

PARTICLE-REINFORCED CARBON-FREE PRECIPITATIONHARDENING HIGH-SPEED STEELS

Herbert Danninger*, Christian Harold**, Fardin Rouzbahani**, Helmut Ponemayr***, Manfred Daxelmüller****, Frantisek Simančík*****, and Karol Iždinský ******

INTRODUCTION For complex-shaped machining tools that combine hardness and toughness, HSS are commonly used.1 Compared with hardmetals or cermets, HSS offer the advantage of availability in a relatively soft (annealed) fabricated condition. This enhances machinability by standard techniques, and the hard condition is then attained by heat treatment, usually consisting of austenitizing, quenching, and triple tempering.2 The high hot hardness of these materials that enables high cutting speeds is not attained by the martensitic transformation but by strengthening through fine secondary carbides precipitated during tempering. Thus the material retains high in-service hardness, at least up to the tempering temperature. Solid-solution strengthening by cobalt is frequently utilized and micron-size primary carbides (M6C and VC) offer high abrasion resistance.3 In common with other precipitation-strengthened alloys, HSS are sensitive to thermal overloading, which results in overaging (coarsening of secondary carbides), with an irreversible loss of hardness. This occurs within a narrow temperature interval above t he optimum tempering range. HSS are commonly intentionally overaged to some degree, to obtain sufficient toughness, but service temperatures >600°C are usually sufficient to result in an appreciable loss of hardness. There have been numerous approaches to improve the cutting performance of HSS by the introduction of hard phases,4,5 PM being particularly well suited for this purpose. This approach improves abrasion resistance but not hot hardness since the matrix softens and the hard phases tend to flow with the matrix and are unable to support the cutting edge.

Carbon-free iron–cobalt– molybdenum and iron– cobalt–tungsten–molybdenum high-speed steels (HSS) offer high hot hardness and temper resistance. While this results in excellent cutting performance when machining stainless steels or titanium alloys, resistance to abrasive wear is limited. In this work, ceramic-particle reinforcement was studied as a possible approach to counter this deficiency utilizing pressand-sinter powder metallurgy (PM) processing. It was shown that carbidic phases are unstable during sintering and decompose. In contrast, alumina is thermodynamically stable during sintering up to 1,400°C, and a homogeneous distribution can be achieved if fine metallic and ceramic powders are used. Fine alumina particles also lower the tendency for matrix grain coarsening during solution treatment, so that the material is less sensitive to overheating. Transverse rupture strength (TRS) is lowered slightly by particle reinforcement but hardness and cutting performance are significantly improved. Presented at the PM2008 World Congress and published in Advances in Powder Metallurgy & Particulate Materials—2008, Proceedings of the 2008 World Congress on Powder Metallurgy & Particulate Materials, which are available from the Publications Department of MPIF (www.mpif.org).

*Professor, **PhD student, Vienna University of Technology, Getreidemarkt 9/7, A-1060 Vienna, Austria, and Materials Center Leoben GmbH, A8700 Leoben, Austria; E-mail: [email protected], ***Managing Director, Böhler Ybbstal Profil GmbH, A-3333 Böhlerwerk, Austria, ****Quality Manager, Böhler Uddeholm Precision Strip GmbH & Co.KG, A-3333 Böhlerwerk, Austria, *****Director, ******* Deputy Director, Institute of Materials and Machine Mechanics, Slovak Academy of Sciences, SK-831 05 Bratislava, Slovakia

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PARTICLE-REINFORCED CARBON-FREE PRECIPITATION-HARDENING HIGH-SPEED STEELS

As early as 1930 it was recognized that carbonfree tool steels with high hot hardness can be obtained by using intermetallic phases for strengthening. Köster and Tonn6–8 studied the iron–cobalt–molybdenum and iron–cobalt–tungsten systems and obtained hardness levels comparable with those of carbidic HSS and with low thermal softening at higher temperatures; thus these materials can be used at higher cutting speeds. In the 1960s, Geller 9–11 studied many carbon-free tool steel grades and found that they were well suited for the machining of titanium alloys. Karpov et al.12 evaluated the PM approach, manufacturing iron–cobalt–tungsten–molybdenum alloys from hydrometallurgically produced powders, obtained by coprecipitation and coreduction followed by pressing and sintering. Hardness levels up to 70 HRC have been reported. Danninger et al.13–15 produced iron–cobalt–molybdenum and iron–cobalt–tungsten–molybdenum steel grades from mixed elemental powders via the press-and-sinter route, and obtained excellent combinations of hardness and TRS after sintering in the range 1,350°C–1,400°C. It was also shown that these alloys, in particular the iron–cobalt– molybdenum grades, were relatively soft in the asquenched condition, enabling machining and cold working, followed by isothermal aging at 550°C–650°C (i.e., without a phase transformation14), with excellent hardness and geometrical precision. This made the materials particularly well suited for precision tools. Machining tests showed that these steels are particularly effective in machining operations where thermal softening is the main reason for tool failure, for example, when cutting materials such as austenitic stainless steels or titanium alloys. In the cutting of abrasive workpiece materials, for example, cold-worked DZ type tool steels of the D2 type with coarse chromium carbides,2 the cutting performance is no better than that of carbidic HSS.16 This was attributed to the lack of micron-size hard phases in the tool steels, since the micron-size µ phases of the type (Fe,Co)7Mo6 and (Fe,Co)7(W,Mo)6 were markedly softer than the M6C and MC carbides present in standard HSS. In this work, an attempt was made to remedy this deficiency by introducing ceramic particles into carbon-free tool steels. The rationale for this approach was that, in contrast to previous

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approaches,4 a deficiency in wear resistance, not in the hot hardness of the matrix, should be remedied by particle reinforcement. EXPERIMENTAL PROCEDURE Starting powders were carbonyl iron (BASF CN), cobalt with d50 = 3.5 µm (Pechiney), molybdenum <32 µm (Plansee) and tungsten with d50 = 3.05 µm (WOLFRAM GmbH). WC, VC, NbC, SiC, and TiC (all <10 µm) were used as the reinforcing phases as well as fused Al 2 O 3 (Treibacher Schleifmittel, grades Alodur F320, F800, and F1000) where 320, 800, and 1000 refer to mesh sizes, which correspond to particle sizes of 45 µm, 6.4 µm, and 4.4 µm, respectively. The elemental and compound powders were dry mixed for 60 min in a tumbling blender and compacted at 400 MPa in tools with a floating die to bars ~100 mm × 12 mm × 14 mm. Die-wall lubrication was applied to avoid contamination of the compacts with carbon. Sintering was carried out in a pusher furnace with molybdenum heating elements (Degussa “Baby”) in flowing technical purity hydrogen for 2 h at 1,370°C, These conditions had previously been found to optimize the homogeneity of the alloying elements.14 After sintering, the bars were hot rolled in a laboratory mill to eliminate any residual porosity. Dilatometric runs were carried out in vacuum using a pushrod dilatometer (Baehr model 801) with an alumina measuring system. Heat treatment was performed by solution treatment in a high-purity nitrogen atmosphere, oil quenching, and aging for 60 min in nitrogen at different temperatures. The bars were ground on all faces, and transverse rupture (TR) tests were carried out in 3-point bending. Rockwell hardness HRC was measured on metallographically prepared cross sections and, for comparison, on the faces. The two sets of results were in excellent agreement. Sections were prepared using standard metallographic techniques, care being taken during cutting to avoid thermal loading of the specimens. Polishing and etching was repeated several times to eliminate any deformed layers. Cutting tests were performed in turning utilizing tools with the appropriate geometry ground from the heat-treated bars. Turning was evaluated against different workpiece materials and tool life was assessed in all cases.

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RESULTS AND INTERPRETATION Addition of Carbide Particles Carbides are the common reinforcing phase for steels since they are well bonded to the matrix in contrast to oxides, as shown by gigacycle fatigue tests.17 The latter showed that carbides tend to crack while oxides decohere from the matrix. Carbides are significantly more reactive than oxides and, in a carbon-free matrix, undesirable reactions could not be excluded. In consequence, specimens were based on a matrix of compositions Fe-25 w/o Co-14 w/o W7.5 w/o Mo, with 10 w/o WC. Sintering was done at temperatures between 1,300°C and 1,370°C. It was observed that the sintered density increased up to ~1,350°C, with a maximum density ~8.5 g/cm3. This compares with a pore-free density of 9.4 g/cm3, or ~10 v/o porosity. Above this temperature density decreased significantly, to a total porosity level ~30 v/o. This effect was particularly visible when sintering in the dilatometer. Figure 1 depicts dilatometric graphs from several runs in which green compacts of the cited composition were heated to different temperatures between 1,300°C and 1,360°C. The expansion encountered in the run-up to the highest temperature is clearly visible. This phenomenon is frequently related to the formation of liquid phases.18,19 Metallographic sections revealed that the WC

Figure 1. Dilatometric graphs for Fe-25 w/o Co-14 w/o W-7.5 w/o Mo containing 10 w/o admixed WC. Heating rate 10 K.min-1 to varying temperatures, 1 h isothermal, in vacuum

Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

particles had agglomerated significantly, which effect is detrimental to the mechanical properties. Sintering at temperatures <1,350°C left some pores (Figures 2(a) and 2(b)) and showed that carbon transfer had occurred from the WC into the matrix. The original WC, having been transformed into eta-carbides, and eta-carbides having been precipitated from the matrix, results in decarburization of WC and carburization of the matrix. At temperatures >1,350°C (Figures 2(c) and 2(d)), this carbon transfer is observed and large pore clusters are evident. The latter have been found in other sintered steels if sintering with a transient liquid phase is followed by a persistent liquid phase.19 This shows that the stability of WC is not sufficiently high to avoid decomposition during sintering. MC-type carbides such as VC or NbC are thermodynamically more stable than WC, and also exhibit lower solubility in the matrix. Thus similar experiments were carried out using 3 w/o

(a) 3.2 h at 1,340°C

(b) 3.2 h at 1,340°C

(c) 1 h at 1,360°C

(d) 1 h at 1,360°C

Figure 2. Microstructures of Fe-25 w/o Co-14 w/o W-7.5 w/o Mo-10 w/o WC alloy sintered at different temperatures. Optical micrographs, Murakami’s etch

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admixed VC and equivalent amounts (in mole %) of other carbides in the matrix. In this case the tungsten-free content iron–cobalt–molybdenum variant was selected. The results were, in part, less than satisfactory, since large amounts of liquid phase were present >1,330°C. In Figure 3, metallographic sections are shown depicting the iron–cobalt–molybdenum alloy reinforced with different admixed carbides. In all cases, the formation of a liquid phase is evident, as is the presence of eta-carbides, which were confirmed by etching with Murakami’s reagent. The results confirm that carbides are not suitable for particle reinforcement in carbon-free tool steels, being too reactive during the sintering process. They might be used if consolidated by hot isostatic pressing (HIP) since lower temperatures can be utilized,20 but some carbon transfer can be expected to take place. Furthermore, when using gas-atomized prealloyed powders the problem of an inhomogeneous distribution of the

(a) 9.3 w/o WC

(b) 3 w/o VC

(c) 5 w/o NbC

(d) 2.9 w/o TiC

(e) 1.9 w/o SiC

(f) 1.9 w/o SiC

Figure 3. Fe-25 w/o Co-7.5 w/o Mo-X w/o hard phase alloy compacted at 400 MPa, and sintered 30 min at 1,340°C in hydrogen. Optical micrographs, Murakami’s etch

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reinforcement arises. Combining micron-size reinforcing phases with 60–100 µm matrix particles resulted in weakening of the particle boundaries, as demonstrated in fiber reinforced HSS.21 Fine-Alumina-Particle Reinforcement Of the other hard phases, nitrides can be regarded as reactive as carbides. Borides are hampered by the low temperature of the Fe-B eutectic, giving rise to decomposition during sintering. Thus oxides remain as the compounds that are sufficiently stable. Alumina was chosen since it is stable, hard, and available in the form of a high-quality product at reasonable cost.22,23 Fused high-purity alumina was selected since it consists exclusively of α-Al2O3. The weak bonding of alumina to steel matrices means that this reinforcing phase must be considered as an inclusion (i.e., a defect), in relation to mechanical behavior. A significant loss of strength can be avoided if the ceramic particles are sufficiently fine and evenly distributed. Therefore, 1,000 mesh Al 2O 3 was used as the standard reinforcing material, coarser grades being used for reference purposes only. In the first test run, the iron–cobalt–molybdenum matrix was reinforced with 2.5 w/o and 5.0 w/o Al2O3, respectively (~5 v/o and 10 v/o, grade F1000). It was confirmed that a sintering temperature >1,350°C resulted in essentially pore-free density. Thus densification is not retarded by the presence of the oxides compared with the monolithic variant.14 The lower density of the particlereinforced alloy (Figure 4) is a consequence of the lower density of Al2O3 which, at 2.5 w/o, lowers

Figure 4. Density of Fe-25 w/o Co-7.5 w/o Mo-2.5 w/o Al2O3 alloy as a function of sintering temperature. Isothermal sintering time 2 h in hydrogen

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the pore-free density by ~0.2 g/cm3. The as-sintered microstructure of the reinforced alloy is shown in Figure 5 compared with the non-reinforced matrix. It is evident that the reinforced alloy exhibits a markedly finer grain structure; apparently grain growth during sintering is significantly inhibited by the presence of the ceramic inclusions. In the as-sintered condition, this is not a problem but, as will be discussed, the grain-refining effect has a pronounced influence on age hardening. In order to assess the hot deformability of the ceramic-reinforced material, hot compression tests were carried out using a GLEEBLE. Cylindrical specimens machined from the sintered bars were tested at 950°C and 1,150°C, respectively, at deformation rates of 1 s-1 and 10 s-1. The flow stress curves showed that the presence of the ceramic reinforcement, at least in concentrations up to 5 w/o, did not markedly affect deformability. Thus hot working should be possible for the composites as well as for the base alloy. In fact, hot rolling of the composite alloy proved to be possible, although the preheating temperatures necessary to avoid crack formation were slightly higher in the reinforced alloy than with the matrix alloy. The optimum temperature window was found to be smaller at the higher Al2O3 content. In consequence, subsequent investigations focused on the alloy containing 2.5 w/o (5 v/o) of the ceramic-reinforcing phase. Both the sintered and the hot-rolled specimens were heat treated: solution treat 30 min at 1,150°C, oil quench, and age 60 min at 600°C.

Figure 5. As-sintered microstructures (a) Fe-25 w/o Co-7.5 w/o Mo monolithic alloy, and (b) alumina-reinforced alloy. Optical micrographs, FeCl3 etch

Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

TABLE I. PROPERTIES OF Fe-25 w/o Co-7.5 w/o Mo-2.5 w/o Al2O3 (F1000) COMPACTED AT 400 MPa & SINTERED 2 h AT 1,370°C Reinforcement None None 2.5 w/o Al2O3 2.5 w/o Al2O3

Deformation Mode/ Temperature Density (°C) (g/cm3) Rolling/1,150 Rolling/1,150

8.06 ± 0.02 8.16 ± 0.03 7.98 ± 0.03 7.96 ± 0.04

TRS (MPa)

Apparent Hardness (HRC)

1,804 ± 130 2,318 ± 158 1,549 ± 186 2,070 ± 146

63.1 ± 0.6 66.7 ± 0.2 66.6 ± 0.3 65.4 ± 0.5

The resulting TRS and apparent hardness in the peak aged condition are listed in Table I. It is clear that hot rolling did not significantly affect the sintered density, confirming that the level of porosity after sintering is low. However, the TRS is increased significantly which indicates that a low level of residual porosity has a detrimental effect on strength. Addition of the ceramic particles does not affect the TRS significantly. Apparently the alumina particles are sufficiently small that they do not act as crack-initiating defects. As noted previously, the precipitation-hardened tool steels are relatively soft in the as-quenched condition, and the high level of hardness needed is attained only after an isothermal aging treatment. For the monolithic alloys, aging temperatures in the range 550°C to 650°C have been shown to be optimal. 14,24 For the reinforced grades, some effect of the ceramic particles on hardening behavior is unlikely, but cannot be completely excluded. Therefore, specimens of the monolithic and the reinforced HSS grades were prepared by the standard press-and-sinter PM route, followed by heat treatment (solution treated 30 min at 1,200°C, oil quenched, and tempered for 60 min at different temperatures). The resulting apparent hardness data are given in Figure 6 from which it can be seen that the hardening response of the reinforced alloy is similar to that of the monolithic alloy: secondary hardening is not affected by the ceramic particles, but the hardness level is higher by about 2–4 HRC. Only in the overaged condition does this difference disappear. Thus the high thermal stability of the (Fe,Co) 7Mo 6 secondary precipitates is evident. Even after aging for 60 min at 700°C hardness levels >60 HRC were obtained in the aluminareinforced alloy. It has been found that the properties of the

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TRS are cited in Table II. The properties of the alloys prepared from F1000 alumina and F800 alumina do not differ significantly; this is expected in light of the marginal difference in the mean particle size of the alumina. The coarse grade of alumina results in a significantly lower TRS. Again, this is expected since large ceramic inclusions act

Figure 6. Apparent hardness vs. aging temperature for unreinforced and reinforced iron–cobalt–molybdenum HSS

carbon-free tool steels, in particular the iron– cobalt–molybdenum grades, are significantly affected by the solution-treatment temperature.14 This process should be performed at a temperature at which most of the micron-size µ phases are dissolved in order to produce a solid solution of molybdenum to generate a pronounced secondary hardening effect during aging. However, overheating, at a temperature at which all the coarser µ phases are dissolved (typically >1,200°C for iron–cobalt–molybdenum), is undesirable, resulting in excessive grain growth and a loss of TRS. Apparently the coarse µ phase stabilizes the grain structure in a way similar to that of the primary carbides in standard carbidic tool steels. In the present study, it has been observed that the as-sintered microstructure of iron–cobalt– molybdenum steels is refined significantly if fine alumina particles are present. A similar effect can be expected for solution treatment, since this is carried out at temperatures below the sintering temperatures. Therefore, heat treatment was carried out utilizing solution treatment temperatures up to 1,300°C. For the unreinforced alloys the expected grain growth was observed (Figures 7(a) and 7(b)) while the variants containing alumina showed a much finer microstructure (Figures 7(c) and 7(d)), thus confirming that the reinforcing phase acts as a grain-growth inhibitor during the aging treatment. In conclusion, the particle-reinforced iron–cobalt–molybdenum alloy variants are less sensitive to the solution-treatment temperature, overheating being much less of a problem than for the respective monolithic alloys. For purposes of comparison, other grades of alumina were evaluated. Apparent hardness and

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(a) No reinforcement

(b) No reinforcement

(c) Reinforced

(d) Reinforced

Figure 7. Microstructures of Fe-25 w/o Co-7.5 w/o Mo-2.5 w/o Al2O3 alloy sintered for 2 h at 1,370°C, solution treated for 1 h at 1,300°C, oil quenched, and aged 600°C for 1 h. Optical micrographs, Nital etch

TABLE II. EFFECT OF SOLUTION TREATMENT ON MECHANICAL PROPERTIES OF Fe-25 w/o Co-7.5 w/o Mo-2.5 w/o Al2O3 COMPACTED AT 400 MPa, SINTERED 2 h AT 1,370°C IN HYDROGEN & HOT ROLLED 30 min at 1,100°C + Oil 5 min at 1,200°C + Oil Quench and Aged 1 h 600°C Quench and Aged 1 h 600°C Al2O3 Size (μm)/Grade

Apparent Hardness (HRC)

TRS (MPa)

4.4/F1000 6.4/F800 45/F320, ZESK

64.2 63.7 64.1

1,853 1,986 1,425

Apparent Hardness TRS (HRC) (MPa) 64.9 65.2 65.9

1,927 2,004 1,443

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as defects under load. The difference in alumina particle size is clearly discernible in Figure 8. Also inhibition of grain growth is less effective, but this should not be a major factor if suitable solutiontreatment conditions are identified. Cutting Tests Machining performance of the alloys was evaluated in turning, utilizing a short-time turning test.25 Cylindrical workpiece specimens (40 mm dia.) were machined at varying cutting speeds at a length of 10 mm, with the cutting speed increasing up to tool failure. Austenitic stainless steel (AISI 304L/1.4301) was used as the workpiece material. The cutting conditions were: feed 0.103 mm/rev., depth 0.5 mm, dry cut. Standard carbidic high-speed steel M42 was used as a reference material. In Figure 9, cutting performance, in terms of the maximum cutting speed, is plotted for several monolithic and alumina-reinforced carbon-free tool-steel grades, and for M42.

(a) F800 Al2O3

(b) F320 Al2O3

Figure 8. Metallographic sections of Fe-25 w/o Co-7.5 w/o Mo-2.5 w/o Al2O3 compacted at 400 MPa and sintered 2 h at 1,370°C. Optical micrographs, FeCl3 etch

From the results in Figure 9, the enhanced performance of the carbon-free grades compared with M42, for the selected machining operation, is evident. Furthermore, it was shown that adding 2.5 w/o of fine alumina particles further improves the cutting performance of the Fe-25 w/o Co-7.5 w/o Mo 2.5 w/o tool steel. Adding phases that have significantly higher hardness than the µ phases present in the monolithic material appears to be a viable way to combine a temper-resistant matrix with hard “micro-cutting edges.” CONCLUSIONS Carbon-free precipitation-hardened tool steels of the type iron–cobalt–molybdenum and iron–cobalt–molybdenum–tungsten are compatible as machining tools for austenitic stainless steels or titanium alloys, in which thermal softening is the dominating mechanism of tool wear. These grades are, however, less resistant in abrasive wear. In the present study, an attempt has been made to remedy this deficiency by reinforcing the steels with hard phases by admixing fine hard phase powder to elemental powder mixes. Carbides have been shown to be unsuitable, at least, as fabricated by pressing and sintering, since they decompose during sintering and form a liquid phase. Fine alumina particles remain stable during sintering and do not adversely affect the densification process and refine the as-sintered microstructure. This effect is most pronounced during solution treatment: overheating results in pronounced grain coarsening in monolithic alloys, with a resulting loss of toughness, while the alumina reinforced grades are much less sensitive. After optimum heat treatment, the reinforced grades exhibit slightly lower TRS than the monolithic alloy variants but a higher apparent hardness and enhanced cutting performance. ACKNOWLEDGEMENT This project was supported financially by the Austrian Federal Government through the Kplus initiative. REFERENCES

Figure 9. Cutting performance (maximum tolerable cutting speed) of iron–cobalt– molybdenum–tungsten alloys (turning wrought 304L compared with M42)

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1. ASM Handbook—Properties and Selection: Irons, Steels, and High-Performance Alloys, Vol. 1, 10th edition, 1990. ASM International, Materials Park, OH. 2. G. Roberts, G. Krauss and R. Kennedy, Tool Steels, 1998, ASM, Materials Park, OH. 3. H.F. Fischmeister and S. Karagöz, “Cutting Performance

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4.

5.

6.

7.

8.

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and Microstructure of High Speed Steels: Contributions of Matrix Strengthening and Undissolved Carbides”, Met. Trans., 1998, vol. 29A, p. 205–216. J.M. Torralba and E. Gordo, “PM High Speed Steel Matrix Composites”, Powder Metall. Progress, Inst. Materials Research, Košice, Slovak Republic, 2002, vol. 2, no. 1, pp. 1–9. J.D. Bolton, A.J. Gant, F.L. Jagger, W.J.C. Price, M. Youseffi, M.M. Oliveira and H. Carvalhinhos, “The Structure and Properties of Sintered Metal Matrix Composites Based on Mixtures of High Speed Steel and Titanium Carbide”, Advances in Powder Metallurgy and Particulate Materials., compiled by J.M. Capus and R.M. German, Metal Powder Industries Federation, Princeton, NJ, 1992, vol. 8, pp. 97–110. W. Köster and W. Tonn, “Das System Eisen-KobaltWolfram”, Arch. Eisenhüttenwesen, 1932, vol. 5, no. 8, pp. 431–440 (in German). W. Köster and W. Tonn, “Das System Eisen-KobaltMolybdän”, Arch. Eisenhüttenwesen, 1932, vol. 5, no. 12, pp. 627–630 (in German). W. Köster, “Mechanische und magnetische Ausscheidungshärtung der Eisen-Kobalt-Wolfram und Eisen-Kobalt-Molybdän-Legierungen”, Arch. Eisen hüttenwesen, 1932, vol. 6, no. 1, pp. 17–23 (in German).

9. J.A. Geller, Instrumentalniye Staly, 1983, Metallurgia Publ., Moscow (in Russian). 10. W.A. Brostrem and J.A. Geller, “Transformation and Properties of High Speed Cutting Alloys Strengthened by Intermetallic Phases”, Metallowedeniye i termicheskaya obrabotka metallov, 1966, no. 11, pp. 35–39 (in Russian). 11. W.A. Brostrem and J.A. Geller: “High Speed Cutting Alloy Strengthened by Intermetallics”, Metallowedeniye i termicheskaya obrabotka metallov, 1970, no. 1, pp. 35–39 (in Russian). 12. M.I. Karpov, V.I. Wnukov, N.W. Medwed and H. Danninger, “Powder Metallurgy Tool Steels with Intermetallic Hardening”, Proc. Powder Metall. World Congress 1998, edited by V. Arnhold and A. Romero, European Powder Metallurgy Association, Shrewsbury, UK, 1998, vol. 3, pp. 519–524. 13. H. Danninger, F. Rouzbahani, C. Harold, H. Ponemayr, M. Daxelmüller, F. Simančik, and K. Iždinský, “Precipitation Hardened Carbon Free Tool Steels with High Temper Resistance”, Proc. 16th Int. Plansee Seminar, edited by G. Kneringer, P. Rödhammer and H. Wildner, Reutte, Austria, 2005, vol. 1, pp. 558–570. 14. H. Danninger, F. Rouzbahani, C. Harold, H. Ponemayr, M. Daxelmüller, F. Simančik, and K. Iždinský, “Heat Treatment and Properties of Precipitation Hardened

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PARTICLE-REINFORCED CARBON-FREE PRECIPITATION-HARDENING HIGH-SPEED STEELS

15.

16.

17.

18.

19.

20.

21.

22.

23.

24.

25.

Carbon-Free PM Tool Steels”, Powder Metall. Progress, Inst. Materials Research, Košice, Slovak Republic, 2005, vol. 5, no. 2, pp. 92–103. C. Gierl, F. Rouzbahani, C. Harold, H. Danninger, H. Ponemayr, and M. Daxelmüller, “Heat Treatment and Microstructure of Carbon-Free Precipitation Hardened PM Tool Steel Fe-Co-Mo/W”, Proc.15th IFHTSE / International Federation for Heat Treatment and Surface Engineering Congress, Vienna, edited by R. Schneider, Austrian Society for Metallurgy and Materials (ASMET), Leoben, Austria, 2006, pp.106–111. R. Florek, “Wear Behaviour of Carbon Free and Carbon Containing HSS in Machining of Ferrous Materials”, 2003, PhD Thesis, Technical University Wien, Vienna, Austria, (in German). Y. Furuya, S. Matsuoka and T. Abe, “A Novel Inspection Method Employing 20 kHz Fatigue Testing”, Mat. Trans., 2003, vol. 34A, pp. 2,517–2,526. N. Dautzenberg and H.J. Dorweiler, “Dimensional Behaviour of Copper-Carbon Sintered Steels”, Powder Metall. Int., 1985, vol. 17, no. 6, pp. 279–282. H. Danninger, “Sintering of Mo Alloyed P/M Steels Prepared from Elemental Powders—II. Mo Homogenization and Dimensional Behaviour”, Powder Metall. Int., 1992, vol. 24, no. 3, pp. 163–168. C. Harold, “Manufacturing Routes, Processing and Application of Precipitation Hardened Carbon Free Tool Steels”, 2002, PhD Thesis, Technical University Wien, Vienna, Austria, (in German). A. Falahati, B. Kriszt, H.P. Degischer, K. Iždinský, F. Simančik, M. Daxelmüller and H. Ponemayr, “Studie zur Faserverstärkung von Schnellarbeitsstahl”, Proc. “Verbundwerkstoffe und Werkstoffverbunde”, John Wiley & Sons, Inc., Hoboken, NJ, 2001, p. 152 (in German). R.A. Queeney, R.E. Masters, R.J. Bleltz and J.D. Dankoff, “Wear Resistance of Al2O3-Reinforced High Speed Tool Steel”, Modern Dev. in Powder Metall., compiled by P.U. Gummeson and D.A. Gustafson, Metal Powder Industries Federation, Princeton, NJ, 1988, vol. 20, pp. 409–419. R.A. Queeney, “Reinforced High Speed Steels as Metal Matrix Composites”, Advances in Powder Metallurgy and Particulate Materials, compiled by J.M. Capus and R.M. German, Metal Powder Industries Federation, Princeton, NJ, 1992, vol. 8, pp. 89–96. E. Eidenberger, E. Stergar, H. Leitner, C. Scheu, P. Staron and H. Clemens, “Precipitates in a Fe-Co-Mo Alloy Characterized by Complementary Methods”, Berg- und Hüttenmännische Monatshefte, 2008, vol. 153, pp. 247–252. A. Salak, K. Vasilko, M. Selecka and H. Danninger, “New Short Time Face Turning Method for Testing the Machinability of PM Steels”, J. Mater. Processing Technol., 2006, vol. 176, pp. 62–69. ijpm

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A V A I L A B L E

MPIF Standard 35 MATERIALS STANDARDS FOR PM SELF-LUBRICATING BEARINGS, 2010 Edition ISBN #: 978-0-9819496-3-5 ~ 28 pages

The first revision to the standard in over 10 years. The 2010 edition includes: NEW Material Section (data property table) Diffusion-Alloyed Iron-Bronze Bearings—FDCT-1802K Revised footnotes for Bronze Bearings and new footnote for the CTG-1004-K10 material Data-table-column heading revisions Data tables now listed in alphabetical order by material system Revised verbiage throughout the standard including a new section under EXPLANATORY NOTES on Oil Impregnation Efficiency ENGINEERING INFORMATION (Inch–Pound and SI Units)— Verbiage and data table modifications New edition includes a 2-part index displaying alphabetical listings & guides to material systems and designation codes used in MPIF Standard 35 • Part 1: for the MPIF Standard 35, Materials Standards for PM Self-Lubricating Bearings document • Part 2: for the other MPIF Standard 35 publications

METAL POWDER INDUSTRIES FEDERATION 105 College Road East, Princeton, NJ 08540-6692 Phone: (609) 945-0009 Fax: (609) 987-8523 Order Online: www.mpif.org

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RESEARCH & DEVELOPMENT

DIAMOND CUTTING TOOLS WITH A Ni3Al + COPPER MATRIX Kuen-Shyang Hwang* and Tsung-Hsien Yang**

INTRODUCTION To attain satisfactory cutting performance with diamond tools, the matrix material must meet several requirements: strong bonding with the diamond particles to avoid diamond pullout; adequate wear resistance to avoid overexposure of the diamond particles when the matrix is too soft or rounding of the sharp diamond edges when it is too hard;1,2 and high hot strength and toughness to overcome the heat and interrupted vibrations generated during cutting.3,4,5 Cobalt has been shown to satisfy these requirements and is widely used. 4–7 However, cobalt is a costly strategic material. Thus the replacement of cobalt has become important and the focus of R&D. Some intermetallic compounds, such as boron-doped nickel aluminide (Ni3Al), are attractive in high-temperature applications. These materials are unique in that their strength increases as the temperature increases.8,9 When used as a matrix material for cutting tools, they improve retention of the dispersed refractory particles that are heated during cutting. Nickel aluminides have been used as a matrix material in WC, TiC, Al2O3, and diamond composites.10–12 To fabricate these composites, gas-atomized Ni3Al powders and ceramic or diamond particles were sintered at temperatures between 1,150°C and 1,500°C for 15–120 min, with or without pressure. Unfortunately, the high temperatures and long sintering times required to densify the prealloyed Ni3Al powders were harmful to the diamond particles, particularly when regular-grade diamond particles were used, as serious graphitization and cracking occurred. To eliminate the problem of diamond degradation, several attempts have been made to produce diamond-containing materials using reaction-sintered titanium diboride (TiB2) and nickel aluminides (NiAl and Ni3Al) as the matrix.13–15 The reaction-sintering process shortens the sintering time at high temperatures. Efforts have also been made to further reduce diamond degradation during the fabrication of diamond tools by adding titanium hydride, which, when dehydrated, creates a protective reducing atmosphere around the diamond particles, preventing oxidation and graphitization.13 Other approaches include the use of bilayer structures to change the combustion temperature.13 The

Diamond cutting tools with a nickel aluminide (Ni3Al) matrix prepared by reaction pseudo–hot isostatic pressing (HIPing) have shown improved stonecutting performance. This is attributed to the high hot strength of the matrix and the short processing time at a high temperature, which minimizes degradation of the diamond particles. It has been shown that graphitization of the diamond particles occurs due to the exothermic heat released during reaction between the nickel and aluminum powders. To minimize this deleterious effect, the addition of copper, iron, and nickel powders has been evaluated. The results show that the addition of 10 w/o Cu decreases the peak temperature of the reaction with retention of the hardness, density, and strength of the matrix. No separate copper phase was found in the matrix, as demonstrated by X-ray diffraction. With a lower reaction temperature and reduced damage to the diamond particles, the performance of the new Ni3Al+10 w/o Cu matrix was superior to that of Ni3Al in dry cutting.

*Professor, Department of Materials Science and Engineering, National Taiwan University, 1, Roosevelt Road, Sec. 4, Taipei, 106, Taiwan, R.O.C.; E-mail: [email protected], **Senior Engineer, Taiwan Semiconductor Manufacturing Co., Park Avenue, Hsin-Chu Science-Based Industrial Park, Hsin-Chu, Taiwan, R.O.C.

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DIAMOND CUTTING TOOLS WITH A Ni3Al + COPPER MATRIX

reaction sintering of nickel–aluminum–copper– titanium has also been reported to reduce the magnitude of the strength decrease of the diamond particles after the reaction. 14 Although these results confirm a reduction in the degree of diamond graphitization and strength degradation, previous studies have not evaluated the mechanical properties or the cutting performance of the tools thus prepared. Recently, Hwang, Yang and Hu16 prepared a new diamond tool with a Ni3Al matrix produced by pseudo-HIPing mixtures of elemental aluminum, boron, nickel, and diamond particles. This new matrix showed an improved performance of over 110% compared with a conventional cobalt matrix during dry cutting of supreme black granites, in which the unique high-temperature properties of the Ni 3 Al were fully utilized. However, some diamond particles were degraded due to the heat of the exothermic reaction with attendant graphitization, in particular when a lower-grade diamond grit was used. In this study, elemental copper, iron, and nickel powders, which can form intermetallic compounds with aluminum, were added to a mixture of elemental aluminum, boron, and nickel powders that was used to prepare the Ni3Al matrix material. The intent was to reduce the amount of exothermic heat that compromised the integrity of the diamonds. The addition of copper was based on the expectation that the heat generated by the exothermic reaction between aluminum and copper would be lower than that between aluminum and nickel, and that some of the reaction heat might be absorbed due to melting of the copper. The nickel was added since an excess of nickel has been shown to slow down the reaction.17,18

The addition of iron may also change the Ni-to-Al ratio, since the iron powder could also react with the aluminum powder, forming iron aluminide, with a reduction in the amount of heat released. EXPERIMENTAL PROCEDURE The powders selected in this study included elemental aluminum, boron, cobalt, copper, iron, and nickel. The characteristics of these powders are listed in Table I. Two grades of diamond grit, SDB1000 and SDA85+ (De Beers, Shannon, Ireland), were selected. SDB1100 is a superior grade with enhanced high-temperature strength. The particle sizes of these two diamond grits were between 300 µm (50 mesh) and 425 µm (40 mesh). In the first phase of the study, Ni3Al specimens without diamond particles were produced so that the effect of the addition of copper, iron, and nickel on the intrinsic properties could be measured. Based on previous results,17,18 the matrix material consisted of elemental nickel and aluminum powders at an atomic ratio of 76 to 24 along with 0.1 w/o boron and various levels of copper, iron, or nickel powders. The powders were then mixed in a V-cone blender for 1 h, followed by compaction at a pressure of 400 MPa into plates 40 mm × 8 mm × 3 mm at 70% of the porefree density. The green compacts were then placed in a graphite die and surrounded by ~2 mm thick alumina powder on all sides. The alumina powder served as the pressurizing medium and also as a heat insulator so that when the exothermic reaction occurred, the heat generated was not lost through the thermally conductive graphite mold. The compacts and the mold were heated at a rate of 30°C/min to 700°C and held for 5 min under a pressure of 20 MPa. To check whether or not

TABLE I. CHARACTERISTICS OF MATRIX POWDERS Ni

Al

B

Cu

Co

Fe

Ni-123

Al-1182

B-1121

635

DiamondGrade P

CIPS-1641

Inco

Cerac

Cerac

ACuPowder

Viridian

ISP

D10 = 3.8 D50 = 10.2 D90 = 24.8

D10 = 15.9 D50 = 56.1 D90 = 147.8

D10 = 4.1 D50 = 18.2 D90 = 39.0

D10 = 8.0 D50 = 14.7 D90 = 26.2

D10 = 0.2 D50 = 4.7 D90 = 15.7

D10 = 2.23 D50 = 4.2 D90 = 7.5

Pycnometer Density, (g/cm3)

8.89

2.69

3.51

8.71

8.87

7.54

Carbon Content (w/o) Oxygen Content (w/o) Nitrogen Content (w/o)

0.081 0.196 0.002

0.025 0.330 0.014

0.068 0.313 0.008

0.267 0.413 0.009

0.220 0.592 0.010

0.716 0.687 0.738

Designation Supplier Particle Size (Laser Scattering) (μm)

38

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DIAMOND CUTTING TOOLS WITH A Ni3Al + COPPER MATRIX

reactive sintering was complete, high-resolution X-ray diffractometry was performed (Phillips model PW1830 unit) with Cu Kα radiation, a 2θ step of 0.01°, and a slow scanning rate of 2°/min in order to enhance separation of possible overlapping diffraction lines of the different phases. To compare the mechanical properties of the new matrix material with those of cobalt, powder compacts of cobalt were also prepared using the same pseudo-HIPing process. The specimens were heated at a rate of 30°C/min. Sintering was carried out at 820°C for 15 min under the same pressure (20 MPa) as that used for the nickel aluminides. After the screening test for the added elements (copper, iron, and nickel), copper was selected and evaluated for cutting performance because it did not lower the sintered density and strength of the Ni3Al matrix. Sintered cutting-insert compacts that included diamond grit were prepared. Diamond grit, at a concentration of 0.88 Karat/cm3 (5.0 v/o/2.4 w/o), was added to the metal powder mixture, precompacted and pseudoHIPed following the same procedure as described previously. To assess a possible change in strength after sintering, diamond grit was extracted from the matrix by immersion in aqua regia. The grit was then sandwiched between two alumina plates and crushed at a crosshead speed of 0.5 mm/min. Crushing-strength data reported are the averages of 20 grits. One hundred extracted diamond grits were also examined utilizing Raman spectroscopy

to check if any graphitization had occurred during sintering. To evaluate the cutting performance of the diamond tools, two cutting inserts were mounted on the opposite ends of a steel disc with an outside diameter of 180 mm. The test was performed against a supreme black granite block using a speed of 30 m/s, at a feed rate of 0.2 mm per cut. The cutting performance was assessed in terms of the grinding ratio (volume removed from the supreme black granite block divided by volume removed (worn) from the diamond tool). RESULTS Nickel Aluminide Matrix Properties The aim of this study was to replace the cobalt matrix with a Ni 3 Al+(copper, iron, or nickel) matrix in diamond cutting tools. Thus the density, apparent hardness, and bending strength of the pseudo-HIPed cobalt were measured; these properties were 8.54 g/cm3 (96% of the pore-free density), HRB 103, and 1,180 MPa, respectively. In terms of the properties of the new matrix, Figure 1 shows that as the amount of copper or nickel increased to 20 w/o (79.9 w/o Ni3Al+0.1 w/o B + 20 w/o Cu (or 20 w/o Ni)), the density remained above 96% of the pore-free level. In contrast, the addition of iron resulted in a significant density loss. Thus no further tests were conducted utilizing iron powders. Figure 2 shows that the apparent hardness remains above HRB 100, which is equivalent to that of the cobalt matrix, with up to 20 w/o Cu or

Figure 1. Effect of copper, iron, and nickel on the sintered density of the Ni3Al matrix: (a) absolute density, (b) relative density Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

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Figure 2. Effect of copper and nickel on apparent hardness of Ni3Al matrix

20 w/o Ni. The bending strength (Figure 3) also remained above 1,300 MPa when up to 20 w/o Ni was added. The bending strength remained at ~1,300 MPa with a copper addition of 10 w/o or less. With more than 10 w/o Cu, a significant decrease in bending strength was observed. Figure 4 compares the effect of the level of copper and nickel on the peak temperature during sintering. Copper and nickel have a similar effect. Without the copper addition, the peak temperature of the matrix occurred at 1,300°C, and this decreased to 1,230°C when ~20 w/o Cu or ~20 w/o Ni was added. These results suggest that in the case of copper or nickel, additions <20 w/o decrease the peak reaction temperature without degrading the mechanical properties of the Ni3Al matrix. Thus diamond tools with a Ni3Al matrix containing 10 w/o Cu, 20 w/o Cu, and 18 w/o Ni (atomic ratio Ni:Al ~79:19) were evaluated further. Table II includes the sintered density, apparent hardness, and bending strength of the specimens that contain diamond particles. The data show that when 20 w/o Cu and 18 w/o Ni were added to the Ni3Al matrix, the sintered density, apparent

Figure 3. Effect of copper and nickel on bending strength of Ni3Al matrix

Figure 4. Effect of copper and nickel on peak temperature of Ni3Al matrix during reaction pseudo-HIPing

hardness, and bending strength decreased significantly. Only the specimen with a 10 w/o Cu addition showed properties comparable with those of the cobalt matrix. These mechanical properties were, in general, slightly lower than those of the diamond-free Ni3Al matrices shown in Figures 1,

TABLE II. COMPARISON OF SINTERED DENSITY, APPARENT HARDNESS, AND BENDING STRENGTH OF DIAMOND TOOLS: SDB1100 DIAMOND PARTICLES EMBEDDED IN A COBALT MATRIX AND Ni3Al MATRIX WITH COPPER AND NICKEL Material

Ni3Al

Ni3Al+10 w/o Cu

Ni3Al+20 w/o Cu

Ni3Al+18 w/o Ni

Co

Sintered Density (g/cm3)

7.23 (99%)*

7.23 (97%)*

7.05 (94%)*

6.86 (91%)*

8.17 (95%)*

Apparent Hardness (HRB)

102

106

94

93

103

1,164

1,071

502

906

1,041

Bending Strength (MPa)

*“Rule of Mixtures” calculation using theoretical densities of Ni3Al, copper, and diamond of 7.50, 8.96, and 3.51 g/cm3, respectively

40

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DIAMOND CUTTING TOOLS WITH A Ni3Al + COPPER MATRIX

2, and 3. This suggests that inferior bonding exists between the diamond particles and the matrix, and serves as the source of crack initiation sites in mechanical testing. Nonetheless, the hardness and bending strength were still comparable with those of the cobalt matrix. In addition to the density and mechanical property measurements, graphitization of the diamond particles was examined for compacts with 10 w/o Cu and 20 w/o Cu and 18 w/o Ni additions using a Raman spectroscope. Figure 5 shows that the degree of graphitization of the SDB1100 diamond grit decreased as the amount of copper increased, probably due to a decrease in the peak reaction temperature. The crushing strength of the extracted diamond particles also improved as the copper content increased, as shown in Figure 6.

Figure 5. Comparison of degree of graphitization between cobalt and Ni3Al matrices with 18 w/o Ni and 0, 10, and 20 w/o Cu additions

Figure 6. Comparison of crushing strength of diamond particles between cobalt and Ni3Al matrices with 18 w/o Ni and 0, 10, and 20 w/o Cu additions Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

However, the addition of nickel did not result in any improvement in reducing graphitization or increasing crushing strength, despite the fact that the peak temperature decreased. This is probably because nickel is a good catalyst for the graphitization of the diamond particles. With excessive nickel present, degradation of the diamond particles occurred. Combining the properties shown in Figure 1 through Figure 6, particularly the bending strength of the matrix and the degree of graphitization and crushing strength of the diamond particles, 10 w/o Cu results in properties most comparable with the matrix prepared from cobalt powder. Thus this composition was used to prepare diamond tools for subsequent cutting-performance tests. Cutting Performance Figure 7 compares the grinding ratios of diamond cutting inserts containing SDA85+ and SDB1100 diamond grit with a matrix of cobalt, Ni3Al, and Ni3Al+10 w/o Cu when cutting was performed in the presence of a cooling fluid. In the case of SDA85+, the cobalt matrix exhibited a grinding ratio of 2380, better than the grinding ratio of 1780 for the Ni3Al matrix. This was primarily because the SDA85+ diamond particles were degraded during reaction pseudo-HIPing. When 10 w/o Cu was added, the grinding ratio improved to 2119. When the superior-grade diamond grit (SDB1100) was used, the grinding ratios were comparable with those of the cobalt matrix and the Ni3Al matrix, at ~4000. When 10 w/o Cu was added to the Ni3Al matrix, the grind-

Figure 7. Grinding ratio of diamond cutting inserts with Ni3Al +10 w/o Cu matrix is slightly better than that of cobalt and Ni3Al matrices in wet cutting

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Figure 8. Grinding ratio of diamond cutting inserts with Ni3Al+10 w/o Cu matrix is ~8% better than that of Ni3Al matrix and superior to that of cobalt matrix in dry cutting

ing ratio improved to 4255, an increase ~6%. The wet-cutting performance showed only a small advantage in using a Ni 3 Al+10 w/o Cu matrix because the operating temperature was lower due to the presence of cooling fluids. However, when the cooling fluid was shut off, the temperature of the cutting insert increased and the benefits of using Ni 3Al+10 w/o Cu as the matrix became obvious because of its superior hot strength. As shown in Figure 8, when SDB 1100 diamond grit was used, the grinding ratio of the Ni3Al matrix was 1850, more than twice that of the cobalt matrix, namely, 850. When 10 w/o Cu was added to the Ni3Al matrix, the grinding ratio further improved to 2000, which is equivalent to an 8% increase. In the case of SDA85+ diamond grit, the grinding ratio of the Ni 3Al matrix was 714, slightly lower than the 780 grinding ratio of the cobalt matrix due to the deleterious effect resulting from the exothermic reaction. With a 10 w/o Cu addition, the grinding ratio of the Ni 3 Al matrix improved to 770, which is still lower than that of the cobalt matrix. DISCUSSION The Ni3Al+10 w/o Cu matrix exhibited the best results among the matrices examined. Figure 9 shows the microstructure of the matrix. The energy-dispersive analysis confirmed that all the grains had a similar composition. The measured contents of nickel, aluminum, and copper were 77.1–79.2 w/o (68.5–70.0 a/o), 11.3–11.7 w/o (22.0–22.5 a/o), and 9.1–11.6 w/o (7.4–9.5 a/o), respectively. In addition, the microindentation-

42

hardness measurements of the matrix were all above HV200, which is much higher than the typical hardness of sintered copper of about HV40. This suggested that there was no separate copper phase. The X-ray diffraction pattern (Figure 10) also confirms that no separate copper, CuAl2, or other copper aluminides were present. Since there were no copper aluminides and the atomic ratio of (Ni,Cu)/Al in the matrix was ~78/22, which is close to the 3/1 ratio of Ni3Al, it is concluded that the copper either dissolved into the matrix or replaced some of the nickel atoms in Ni3Al. A closer examination showed that the peaks were broadened but not shifted. A new small NiAl peak was also found. This also suggested that copper atoms substituted for nickel atoms and caused non-uniform lattice distortion. In contrast, the addition of iron and nickel did

Figure 9. Representative microstructure of Ni3Al+10 w/o Cu showing uniform grains and the absence of copper phases. SEM

Figure 10. X-ray diffraction pattern of Ni3Al+0.1 w/o B+10 w/o Cu showing no separate copper phase Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

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not improve the properties of the Ni3Al matrix containing diamond grit. The addition of iron decreased the sintered density of Ni3Al significantly. Sintered density, apparent hardness, and bending strength were maintained with the addition of nickel (Figures 1, 2, and 3), but significant degradation of the diamond particles occurred when the amount of nickel was greater than that required to maintain the Ni3Al stoichiometry; thus free nickel may be present. Since nickel is an effective catalyst in promoting graphitization, most of the diamond particles embedded in the Ni3Al+18 w/o Ni matrix suffered graphitization, as shown in Figure 5. With graphitized diamond surfaces, the interfacial bonding strength between the diamond particles and the matrix was reduced and the crushing strength of the diamond particles decreased, as shown in Figure 6. CONCLUSIONS An attempt has been made to decrease the extent of graphitization and enhance high-temperature retention of diamond grit by adding copper, nickel, and iron powders to the matrix. The results show that both copper and nickel can lower the peak temperature while maintaining the required strength, apparent hardness, and sintered density. However, the crushing strength of the diamond particles retrieved from the matrix decreases when nickel is added. This is attributed to severe graphitization since free nickel serves as a catalyst. The optimal result is attained by adding 10 w/o Cu. During wet cutting, the performance of the new matrix is similar to that of the conventional cobalt matrix. In the dry cutting, which is a high-temperature operation, the benefits of using the new Ni3Al+10 w/o Cu matrix with high hot strength are demonstrated. When SDB1100 diamond particles are used, the grinding ratio improves to 2000, an 8% increase over the value of 1850 for the Ni3Al matrix, while the cobalt matrix exhibits a grinding ratio of only 850. REFERENCES 1. J. Konstanty, “Production Parameters and Materials Selection of Powder Metallurgy Diamond Tools”, Powder Metall., 2006, vol. 49, pp. 299–306. 2. Y.S. Liao and S.Y. Luo, “Wear Characteristics of Sintered Diamond Composite During Circular Cutting”, Wear, 1992, vol. 157, pp. 325–37. 3. J. Konstanty, “Factors Affecting Diamond Retention in Stone Sawblade Segments”, Key Engineering Materials, 2003, vol. 250, pp. 13–20.

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4. S.W. Webb, “Diamond Retention in Sintered Cobalt Bonds for Stone Cutting and Drilling”, Diamond and Related Materials, 1999, vol. 8, no. 11, pp. 2,043–2,052. 5. Y.S. Liao and S.Y. Luo, “Effects of Matrix Characteristics on Diamond Composites”, J. Mat. Sci., 1993, vol. 28, pp. 1,245–1,251. 6. A. Romanski, H. Frydrych and J. Konstanty, ” Mechanical Properties of Cobalt-Base Materials for Diamond Impregnated Tools”, Proc. International Workshop on Diamond Tool Production, Turin, Italy, 1999, European Powder Metallurgy Association, Shrewsbury, UK, pp. 191–196. 7. D.A. Akyüz and H. Hofmann, “Interface Aspects in Cobaltbased Diamond Cutting Tool Segments”, Proc. 1998 PM World Congress, Hard Materials, European Powder Metallurgy Association, Shrewsbury, UK, vol. 2, pp. 158–163. 8. A. Aoki and O. Izumi, “Improvement in Room Temperature Ductility of the L1 2 Type Intermetallic Compound Ni 3 Al by Boron Addition”, J. Japan Inst. Metals, 1979, vol. 43, pp. 1,190–1,194. 9. P.H. Thornton, R.G. Davies and T.L. Johnston, “The Temperature Dependence of the Flow Stress of the γ’ Phase Based upon Ni3Al”, Metall. Trans., 1970, vol. 1, pp. 207–218 10. T.N. Tiegs, K.B. Alexander, K.P. Plucknett, P.A. Menchhofer, P.F. Becher and S.B. Waters, “Ceramic Composites with a Ductile Ni3Al Binder Phase”, Materials Science and Engineering, 1996, vol. 209A, pp. 243–247. 11. V. Weihnacht, W.D. Fan, K. Jagannadham, J. Narayan and C.T. Liu, “A New Design of Tungsten Carbide Tools with Diamond Coatings“, J. Mater. Res., 1996, vol. 11, no. 9, pp. 2,220–2,230. 12. D.E. Wittmer and P. Philip, “Intermetallic Bonded Diamond Composite Composition and Methods of Forming Articles from Same”, U.S. Patent Application 20060280638, December 14, 2006. 13. E.A. Levashov, I. Borovinskaya, A. Rogachov, M. Koizumi, M. Ohyanagi and S. Hosomi, “A New Method for Production of Diamond-Containing Ceramics”, Int. J. of Self-Propagating High-Temperature Synthesis, 1993, vol. 2, no. 2, pp. 189–201. 14. K.L. Padyukov and E.A. Levashov, “Self-Propagating HighTemperature Synthesis: a New Method for the Production of Diamond-Containing Materials”, Diamond and Related Materials, 1993, vol. 2, pp. 207–210. 15. S.C. Hu and K.S. Hwang, “Diamond Tools with Reaction Sintered Ni3Al Matrix”, Proc. 1998 PM World Congress, Reactive Synthesis, European Powder Metallurgy Association, Shrewsbury, UK, vol. 3, pp. 468–473. 16. K.S. Hwang, T.H. Yang and S.C. Hu, “Diamond Cutting Tools with a Ni3Al Matrix Processed by Reaction PseudoHipping”, Metall. And Mater. Trans. A, 2005, vol. 36A, pp. 2,801–2,806. 17. W. Misiolek and R.M. German, “Reactive Sintering and Reactive Hot Isostatic Compaction of Aluminide Matrix Composites”, Materials Science and Engineering A., 1991, vol. 144, pp. 1–10. 18. K.S. Hwang and Y.C. Lu, “Reaction Sintering of 0.1% B Added Ni3Al”, Powder Metall. Int., 1992, vol. 24, no. 5, pp. 1–10. ijpm

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Advances in Powder Metallurgy & Particulate Materials—2009 Proceedings of the 2009 International Conference on Powder Metallurgy & Particulate Materials Now available on a fully searchable CD-ROM in a format that preserves the original color of all figures—contains 110 technical papers encompassing over 1,200 pages

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RESEARCH & DEVELOPMENT

CORROSION BEHAVIOR OF SINTERED SURFACETREATED Nd-Fe-B MAGNETS

Emerson Alves Martins*, Jesualdo Luiz Rossi**, Mara Cristina Lopes de Oliveira***, Isolda Costa**, and Hercílio Gomes de Melo**** INTRODUCTION The outstanding magnetic properties of neodymium–iron–boron magnets are derived from the magnetic Nd2Fe14B intermetallic phase. Magnets based on this phase are of commercial interest since their energy product significantly exceeds that of samarium–cobalt magnets.1 They also exhibit high coercivity and are suitable for permanent magnet applications.2 Due to their excellent magnetic properties, they have been the subject of numerous studies and find many applications.3 However, these magnets exhibit low corrosion resistance4–8 due primarily to the fact that the rare-earth elements are electrochemically active, as well as to their complex microstructures. The microstructure of neodymium–iron–boron magnets is composed of two primary phases, the magnetic ϕ phase (Nd2Fe14B),9 and a neodymium-rich phase.10 The electrochemical potential differences among the various phases is significant, generating galvanic cells and attendant preferential dissolution of the most active phases. The most active is the neodymiumrich phase which surrounds the ϕ phase, leading to intergranular corrosion.11 These powder metallurgy (PM) sintered magnets exhibit intrinsic porosity, which further decreases their corrosion resistance. This class of magnets is useful in many applications, including small parts in the electro–electronics industry and larger parts in industrial machinery and dentistry. Magnets have been used in dentistry since the 1950s to improve the retention and stability of dental prostheses.12 In the past, however, the large size required to produce adequate forces limited their use.13–16 Since the introduction of rare-earth magnets, it is now possible to produce magnets with small dimensions for use in dental applications Presented at PowderMet2009 and published in Advances in Powder Metallurgy & Particulate Materials—2009, Proceedings of the 2009 Conference on Powder Metallurgy & Particulate Materials, which are available from the Publications Department of MPIF (www.mpif.org).

Permanent sintered neodymium–iron–boron magnets find many applications due to their excellent magnetic properties. However, these materials are highly susceptible to corrosion due primarily to their complex microstructure and inherent porosity. For applications in corrosive environments they are usually surface treated and coated or encapsulated. One of the surface treatments used is chromating, which, although effective, generates toxic Cr(VI) and carcinogenic residues. Thus it is important to find environmentally friendly treatments to replace those based on Cr(VI). The aim of the present work was to evaluate and compare the effectiveness of the corrosion protection afforded by two different surface treatments: a NaH2PO4-based conversion treatment and chromating with Cr(III) compounds, and to compare the corrosion response with that resulting from Cr(VI) chromating. The corrosion resistance of untreated and surfacetreated magnets was evaluated by electrochemical tests performed in a phosphatebuffered solution (PBS) at a neutral pH to simulate body fluids. Surface treatments with Cr(III) improved the corrosion resistance but phosphating provided superior corrosion resistance to chromating. Microstructural characterization showed that immersion in the phosphating solution resulted in a selective attack of the magnetic phase in the neighborhood of the neodymium-rich phase. There is also evidence that increased corrosion resistance is provided by the formation of a thin layer of phosphate on the magnetic phase.

*PhD Student, **Senior Researcher, ***Researcher, IPEN–CNEN/SP, CCTM, Av. Lineu Prestes, 2242, CEP 05508-000–São Paulo, Brazil; E-mail: [email protected], ****Lecturer, EPUSP–Departamento de Engenharia Química, Caixa Postal 61548, CEP 05424-970–São Paulo, Brazil

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CORROSION BEHAVIOR OF SINTERED SURFACE-TREATED Nd-Fe-B MAGNETS

as retentive devices for overdentures, mainly due to their strength and compactness.2,17,18 These magnets exhibit improvements in their maximum energy product compared with older types, leading to a significant reduction in the size required to generate the necessary magnetic flux.19 Dental materials mandate a high resistance to corrosion and must be inert to human tissues; however, neodymium–iron–boron magnets are highly susceptible to corrosion. One of the main problems associated with the clinical use of rareearth–iron–boron magnets is corrosion due to their low corrosion resistance in aqueous media. For dental applications these magnets are usually encapsulated in a stainless steel or titanium can. However, due to wear of the can or failure of the laser weld, saliva can leak into the can and lead to corrosion of the magnet. Due to their low corrosion resistance, neodymium–iron–boron permanent magnets need corrosion control methods to improve their viability as an engineering material and attempts have been made to improve their corrosion resistance by metallic coatings,20 by organic coatings,21 or by alloying.7 The corrosion control treatments must not decrease magnetic performance. Phosphating treatments have been used to prepare the surface of ferrous alloys for subsequent organic coating. However, only limited research has been reported on the effect of phosphating on the corrosion resistance of neodymium–iron–boron sintered magnets. Results suggest a significant improvement in their corrosion resistance compared with the untreated condition.22–25 Phosphating in a NaH2PO4 solution at room temperature has been shown to result in a substantial improvement in corrosion resistance. 24 The aim of the present work was to evaluate the effect of phosphating and chromating with Cr(III) baths on the corrosion resistance of sintered neodymium–iron–boron magnets as environmental friendly alternatives to surface treatment with solutions containing Cr(VI). MATERIAL AND METHODS Neodymium–iron–boron magnets produced by the Crucible Materials Corporation were used in this investigation. The composition (Table I), was

determined by X-ray fluorescence analysis and atomic absorption. The corrosion resistance of the magnets was investigated by electrochemical measurements, specifically, potentiodynamic polarization curves and electrochemical impedance spectroscopy (EIS) in a PBS of composition NaCl 8.77 g/L, Na2HPO4 1.42 g/L and KH2PO4 2.72 g/L at a pH 7, to simulate body fluids. Immersion tests were also carried out in sodium chloride electrolytes, specifically Hank’s solution and 3.5 w/o NaCl solution, since these simulate aggressive service applications. The neodymium– iron–boron magnets were tested in the demagnetized state. Scanning electron microscopy (SEM) was used for surface characterization following the corrosion tests. Specimen Preparation: For the electrochemical measurements, electrodes with an area approximately 130 mm2 were prepared by cold resin mounting. The electrode surface for exposure to the electrolyte was prepared by grinding on silicon carbide paper up to grade #600, degreasing with acetone, rinsing in deionized water, and drying under a hot-air stream. Some specimens were phosphated in a 10 g/L NaH2PO4 solution (pH 3.8) for times up to 4 h; others were chromated in two commercial solutions (SurTec 652 and SurTec 650) for Cr(VI) and Cr(III), respectively, at a concentration of 20 g/L (pH 3.7) for 3 min at 40°C. Experimental Set-Up: A three-electrode cell arrangement was used for the electrochemical measurements, with a platinum wire and a saturated calomel electrode (SCE) as the counter and reference electrodes, respectively. EIS measurements were made with a 1255 Solartron frequency response analyzer coupled to an EG&G 273A potentiostat. All the EIS measurements were performed in the potentiostatic mode at an open circuit potential, Eocp. The amplitude of the perturbation signal was 10 mV, and the frequency range studied was from 105 to 10-2 Hz, with 6 points per decade. The test medium was a PBS, naturally aerated at 23 ± 2°C. The potentiodynamic polarization measurements were

TABLE I. CHEMICAL COMPOSITION OF NEODYMIUM–IRON–BORON MAGNET (w/o)

46

Fe

Nd

B

Dy

Al

Co

Si

Cu

Nb

Na

Ca

S

60.59

28.31

1.00

2.09

3.73

1.28

1.39

0.18

0.66

0.41

0.15

0.16

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carried out with the potentiostat coupled to a computer with a scan rate of 1 mV/s. All the corrosion studies were carried out in a PBS at 25°C. RESULTS AND DISCUSSION Figure 1 shows the surface of the magnets after

Figure 2. Representative micrograph of neodymium–iron–boron magnet phosphated for 1 h and subsequently immersed for 4 h in PBS. SEM/ backscattered electron image

Figure 1. Representative surface images of neodymium–iron–boron magnets after: (a) 10-day immersion in Hank’s solution; (b) 30-day immersion in 3.5 w/o NaCl; (c) 4 h immersion in PBS. Markers 1 and 2 correspond to the neodymium-rich phase, 3 and 5 to pores, and 4 to the magnetic phase. SEM/secondary electron images Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

different periods of immersion in the various solutions. As expected, all the samples exhibited corrosion, independent of the electrolyte used, but the form of corrosion was dependent on the aggressiveness of the test solution. Intergranular corrosion was observed on the samples that were immersed in the PBS for short times (4 h). After this period of immersion, many areas of the neodymium-rich phase remained on the specimen surface but the areas surrounding this phase were generally attacked. In previous studies26 the same magnet demonstrated pitting corrosion after immersion in a cell culture medium for 10 days. Chloride was detected on the corroded area, showing that it must have initiated the corrosion process. The pitting corrosion found in the samples immersed in the culture medium also demonstrated passivation of some regions on the surface. In Hank’s solution a passive state was absent and corrosion was of a generalized type. These results are related to the increased aggressiveness of Hank’s solution compared with the cell culture medium. The combination of a porous sintered structure, a complex microstructure (mixture of phases of different electrical potential), and the presence of chlorides in the test medium are the primary causes of corrosion in these magnets. Figure 2 shows a representative SEM image of the surface of a phosphated magnet after 4 h of immersion in the PBS. After this period, only limited signs of corrosion are visible, confirming the beneficial effect of the phosphating treatment on corrosion resistance. The effect of chromatizing the neodymium–iron–

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Figure 3. Representative micrographs of the surface of neodymium–iron–boron magnets after: (a) immersion in Cr(III) and (b) immersion in Cr(VI)-containing baths. SEM/backscattered electron images

boron magnet in a Cr(III)-containing bath on its corrosion resistance was also investigated as another alternative to the toxic Cr(VI)-containing baths. Figure 3 shows SEM images of the surface of specimens treated with Cr(III) or Cr(VI)-containing baths. The surface treatments localized the attack primarily at the pores and the boundaries between the neodymium-rich phase and the magnetic phase. However, the Cr(VI) bath was more aggressive and increased localized corrosion was identi-

fied in the SEM image. The formation of chromate layers from the Cr (III) or Cr (VI) baths is due to oxidation of either the HCrO4- (chromate) or the Cr2O72- (dichromate), both of which are known to exist at high CrO3 concentrations.27 These species are involved in a redox process leading to precipitation of the chromate layer (Cr(III)). It is possible that the heterogeneous nature of the magnet surface leads to precipitation of the conversion layer in preferred sites. Accordingly, oxidation of the substrate takes place only at the more active sites

Figure 4. (a) Nyquist and (b) Bode phase-angle diagrams of neodymium–iron–boron magnets untreated and surface treated by phosphating and chromating in (Cr(III) and Cr(VI)) baths

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leading to the localized corrosion depicted in Figure 3(b). The EIS diagrams for the neodymium–iron– boron untreated magnet, and for the phosphated and chromated (Cr(III) and Cr(VI))-treated samples are shown in Figure 4 as Nyquist and Bode (phaseangle) diagrams. The Nyquist diagrams exhibit a depressed capacitance loop at medium to low frequencies. The flattened appearance of this time constant is related to the complex electrochemical processes taking place on the heterogeneous magnet surface and is indicative of the superposition of several time constants. For samples tested under the same conditions, a larger diameter corresponds to a higher corrosion resistance. The Bode phase-angle diagrams for all surface treated specimens show a peak at high frequencies associated with the conversion coating, indicating the presence of a surface layer independent of the corrosion protection afforded; this feature is absent in the untreated sample. For the two chromate layers, the increased high-frequency phase angle of the sample protected with the layer formed in the Cr(III) bath is indicative of its superior performance, as verified in the Nyquist diagrams. Conversely, the corrosion resistance of the sample treated in the Cr(VI) bath is inferior to that exhibited by the uncoated sample, Figure 5. The phosphate-treated samples exhibit a broader phase angle (indicative of the superposition of several phenomena) and the 45° angle between the Nyquist diagram and the real axis is indicative

Figure 5. Nyquist diagrams of untreated and chromated Cr(VI) neodymium–iron–boron magnets in PBS Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

of a porous electrode response, consistent with de Levie theory.28 The results presented in Figure 4 allow for a ranking of the corrosion protection afforded by the surface treatments as follows: phosphate > Cr(III) > Cr(VI). While the two former treatments improve corrosion resistance, the latter treatment results in decreased corrosion resistance. With phosphating, the solution is acidic (pH 3.8) and will have a small corrosive effect on the magnet surfaces during treatment. Commercial phosphating baths with a lower pH (2.7) have also been evaluated but were too aggressive. Consequently, phosphating solutions must be specifically developed for use with these magnets. The solution adopted in the present study had a higher pH than commercial baths and allowed the formation of a thin and adherent phosphate layer on essentially all the magnet surfaces. These results show that the corrosiveness of the phosphate solution must be precisely controlled to avoid attack of the magnets surfaces. Previous work carried out with the same magnet and a phosphated solution of similar composition (but with the pH adjusted to 2) exhibited severe attack of the substrate accompanied by formation of a white corrosion products after 1 day of immersion. During acid attack in the first step of phosphating (formation of cations for primary phosphate formation), NdH2 tends to form, which may result in the following reactions, depending on the pH:29 NdH2 + 3H+→ Nd3++ 5/2H2 (pH ≤ 3.5)

(1)

NdH2 + 3H2O → Nd(OH)3 + 5/2H2 (pH range from 3.5 to 5)

(2)

Figure 6. Micrograph of neodymium–iron–boron magnet phosphated 1 h in 10 g/L NaH2PO4. SEM/backscattered electron image

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Local changes in pH can favor the first reaction leading to localized corrosion processes. However, under normal conditions at a higher pH, the formation of a protective hydroxide or phosphate occurs, depending on the solution chemistry. Upon immersion in the NaH 2 PO 4 solution, phosphating of the neodymium–iron–boron magnets occurred resulting in a thin surface film with interference colors. Some surface attack occurred during this treatment, as shown by the SEM observations, Figure 6. As documented in the images, areas attacked corresponded to the magnetic (φ) phase surrounded by the neodymium-rich phase. Unpredictably, this phase was not attacked, whereas the nobler φ phase around it was corroded. Thus a polarity inversion must have resulted from the phosphating treatment. This is attributed to the rapid formation of a protective phosphate film on the neodymium-rich phase, leading to its ennoblement, and to the attack of surrounding areas of the magnetic phase. The extremely low solubility of neodymium phosphate explains this behavior. Phosphating of the magnetic phase also occurs, but the kinetics are slower than that of the neodymium-rich phase. EIS tests carried out in the PBS, for sample surfaces treated in a 10 gL-1 NaH2PO4 solution (pH 3.8) for varying times, showed an increase in the impedance with time. The EIS results showed the presence of a high frequency (HF) time constant that can be ascribed to a thin phosphate layer on the surface (confirmed by EDS analysis

and visual observations through interference colors). The EIS results also showed that the neodymium–iron–boron behaves as a porous electrode, and the phosphate layer is also formed on the pore walls, hindering the corrosion reaction. The results of the present study also determined that the phosphate treatment time is an important parameter. Even though the EIS results indicated an increase in the impedance of the magnets with phosphating time from 1 h to 4 h, SEM observations of the surfaces revealed that the degree of attack also increased (Figure 7). The preferential attack of the matrix occurred at the vicinity of the neodymium-rich phase, due probably to a combination of surface defects and passivating phosphate film on this phase. This result indicates that the phosphating treatment should not be carried out for long periods of time. CONCLUSIONS Surface treatment of neodymium–iron–boron sintered magnets in Cr(III) solutions improved corrosion resistance, but phosphating provided better corrosion resistance than chromating, as demonstrated by EIS. The solution pH and time of treatment are important variables that must be controlled in order to optimize anticorrosion performance. SEM characterization of the magnets after surface treatment showed that immersion in the phosphating solution resulted in selective attack of the magnetic phase at the boundaries with the neodymium-rich phase, whereas this

Figure 7. Micrographs of phosphated neodymium–iron–boron magnets (a) 1 h, and (b) 4 h in 10 g/L NaH2PO4. SEM/backscattered electron images

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phase was not unattacked. This is attributed to the formation of a protective phosphate film on the neodymium-rich phase, due to the low solubility of neodymium phosphates. After the formation of this film, a polarity inversion occurred causing attack of the magnetic phase surrounding the neodymium-rich one. There was also evidence that a thin phosphate layer formed on the magnetic phase, confirmed by EDS analysis and visual observations through interference colors. This contributed to the increase in the corrosion resistance of the phosphated magnets. REFERENCES 1. J.F. Herbst and J.J. Croat, “Neodymium Iron Boron Permanent-Magnets”, J. Mag. Magn. Mat., 1991, vol. 100, no. 1–3, pp. 57–78. 2. M. Sagawa, S. Fujimura, N. Togawa, H. Yamamoto and Y. Matsuura, “New Material for Permanent-Magnets on a Base of Nd and Fe”, J. Appl. Phys., 1984, vol. 55, no.. 6, pp. 2,083–2,087. 3. K.H.J. Buschow, “New Permanent Magnet Materials”, Mater. Sci. Report, 1986, vol. 1, no. 1, pp. 1–66. 4. T.S. Chin, R.T. Chang, W.T. Tsai and M.P. Hung, “Electrochemical-Behavior of Rare-Earth Magnet Alloys in Various Solutions” IEEE Trans. Mag., 1988, vol. 24, no. 2, part 2, pp. 1,927–1,929. 5. J. Jacobson and A. Kim, “Oxidation Behavior of Nd-Fe-B Magnets”, J. Appl. Phys., 1987, vol. 61, no. 8, part 2A, pp. 3,763–3,765. 6. H. Bala, S. Szymura and J.J. Wyslocki, “Electrochemical Corrosion-Resistance of Fe-Nd-B Permanent-Magnets” J. Mater. Sci., 1990, vol. 25, no. 1B, pp. 571–574. 7. H. Bala, G. Pawlowska, S. Szymura and Y.M. Rabinovich, “Corrosion Characteristics of Nd-Fe-B Sintered Magnets Containing Various Alloying Elements”, J. Mag. Magnetic Mat., 1990, vol. 87, no. 3, pp. L255–L259. 8. H. Bala and S. Szymura, “An Electrochemical Investigation of Dissolution of Nd-Fe-B Magnets in AcidSolution Under Cathodic Polarization”, Corr. Sci., 1991, vol. 32, no. 9, pp. 953–963. 9. R.H.J. Fastenau and E.J. vanLoenen, “Applications of Rare Earth Permanent Magnets”, J. Mag. Magn. Mat., 1996, vol. 158, pp. 1–6. 10. J.M. Jacobson and A.S. Kim, “Oxidation Behavior of NdFe-B Magnets”, J. App. Phys., 1987, vol. 61, no. 8, pp. 3,763–3,765. 11. D.W Scott, B.M. Ma, Y.L. Liang and C.O. Bounds, “The Effects of Average Grain Size on the Magnetic Properties and Corrosion Resistance of NdFeB Sintered Magnets”, J. Appl. Phys., 1996, vol. 79, no. 8, part 2A, pp. 5,501–5,503. 12. H. Freedman, “Magnets to Stabilize Dentures”, J. Am. Dent. Assoc., 1953, vol. 47, no. 3, pp. 288–297. 13. B.K. Moghadam and F.R. Scandrett, “Magnetic Retention for Overdentures”, J. Prosthet. Dent., 1979, vol. 41, no. 1, pp. 26–29.

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14. S.J. Behrman, “Magnets Implanted in Mandible—Aid to Denture Retention”, J. Am. Dent. Associ., 1964, vol. 68, no. 2, pp. 206–215. 15. N. Javid, “Use of Magnets in a Maxillofacial Prosthesis”, J. Prosthet. Dent., 1971, vol. 25, no. 3, pp. 334–341. 16. M. Wilson, H. Kpendema, J.H. Noar, N. Hunt and N.J. Mordan, “Corrosion of Intraoral Magnets in the Presence and Absence of Biofilms of Streptococcus-Sanguis”, Biomaterials, 1995, vol. 16, no. 9, pp. 721–725. 17. J.J. Becker, “Permanent Magnets”, Scientific American, 1970, vol. 223, no. 6, pp. 92–100. 18. H. Tsutsui, Y. Kinouchi, H. Sasaki, M. Shiota and T. Ushita, “Studies on the Sm-Co Magnet as a Dental Material”, J. Dent. Res., 1979, vol. 58, no. 6, pp. 1,597–1,606. 19. I.R. Harris, “Hard Magnets”, Mater Sci. Technol., 1990, vol. 6, no. 10, pp. 962–966. 20. A. Ahmad, P.J. McGuiness and I.R. Harris, “A Study of the Microstructures and the Effects of Coating on Nd2Fe14B Alloys”, IEEE Trans Magn., 1990, vol. 26, no. 5, pp. 2,625–2,627. 21. C.W. Cheng, H.C. Man and F.T. Cheng, “Magnetic and Corrosion Characteristics of Nd-Fe-B Magnet with Various Surface Coatings”, IEEE Trans Magn., 1997, vol. 33, no. 5, pp. 3,910–3,912. 22. A.M. Saliba-Silva and I. Costa, “Corrosion Protection of a Commercial NdFeB Magnet by Phosphating”, Key Eng. Mat., 2001, vol. 189-1, pp. 363–368. 23. A.M. Saliba-Silva, H.G. de Melo, M.A. Baker, A.M. Brown and I. Costa, “Characterization of Sintered NdFeB Magnets after Phosphating in Alkaline and Acidic Environments”, Mater. Sci. Forum, 2003, vol. 416-4, pp. 54–59. 24. A.M. Saliba-Silva, R.N. Faria, M.A. Baker and I. Costa, “Improving the Corrosion Resistance of NdFeB Magnets: an Electrochemical and Surface Analytical Study”, Surf. Coat. Technol., 2004, vol. 185, no. 2–3, pp. 321–328. 25. S.M. Tamborim Takeuchi, D.S. Azambuja, A.M. SalibaSilva and I. Costa, “Corrosion Protection of NdFeB Magnets by Phosphating with Tungstate Incorporation”, Surf. Coat. Technol., 2006, vol. 200, no. 24, pp. 6,826–6,831. 26. I. Costa, M.C.L. Oliveira, H. Takiishi, M. Saiki and R.N. Faria, “Corrosion Behaviour of Commercial NdFeB Magnets – the Effect of Magnetization”, Key Eng. Mat., 2001, vol. 189-1, pp. 340–345. 27. P. Campestrini, G. Goeminne, H. Terryn, J. Vereecken and J.H.W. de Wit, “Chromate Conversion Coating on Aluminum Alloys - I. Formation Mechanism”, J. Electrochem. Soc., 2004, vol. 151, no. 2, pp. B59–B70. 28. R. de Levie, Adv. Electrochemistry, 1967, vol. 6, pp. 329. In: O.E. Barcia, E. D’Elia, I. Frateur, O.R. Mattos, N. Pébère and B. Tribollet, “Application of the Impedance Model of de Levie for the Characterization of Porous Electrodes”, Electrochimica Acta, vol. 47, no. 13–14, pp. 2,109–2,116. 29. B. Rupp, A. Resnik, D. Shaltiel and P. Rogl, “PhaseRelations and Hydrogen Absorption of Neodymium Iron (boron) Alloys”, J. Mat. Science, 1988, vol. 23, no. 6, pp. 2,133–2,141. ijpm

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MEETINGS AND CONFERENCES

2010 INTERNATIONAL CONFERENCE & EXHIBITION ON POWDER METALLURGY PROCESSING OF PARTICULATE MATERIALS AND PRODUCTS & 36TH ANNUAL TECHNICAL MEETING January 28–30 Rajasthan, India www.pmai.in MIM2010: INTERNATIONAL CONFERENCE ON INJECTION MOLDING OF METAL, CERAMICS AND CARBIDES MARCH 29–31 Long Beach, CA MPIF* POWTECH 2010 April 27–29 Nuremberg, Germany www.powtech.de

PRICM 7 7th PACIFIC RIM INTERNATIONAL CONFERENCE ON ADVANCED MATERIALS AND PROCESSING August 1–5 Cairns, Australia www.materialsaustralia.com. au/scripts/cgiip.exe/WServ ice=MA/ccms.r?PageID=190 70 ILASS 2010 23rd ANNUAL CONFERENCE ON LIQUID ATOMIZATION AND SPRAY SYSTEMS September 6–8 Brno, Czech Republic www.ilasseurope2010.org 7th INTERNATIONAL SYMPOSIUM ON ALLOY 718 & DERIVATIVES September 10–13 Pittsburgh, PA www.tms.org

PowderMet2010: MPIF/APMI INTERNATIONAL CONFERENCE ON POWDER METALLURGY & PARTICULATE MATERIALS June 27–30 Hollywood (Ft. Lauderdale), FL MPIF*

PM SINTERING SEMINAR September TBA MPIF*

BASIC PM SHORT COURSE July 25–28 State College, PA MPIF*

PM2010 WORLD CONGRESS October 10–14 Florence, Italy www.epma.com/pm2010

2011 PowderMet2011: MPIF/APMI INTERNATIONAL CONFERENCE ON POWDER METALLURGY & PARTICULATE MATERIALS June 19–22 Chicago, IL 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

TITANIUM 2010 October 3–5 Orlando, FL www.titanium.org

*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

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MEMBERSHIP APPLICATION Please type or print legibly.

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YEARLY CONTENTS INTERNATIONAL JOURNAL OF POWDER METALLURGY TABLE OF CONTENTS FOR VOLUME 45, NUMBERS 1–6, 2009

45/1 JANUARY/ FEBRUARY 2009 2 4 6 9

Editor’s Note PM Industry News in Review PMT Spotlight On …Zachary Z. Zebrovious Consultants’ Corner David Whittaker

ENGINEERING & TECHNOLOGY 13 Engineering the Green State of Powder Products D. Whittaker RESEARCH & DEVELOPMENT 19 Optimization of Metal Powder-Mixing Parameters for Chemical Homogeneity and Agglomeration N. Vlachos and I.T.H. Chang 29 Effect of Axial and Radial Metal Powder Mixing on Chemical Homogeneity and Agglomeration N. Vlachos and I.T.H. Chang OUTSTANDING TECHNICAL PAPER: PM2008 WORLD CONGRESS 38 Development of a Dual-Phase PrecipitationHardening PM Stainless Steel C.T. Schade, T.F. Murphy, A. Lawley and R.D. Doherty HISTORICAL PROFILE 47 The Origin and Role of APMI International in North America’s PM Industry K.H. Roll DEPARTMENTS 57 APMI Membership Application 58 Web Site Directory 64 Advertisers’ Index

45/2 MARCH/APRIL 2009 2 4 7 11 13

FOCUS: PM Metallography 17 Cutting-Edge PM Metallography T.F. Murphy, FAPMI 19 Three-Dimensional Characterization and Modeling of Porosity in PM Steels N. Chawla, J.J. Williams, X. Deng, C. McClimon, L. Hunter and S.H. Lau 29 Characterization of Powders and PM Components Utilizing Transmission Electron Microscopy C. Blais, G. L’Espérance, FAPMI, and P. Plamondon 39 Porosity Statistics and Fatigue Strength of Sintered Iron P. Beiss, FAPMI, and S. Lindlohr 49 Evaluation of PM Fracture Surfaces Using Quantitative Fractography T.F. Murphy, FAPMI DEPARTMENTS 62 Meetings and Conferences 63 PM Bookshelf 64 Advertisers’ Index

45/4 JULY/AUGUST 2009

45/3 MAY/JUNE 2009 2 4 7 11

Editor's Note PM Industry News in Review Consultants’ Corner James G. Marsden Technology Investments Key To PM’s Future Peter K. Johnson 17 Exhibitor Showcase: PowderMet2009 RESEARCH & DEVELOPMENT 25 Powder Injection Molding of Metal and Ceramic Hip Implants J. Song, T. Barriere, J-C. Gelin and B. Liu 36 Iron-Base PM Matrix Alloys for Diamond-Impregnated Tools M. Zak-Szwed, J. Konstanty and A. Zielinska-Lipiec 45 Processing of Bulk Fe-Zn Alloys Using Explosive Compaction R.P. Corson, S. Guruswamy, M.K. McCarter and C-L. Lin ENGINEERING & TECHNOLOGY 55 Improvement in Fatigue Performance of PowderForged Connecting Rods by Shot Peening E. Ilia, R.A. Chernenkoff and K.T. Tutton DEPARTMENTS 62 Meetings and Conferences 63 PM Bookshelf 64 Advertisers’ Index

54

Editor's Note PM Industry News in Review Company Profile Magnesium Elektron Powders PMT Spotlight On …Stan Cuthbert Consultants’ Corner Joseph Tunick Strauss

2 4 7 9 11 12 15

Editor's Note PM Industry News in Review PMT Spotlight On …Shawn Metcalfe Consultants’ Corner O. Grinder 2009 APMI Fellow Award 2009 Poster Awards 2009 PM Design Excellence Award Competition Winners 20 Axel Madsen/CPMT Scholar Reports ENGINEERING & TECHNOLOGY 23 State of the PM Industry in North America—2009 M. Paullin RESEARCH & DEVELOPMENT 27 Processing and Behavior of Fe-Based Metallic Glass Components via Laser-Engineered Net Shaping B. Zheng, Y. Zhou, J.E. Smugeresky and E.J. Lavernia 40 Exothermic Reactions During the Sintering of Elemental Iron and Aluminum Powder Mixes S. Jó´zwiak, K. Karczewski and Z. Bojar DEPARTMENTS 45 Meetings and Conferences 46 Instructions for Authors 48 Advertisers’ Index

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YEARLY CONTENTS 45/5 SEPTEMBER/OCTOBER 2009 2 4 7 9

Editor’s Note PM Industry News in Review Newsmaker Richard Pfingstler Consultants’ Corner John A. Shields, Jr.

FOCUS: Precious Metals 15 Precious Metals: A Valuable Powder Metallurgy Player P.W. Taubenblat, FAPMI 21 Applications for Precious Metal Powders J.T. Strauss 29 The Manufacture of Platinum, Gold, and Palladium Powders H.D. Glicksman 37 Precious Metal Powder Precipitation and Processing S. Frink and P. Connor 43 Additive Manufacturing of Precious Metal Dental Restorations A.L. Hancox and J.A. McDaniel DEPARTMENTS 53 Meetings and Conferences 55 APMI Membership Application 56 Advertisers’ Index

Volume 45, Issue 6, 2009 International Journal of Powder Metallurgy

45/6 NOVEMBER/DECEMBER 2009 Editor’s Note PM Industry News in Review PMT Spotlight On …Leander F. Pease III, FAPMI Excellence in Metallography Awards Consultants’ Corner Myron I. (Mike) Jaffe ENGINEERING & TECHNOLOGY Particle-Reinforced Carbon-Free PrecipitationHardening High-Speed Steels H. Danninger, C. Harold, F. Rouzbahani, H. Ponemayr, M. Daxelmüller, F. Simancik and K. Izdinsky RESEARCH & DEVELOPMENT Diamond Cutting Tools with a Ni3Al+ Copper Matrix K.S. Hwang and T.H. Yang Corrosion Behavior of Sintered Surface-Treated Nd-Fe-B Magnets E.A. Martins, J.L. Rossi, M.C.L. de Oliveira, I. Costa and H.G. de Melo DEPARTMENTS Meetings and Conferences APMI Membership Application Instructions for Authors PM Bookshelf Table of Contents: Volume 45, Numbers 1–6, 2009 Advertisers’ Index

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ADVERTISERS’ INDEX

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ACE IRON & METAL CO. INC._________(269) 342-0185 ______________________________________________________5 ACUPOWDER INTERNATIONAL, LLC ___(908) 851-4597 ________www.acupowder.com ___________________________35 ASBURY CARBONS __________________(908) 537-2908 _________www.asbury.com _________________________________20 ELNIK SYSTEMS ____________________(973) 239-6066 _________www.elnik.com ___________________________________6 GLOBAL TITANIUM _________________(313) 366-5305 ________www.globaltitanium.com ________________________34 HOEGANAES CORPORATION _________(856) 786-2574 ________www.hoeganaes.com ___________INSIDE FRONT COVER QMP ____________________________(734) 953-0082 ________www.qmp-powders.com ________________BACK COVER SCM METAL PRODUCTS, INC. ________(919) 544-7996 ________www.scmmetals.com ____________INSIDE BACK COVER

ADVERTISER’S REQUEST FOR INFORMATION FAX FORM Need more information on products or services seen in this issue?

Complete the form below and fax to the advertiser(s) of your choice. Fax numbers are listed in the advertisers’ index above.

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 45, Issue 6, 2009 International Journal of Powder Metallurgy

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