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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] President of APMI International Nicholas T. Mares
[email protected] Executive Director/CEO, APMI International C. James Trombino, CAE
[email protected]
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international journal of
powder metallurgy Contents 2 4 7 11 17
45/3 May/June 2009
Editor's Note PM Industry News in Review Consultants’ Corner James G. Marsden Technology Investments Key To PM’s Future Peter K. Johnson 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 Powder-Forged 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 Cover: TEM micrograph of Fe-Cu-Sn compact hot-pressed at 900°C. Photo courtesy Janusz Konstanty, AGH–University of Science & Technology.
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 by Cadmus Communications Corporation, P.O. Box 27367, Richmond, Virginia 23261-7367. 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 INTERNATIONAL 105 College Road East, Princeton, New Jersey 08540-6692 USA Tel: (609) 452-7700 • Fax: (609) 987-8523 • E-mail:
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EDITOR’S NOTE
I
t’s “Show Time” in Las Vegas again! The technical program for PowderMet2009 embraces all facets of PM science and technology—from basic research to parts fabrication. Attendees are expected from over 30 countries at this “window” on new developments, trends, and future prospects for the PM industry. The Exhibitor Showcase in our Show Issue includes profiles of the participating companies. From Peter Johnson’s annual technology review of the PM industry, based on input from MPIF-member companies, it is evident that, notwithstanding the current economic climate, investment in new PM technologies has been sustained. This is viewed as the key to the future health and viability of the powder producers, equipment manufacturers, and the parts fabricators. Consistent with the MPIF-member company input, the Wall Street Journal recently reported that many large companies are maintaining the level of their R&D spending in the face of falling revenues. If you have a sintering problem, particularly in the context of carbon control in ferrous alloys, Jim Marsden’s “Consultants’ Corner” should prove to be invaluable. His column also discusses practical approaches to improving the environmental performance of PM parts-manufacturing plants. In the “Engineering & Technology” section, Ilia, Chernenkoff, and Tutton quantify and interpret the beneficial effect of shot peening on the fatigue performance of powder-forged connecting rods. Three diverse topics make up the content of the “Research & Development” section: Song et al. detail the results of a dual experimental and modeling study of the fabrication of metal and ceramic implants by powder injection molding; The viability of using iron-base PM alloys in place of cobalt-base materials as a matrix in diamond-impregnated tools is demonstrated by Zak-Szwed, Konstanty, and Zielinska-Lipiec. The front cover shows the microstructure of a iron–copper–tin matrix alloy, as seen by transmission electron microscopy; To compare magnetostriction in iron–zinc alloys with that of iron–gallium, Corson et al. describe the fabrication of bulk iron–zinc alloys with a [100] texture by explosive compaction in the presence of a magnetic field. Unlike with gallium, no significant change in magnetostriction is observed when zinc is substituted for iron.
Alan Lawley Editor-in-Chief
The new U.S. administration, as well as the global economic slowdown and the stimulus packages it gave rise to, are expected to have a major impact on a number of key issues in academe: • With a renewed focus on energy research, the nation’s major research universities are likely to be the primary beneficiaries. • The National Institutes of Health and the National Science Foundation, the leading providers of federal science money to universities, apparently are taking diverse approaches to the use of stimulus funds. The former will modify its distribution guidelines to ensure a measure of geographic parity while the latter will not. • The economic downturn is reflected in a tight job market for students enrolled in cooperative education programs. In the longer term, however, the recession may well give new life to cooperative education since, increasingly, employees are recognizing the benefit of hiring students with experience, without the cost of a benefit package. • The New York Times notes that with the new administration avowedly committed to science, now is the time to attract, and retain, more women. The article cites a number of compelling historical reasons why this has not been the case to date, and offers potential solutions. Yes, indeed, the times they are a-changing!
2
Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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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.”
Plansee Makes Massive HIP Cylinder Plansee Metall, Reutte, Austria, has made a doped molybdenum cylinder 16.5 feet long and more than 7.26 feet in diameter for Avure Technologies, Inc., builder of hot isostatic pressing (HIP) units. Plansee claims its manufacturing feat represents the largest HIP unit ever built. New MIM Materials Standard Released MPIF has released a new, MIM standardized material designation code, information, and property data as an addendum to the 2007 edition of MPIF Standard 35, “Materials Standards for Metal Injection Molded Parts.” The new Low-Alloy Steel standard includes chemical composition and data tables (both inch–pound and SI units) for the MIM-4140 quenched-and-tempered material. The new standard is posted on the MPIF Web site as a free-access document until it is included in the next printed edition of the standard publication. New Magnesium Powder Applications Magnesium powders are used in a variety of applications, reports Magnesium Elektron Powders, Manchester, N.J. Significant markets include defense, chemical, and pharmaceutical markets. Laser Sintering Dental Crowns and Bridges Dental implants manufactured by
4
direct laser sintering represent a growth market, reports EOS GmbH Electro Optical Systems, Krailling bei München, Germany. The company’s EOSINT M 270 system uses CAD data to produce complex parts from a CobaltChrome SP2 powder alloy it developed for dental crowns and bridges. Energy-Efficient Injection Molding Arburg GmbH + Co KG, Lossburg, Germany, introduces the new Allrounder H hybrid-drive concept injection molding machine that combines servo-electric and hydraulicmovement axes. The machine combines high performance with energy-efficient drive technology. Powder Maker Pauses Production Qit-Fer et Titane Inc., a wholly owned subsidiary of Rio Tinto, announced a temporary closure this summer of its smelter, upgraded slag division, and steel billet plant in Sorel-Tracy, Québec, for eight weeks from July 12 to September 8. Because of this decision, the QMP steel powder plant will have limited operations but will stockpile material to meet ongoing demand of customers. New Tungsten Investment Malaga Inc., Montréal, Québec, a tungsten mining company, has closed the first tranche (portion) of a private placement by issuing a fiveyear $1.3 million senior secured convertible promissory note to Global Tungsten & Powders Corp. (GTP),
Towanda, Pa. The private placement will be conducted in four tranches to be closed within the next four months, for a total of $3.8 million. PM Competes in Automotive Awards Program A camshaft-phasing system developed by BorgWarner Morse TEC Inc. is one of 25 finalists in the Automotive News annual PACE awards honoring innovative suppliers. The variable valve timing (VVT) product contains three PM steel parts made at the company’s Cortland, N.Y., plant. Nanoparticles Transmit Energy in Solar Cells Clarkson University’s Center for Advanced Materials Processing (CAMP), Potsdam, N.Y., is developing nanomaterials for advanced siliconbased solar cells. The CAMP group headed by Professor Dan Goia uses inkjet printing to deposit silver and gold nanoparticles on silicon wafers. PM Company Receives SmallBusiness Loan The state of Pennsylvania Department of Community and Economic Development (DCED) gave a $200,000 Small Business First loan to SinterFire Inc., Kersey, Pa. It is the company’s fourth such loan and will be used to buy equipment to increase manufacturing capacity by 20 percent and retain 25 jobs. ijpm
Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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Conversion Kit for Laboratory Batch Mills Union Process, Inc., Akron, Ohio, offers a conversion kit to retrofit its 1-S attritor model from a laboratory batch mill to a circulation mill. The kit contains an upper discharge chamber, grid plate assembly, impeller, and agitator shaft/arm assembly.
Chinese PM Industry Results PM parts production in China declined five percent to 102,048 short tons last year, reports the PM Association of China, Beijing. Iron-base parts production declined to 95,706 short tons while copper parts production weakened slightly to 6,342 short tons.
Höganäs Powder Sales Sag Swedish metal powder producer Höganäs AB reports first quarter 2009 sales declined 42 percent to MSEK 916 (about $112 million). Production volumes fell sharply in all regions.
New Resource for Metal Injection Molding Information The Metal Injection Molding Association (MIMA), one of the six federated trade associations of the Metal Powder Industries Federation (MPIF), has launched
a new industry-funded Web site, mimaweb.org, to promote the benefits of metal injection molding (MIM) as a part-manufacturing technology. Miba Advances amid Softening Sales Miba AG, Laakirchen, Austria, announced a 2.1 percent increase of fiscal year 2008–09 sales to 374.6 million (about $508 million), despite a sharply declining fourth quarter. Based on the slowdown in the automotive market, the Sinter (PM parts) Group reported a 15 percent sales drop to 135.4 million (about $183 million). ijpm
PURCHASER & PROCESSOR
Powder Metal Scrap (800) 313-9672 Since 1946
Ferrous & Non-Ferrous Metals Green, Sintered, Floor Sweeps, Furnace & Maintenance Scrap
1403 Fourth St. • Kalamazoo, MI 49048 • Tel: 269-342-0183 • Fax: 269-342-0185 Robert Lando E-mail:
[email protected] Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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Focusing on Solutions
At ACuPowder, we’ve created a unique “focused approach” designed to provide innovative solutions to our customers’ problems. We’re more than just a supplier of goods, we’re a provider of ideas. We work closely with customers to assess needs and create workable responses, tailored exactly to meet their objectives. Our knowledgeable support staff looks beyond the ordinary to develop programs that deliver extraordinary results. With more than 90 years industry experience, ACuPowder welcomes even the toughest assignments. Put us to the test. You’ll quickly learn that we are totally focused on you. So depend on ACuPowder as your “one-stop source” for Copper, Tin, Bronze, Brass, Copper Powder Infiltrant, Bronze Premixes, Antimony, Bismuth, Manganese, MnS+ Nickel, Silicon, Graphite and P/M Lubricants. New products include powders for MIM, Thermal Management, “Green” Bullets, Lead Free Solders, Plastic Fillers, Cold Casting and most recently Ultra Fine/Ultra Pure Copper Powders for the electronics industry and ULTRA INFILTRANT the wrought/wire infiltration solution.
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[email protected] • web: www.acupowder.com
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CONSULTANTS’ CORNER
JAMES G. MARSDEN, FAPMI* Q
Apart from eliminating metallic stearates as PM lubricants, what can PM part manufacturing plants do to improve environmental performance? A critical step in improving the environment is the reduction of smoke, odor, and particulate material emitted from the furnace stack during the lubricant burn-off cycle. Even with wax-base lubricants, such as Acrawax, the percentage of carbon is high and is a major player in causing not only environmental problems but also premature deterioration of the furnace components. Most protective atmospheres used today are oxygen-free atmospheres such as nitrogen–hydrogen or nitrogen-diluted atmosphere blends that normally exhibit a low dew point. The dew point can range anywhere from -21°C to -26°C (-5°F to -15°F) for the nitrogen– endothermic gas system to -40°C to -51°C (-40°F to -60°F) for the nitrogen–dissociated ammonia and nitrogen–hydrogen systems. These low dew points are excellent for carbon control in the high-heat zone but devastating for the preheat zone of the sintering furnace. Not only do these carbonaceous vapors deposit as stalactites and stalagmites in both the preheat and high-heat zones of the furnace, but they also tend to carburize both the belts and muffles, which causes premature failure of the furnace components. In addition, the vapors that are removed from the furnace by the atmosphere will tend to pollute the environment, as well as accumulate in the furnace stack. The accumulation of these carbonaceous vapors can, in some cases, result in stack fires, which can prove dangerous as well as be costly to the PM parts manufacturer. To remove these vapors from the furnace before they can accumulate, there must be an oxidant added to the atmosphere to attack these vapors and form gaseous compounds of CO, CO2, and some hydrocarbons. Once the gaseous compounds are formed they can easily be removed from the furnace by the furnace atmosphere. To my knowledge there are only two hydration systems available that are
A
designed exclusively for this purpose. One is the bubbler system and the other is the steam-injection system. Although both systems are designed to inject moisture into the preheat zone, using nitrogen as a carrier gas, there is one major difference. The bubbler system offers an intermittent dew point and moisturized nitrogen at temperatures ~71°C to 82°C (160°F to 180°F). With this system, nitrogen is passed through a tank of water and the amount of moisture picked up by the nitrogen is dependent on the water temperature. The steaminjection system maintains a constant dew point with the moisture injected at temperatures between 121°C and 177°C (between 250°F and 350°F) for more precise and continuous lubricant removal. The steam-injection (moisture-control) system injects a meter-controlled amount of water into a stream of nitrogen. The mixture of water and nitrogen is then passed through a heater to produce steam. The steam is then injected into the preheat zone of the furnace and directed toward the entrance of the furnace. With this system the moisture level in the atmosphere (preheat zone) can be adjusted to produce any dew point the operator desires. However, the amount of moisture distributed by the bubbler is dependent solely on the water temperature. Since the oxidant (H2O) attacks the carbonaceous vapors as they are emitted from the metal compact, the furnace will remain relatively free of high-carbon stalactites and stalagmites that attack the furnace components. It has also proven to reduce stack particulate material by as much as 74%, as established from stack analyses conducted by an independent company. It has also reduced, and in some cases eliminated, both smoke and odor created by the stack emissions. The removal of these vapors from the stack has also been reported to eliminate stack
*Consultant, Furnace & Atmosphere Service Technology, Inc. (F.A.S.T., Inc.), P.O. Box 43, Big Run, Pennsylvania 15715-0043, USA; Phone: 814-427-2228; E-mail:
[email protected]
Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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CONSULTANTS’ CORNER
TABLE I. ANALYSIS OF STACK EMISSIONS* Furnace
CO2
CO + N2
Particulates g/ft.3 (g/m3)
A B
0.00 0.11
83.00 83.88
0.0194 (0.685) 0.0677 (2.39)
*Average figures for three 1 h tests on both furnaces
fires in one manufacturing facility. Two similar continuous-belt furnaces operating side by side were used for the stack analysis tests. One furnace had the moisture-control system operating for one year and the other furnace never contained a lubricant removal system. Identical parts and work load were used in each furnace during the stack testing. There were three 1 h tests conducted on each furnace. Sample results from these tests, which include CO2 and CO + N2 (stack gas % emissions), as well as total particulate concentrations for each furnace, are shown in Table I. Furnace (A) has the moisture-control system and furnace (B) is without a delubing system. Figure 1 shows a comparison of particulate material gathered during stack emission tests with the moisture-control system in operation and with the moisture-control system turned off. The exhaust system in this customer’s facility had three continuous-belt furnaces connected to one main stack. The moisture system was in operation on one furnace, whereas the other two furnaces did not contain any lubricant removal system. All three furnaces were operating under production conditions but not using identical parts or loading parameters. The moisture system was operating for approximately 24 h before the stack emissions test was performed. The particulate material from this test is shown in Figure 1(a). The system was then turned off for 1 h and additional tests run. The results of these tests are shown in Figure 1(b). Even though the unit was only on one furnace the reduction in particulate material is
(a)
(b)
Figure 1. (a) Filter with moisture-control system, (b) filter without moisture-control system
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Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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CONSULTANTS’ CORNER
significant. By looking at the difference between the two filters, it is readily seen that while operating on one furnace with one emissions stack, the addition of moisture could easily produce a 74% reduction in particulate emissions. It should also be noted that all the testing was conducted in production facilities, not in a laboratory. There are testimonials that show a significant increase in belt life (25%) and high-heat muffle life (33%) which are the direct result of maintaining a soot-free furnace environment. More information and testimonials are available at www.fast-incpmt.net. What is the current state of the art with respect to the degree of precision in carbon control during the sintering of steels? With regard to atmosphere control, what technologies need to be developed in order to improve carbon-control precision? I have worked for an industrial gas supplier serving the PM industry for more than 18 years, as well as with PM suppliers for over 20 years. I can assure you that atmosphere technology for precise carbon control is already in existence. What the industrial gas industry has to understand is that there are more factors that affect carbon control than just the sintering atmosphere. For instance, the company providing the powder blend has to know the oxygen content of the base metal as well as any alloying elements added to the powder to complete the mix. Once the oxygen content of the mix is established, the graphite level should be adjusted so that the final carbon level in the PM part is obtained after sintering. This is critical since the oxides in the powder will be reduced by the graphite in the premix and not by the atmosphere, as some might think. Another critical point is to have a tight, leak-free furnace that includes leak-free gas-line piping, and flanges, as well as to control air ingression through the entrance or exit to the furnace. Any oxygen entering the furnace will result in decarburization of the sintered compact and poor carbon control. All these factors must be taken into consideration to obtain precise carbon control. The best atmosphere for carbon control is a mixture of two pure cryogenic gases: nitrogen and hydrogen. Both gases should have an oxygen content ~ 2 to 5 ppm with an oxygen level <10 ppm in the high-heat zone of the furnace. The dew point in the high-heat zone should be between -40°C and 51°C (between -40°F and -60°F) with an atmosphere
Q A
Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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CONSULTANTS’ CORNER
composition approximately 95 v/o nitrogen/5 v/o hydrogen. The atmosphere distribution is shown in Figure 2. This system will require a low-flow nitrogen curtain, as well as layers of thin cloth curtains (fiber fax) at the exit end of the furnace. With a dew point in the -40°C to -51°C (-40°F to -60°F) range the system will also require an oxidant to be added to the preheat zone of the furnace to remove the carbonaceous vapors emitted during lubricant removal from the powder compact. The oxidant can be in the form of air or water but research has proven the latter to be the more stable of the two. The recommended injection points are shown in Figure 3. I have installed these systems in many PM parts manufacturing facilities and, from my experience, the best atmosphere system for precise carbon control is pure nitrogen/hydrogen gas with a hydration system that will raise the preheat-zone dew point to between 2°C and 7°C (between 35°F and 45°F). This system will not only give precise carbon control but will also remove the carbonaceous vapors emitted during lubricant burn-off by turning them into gaseous compounds of CO, CO2, and some hydro-
Figure 2. Furnace-atmosphere distribution—schematic
Figure 3. Furnace-atmosphere-injection points—schematic
10
carbons. This system has been proven to retain good carbon control within the sintered compact, as well as extend the life of the furnace components, such as the belt and muffle, by eliminating build-up of the high-carbon vapors in the preheat and high-heat muffles.
Q A
What would be the implications if the precision in carbon control was as low as ±0.01 w/o or at least ±0.05 w/o carbon? The question does not disclose the composition of the material being sintered, the application(s) of the finished parts, and/or what is the ultimate goal. My answer assumes different scenarios that I hope will answer the question. If the material is austenitic stainless steel or pure iron and is being sintered for magnetic applications, the presence of any carbon level could be significant in lowering the magnetic properties. It would also depend on whether or not this was a high-quality magnetic part. If it is a low-quality magnetic application using iron, it probably would not show much of a difference. With austenitic stainless steel and a highquality application it could be significant. This is why it is always recommended that austenitic stainless steel parts be sintered in a clean (carbon-free) furnace using a pure hydrogen atmosphere. Another scenario would be sintering a carbon-bearing part. When sintering iron without carbon additions, I have frequently encountered the formation of small, low-carbon grain-boundary carbides scattered throughout the cross section of the sintered part. This condition has always been attributed to slow cooling during the cooling cycle but I could not honestly say for sure that it was what caused these precipitates to migrate to the grain boundaries. Although there was never a problem with transverse rupture strength or hardness, it may have a negative effect on the impact properties of the finished part. There is also the question of atmosphere composition. If this, as well as the material composition and application were made available, I may be able to offer more help, not only answering the question, but in providing guidance on solving the problem if indeed there is one. I would certainly be open to discussing this with the person submitting the question if they care to contact me at (814) 427-2228 or
[email protected]. 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] ijpm Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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ANNUAL TECHNOLOGY REVIEW
TECHNOLOGY INVESTMENTS KEY TO PM’S FUTURE Peter K. Johnson*
While the PM industry may be struggling because of the sagging economy and plunging vehicle production in North America, companies are still investing in new technology. When the marketplace returns to health, these investments in the development of new metal powders, equipment, and parts fabrication will reap rewards.
METAL POWDER DEVELOPMENTS Hoeganaes Corporation, Cinnaminson, New Jersey, has developed several new powder products, reports K.S. (Sim) Narasimhan, vice president and chief technology officer. Ancorsteel 30 HP offers PM parts makers a way of reducing costs by replacing higher-molybdenum-containing prealloys with a lower-molybdenum-containing grade, especially in the quenchand-tempered condition. A lower-cost sinter-hardening grade, Ancorsteel 721 SH, is now available for sinter hardening or heat treating. Densities approaching 7.55 g/cm3 by single pressing and sintering at 1,200°C (2,190°F) can be achieved with Ancordense 450, an improvement over Ancormax 200. Ancorsteel AMH is a low-apparent-density atomized powder that can replace sponge iron powder for PM parts. Examples of the compaction response of these powders are illustrated in Figure 1. Additional new powders include Ancorlam for electronic applications, and nickel-free steel powder. Ian Howe, director of application and product development, North American Höganäs, Inc. (NAH), Hollsopple, Pennsylvania, says his company is focusing on cost-effective alloys, improved machinability additives, enhanced bonded mixes, and a warm-die-compaction lubricant system. Some cost-effective NAH alloys include lean-chromium-containing materials such as Astaloy CrM and Astaloy CrL, as well as new materials Astaloy LH and D.L.H. In addition, the company introduced CMN, a lean
Figure 1. Compaction response of new ferrous powders. Base mix FL 4400 (0.35 w/o graphite) *Contributing Editor, International Journal of Powder Metallurgy, APMI International, 105 College Road East, Princeton, New Jersey 08501-6692, USA; E-mail:
[email protected].
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TECHNOLOGY INVESTMENTS KEY TO PM’S FUTURE
alloy for PM parts requiring heat treatment to develop properties. In 2008 NAH developed machining additives to greatly improve machinability relative to MnS for pearlitic structures (HRB <75), and for martensitic structures (>HRC 35). Focusing on increasing green densities, the company launched Intralube E, a new lubricant for warm-die compaction for applications in the green-density range of 7.20–7.35 g/cm3. A new prototyping center incorporates CAD/FEA and compacting presses including an 800 mt (881 st) hydraulic press. Sintering and heat treating, and CNC machining, are also available for producing prototype parts. Rio Tinto’s powder plant (formerly Quebec Metal Powders Ltd.), Sorel-Tracy, Québec, Canada, is concentrating R&D on lower-cost alternative materials to replace diffusion-bonded powders and high-nickel-containing grades, reports Francois Chagnon, principal scientist, technology center. This will be achieved through new organic bonded powders and by redesigning alloys and/or mix formulations with elements that are less sensitive to price volatility. Reaching a density of 7.5 g/cm3 with single pressing and sintering is a very important goal for the PM industry. This will be achieved by developing high-compressibility powders and redesigning compacting presses and tooling to attain pressures up to 1,100 MPa (80 tsi). Carpenter Powder Products Inc., Bridgeville, Pennsylvania, is developing improved powder-making and processing procedures to provide cleaner prealloyed iron, nickel, and cobalt-based powders for high-performance parts, says Louis W. Lherbier, director of technical market development. The company has also increased its capacity to produce metal injection molding (MIM) powders by acquiring Ultrafine Metals Powders. It sees increased demand for stainless steel, nickel, and cobalt-based alloys for hot isostatic pressing (HIPing). Among future trends, Carpenter sees the increasing use of laser technology to make free-form near-net-shape parts from specialty alloy powders. Edul Daver, president of ACuPowder International, LLC, Union, New Jersey, reports on nonferrous powder trends. Extra-fine spherical copper powders are finding applications in electronics. Specialty bronze powders, which have high strength and high hardness, are being used to make bronze structural parts. Magnesium Elektron Powders, Manchester, New Jersey, will complete an expansion of its Hart Metals facility in Tamaqua, Pennsylvania, says Deepak Madan, vice president of technology & new
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product development. The investment will increase the company’s total magnesium atomization capacity by more than 150 percent, Figure 2. EQUIPMENT INNOVATIONS Uwe Haupt, sales representative at ARBURG GmbH + Co KG, Lossburg, Germany, cites the Allrounder A series injection molding machines with electrically driven motion axes for achieving up to 50 percent in energy savings in producing MIM parts, Figure 3. Dorst will continue investing in new compacting equipment that improves the performance and productivity of PM parts manufacturing, says Greg Wallis, CEO of Dorst America, Inc., Bethlehem, Pennsylvania. Capable and reliable equipment supported both locally and remotely will be in demand, as PM parts makers will not be able to meet higherperformance requirements demanded by customers with outdated equipment. Ingo Cremer, Cremer Thermoprozessanlagen GmbH, Düren, Germany, cites a growing demand for sintering at up to 1,200°C (2,190°F). The company has designed a range of conveyor furnaces that can use conventional belts at up to 1,200°C (2,190°F), which can generate significant savings. Elnik Systems, Div. of PVA MIMtech, LLC, Cedar Grove, New Jersey, offers a new furnace, Figure 4,
Figure 2. Expanded magnesium atomizing plant
Figure 3. Allrounder A series MIM presses save energy
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TECHNOLOGY INVESTMENTS KEY TO PM’S FUTURE
for debinding the secondary binder in sintering MIM parts, reports Claus J. Joens, president. Debinding is achieved using a plasma at 1 to 7 mbar pressure and sintering with a laminar gas flow at 400 mbar. Impco, Inc., East Providence, Rhode Island, has improved PM impregnation technology substantially through a combination of material and process developments, reports Terry Chwalk, vice president. These improvements include improved pore filling, which results in higher-quality plating and moreconsistent machining of PM parts. Minox-Elcan Industries, Inc., Mamaroneck, New York, offers Kroosh multi-frequency vibratory screeners, Figure 5, capable of screening particulate materials with efficiencies of more than 95 percent, reports Bob Grotto, president. These machines have increased throughputs by a factor of 15 to 20 over conventional screeners, the company reports.
Figure 4. Improved debinding furnace
Figure 5. Multi-frequency vibratory screeners increase productivity
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TECHNOLOGY INVESTMENTS KEY TO PM’S FUTURE
Dave Quilter, manager of refractory systems, Pyrotek, Inc., Canastota, New York, reports that refractory castables, shape-casting techniques, mold making, and curing processing improvements have significantly impacted the cost effectiveness of refractory shapes. Blending traditional brick and cast-in-place techniques with precast technology has provided an enhanced finely tuned rebuild package. Rebuilding a high-temperature pusher furnace with all precast cured refractory shapes can cut 75 percent of the labor hours needed to rebuild a traditional brick design. UTRON Kinetics, UTRON, Inc., Manassas, Virginia, has developed the combustion-driven compaction (CDC) process for making high-density PM and ceramic parts with improved properties, says Dennis Massey, CEO. The company has successfully made cost-effective parts for defense applications. The CDC process uses the energy generated from the controlled combustion of a refined-gasand-air mixture to power press motions, Figure 6. The energy source can generate extreme forces >91 × 104 mt (106 st) and more moderate forces of 91 to 2,723 mt (100 to 3,000 st). When combined with powder and lubricant developments, the process can attain full or near-full density with a single stroke, the company claims. This is achieved through incorporating press stroke speed, impact control, and tooling preloading not attainable with current mechanical and hydraulic compacting presses. Tooling can withstand a 100-to-200 percent increase in applied compaction pressure without failure. PM PARTS TRENDS C. L. Chu, general manager, Porite Taiwan Co.,
Figure 6. CDC press concept
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Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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TECHNOLOGY INVESTMENTS KEY TO PM’S FUTURE
Figure 7. Complex high-speed tool steel applications
Figure 8. Tungsten carbide MIM part
Ltd., Taiwan, reports on new projects developing high-nickel and high-chromium alloys for electronics and energy applications. New automotive VVT/VCT and high-pressure pumps have passed validation tests and Porite intends to supply more automotive parts to the U.S. and Europe. In addition, the company is leveraging its 13 years of warm compaction experience for higher -density compaction. PSM Industries, Los Angeles, California, is using PM techniques to create new engineered materials for applications requiring extreme wear resistance and toughness, reports Craig Paullin, president. Capturing net or near-net shape lowers manufac-
turing costs. The company is working with powder suppliers to provide more compressible materials in alloy-rich compositions. Its PM Krupp division uses a proprietary sintering process to make fully dense high-speed tool steels in complex shapes with hardness up to HRC 67, Figure 7. The Ferro-Tic division supplies steel-bonded titanium carbide which PSM says is 50 percent harder than tungsten carbide. PSM’s Yillik Precision Carbide division has recently developed a material containing tungsten carbide and titanium carbide sub-micron particles in a cobalt/nickel base. Yillik and the PolyAlloys division are applying MIM technology to make complex fully dense tungsten carbide parts, Figure 8. ijpm
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Experts from leading PM and particulate materials companies will answer questions about the latest trends in powders, production equipment, process technologies, testing, and QC equipment and products. The exhibition features process equipment and provides a valuable opportunity to meet with current or new suppliers. Receive immediate help with production and materials questions. Arrange appointments now with the companies you want to visit and arrive with your list of technical issues for one-on-one discussions. Take advantage of this valuable opportunity to gain new information from major suppliers and network with industry technical leaders.
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Exhibitor Showcase
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Exhibitors Listing as of April 3, 2009 ABBOTT FURNACE COMPANY St. Marys, PA Abbott specializes in continuous furnaces for sintering, steam treating, quenching, annealing, tempering, and brazing. Silicon Carbide muffles, a Quality Delube Processor, and VariCool are all popular options. Pusher furnaces and ceramic belt models are suitable for higher-temperature applications. Spare parts, fabrications, repairs, and calibrations are offered. ISO/IEC 17025 Accredited. AC COMPACTING LLC North Brunswick, NJ AC Compacting LLC carries a line of small parts sorters from CI electronics. The units will sort pieces by weight up to 10 grams with accuracies down to 0.5 mg and speeds to 85 pieces per minute. AC also carries from 10 ton to 60 ton rotary presses and an instrumented compaction research press and press simulator. ACUPOWDER INTERNATIONAL, LLC. Union, NJ/Greenback TN ACuPowder, with plant in NJ & TN, is a major U.S. producer of metal powders. Products include: Antimony, Bismuth, Brass, Bronze, Bronze Premixes, Copper, Copper Alloys, Copper Oxide, Copper Premixes, Diluted Bronze Premixes, Graphite, High Strength Bronze, Cu Infiltrant , Manganese, MnS+, Nickel, Phos Copper, Silicon, Silver, Tin, Tin Alloys and PM Lubricants. New products include powders for MIM, Thermal Management, "Green" Bullets, Lead Free Solders, Plastic Fillers, Cold Casting and most recently Ultra Fine/Ultra Pure Copper Powders for the electronics industry and ULTRA INFILTRANT the wrought/wire infiltration solution. AMERICAN CHEMET CORPORATION East Helena, MT & Deerfield, IL American Chemet, est. in 1946, manufactures copper powders, dispersion strengthened Cu, and copper and zinc oxides. Chemet’s oxide reduction process allows a high degree of control over particle size and shape in powders ranging from molding grade (150 mesh) to 12 micron median size. AMETEK, INC. Eighty Four, PA Ametek and Reading Alloys produce specialty powders primarily for aerospace, automotive, electronic, hardware and med-
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ical industries. Products include Ultra 300 and 400 series stainless powders for PM, MIM, and filter markets, nickel-base thermal spray powders and specialty alloys such as titanium CP and Ti 6/4 powders. Ametek/ Reading Alloys is a world leader in research, development and manufacture of high-grade aerospace master alloys, specialty metals, and coatings materials. APMI INTERNATIONAL Princeton, NJ Celebrating its 50th anniversary (1959–2009) APMI International is the professional society for individuals involved in powder metallurgy and particulate materials. Members include metallurgists, engineers, teachers, students and business people. Some of the many benefits include: International Journal of Powder Metallurgy, Who's Who in PM membership directory, full access to PM NEWSBYTES and monthly PM Industry NewsLine. Stop by our booth and learn how APMI can be your professional resource. ARBURG GMBH + CO KG Lossburg, Germany & Newington, CT Joining PowderMet2009 in Las Vegas, ARBURG offers intensive individual consulting on site when it comes to the PIM sector. The ARBURG machines for processing metal and ceramic powders using the PIM process are based on the current ALLROUNDER machine series. In Las Vegas, the ARBURG PIM specialists look forward to seeing you at booth 430. ASBURY-SOUTHWESTERN GRAPHITE Asbury, NJ For over 100 years the worldwide leader in graphites and carbons for the Powder Metal industry. Our complete line of natural and synthetic graphites for conventional PM applications, specialty materials for forging, bearing, and hard metal applications will be presented. Asbury also supplies a complete line of graphite sintering trays and graphite lubricants to the industry. Metal sulphides and metal alloy powders are also available from Asbury. BASF CORPORATION Evans City, PA Catamold® is BASF's ready-to-mold feedstock for MIM and CIM, available in a wide range of options for standardized steel, stainless steel and ceramics. BASF can
help your company get started in PIM with samples, training and technical support based on our experience as the leading feedstock supplier worldwide. Contact BASF: 724-538-1363, or
[email protected]. BODYCOTE-HIP Andover, MA Bodycote–HIP Powder Metal (PM) technology has the capability to provide a unique combination of properties for demanding applications. Unlike traditional press and sinter PM technology the HIP process is without die friction forces that limit product size and density. We routinely make 100% dense parts as large as 25,000 lbs. in weight. BRONSON & BRATTON, INC. Burr Ridge, IL Bronson & Bratton, Inc. has been in the Tool & Die business since 1948, and has been building PM Tooling since 1970. We have the Design (CAD), Manufacturing (CAM), and the experience to design and build the Tooling/Adapters required to fit your existing Compacting/Sizing Presses. We are ISO 9001:2000 certified. C.I. HAYES INC. A SUBSIDIARY OF GASBARRE PRODUCTS, INC. Cranston, RI Manufacturers of custom-designed sintering and heat-treating furnaces with temperatures to 3,000ºF. Hayes' atmosphere furnace designs include, belt, pusher, walking beam. Vacuum furnaces in batch or continuous and feature isolated heating and quenching chambers. Continuous vacuum carburizing. Endothermic, exothermic, and DA generators. Full line of replacement parts. CARPENTER POWDER PRODUCTS INC. Bridgeville, PA Provides prealloyed powders that are tailored to meet customer requirements for thermal surfacing processes, metal injection molding, near net shape hot consolidation technologies, and mill form products (billet, bar, wire, plate, sheet, and strip). Our manufacturing versatility and technical knowledge enable us to provide you with consistent high quality products.
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CENTER FOR POWDER METALLURGY TECHNOLOGY (CPMT) Princeton, NJ The Center for Powder Metallurgy Technology (CPMT) is a not-for-profit foundation established by members from the PM community. CPMT funds cooperative technology programs focusing on R&D that bring together the corporate, academic, and research organizations to advance PM technology. Center members benefit from periodic research reports and guide the direction of research activities. Other activities include scholarships and grants provided to industry students. CENTORR/VACUUM INDUSTRIES, INC. Nashua, NH High-performance Metal Injection Molding Furnaces for alloy steels, stainless steel, tool steel, hardmetals and ceramics. Laboratory to production size. Temperatures to 2,300ºC in vacuum, inert, or hydrogen gas from 10–750 torr. Graphite or refractory metal hot zones with proprietary Sweepgas™ binder removal systems for injection molded parts.
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CINCINNATI INCORPORATED Cincinnati, OH CINCINNATI INCORPORATED manufactures PM Compacting and Restrike Presses. All presses are backed by extensive support services including a ReManufacturing Facility for Reconditioning and Up-grading existing equipment to ensure maximum performance and productivity. Video and photographs will be shown highlighting various products and services available. CM FURNACES, INC. Bloomfield, NJ Fully automated high-temperature continuous pusher furnaces for both traditional powder metal and metal injection molding with inline debinding. These furnaces operate in a hydrogen/nitrogen atmosphere up to 3,100°F with extremely low dew points. Also being displayed will be our line of high-temperature hydrogen batch furnaces. DIAPAC Houston, TX Decades of extensive industry experience
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and a proven commitment to superior service makes DiaPac an internationally recognized leader in high-quality, solutionsdriven products. From Research and Development to expertise in restoring value from used materials, as well as offering ready-to-use diamond and tungsten carbide powders, DiaPac ensures you get the job done right. DORST AMERICA, INC. Bethlehem, PA Continuous innovation, leading technology and outstanding customer service have made Dorst the market leader for CNC hydraulic presses in the PM and related industries. Our all-encompassing approach, ranging from products to technological support and after-sales service, enables customers to optimize the most demanding jobs and perform with exceptional capability and productivity. ECKA GRANULES OF AMERICA LLC Orangeburg, SC ECKA Granules is the leading manufacturer of non-ferrous metal powders. The prod-
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Exhibitors uct range includes Aluminium, Magnesium, Copper, Calcium, Tin, Lead, Zinc, Silicon and their alloys as well as ready-to-press premixes. Production techniques include milling and grinding, electro-deposition, air, water and gas atomization, granulation and melting for recycling. ELMCO ENGINEERING INC. Indianapolis, IN ELMCO Engineering Inc. is a leading manufacturer of new and rebuilt PM equipment of all makes and sizes. We service all makes of presses, and have an extensive parts inventory. We are North American Representatives for Yoshizuka presses. ELMCO also offers custom engineering for special applications. Visit us in Booth #301. ELNIK SYSTEMS (Division of PVA MIMtech, LLC.) Cedar Grove, NJ Elnik will introduce a new furnace called "PLASMIM" incorporating a plasma source providing clean debinding and reduced processing time. The “ONE SOURCE MIM” equipment is also featured. Let DSH
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Technologies, our affiliate, prove the feasibility of the process parameters for your MIM parts before you invest in expensive capital equipment. ERASTEEL—DIVISION OF ERAMET Paris, France Erasteel: Your flexible powder source With atomization units based in Sweden and 40 years of experience, Erasteel is the world leading producer of high-quality gasatomized metal powders for tooling and components. Alloy types include highspeed steels, tool steels, stainless steels and other alloys. Contact us at
[email protected] EROWA TECHNOLOGY, INC. Arlington Heights, IL “Pulverizing Set Up Times”—EROWA Technology (Arlington Heights, IL) is the world leader in palletization and automation solutions for the manufacturing industry. EROWA’S PM Tooling System palletizes the punches as well as the die/mold; enabling press resetting in less than 3 minutes. The 0.002 mm repeatabili-
ty eliminates punch damage during press set-up. See us at the PowderMet2009 show in booth #324! EVANS ANALYTICAL GROUP LLC Sunnyvale, CA Evans Analytical Group (EAG), is a global leader of Materials Characterization and microelectronic Release to Production testing services, including Electron Microscopy, Failure Analysis, Burn-In, FIB, ESD and ATE services. EAG provides fast turnaround time, superior data quality and excellent results. EAG has over 20 locations in Asia, Europe and the U.S. GASBARRE PRESS DIVISION– GASBARRE PRODUCTS, INC. DuBois, PA Designers and manufacturers of singlelevel and multi-level Mechanical and CNC Hydraulic Presses—5 to 1,200 Tons for compacting and sizing of structural P/M parts. Removable die set presses are available. Hydropulsor High Velocity Compacting Presses to 2,000 Tons. TOPS Powder Heating Systems, Die Wall
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Lubrication Units, Fluidized Filler Shoes, Parts Automation, and Powder Handling Systems. Extensive rebuild services. GLOBAL TUNGSTEN & POWDERS CORPORATION Towanda, PA Global Tungsten & Powders Corp. (GTP) located in Towanda, Pennsylvania, and GTP BRUNTAL, located in Bruntal, Czech Republic, combine to create a world leader in the production of tungsten, tungsten carbide, molybdenum, cobalt, and tantalum powder products. GTP features its tungsten carbide powders, POWDER PERFECT™ thermal spray powders, and high green strength tungsten powders for a number of applications and manufacturing processes including MIM. We are also a major producer of tungsten and molybdenum ingots, billets, plate, sheet and wire. We service the hard materials, energy, automotive, defense, electronics, medical, and aerospace markets. H.C. STARCK, INC. Newton, MA H.C. Starck ranks among the world’s leading manufacturers of refractory metals such as tungsten, molybdenum, tantalum, niobium, and rhenium; electronic chemicals and ceramic powders. H.C. Starck continues to strive to further strengthen its ability to bring material solutions to the market. Please visit our booth for more details. HOEGANAES CORPORATION Cinnaminson, NJ Hoeganaes Corporation, world leader in ferrous powder production, has been a driving force within the PM industry’s growth for over 50 years. It has seven manufacturing facilities in the United States and Europe to meet customers’ needs worldwide. It holds these certifications: ISO 14001, ISO/TS 16949, and QS 9000. INCO SPECIAL PRODUCTS See Vale Inco Americas INDUSTRIAL HEATING MAGAZINE Pittsburgh, PA The metal-powder industry's only fully audited monthly trade journal for metalpowder engineers, part designers, applications engineers, equipment manufacturers, powder producers and suppliers.
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KITTYHAWK PRODUCTS Garden Grove, CA Kittyhawk Products—qualified experts in the field of Hot Isostatic Processing—HIP is a process of unique benefit in solving complex design problems while increasing the strength of properties. Through our sister company, Synertech P/M, Inc., we offer unmatched net-shape capabilities with powder metal parts design and manufacture. Kittyhawk Products holds ISO9001 and AS9100 certification. KOMAGE GELLNER MASCHINENFABRIK KG Kell am See Germany KOMAGE manufactures machinery, tooling and handling system for the metal powder and ceramic industry. There are mechanical presses in a range of 5–50 metric tons, hydraulic presses from 20–1,200 metric tons with fixed or floating die table. and high speed hybrid presses from 20–250 metric tons. All presses are offered with multi plates adapter and closed loop controlled axes. LABORATORY TESTING INC. Hatfield, PA Laboratory Testing Inc. is an independent, accredited laboratory performing materials testing, nondestructive testing, failure analysis, specimen machining, dimensional inspection and NIST-traceable calibration services at one convenient location. LTI tests and analyzes metals, powdered metals, ores, ferroalloys, composites, ceramics, aerospace materials and nuclear materials. Certified reports include detailed results. LASCO ENGINEERING SERVICES, L.L.C. Detroit, MI LASCO Engineering Services is the US arm of LASCO Umformtechnik in Coburg, Germany. LASCO is a 135-year-old company producing metal forming machines for export around the world. LASCO Engineering Services will be presenting their new line of powder metal compacting presses along with coining and powder metal forging equipment. LINDE AG, GAS DIVISION HEADQUARTERS Pullach, Germany With its innovative solutions, Linde Gas plays a pioneering role in the global gases
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market. We offer more than just high-quality gases to customers in the PM industry. No matter which processes comprise your daily business—from the automation, hot isostatic pressing or actual sintering—our experts can always provide you with the right solution. MAGNESIUM ELEKTRON POWDERS Manchester, NJ Magnesium Elektron Powders is a producer of magnesium powders and specialty niche alloy powders. It has three facilities in North America, producing various types of powders. The company manufactures a wide range of atomized and ground powders to military specification. The company also manufactures powders for steel desulphurization, chemical synthesis, welding applications, powder metallurgy, specialty pyrotechnics, and flameless ration-heater pads. MASRIA FOR METALLURGICAL POWDER INDUSTRY (MPI) Cairo, Egypt MPI Ltd is Specialist manufacturers of water-atomized metal powders. The product range includes Atomized copper, Bronze, Brass, Aluminum, Zinc, Lead, and Tin powder. The mission of MPI is to become a world-class company which will allow satisfying our customers. METAL POWDER INDUSTRIES FEDERATION Princeton, NJ Stop by to learn about membership benefits, programs, association committee activities, and any other topic of interest to you. If you have comments or ideas regarding the Federation and its services, let us know when visiting us at our booth. If your company is not a member of MPIF, you can discuss membership opportunities and benefits with someone from headquarters staff. METAL POWDER REPORT Oxford, UK England Metal Powder Report has charted the expansion of the powder metallurgy business over the past 50 years. It is the premier international and independent magazine for the powder metallurgy industry reporting on technical trends in the manufacture, research and use of metal powders. Sample copies will be available at
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Exhibitors the booth or log on to http://www.metalpowder.net/sample.asp MINOX-ELCAN, INC. Mamaroneck, NY Minox-Elcan specializes in advanced screening systems used in the powdered metals industry. Kroosh technology provides high energy/high amplitude multifrequency energy directly to the screen surface. The value of our performance improvement separating powdered metals is significant. Our testing and toll processing facilities annually process a variety of powdered metals down to 10 micron screen size. We help develop advanced high value products as well as maximize current product profitability by improving screen efficiency. NORTH AMERICAN HÖGANÄS, INC. Hollsopple, PA North American Höganäs, Inc., offers metal powder solutions that create new business and profitable growth for partners and customers. Metal powder range includes: Plain Iron, Prealloyed Steel, Diffusion Alloyed, Stainless Steel, Tool Steel, Gas Atomized and Electrolytic Iron. Premixed and bonded Starmix materials. ORTON CERAMIC FOUNDATION Westerville, Ohio The Orton Ceramic Foundation has produced devices that measure thermal energy (temperature and time at temperature) for well over 100 years. They have recently introduced a new, easy-to-use product TempTab, to help sintering and heat treating operations benchmark their processes and verify temperature uniformity inside their furnaces, without interrupting production. OSTERWALDER AG Lyss, Switzerland Switzerland OSTERWALDER AG, the leading powder press manufacturer, presents our newest developments in our Hydraulic Powder Presses CA-SP, CA-MP, CA-NC II, UPP as well as the Mechanical Hydraulic Powder Press KPP. These developments bring surpassing savings in set-up time and unrivaled benefits in the overall quality and productivity of your compacts.
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POMETON POWDER Venice, Italy Founded in 1940 and based in Venice, Italy, Pometon supplies its range of ferrous and non-ferrous powders to PM and other industrial clients in over 40 countries worldwide. We produce pure powders such as iron, copper (both electrolytic and atomised), bronze, brass, tin and zinc, and press-ready iron and bronze premixes. POWDER INJECTION MOULDING INTERNATIONAL Shrewsbury, England Powder Injection Moulding International is a quarterly magazine that offers in-depth industry coverage of the MIM, CIM and carbide injection moulding industries. Each issue features industry news, company reports, exclusive commissioned features and leading technical papers. The publisher, Inovar Communications, will also be promoting the forthcoming 14th Edition “International Powder Metallurgy Directory 2010-2011.” PTX-PENTRONIX, INC. A SUBSIDIARY OF GASBARRE PRODUCTS, INC. Lincoln Park, MI Designers and manufacturers of highspeed, mechanical compacting presses, 2 tons to 35 tons. Anvil and opposed-ram designs available. With speeds up to 300 pcs/min, and multiple cavity capabilities, extremely high production and high precision are achieved on PTX presses. PTXPentronix also manufactures automatic, high-speed parts handling and robotic parts-palletizing systems. Distributors for Simac Isostatic Dry Bag Presses. QMP See Rio Tinto, Metal Powders RIO TINTO, METAL POWDERS (QMP) Sorel-Tracy, Canada Rio Tinto, Metal Powders (QMP): registered ISO 9001, ISO 14001, ISO/TS 16949; manufactures a full product line of iron and steel powders including ATOMET standard grades, prealloys, binder treated FLOMET™ mixes, diffusion bonded ATOMET DB, machinable (sulphur-free) grades, sinter-hardening grades, and soft magnetic composite materials for customers worldwide.
RUSSELL FINEX, INC. Pineville, NC Russell vibratory screeners and separators improve particle size control and ensure that your products meet precise specification. The Compact screener is suitable for high-capacity check-screening and grading metal powders. The Vibrasonic deblinding system eliminates mesh blinding and increases screening efficiency, allowing metal powders to be accurately screened down to 20 microns. RYER, INC. Temecula, CA Ryer, Inc., is a Manufacturer, Developer and Supplier of Custom and Standard Feedstocks for the Metal Injection Molding Industry. Ryer manufactures a variety of Standard Feedstocks in addition to our Custom- Formulated Feedstocks to match your current material shrink specifications. For more information visit us on the Web at www.ryerinc.com. SANDVIK OSPREY LTD. (Powder Group) Neath, United Kingdom Specialist manufacturer of Gas Atomized powders with a size range from 1 to 250 microns. Our alloy range, already the largest in the world for MIM applications, also includes Thermal Spray, Rapid Prototyping, HIPping and Brazing powders. Accredited to ISO 9001:2000 and ISO 14001, Osprey is a Sandvik Materials Technology company. SCM METAL PRODUCTS, INC. Research Triangle Park, NC & Suzhou, China SCM Metal Products is a leading manufacturer of metal powders, pastes, flakes and infiltrating and brazing preforms with manufacturing facilities in the U.S. and China. Our metal powders include copper, bronze, brass, infiltration, friction copper, copper oxide, tin and lead. SCM also produces a line of specialty paste products for infiltrating and sinterbrazing PM parts as well as for furnace brazing of steel components. SINTERITE FURNACE DIVISION– GASBARRE PRODUCTS, INC. St. Marys, PA Sinterite designs and manufactures continuous-belt and batch furnaces for sintering, steam-treating, annealing, brazing, and Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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heat-treating applications. High-Temperature Pusher Furnaces (over 350 manufactured) in several designs for iron and stainless steel parts (to 3,000°F). VersaCool inline cooling systems for sinter-hardening; Accelerated De-lubrication Systems (ADS). Alloy or Ceramic muffles available. Replacement muffles, powder-handling equipment, and fabrication products. SMS MEER– A COMPANY OF THE SMS GROUP Moenchengladbach, Germany & Pittsburgh, PA In addition to equipment for pipe and long product rolling mills, forging presses and the NF metal industries, we design and build hydraulic and mechanical powder presses of which we have already sold more than 1,800. For over 50 years, we have been the competent partner for the metal powder, ceramics and tungsten carbide industry. SOLAR ATMOSPHERES INC. Souderton, PA Solar Atmospheres, vacuum heat treating specialists, provides vacuum sintering, degassing, low-temperature drying, hightemperature purification, high-temperature compound formation, carburizing, and nitriding for the powder metal industry. Capabilities include over 40 vacuum furnaces, from laboratory, for cycle development, up to 36 feet long for production services. AS9100:2004 SUPERIOR GRAPHITE Chicago, IL Superior Graphite specializes in thermal purification, advanced sizing, blending and
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coating technologies, providing value added graphite and carbon-based solutions globally. Combining 90 years of experience and advanced technologies into every facet of the organization, a wide range of markets are served such as; advanced ceramics, agriculture, battery/fuel cells, ceramic armor, carbon parts, ferrous/nonferrous metallurgy, friction management, hot metal forming, polymer/composites, powder metals, lubricants and performance drilling additives. North and South America contact: CustomerServiceUSA@SuperiorGraphite. com or in Europe/Africa/Asia/Australia contact: CustomerServiceEU@Superior Graphite.com. THE ALLOY ENGINEERING COMPANY Berea, OH The Alloy Engineering Co. has been recognized for its expertise in the design and fabrication of products utilized in high-temperature and corrosive environments since 1943. We have also acquired two major high-temperature fabricators—TEI/Rolock and Walmil—that have strengthened our engineering and fabricated-product offering to the powder metal industry. In addition to a variety of products including fabricated muffles and high-temperature fans, Alloy Engineering offers extensive alloy-materials expertise, design know-how, and fabrication capabilities. THE MODAL SHOP, INC. Cincinnati, OH The Modal Shop’s NDT-RAM systems are designed to help you deliver fully inspected parts, on time, giving you and your customer confidence in the quality of your
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parts. NDT-RAM systems detect cracks, voids, variances in dimension, geometry, weight, density, porosity, bonding, brazing, and machine process. A free parts test will determine if your part is a good candidate for NDT-RAM. Contact TMS at 513-3519919 or www.ndt-ram.com. THE WIRE MESH BELT COMPANY Brampton, Ontario, Canada Manufacturing top-quality mesh belting for use in high-temperature furnaces for 40 years. Specializing in custom designed Double Balanced & Balanced Flat Spiral (BFS) belting used in sintering, brazing and annealing operations in temperatures to 2,300°F. Our flexibility and service will eliminate costly downtime with delivery in days. THERMAL TECHNOLOGY LLC Santa Rosa, CA Thermal Technology LLC is a high-temperature equipment manufacturing company whose broad line of equipment includes: spark plasma sintering (SPS), crystal growing systems, arc furnaces, and high-temperature vacuum and controlled-atmosphere furnaces. This incredible product line with its associated engineering and applications skills make Thermal Technology LLC one of your best resources for thermal processing. THINK “SOLUTIONS!” TIMCAL GROUP Westlake, OH Timcal Graphite and Carbon, a member of Imerys a global leader in adding value to minerals, produces a full line of graphites designed specifically for the PM industry: including high-performance primary syn-
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Exhibitors thetics and custom-sized natural flakes using raw material sourced from our 100% owned North American mining operation. ULTRA INFILTRANT Carmel, IN The patented Ultra Infiltrant Wrought/Wire Copper Infiltration Technology has raised the bar. The surface erosion, adherent residue, high cost and hassle factor associated with pressed copper powder infiltrants are things of the past. UI delivers superior mechanical and metallurgical results that far exceed the MPIF Standard 35. UI also adds solid performance to your bottom line—Custom made preform parts manufactured to your exacting specifications are delivered ready for assembly with your green parts—All the non-value-added process is removed. Now that’s a Solid Line of Thinking!!! Come visit Ultra Infiltrant in booth #412. UNION PROCESS, INC. Akron, OH Attritor mills for fine grinding, flaking or mechanical alloying of metal powders are
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displayed. Attritors are ruggedly constructed and designed with interchangeable components to meet a variety of processing requirements, wet or dry. Sizes range from research to production-sized mills. Systems for grinding under inert gases or cryogenic grinding and metal-free grinding are offered. UNITED STATES METAL POWDERS, INCORPORATED Flemington, NJ Major global producer of non-ferrous metal powders and flakes, including aluminum, aluminum premixes, copper, copper alloys, bronze premixes, nickel silver, infiltrants, and tin. Subsidiaries are AMPAL, Inc., Palmerton, PA; Makin Metal Powders, Ltd., United Kingdom; and Poudres Hermillon, France. UTRON KINETICS, LLC. Manassas, VA UTRON Kinetics, LLC., is an award winning R&D company with an exemplary history of providing advanced technological innovations to NASA, DOE, NSF, the Army,
the Navy, and other organizations. We have pioneered the development and application of Combustion Driven Compaction and developed a set of globally unique technologies that are providing revolutionary improvements in materials and materials processing. VALE INCO AMERICAS Wyckoff, NJ Vale Inco Limited produces nickel, copper, cobalt and precious metals. Carbonyl refineries in North America and Europe make nickel powders of various sizes and shapes to the ISO9002 standard. Nickel powder products supplied to the PM and MIM industries include: T123 PM, T110 D, T255, T287, T210 and Novamet 4SP-10. ZIRCAR CERAMICS, INC. Florida, NY High Alumina purity porous sintering setters & custom machined sintering fixtures. Furnace insulation, molydisilicide heating elements, alumina-silica papers & blankets. ijpm
Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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RESEARCH & DEVELOPMENT
POWDER INJECTION MOLDING OF METALLIC AND CERAMIC HIP IMPLANTS Jiupeng Song*, Thierry Barriere**, Jean-Claude Gelin***, Baosheng Liu****
INTRODUCTION PIM is a relatively new processing technology for manufacturing small and intricate metallic or ceramic components in large batches.1 The four basic processing steps in PIM are: powder and binder mixing, injection molding, debinding, and sintering. Each stage influences the physical and mechanical properties of the final component. The mixing process prepares the feedstock for injection molding, and optimization involves the particle characteristics, binder formulation, powder-volume fraction, and the mode of mixing. Research has focused on the injection molding stage in which the green parts are shaped and it has been demonstrated that injection pressure, velocity, and temperature are the primary processing parameters in relation to dimensional accuracy.2 Concurrently, numerical simulations have been performed to analyze the PIM process. One method is to utilize commercially available software for the injection molding of plastics or make use of casting simulations, incorporating the rheological properties of the PIM feedstock.3–5 An alternative approach is to develop dedicated software for PIM such as PIMSolver,6 in which the effect of slip between the feedstock and the mold cavity is taken into account. Powder–binder phase separation (phase segregation) can be induced during injection due to flow at high speeds and high pressures. This is a natural phenomenon because of the difference in the density of the metallic or ceramic powders and the organic binder. Inhomogeneity in the green component is the result of segregation and its effect is normally amplified in the subsequent debinding and sintering steps.7,8 Our previous work resulted in a biphasic model for the prediction of segregation in the injection stage.9,10 A new and efficient explicit algorithm was developed by the authors,9,10 and implemented with in-house finite element software (FEAPIM). Debinding removes a majority of the binder in the molded part employing solvent, catalyst, or other techniques. Physical and numerical models have been proposed for thermal and solvent debinding in order to improve their efficiency. 11,12 In the sintering stage, the
Design of the powder injection molding (PIM) process for the fabrication of prototype hip implants is considered. Experimental studies were carried out using metallic (316L stainless steel) and ceramic (alumina) powders, focusing on the effect of primary processing parameters on defect formation and attendant properties. Simulation of the injection and sintering steps was also included. A biphase injection model gave a reliable prediction of segregation defect control. The model for sintering resulted in good agreement with experimental results and observations on density, shrinkage, and strength.
*Post-doctoral researcher, **Associate professor, *** Professor, FEMTO-ST Institute, Department of Applied Mechanics, ENSMM Besançon, 26 Rue de l’Epitaphe, 25030 Besançon, France, E-mail:
[email protected], ****Professor, Southwest Jiaotong University, Department of Applied Mechanics and Engineering, 610031 Chengdu, China
Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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POWDER INJECTION MOLDING OF METALLIC AND CERAMIC HIP IMPLANTS
debound components are subjected to a temperature close to, but below, the melting point of the main constituent. The powder particles are then sintered to achieve their final density by solidstate diffusion. The sintering step13 dictates the final dimensional accuracy and the physical and mechanical properties. In principle, the sintering process for PIM parts does not differ from that utilized for conventional die-pressed compacts. However, shrinkage (10–15 v/o) is much larger in PIM and the sintering kinetics are higher due to the high level of porosity after debinding. Both experimental and modeling investigations have been performed on the sintering of PIM parts.14,15 In the present study, experimental work on PIM to produce a complex hip-implant prototype is presented using two common PIM materials, 316L stainless steel and alumina. The sintering stage is further analyzed, because of its important influence on density, microstructure, mechanical properties, and defects in the final component. The biphasic injection simulation based on our previous work,9,10 has recently been extended to three-dimensional (3D) components of complex shapes and it is applied here to hip implants. Macroscopic finite element simulations of the sintering process have also been performed to predict the final density and mechanical properties of the sintered component. EXPERIMENTAL INVESTIGATION Materials and Procedures Two commercially available PIM feedstocks, in the shape of pellets, were used in the experiments. ADVAMET® 316L stainless steel feedstock (Advanced Metalworking Practices, Inc., U.S.) consists of gas-atomized powder and a thermoplastic binder. The powder content is ~62 v/o with a particle size D80 of 16 µm. ELUTEC® A-99-S alumina feedstock (Zschimmer & Schwarz, Germany) consists of a mixture of 81.5 w/o alumina powders and 18.5 w/o thermoplastic binder. The purity of the alumina powder is 99.8 w/o with a particle size D50 of 0.7 µm. The binder density in the alumina feedstock is 1.2 g/cm3 and its water solubility is ~65 w/o at 20°C. Hip implants are one of the most widely used medical prostheses. In our laboratory, various manufacturing processes have been investigated to manufacture hip implants, including polymer injection molding, PIM, 3D printing and highspeed machining. These processes involve materi-
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TABLE I. INJECTION MOLDING PARAMETERS FOR HIP IMPLANT PROTOTYPES Parameter
316L Stainless Steel
Alumina
Melt Temperature Mold Temperature Injection Pressure Injection Velocity Packing Pressure Packing Time Cooling Time
170°C 60°C 14 MPa 160 mm/s 6 MPa 5s 20 s
190°C 40°C 14 MPa 160 mm/s 3 MPa 5s 20 s
als such as polyetheretherketone (PEEK), 316L stainless steel, titanium, and alumina. In the present study, the hip implant prototype was selected to demonstrate the viability of PIM in fabricating intricate 3D components. In general, neither 316L stainless steel nor alumina is an ideal candidate material for a hip implant. The hip implant was chosen for purposes of demonstration and investigation and was molded in an injection machine (Boy 22M/D, Germany). Both feedstocks were evaluated in relation to power–binder segregation and the formation of jetting defects.16–18 The injection molding parameters used for the hip implant are summarized in Table I. The debinding process was dependent on the constituents of the binder. A wax-based binder was used in the 316L stainless steel feedstock. The wax component in the binder was removed by thermal debinding using a low heating rate to avoid distortion and the formation of defects during the debinding process. A pre-sintering step was carried out to eliminate the remaining binder, Figure 1(a). The binder in the alumina feedstock was a modified polyalcohol, with a high water solubility (~65 w/o at 20°C). The alumina hip implant was solvent debound in water at 70°C for 24 h, and then thermally debound at 600°C for 1 h, Figure 1(b). Tests in a horizontal dilatometer (Netzch 402C, Germany) were carried out to monitor the sintering behavior of the two materials.14 The results provided reference values and parameters in the design of the sintering process for hip implants, Table II. 316L stainless steel tensile specimens were prepared for the determination of mechanical properties of the sintered parts. The design of the mold and injection stage for these test specimens is described elsewhere.9,10 The specimens were sintered at various temperatures to investigate the effect of this process parameter on the mechanical Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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Figure 1. Process parameters: (a) thermal debinding and pre-sintering for 316L stainless steel PIM parts, (b) solvent and thermal debinding for alumina PIM parts
TABLE II. SINTERING PARAMETERS FOR HIP IMPLANT PROTOTYPES Parameter
316L Stainless Steel
Alumina
Temperature Heating Rate Holding Time Cooling Mode
1,360°C 5°C/min 2h Natural Cooling in Furnace Vacuum
1,550°C 2°C, 5°C, 10°C/min 1h Natural Cooling in Furnace Air
Atmosphere
properties of the final parts. For sintering of the alumina hip implants, crack formation was the primary defect. The implant is sensitive to the heating rate used in sintering. In order to avoid distortion of the hip implants during thermal debinding and sintering, supports were designed to match the shrinkage of the sintered hip implants. The supports were made by PIM using the same material as the hip implant. For the area in contact, the surface of the support that carries the hip implant was designed to be the same size and same shape as the implant. For the area supported, the convex profile of the hip implant matched the concave surface of the support. RESULTS AND DISCUSSION The PIM hip implant prostheses are shown in Figure 2, confirming the capability of PIM to fabricate complex shaped components. The shape distortion in the debinding and sintering steps is Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
controlled by the supports, as illustrated in Figure 3. As the debound samples are fragile, it is necessary to handle them with care prior to sintering. Complete debinding involves two steps: primary debinding and pre-sintering. The pre-sintering step is designed to eliminate any residual binder and to develop the necessary strength for handling. The reductions in mass and pycnometer volume of the cylindrical 316L stainless steel specimens were measured, Figure 4(a): 77.4 w/o (85.4 v/o) and 95.9 w/o (99.7 v/o) of the binder was removed in the debinding stages I and II, respectively. The pre-sintered part had an ultimate tensile strength (UTS) of 112 MPa. Compared with the injection molded parts, the pre-sintered parts exhibited a shrinkage ~1%. The density of the pre-sintered parts, measured by pycnometer, was 7.92 g/cm3 and this is taken as the pore-free density of 316L stainless steel in the present study. The relative density of the pre-sintered samples was 64%. Binder removal as a function of time during solvent debinding of the alumina samples is shown in Figure 4(b). Solvent debinding for 24 h removed 88 v/o of the binder. The remaining binder was eliminated by subsequent thermal debinding. The measured pycnometer density of the debound samples was 3.91 g/cm3, corresponding to a relative density of 58%. Surface cracking was the primary defect in the debound alumina hip implants,
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Figure 2: PM hip implant prototypes: (a) 316L stainless steel, (b) alumina
Figure 3. Supports in hip implant prototypes for debinding and sintering: (a) alumina, (b) 316L stainless steel
Figure 5. This defect is dependent on the kinetics of binder removal, which is influenced by the solvent and/or thermal debinding cycles.18–21 Density variations during sintering are shown in Figure 6; these are derived from the in situ shrink-
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age curves in the dilatometric tests. Isotropic shrinkage is assumed. Densification of the 316L stainless steel begins when the temperature reaches 1,050°C. Below this threshold temperature, the material undergoes thermal expansion only. Densification is accomplished mainly in the temperature range from 1,130°C to 1,320°C. The alumina has a wide densification range from 1,000°C to 1,550°C, and densification continues during the holding period. A representative microstructure of the surface of the sintered 316L stainless steel, observed by scanning electronic microscopy (SEM), is given in Figure 7(a). Grain growth is evident and the final grain size is ~150 µm. The fracture surface of a sintered alumina sample shows that the grain size is ~1 µm, Figure 7(b). The mechanical properties of the sintered 316L stainless steel, as a function of sintering temperature, are shown in Figure 8. The strength increases rapidly in the temperature range of 1,250°C–1,300°C, while elongation increases primarily in the sintering range 1,300°C–1,360°C. After sintering at 1,360°C for 1 h the material exhibits a yield strength of 178.1 MPa, a UTS of 483 MPa, and an elongation of 49.9%. The mechanical properties are similar to those reported in the EPMA standard for components obtained by metal injection molding (MIM). 22 Reference values in standards for 316L stainless steel are 140 MPa for yield strength, 450 MPa for UTS, and 40% for elongation. Barriere, Liu, and Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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Figure 4. Reduction in mass and pycnometer volume of PIM parts during debinding and pre-sintering: (a) 316L stainless steel, (b) alumina
Figure 5. Cracks in alumina hip implant after solvent and thermal debinding
Figure 7. Representative SEM images of sintered samples: (a) surface of 316L stainless steel, (b) fracture surface of alumina
Figure 6. Variation in density during sintering stage (dilatometry)
Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
Gelin16 have shown that 316L stainless steel PIM parts sintered at 1,380°C exhibit a lower tensile strength and elongation due to the occurrence of a liquid phase during sintering. The strength of alumina decreases at high temperatures due to thermal softening and the alumina hip implant has low thermal conductivity. The stress and temperature gradients in the rapid sin-
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Figure 8. Tensile strength and strain for 316L stainless steel samples, sintered at different peak temperatures. Heating rate 5°C/min; furnace cooling. No hold time except for 1,260°C (1 h)
controlled by the flow of two phases, namely, the solid phase of the metallic powder and the fluid phase of the polymer binder. These two flows are described by the appropriate Navier–Stokes equations and their coupling is taken into account by momentum exchange. For solution of the biphasic injection model, an explicit algorithm of the finite element simulation had been developed by the authors. 9,10,16 It is completed by an approach that determines the constitutive behavior of each phase, based on capillary tests.23 The 316L stainless steel feedstock results are shown in Figure 10. Based on the process parameters used for the hip implant in 316L stainless steel, the filling process and powder volume fraction obtained by biphasic injection simulations are shown in Figure 11. Powder–binder segregation occurs as expected at corners and at the end position of the hip implant. On completion of the filling step, the powder volume is in the range 60.2–63.8 v/o. This inhomogeneity in green density is then used for the initial value in the sintering simulations. Analysis of Shrinkage During Densification Densification of the materials during sintering is governed by diffusion processes. The associated macroscopic behavior can be regarded as creep deformation, which results in shrinkage and distortion of the sintered part.1 Deformation in sintering is rate dependent and permits the use of a viscoplastic constitutive law based on continuum mechanics:24
Figure 9. Cracks resulting from the sintering of alumina hip implants at high heating rates: (a) 2°C/min, (b) 5°C/min, and (c) 10°C/min
tering cycles can induce cracks. Figure 9 shows that, with low heating rates, surface cracking in alumina is prevented, but the low heating rate may lead to grain growth. NUMERICAL SIMULATIONS IN PIM Analysis of Powder–Binder Segregation in Injection Stage The injection molding of feedstock mixtures is
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Figure 10. Viscosity curves of 316L stainless steel–base feedstock and associated viscous behavior for each modeling phase
Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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Figure 11. Filling state and powder-volume fraction during PIM of 316L stainless steel hip implant; biphasic injection simulation
Figure 12. Inhomogeneous distribution of final relative density in sintered 316L stainless steel hip implant; sintering simulation with initial condition of injection molded segregation
Figure 13. Inhomogeneous shrinkages of sintered 316L stainless steel hip implant in three orthogonal directions; sintering simulation with initial condition of injection molded segregation
σ) dev(σ σm– σs ε· vp = ———— + ———— I 3Kp 2Gp
(1)
where ε· vp is the viscoplastic strain rate, σ is the σ) is the deviatoric Cauchy stress tensor, dev(σ σ)/3 is the mean stress, I is stress tensor, σm = tr(σ the second order identity tensor, Gp and Kp and are the shear and bulk viscosity moduli of the porous material, and σs is the sintering stress that drives the densification process. Song et al.14 used dilatometric and bend tests during sintering to determine the viscosity modulus and sintering stress parameters in equation (1). In order to perVolume 45, Issue 3, 2009 International Journal of Powder Metallurgy
form the sintering simulation, this constitutive law is implemented in Abaqus ® finite element software by a user subroutine. The final density contours obtained by the sintering simulation are shown in Figure 12. The final relative density after sintering is in the range of 97.6%–98.2%, which is smaller than the variations in initial green density. The shrinkage of the sintered hip implant in three orthogonal directions is shown in Figure 13, based on the sintering simulations. The relatively large difference in green density can induce large variations in the dimensions of the sintered components.
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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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Strength of Sintered Parts The yield strength and ultimate tensile strength of a sintered body can be expressed by the equation:25 1–θ σ 1–θ 0 σy = σ0y ——, UTS = σUTS —————— 1+αs(Kc–1)θ Kc
(2)
where the superscript 0 indicates the strength of the wrought material, and the subscripts y and UTS refer to the yield strength and ultimate tensile strength, θ is the porosity factor, Kc is the stress concentration factor, and αs is a constant. α s is set at 1.8; K c is taken from the literature;25–27 for wrought 316L stainless steel, the yield stress σ0y is assigned 261 MPa; and the UTS 0 σUTS is 580 MPa. The evolution of density (or decrease in porosity) during sintering is obtained from the dilatometric test shown in Figure 6. The sintered strength of 316L stainless steel, predicted by the simulation, is shown in Figure 14, and compared with the values obtained from the tensile tests. When the sintering temperature for 316L stainless steel is <1,300°C, the strength predicted by the model is in good agreement with the experimental values. The relative density of the sintered parts is <92% in this case. According to sintering theory,28 this density corresponds to the initial and intermediate stages of sintering. During these two stages, densification occurs rapidly, accompanied by neck formation and growth and the pores
are interconnected. When the temperature exceeds 1,300°C, sintering enters the final stage. The density increases slowly and the pores close or are eliminated. Under these conditions, the strength model (equation (2)) is no longer too accurate in predicting mechanical properties. CONCLUDING REMARKS PIM is a multistep processing technology, and the final density, microstructure, strength, elongation, and fracture stress of the product are affected by the process parameters in each step. Powder–binder segregation induced by injection molding is a typical defect in PIM. The proposed biphasic injection model and related finite element simulations are shown to be reliable in the prediction and control of segregation defects. The inhomogeneous green density at the injection stage is taken as the initial condition for sintering modeling. Simulation of the sintering step gives a prediction of the density and inhomogeneous shrinkage of the sintered part. The strength of the sintered parts can be evaluated by the model and the predictive value is in good agreement with the experimental results. ACKNOWLEDGEMENT Portions of this work were supported by the National Natural Science Foundation of China (10772154) and the alternate PhD Program of the French government. REFERENCES
Figure 14. Experimental and predicted (model) tensile strength of 316L stainless steel samples, as a function of peak sintering temperature. Heating rate 5°C/min; furnace cooling. No hold time except for 1,360°C (1 h)
Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
1. R.M. German and A. Bose, Injection Molding of Metals and Ceramics, 1997, Metal Powder Industries Federation, Princeton, NJ. 2. C.D. Greene and D.F. Heaney, "The PVT Effect on the Final Sintered Dimensions of Powder Injection Molded Components", Mater. Design, 2007, vol. 28, no. 1, pp. 95–100. 3. C. Binet, D.F. Heaney, R. Spina and L. T ricarico, "Experimental and Numerical Analysis of Metal Injection Molded Products", J. Mater. Process. Technol., 2005, vol. 164–165, pp. 1,160–1,166. 4. C. Hinse, R. Zauner, R. Nagel, P. Davies and M. Kearns, "Simulation-Based Design for Powder Injection Moulding", Powder Injection Moulding International, 2007, vol. 1, no. 2, pp. 54–56. 5. V.V. Bilovol, L. Kowalski, J. Duszczyk and L. Katgerman, "The Effects of Constitutive Description of PIM Feedstock Viscosity in Numerical Analysis of the Powder Injection Moulding Process", J. Mater. Process. Technol., 2006, vol. 178, pp.194–199. 6. S.V. Atre, S-J. Park, R. Zauner and R.M. German, "Process Simulation of Powder Injection Moulding: Identification of Significant Parameters during Mould
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7.
8.
9.
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11.
12.
Filling Phase", Powder Metall., 2007, vol. 50, no. 1, pp. 76–85. R.T. Fox and D. Lee, "Analysis of Temperature Effects during Cooling in Powder Injection Moulding", Int. J. Powder Metall., 1994, vol. 30, no. 2, pp. 221–229. R.M. German, "Green Body Homogeneity Effects on Sintered Tolerances", Powder Metall., 2004, vol. 47, no. 2, pp. 157–160. J.C. Gelin, T. Barriere and M. Dutilly, "Experimental and Computational Modelling of Metal Injection Molding for Forming Small Parts", CIRP Ann.—Manuf. Techn., 1999, vol. 48, no. 1, pp. 179–182. T. Barriere, J.C. Gelin and B. Liu, "Experimental and Numerical Analyses of Powder Segregations on the Properties and Quality of Parts Produced by MIM", Powder Metall., 2001, vol. 44, no. 3, pp. 228–234. L.E. Khoong, Y.C. Lam, J.C. Chai, L. Jiang and J. Ma, "Numerical and Experimental Investigation on Thermal Debinding of Polymeric Binder of Powder Injection Molding Compact", Chem. Eng. Sci., 2007, vol. 62, no. 21, pp. 6,927–6,938. V.A. Krauss, A.A.M. Oliveria, A.N. Klein, H.A. Al-Qureshi and M.C. Fredel, "A Model for PEG Removal from Alumina Injection Moulded Parts by Solvent Debinding", J. Mater. Process. Technol., 2007, vol. 182, pp. 268–273.
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CAD + CAM = *CIT
*COMPUTER INTEGRATED TOOLING
13. B. Berginc, Z. Kampus and B. Sustarsic, "Influence of Feedstock Characteristics and Process Parameters on Properties of MIM Parts Made of 316L", 2007, Powder Metall., vol. 50, no. 2, pp. 172–183. 14. J. Song, J.C. Gelin, T. Barriere and B. Liu, "Experiments and Numerical Modelling of Solid State Sintering for 316L Stainless Steel Components", J. Mater. Process. Technol., 2006, vol. 177, pp. 352–355. 15. D.F. Heaney and R. Spina, "Numerical Analysis of Debinding and Sintering of MIM Parts", J. Mater. Process. Technol., 2007, vol. 191, pp. 385–389. 16. T. Barriere, B. Liu and J.C. Gelin, "Determination of the Optimal Process Parameters in Metal Injection Molding from Experiments and Numerical Modeling", J. Mater. Process. Technol., 2003, vol. 143–144, pp. 636–644. 17. P. Dvorak, T. Barriere and J.C. Gelin, "Jetting in Metal Injection Moulding of 316L Stainless Steel", Powder Metall., 2005, vol. 48, no. 3, pp. 254–260. 18. P. Dvorak, T. Barriere and J.C. Gelin, "Direct Observation of Mould Cavity Filling in Ceramic Injection Moulding", J. Eur. Ceram. Soc., 2008, vol. 28, pp. 1,923–1,929. 19. W.J. Tseng and C-K. Hsu, "Cracking Defect and Porosity Evolution during Thermal Debinding in Ceramic Injection Moldings", Ceram. Int., 1999, vol. 25, pp. 461–466. 20. R.V.B. Oliveira, V. Soldi, M.C. Fredel and A.T.N. Pires, "Ceramic Injection Moulding: Influence of Specimens and Temperature on Solvent Debinding Kinetics", J. Mater. Process. Technol., 2005, vol. 160, pp. 213–220. 21. P. Thomas-Vielma, A. Cervera, B. Levenfeld and A. Várez, "Production of Alumina Parts by Powder Injection Molding with a Binder System Based on High Density Polythylene", J. Eur. Ceram. Soc., 2008, vol. 28, pp. 763–771. 22. European Powder Metallurgy Association (EPMA), A Manufacturing Process for Precision Engineering Components—Metal Injection Moulding, 2000, EPMA, Shrewsbury, UK. 23. B. Liu, T. Barriere and J.C. Gelin, "Bi-phasic Simulation of Metal Injection Molding Constitutive Determination", J. Southwest Jiaotong University (English version), 2003, vol. 11, pp. 122–130. 24. E.A. Olevsky, "Theory of Sintering: from Discrete to Continuum", Mater. Sci. Eng. R, 1998, vol. 23, pp. 41–100. 25. P. Suri, D.F. Heaney and R.M. German, "Defect-Free Sintering of Two Material Powder Injection Molded Component Part II-Model", J. Mater. Sci., 2003, vol. 38, no. 24, pp. 4,875–4,881. 26. R.M. Ger man, "Computer Modeling of Sintering Processes", Int. J. Powder Metall., 2002, vol. 38, no. 2, pp. 48–66. 27. E.A. Olevsky, G.A. Shoales and R.M. Ger man, "Temperature Effects on Strength Evolution under Sintering", Mater. Res. Bull., 2001, vol. 36, no. 3, pp. 449–459. 28. R.M. German, Sintering Theory and Practice, 1996, John Wiley, New York, NY. ijpm
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RESEARCH & DEVELOPMENT
IRON-BASE PM MATRIX ALLOYS FOR DIAMONDIMPREGNATED TOOLS Marta Zak-Szwed*, Janusz Konstanty** and Anna Zielinska-Lipiec**
INTRODUCTION The instability of the cobalt market continues and the average price of this metal is at record levels, Figure 1. In consequence, replacement of cobalt-based diamond-impregnated materials with low-cobalt, or cobalt-free compositions, is a major focus in the diamond-tool industry. To date three families of fine prealloyed copper-base and iron-base powders have been developed commercially by Eurotungstene in France, and Umicore in Belgium, as a substitute for cobalt powders in the manufacture of diamond-impregnated tools. These powders (Table I) consist of a combination of at least two elements which are coprecipitated via proprietary manufacturing processes3–5 to yield prealloyed agglomerates of submicron-sized particles. The new alloys show excellent consolidation behavior and mechanical strength comparable with those of cobalt.6–10 Over the past decade toolmakers have benefited significantly from the increasing choice of new powder grades. To meet ever-increasing demands for tool performance, it has become necessary to further refine existing product characteristics and to develop new cobalt-free grades that fulfill complementary needs. In the present study we assess the viability of using iron-based powder metallurgy (PM) alloys as a matrix in diamond-impregnated tools.
The objective of this study was to determine the effects of powder composition and consolidation conditions on the microstructure and mechanical properties of iron–copper and iron–copper– tin alloys for use as a matrix in diamond-impregnated tools. Prealloyed powders were fully densified by hot pressing (between 650°C and 980°C) and the compacts evaluated for density, hardness, and bend properties. Microstructures and alloy constitution were characterized utilizing scanning electron microscopy (SEM), transmission electron microscopy (TEM), and X-ray diffraction (XRD). Consolidation in the stable austenite range resulted in a significant increase in ductility, at the expense of strength, especially in the absence of tin. The alloys exhibited a fine-grain microstructure independent of consolidation conditions. For both alloy powders, a hotpressing temperature was determined to optimize hardness and yield strength, coupled with ductility.
Figure 1. Free-market price quotation for high-grade (>99.8 w/o) refined cobalt1,2
*PhD Student, **Professor, AGH–University of Science & Technology, Faculty of Metals Engineering & Industrial Computer Science, Mickiewicz Avenue 30, 30-059 Krakow, Poland; E-mail:
[email protected]
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TABLE I. CHEMICAL COMPOSITION AND MEAN PARTICLE SIZE OF SUBSTITUTE ALLOYS11,12 Designation Fe
Nominal Composition (w/o) Cu Co Others
Fisher Subsieve Size (µm)
Source
Next 100 Next 200 Next 300 Next 900 Keen 10 Keen 20
25 15 72 80 n/a n/a
50 60 3 20 n/a n/a
25 25 25 25 19
n/a n/a
0.8–1.5 0.8–1.5 ~4 ~3 ~2.5 ~3
Eurotungstene
Cobalite 601 Cobalite HDR Cobalite CNF
70 66 68.4
20 7 26
10 27 -
3Sn; 2W; 0.6Y2O3
~5 6–7 ~2
Umicore
EXPERIMENTAL Two prealloyed iron-base powders were manufactured (H.C. Starck) by mixing aqueous metal–salt solutions with a carboxylic acid solution, separating the precipitation products from the mother liquor and reducing them in a hydrogen-containing atmosphere.13 The powders were analyzed for chemical composition, particle size and morphology, apparent density, and tap densi-
Figure 2. Representative micrographs of experimental powders: (a) iron–copper and (b) iron–copper–tin. SEM
Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
ty. Characterization of the degree of alloying was performed by XRD. The results and observations are shown in Figures 2 and 3, and in Table II. Compacts of both powders were fabricated by hot pressing in a graphite mold at temperatures between 650°C and 980°C in order to evaluate the effect of consolidation temperature on density, hardness, bend properties, phase composition, and microstructure. In each case the powder was held at the peak temperature for 3 min. under a pressure of 35 MPa. Specimens, 25 mm in dia. × 5 mm high, were fabricated for density and hardness determination by the water-displacement method and Rockwell B test, respectively. In the latter case, hardness numbers between 100 and 110 RB were also included in order to comply with quality-control procedures used within the diamond-tool industry.11 Bend test bars were 4 mm high, 5 mm wide, and 40 mm long. The bars were supported by two high-speed steel (HSS) rods 3.5 mm in dia. spaced 30 mm apart, and loaded perpendicular to the hot
Figure 3. X-ray diffraction patterns of as-received experimental powders
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IRON-BASE PM MATRIX ALLOYS FOR DIAMOND-IMPREGNATED TOOLS
TABLE II. COMPOSITION AND PROPERTIES OF POWDERS Designation Fe-Cu Fe-Cu-Sn
Fe
Composition (w/o) Cu Sn Sm
82.5 79.1
15.9 16.5
3.0
0.6
O
Fisher Subsieve Size (µm)
Apparent Density (g/cm3)
Tap Density (g/cm3)
0.48 0.69
3.30 1.63
0.98 0.79
1.50 1.33
pressing axis, using a similar HSS rod positioned midway between the supports. Transverse rupture strength (TRS), 0.2% offset yield strength, and the extent of plastic deformation at failure were calculated. Selected fractured specimens, representative of each alloy, were examined by means of XRD. Their microstructures were also observed on diamond-polished-and-etched transverse sections. RESULTS AND DISCUSSION The dependence of density and hardness on the hot-pressing temperature is presented for both alloys in Figure 4. The data confirm an excellent densification response of the powders. Hot pressing within the temperature range 650°C and 980°C results in near-pore-free densification of both alloys. Hardness remains high at temperatures ≤800°C but decreases markedly at higher temperatures. In Table III the results of the three-point bend test are benchmarked against data obtained for hot pressed Umicore Extrafine Co14, a material commonly used as a matrix in commercial diamond-impregnated tools.11 Selected fractured specimens, representing each alloy, were examined utilizing SEM and XRD for phase compositions as well as for evidence of grain growth. The results are given in Table IV. Representative stress–strain curves, microstruc-
tures, and back-reflection patterns are presented in Figures 5–7. Microstructures of the iron–copper–tin alloys were also observed in thin foils by utilizing TEM. Representative observations are presented in Figures 8 and 9. From the bend-test results (Table III) and Figure 5, it is seen that the iron–copper alloys attain an attractive combination of yield strength and ductility after hot pressing at 820°C. The compacts densified at 700°C fail in a brittle manner, while those hot pressed at 900°C show a marked increase in ductility at the expense of yield strength. The X-ray back-reflection patterns (Figure 6) confirm that the eutectoid reaction occurring in the alloy during the densification cycle imparts ductility. It remains unclear whether austenite was present during hot pressing at nominally 820°C (Figure 10) because the actual temperature may vary slightly between specimens within the resistance-heated graphite mold, and differ from the temperature measured by a pyrometer on the outer surface of the mold. The quantitative XRD analysis (Table IV) shows a similar phase composition for all the iron–copper compacts regardless of the hot-pressing temperature. However, increases in the lattice spacing of the α-iron solid solution with temperature are attributed to an increasing supersaturation of fer-
Figure 4. Effect of powder-consolidation temperature on density and hardness (a) iron–copper and (b) iron–copper–tin
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TABLE III. EFFECT OF COMPOSITION AND HOT-PRESSING TEMPERATURE ON DENSITY, HARDNESS, AND BEND PROPERTIES Material (Hot-Pressing Temperature)
Compact No.
Density (g/cm3)
Hardness (RB)
TRS (MPa)
Yield Strength (MPa)
Plastic Strain (%)
Fe-Cu (700ºC)
1 2 3
7.94 7.95 7.92
96.5 97.2 95.6
1,077 1,219 1,133
1,077 1,219 1,133
0 0 0
Fe-Cu (820ºC)
4 5 6
7.95 7.95 7.95
95.3 94.8 94.1
1,186 1,163 1,114
1,161 1,145 1,103
6 5.5 6
Fe-Cu (900ºC)
7 8 9
7.96 7.96 7.94
86.6 85.9 86.4
not broken
765 760 716
not broken
Fe-Cu-Sn (700ºC)
10 11 12
7.99 7.98 7.94
104.8 104.3 103.9
297 1,116 1,171
297 1,116 1,171
0 0 0
Fe-Cu-Sn (900ºC)
13 14 15
7.88 7.88 7.86
98.4 96.3 97.5
1,233 1,182 1,212
1,156 1,111 1,124
4.5 5 5
Fe-Cu-Sn (980ºC)
16 17 18
7.85 7.85 7.89
97.3 96.9 98.1
1,216 1,157 1,182
1,143 1,139 1,165
4 0.5 0.5
Co (850ºC) *
-
8.74±0.03
105±1.4
1,854±450
1,186±60
11±9.5
* Confidence intervals estimated by t-distribution at 90% confidence level TABLE IV. PHASE ANALYSIS AND LATTICE SPACINGS Material (Hot-Pressing Temperature Compact No.)
Volume Fraction: Cu Solid Solution
Lattice Spacing: Cu Solid Solution (nm)
Fe-Cu (700ºC-2) Fe-Cu (820ºC-6) Fe-Cu (900ºC-7) Fe-Cu-Sn (700ºC-11) Fe-Cu-Sn (900ºC-15) Fe-Cu-Sn (980ºC-18)
0.21 0.22 0.22 0.20 0.10 0.07
0.361727 0.361921 0.361900 0.368859 0.366698 0.364550
Sn in Cu Solid Lattice Spacing: α−Fe Sn in α−Fe Solid Solution* (w/o) Solid Solution (nm) Solution** (w/o) 13.0 9.1 5.6
0.286606 0.286710 0.286786 0.286731 0.286780 0.286715
0.43 0.58 0.39
*Estimated from reference 15 (Figure 114, p. 600), **Estimated from reference 15 (Figure 124, p. 633)
Figure 5. Representative stress–strain curves in bending. (a) iron–copper and (b) iron–copper–tin. For clarity, individual curves are offset by a strain of 0.5%
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Figure 6. Representative micrographs (SEM) and X-ray back-reflection patterns of iron–copper: (a) 700°C-2, (b) 820°C-6, and (c) 900°C-7
rite with copper (Figure 10(a)). The iron–copper–tin alloy displays a high yield strength which appears to be insensitive to the hot-pressing temperature. The X-ray back-reflection patterns (Figure 7) show less extensive grain growth than in the iron–copper alloy consolidated at 900°C, even after hot pressing at 980°C.
Figure 7. Representative micrographs (SEM) and X-ray back-reflection patterns of iron–copper–tin: (a) 700°C-11, (b) 900°C-15, and (c) 980°C-18
The powder compacts consolidated at 700°C are brittle. The TEM micrograph shown in Figure 8(a) suggests the presence of microdiscontinuities which weaken the grain boundaries and cause intercrystalline fracture, Figure 8(b). By increasing the hot-pressing temperature to 900°C it was possible to obtain a reasonable combi-
Figure 8. (a) TEM micrograph and (b) SEM fracture surface of compact 700°C-11
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Figure 9. Represenative TEM micrographs of iron–copper–tin hot-pressed compacts: (a) and (b) 700°C, (c) and (d) 900°C, (e) and (f) 980°C
nation of hardness, yield strength, and ductility. It was found that a further increase in the hot-pressing temperature (to 980°C), restored shortness. The quantitative phase analysis data in Table IV reveal that the volume fraction of the copper solid solution in the iron–copper–tin alloy is inversely related to the consolidation temperature. It has Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
also been found that the content of tin in copper decreases from 13 to 5.6 w/o as the hot-pressing temperature increases from 700°C to 980°C. By analyzing the XRD data in Figure 3 it is evident that the prealloyed iron–copper–tin powder contains a solid solution of tin in copper. The solid solution melts above 798°C, as can be deduced
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Figure 10. Iron–copper, copper–tin and iron–tin phase diagrams16
from the copper–tin phase diagram (Figure 10), to form a liquid which, theoretically, may have a transient nature due to considerable solubility of tin in ferrite at and above 798°C (Figure 10). Calculations based on the lattice spacing data from Table IV and Pearson,15 however, give relatively low contents of tin dissolved in ferrite, namely, 0.58 and 0.39 w/o after hot pressing at 900°C and 980°C, respectively. This may indicate that the missing tin was ejected from the graphite mold as a tin-rich flash, decreasing both the volume fraction of the copper solid solution and the content of tin in the material. As seen in Figure 7(b) and (c), at 900°C and 980°C, the liquid phase penetrates the grain boundaries to form a copper–tin film. The amount of liquid, as well as its tendency for spreading and formation of a grain boundary film, increases with temperature which is likely to cause brittle failure of the over-sintered material. The TEM micrographs in Figure 9 give corroborating evidence for the formation of a copper solid solution film. The microstructure of the material consolidated at 700°C consists primarily of submicron-sized copper solid solution grains and α-iron solid solution grains (Figure 9(a)). The α-iron solid solution is precipitation hardened by copper-solid-solution precipitates, observed at certain grain orientations (Figure 9(b)), mostly <20 nm in size. Such precipitates have been identified as coherent or semi-coherent with the ferrite matrix by high-resolution electron microscopy.17,18 The grains coarsen with the hot-pressing temperature (Figures 9(c)), and likewise coarsen the copper -solid-solution precipitates (Figure 9(d)), which is associated with a loss of
42
coherency.18 On cooling from 980°C, the copper solid solution solidifies as a grain boundary film (Figure 9(e)) or eutectic-like microconstituent (Figure 9(f)). In either case the zone adjacent to the film, or eutectic, is free from precipitates of the copper solid solution. CONCLUSION The experimental work in this study has shown that by correct choice of processing conditions it is possible to obtain iron–copper and iron–copper–tin exhibiting an excellent combination of hardness, yield strength, and ductility. As such, these alloys are suitable as a matrix in diamondimpregnated tools.11 It is important that the powders can be consolidated to essentially pore-free density by hot pressing at temperatures markedly lower than 900°C to ensure minimal degradation of the synthetic diamond.11 The application of hot pressing enables near-pore-free densification of both alloys even at 650ºC. This gives the alloys a significant advantage over cobalt powders which, depending on particle size, need to be hot pressed at temperatures ranging from 750°C to 950°C.19 The optimal combination of mechanical strength and ductility was obtained after hot pressing the iron–copper and iron–copper–tin powder at 820°C and 900°C, respectively. It is noteworthy that the latter alloy is comparable with cobalt in relation to a stable yield strength (~1,140 MPa) over a broad range of hot-pressing temperatures. ACKNOWLEDGEMENTS The authors gratefully acknowledge W. Ratuszek and A. Radziszewska for assistance with the XRD Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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and SEM chracterization, and Benno Gries and Ber nd Mende, H.C. Starck GmbH, Goslar, Germany, for providing the powders. The work was supported by the Polish Ministry of Science & Higher Education, contract 11.11.110.788.
19. J. Konstanty, Cobalt as a Matrix in Diamond Impregnated Tools for Stone Sawing Applications, 2002, AGH Uczelniane Wydawnictwa Naukowo-Dydaktyczne, Krakow, Poland. ijpm
REFERENCES 1. “2004 Production Statistics”, Cobalt News, 2005, vol. 05, no. 2, pp. 3–4. 2. “2007 Production Statistics”, Cobalt News, 2008, vol. 08, no.2, pp. 3–4. 3. R. Standaert, “Pre-alloyed, Copper Containing Powder, and its Use in the Manufacture of Diamond Tools”, US Patent 6,312,497, November 6, 2001. 4. M. Bonneau, S. Chabord and G. Prost, “Micronic Prealloyed Metal Powder Based on Three-dimensional Transition Metal”, US Patent 6,613,122 B1, September 2, 2003. 5. B-J. Kamphuis and J. Peersman, “Pre-alloyed Bond Powders”, US Patent 7,077,883 B2, July 18, 2006. 6. I.E. Clark and B-J. Kamphuis, “Cobalite HDR—a New Prealloyed Matrix Powder for Diamond Construction Tools”, Industrial Diamond Review, 2002, vol. 62, no.3, pp. 177–182. 7. “A New Generation of Powders for the Diamond Tool Industry”, Marmomacchine International, 1997, vol. 18, pp. 156–157. 8. M. Bonneau, “NEXT and NEXT Pre-mixed Powders: a Complete Range of Basis”, Diamante Applicazioni & Tecnologia, 1999, vol. 5, no.18, pp. 45–52. 9. B-J. Kamphuis and A. Serneels, “Cobalt and Nickel Free Bond Powder for Diamond Tools: Cobalite CNF”, Industrial Diamond Review, 2004, vol. 64, no. 1, pp. 26–32. 10. “Keen—a New Concept in Prealloyed Powders”, Industrial Diamond Review, 2005, vol. 65, no. 3, pp. 45–47. 11. J. Konstanty, Powder Metallurgy Diamond Tools, 2005, Elsevier, Oxford, UK. 12. “Next Range Brochure”, Eurotungstene Metal Powders, Grenoble Cedex, France, www.eurotungstene.com, 09/2004. 13. B. Mende, G. Gille, B. Gries, P. Aulich and J. Munchow, “Pre-alloyed Powder”, U.S. Patent 6,554,885 B1, April 29, 2003. 14. A. Romanski, “Engineering Structure and Properties of Diamond-Impregnated, Metal-Bonded Composite Materials”, PhD Thesis, 2000, AGH-University of Science & Technology, Krakow, Poland. In Polish. 15. W.B. Pearson, A Handbook of Lattice Spacings and Structures of Metals and Alloys, 1964, Pergamon Press, Belfast, Northern Ireland. 16. Binary Alloy Phase Diagrams, Second Edition, edited by T.B. Massalski, ASM International, Materials Park, OH, 1990, vol. 2, pp. 1,409, 1,482 & 1,775. 17. H.R. Habibi, “Atomic Structure of the Cu Precipitates in Two Stages Hardening in Maraging Steel”, Materials Letters, 2005, vol. 59, pp. 1,824–1,827. 18. G. Fourlaris, A.J. Baker and G.D. Papadimitriou, “Microscopic Characterisation of e-Cu Interphase Precipitation in Hypereutectoid Fe-C-Cu alloys”, Acta Metall. Mater., 1995, vol. 43, no. 7, pp. 2,589–2,604.
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HOT ISOSTATIC PROCESSING SERVICES FOR PRODUCTION AND RESEARCH PROGRAMS ISO 9001, AS9100 REGISTERED
• CASTING DENSIFICATION • Improved Properties • Reduced Rejection Rate • Reduced Scrape Rate
• POWDER CONSOLIDATION • PRESSURE BRAZING • DIFFUSION BONDING • CERAMICS
KITTYHAWK PRODUCTS
11651 MONARCH ST. • GARDEN GROVE, CA 92841 Tel. (714) 895-5024 Fax (714) 893-8709 www.kittyhawkinc.com E-mail:
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RESEARCH & DEVELOPMENT
PROCESSING OF BULK Fe-Zn ALLOYS USING EXPLOSIVE COMPACTION Robert P. Corson,* Sivaraman Guruswamy,** Michael K. McCarter,*** and Chen-Luh Lin****
INTRODUCTION The phenomenon of magnetostriction refers to changes in the dimensions (strain) of a ferromagnetic material on the application of a magnetic field. It has been shown that the addition of gallium to α-iron results in a large increase in Joule magnetostriction.1–3 While the mechanism for this large increase is not clear, this discovery raises the possibility that other non-rare-earth alloying elements might also lead to large increases in the magnetostriction of α-iron. The phenomenon of magnetostriction is not well understood, and so it is difficult to decide from first principles which elements, when added to iron, will lead to an increase in magnetostriction. As zinc occupies a position adjacent to gallium in the periodic table, an examination of the influence of zinc on the magnetostriction of iron is of particular interest. In its ground state, zinc has a filled d shell and one less valence electron than gallium.4,5 The zinc atoms are also larger than the gallium atoms and lead to a larger dilation of the iron lattice. Zinc also has a large solubility in α-iron at elevated temperatures. Based on these properties, it appears fruitful to examine the magnetostrictive effects of alloying iron with zinc. It is expected that a knowledge of magnetostriction in iron–zinc alloys could lead to an improved understanding of the factors responsible for the large magnetostriction in certain iron alloys compared with that in pure iron. The processing of these alloys in bulk form with the desired crystallographic texture, and in a suitable size, presents a challenge. The boiling point of zinc (1,179K) is much lower than the melting point of iron (1,809K),5 and this limits processing temperatures to below 1,179K. The two metals cannot be melted together, as the zinc vapor pressure in contact with the iron-rich melt would be too high. Generally, iron-rich iron–zinc alloys have been made by either (i) vapor-phase transport and dissolution of zinc in iron, or (ii) mechanical alloying (MA), with product dimensions <100 µm.5–7 In the first approach, zinc is vaporized in a controlled manner at process temperatures below the boiling point of zinc. The zinc vapor is then transported to, and deposited on, solid αiron maintained at a high process temperature (~1,038K). The zinc that
Iron–zinc alloys are of interest in the study of magnetostriction in α-ironbase binary alloys. A novel powder metallurgy (PM) approach to fabricate bulk iron-rich iron–zinc alloys with a [100] texture was examined. The approach embraced magnetic field orientation, liquid-phase sintering (LPS) to form a porous body, explosive compaction, and a homogenization anneal. X-ray tomography (XRT), scanning electron microscopy (SEM), energy dispersive X-ray spectroscopy (EDAX), X-ray diffraction (XRD), orientation imaging microscopy (OIM), and vibrating sample magnetometry (VSM) were utilized to characterize the alloys. The feasibility of obtaining bulk iron–zinc alloys was demonstrated, albeit with a weak [100] texture. The addition of zinc does not change the magnetostriction of iron significantly.
*Formerly with the University of Utah, **Professor, ****Research Professor, Department of Metallurgical Engineering, University of Utah, 135 South 1460 East, Room 412, Salt Lake City, Utah 84112, USA; E-mail:
[email protected], ***Professor, Department of Mining Engineering, University of Utah, 135 South 1460 East, Room 313, Lake City, Utah 84112, USA
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is deposited diffuses into the iron at a rate that is dependent on the processing temperature, forming a series of zinc-rich and iron-rich phases.5–8 A subsequent homogenization anneal, carried out over a month at a temperature below the boiling point of zinc (e.g., 1,158K), results in an iron–zinc alloy of the desired composition. Samples prepared using this approach are typically in the form of small foils or powders with a thickness/size of a few tens of microns6 in order to minimize diffusion distances and homogenization times. The synthesis of iron–zinc alloys using MA7 involves milling a blend of the iron and zinc powders (few tens of microns in size) in an attritor or ball mill over an extended period of several hours. The impacting of the powders between colliding balls leads to repeated cold welding and fracture of particles and mixing of the powder constituents at the microscale or nanoscale levels. Subsequent annealing allows for chemical homogenization and strain relief. The product obtained is a powder with particle sizes in the range of only a few tens of microns. The process is not particularly efficient and suffers from contamination issues. Because the magnetostrictive properties vary significantly with crystallographic direction (in iron the direction of maximum magnetostriction is [100]), a new process that results in [100] textured bulk samples was sought. In the present investigation the samples used for magnetostriction measurements were typically about 10 to 15 mm in length and about 3 to 6 mm in the lateral dimension. This allowed for a reasonably uniform magnetization region in the sample over which strain gauges could be attached. This study addressed the fabrication of [100] textured iron–zinc alloys with zinc contents varying from 5 a/o to 30 a/o in bulk form (several mm size range) suitable for use in magnetostriction studies. The approach used involved magnetic field orientation of the iron component in the powder blend, partial LPS of loose powders, explosive compaction, and a homogenization anneal. SEM, EDAX, XRT, XRD, OIM, and VSM were used to analyze the alloys and to verify if this fabrication approach was viable. Magnetostriction measurements were made on the samples to determine the influence of zinc on the magnetostriction of iron. EXPERIMENTAL PROCEDURE Explosive Compaction Explosive compaction is useful for processing
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materials that are difficult to fabricate utilizing conventional metal-forming operations. Because of the rapidity of the operation in explosive compaction, consolidation is achieved essentially at room temperature and materials in powder form can be consolidated while preserving the structure present in the individual powder particles.9 In the case of iron–zinc alloys, the consolidation step avoids the use of high temperatures at which the zinc vapor pressure is prohibitively high. In this experimental work, sintered samples of pure iron, Fe-5 a/o Zn, Fe-10 a/o Zn, Fe-15 a/o Zn, Fe-20 a/o Zn, Fe-25 a/o Zn, and Fe-30 a/o Zn were explosively compacted to obtain bulk samples suitable for use in magnetic and magnetostriction measurements. The alloys were produced by mixing powders of each of the components in the appropriate amounts. The pure iron and zinc powders were obtained from Alfa Aesar. The iron powder had a purity >99.9% with a size range 75–180 µm (-80 + 200 mesh). The zinc powder had a purity of 99.9% (metals basis) with a size <150 µm (-100 mesh). The mixed powders were slowly dropped into a quartz tube with a SmCo magnet placed in contact with the bottom of the tube. The magnet aligned the easy magnetization direction [100] in the largest grain within each of the ferromagnetic particles in the powder mix (Figure 1(a)) in a direction parallel to the axis of the magnet (the magnetic field). This was expected to assist in developing a [100] texture in the final product. The tube was then placed in a furnace and the powders were sintered to create a compact strong enough for handling. The sintering temperature was well below the Curie temperature of the magnet used for powder alignment and the boiling point of zinc. The sintering sequence used in this experiment was 1 h at 673K, 2 h at 813K, and 3 h at 958K. This sequence involved LPS until the zinc-rich liquid was consumed by dissolution in the iron and by the formation of intermetallic compounds. In LPS, a minor component is present in the liquid state which facilitates rapid mass transport through dissolution, diffusion through the liquid, and redeposition in the neck regions connecting the solid powder (primary-phase) particles.10,11 The compact was then placed inside a coppertube assembly (Figure 1(b)). The tube was sealed at the bottom with a steel plug, and a steel ball bearing or shaped plug was used to seal the top of this copper tube (Figure 1(b)). The fill volume of Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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pressure generated by the explosion) is given by the formula: P = (.25 × 10-6) ρD2
Figure 1. (a) Schematic illustrating geometry used for alignment of easy magnetic direction of iron powder, and (b) schematic of explosive compaction geometry
the copper tube has an ID of 6.25 mm and an OD of 7.9 mm, and the length of the powder-fill volume was about 38 mm. The copper-tube assembly was then positioned and attached to the plastic-explosive-container assembly that housed the explosive and the detonation fuse. The explosive used in this set of experiments was a twocomponent perchlorate–aluminum-based IRECO® 207 X explosive. The reason for the shaped topplug seal was to change the shock-wave profile in the blast.12–15 The shock-wave profile allowed the tube to compress around the sample, and not peel off the sample. During detonation, high pressures are generated at the shock-wave front in the sample. The nominal detonation pressure (the peak Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
(1)
where P is the pressure generated by the explosive in GPa, ρ is the density of the explosive (g/cm3), and D is the detonation velocity (m/s).16 With many explosives detonating in the km/s range, the pressure generated by the blast is high. The explosive used in this set of experiments had a density of 1.16 to 1.25 x103 kg/m3. The detonation velocity measured for this explosive was 3,600 to 3,700 m/s resulting in a pressure of 3.76 GPa as the minimum overpressure generated. Due to the small diameter of the copper tube containing the powder, the shock-wave velocity in the sample was close (negligible lag) to that of the detonation velocity. As the shock wave moves through the compact, it moves in a wavy fashion due to the porous nature of the material. This leads to extremely high temperatures at the particle interfaces due to adiabatic heating, which welds adjacent particles. Additionally, jetting (motion of material as if in a high-velocity-fluid stream) due to high particle velocity occurs, filling voids and joining particles together by friction welding. The material can be considered to be under hydrostatic loading, which is similar to the loading condition existing in conventional PM processes.12–15 In addition to its ability to achieve near -pore-free density in the product, other important benefits of explosive compaction are the preservation of nonequilibrium microstructures in the powder due to the short duration of the pressure and temperature pulses and break-up of oxide films on particle surfaces. This break-up minimizes fracture along prior particle boundaries and allows for intimate contact between adjacent powder particles. The latter feature is particularly important in PM iron–zinc alloys. An important consideration in explosive compaction is the shape of the shock wave as it moves through the powder since it determines the effectiveness of compaction. For example, if pressure decreases too much at the center of the sample, the powder will not be compacted at this location, or if the pressure increases towards the middle of the sample, a central hole (mach tube) will form due to the intense release wave. Ideally, a shock wave with a conical cross section should form since this will give a constant pressure over the cross section.13
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The best way to maximize compaction without leaving a mach tube in the sample is to optimize the explosive mass-to-powder mass ratio (E:M). The amount of explosive used was 30 g and the E:M was approximately 6. As these were proof-ofconcept experiments, no optimization of the E:M for compaction was attempted. After compaction, the copper tubing was machined off the compacts, which were then subjected to a homogenization anneal at 973K for 96 h. At each step of the process, slices were taken from the samples for characterization. One final compact sample, ~30 mm long × 4 mm in dia., was obtained from each of the alloys. Characterization Optical microscopy was used to visually examine the grain size and grain-size distribution in the samples. XRD characterization was performed using a Siemens D5000 X-ray diffractometer and Cu Kα radiation. On a finer scale, SEM analysis was used to determine the microstructure, chemical composition, and crystallographic texture in the iron–zinc samples. SEM analyses were performed using a Hitachi Environment SEM S3000N. The SEM had TSL®-EDAX and TSL®-OIM attachments. The OIM hardware and software enabled the collection and analysis of electron backscatter diffraction patterns (Kikuchi lines) to determine the crystallographic orientations of the grains in the sample. Assessment of the crystallographic texture in the various iron–zinc alloy samples after explosive compaction and annealing was made using the pole figures and inverse pole figures generated during OIM analysis. Surface-area and particlesize estimates of the metal powders were obtained prior to processing utilizing the BET method with a Micromeritics ASAP 2010 unit and SEM images. At the powder-particle sizes in this study, BET is not too accurate for the determination of particle size or surface area. Thus, the assessment of powder-particle size relied primarily on the SEM images of the powder particles. XRT of sections excised from samples after different stages of processing was performed using a Konoscope 40-130 X-ray high-precision microtomography system with a tungsten filament. This provides three-dimensional (3D) images of the internal structure of opaque materials in a nondestructive manner. In this system, a conical beam of tungsten Kα X-ray radiation is generated and
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projected on a two-dimensional (2D) charge-coupled-device (CCD) detector. The sample is placed in the path of the beam and rotated over 360° in small angular steps as the images are collected at a 10-µm resolution. The images formed depend on the density, average atomic number, size, shape, and the spatial distribution of the different phases or regions present in the sample. Aracor® Konoscope Acquisition and Image Reconstruction (AIR) software was used to acquire true 16-bit image data and to process the images to reconstruct a 3D image. Visualization of the reconstructed 3D image was done using Volsuite ® (http://
[email protected]), and image processing was performed using medical image processing, analysis, and visualization (MIPAV) software from the National Institute of Health (NIH) (http://www.mipav.cit.nih.gov). Magnetization measurements were performed on compacted and annealed iron–zinc alloy samples using a Lakeshore Model 7307 vibrating sample magnetometer with Lakeshore Model 735 VSM control electronics and a Lakeshore Model 450 Gaussmeter with a 3 Tesla Hall probe. The sample dimensions were ~2 mm × 2 mm × 2 mm. Saturation magnetostriction measurements were also made on compacted and annealed iron–zinc alloy samples ~10 mm long × 4 mm dia. using strain gauges cemented to the sample and applied fields in the range of -63.64 to +63.64 kA/m (-800 to + 800 Oe). RESULTS AND DISCUSSION Figure 2 shows SEM images of the iron and zinc powders. The iron-powder particles are roughly spherical, whereas the zinc-powder particles are angular. The BET surface areas of the iron and zinc powders were 0.0240 m 2/g and 0.0560 m2/g, respectively. The largest dimensions of both powders were approximately the same (≥100 µm) but the aspect ratio of the zinc powder was larger. Even with differences in the shapes of the powders and the different densities of the metals (7.87 × 103 g/cm3 for iron and 7.14 × 103 g/cm3 for zinc), the powders blended readily. After the initial sintering sequence of the magnetically aligned loose powder mix (1 hr at 673K, 2 hr at 813K, and 3 hr at 958K), the particles were bonded sufficiently to allow the compact to be handled without fragmenting. SEM images of the loose-powder-sintered Fe-30 a/o Zn sample show that all the zinc has alloyed with the iron Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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Figure 2. Representative SEM micrographs of (a) iron and (b) zinc powders
powders (Figure 3(a)) and necks have formed between the particles (Figure 3(b)). The density of this compact is low and the compact is brittle, because the only bonds holding it together are the narrow interconnecting necks between the particles. Figure 4 shows the XRT images obtained from the sample that had been loose-powder sintered. A more detailed description of Figure 4 is given subsequently in this section. Figure 5(a) shows an explosively compacted sample assembly. The sample is encased in the copper tube. Figure 5(b) shows a cross section of a typical sample. The central hole is a mach tube resulting from overpressure during compaction at the center of the sample. After the compression wave passes, a rarefaction wave results, and this leaves an empty region in the center of the sample. Of note is the high density of the compact, except for the central mach hole region. The comVolume 45, Issue 3, 2009 International Journal of Powder Metallurgy
Figure 3. SEM images of porous iron–zinc alloy compact annealed at 673K for 1 h, 813K for 2 h, and 958K for 3 h prior to explosive compaction: (a) particle morphology, and (b) neck region between two particles
pact was then sectioned into slices for further annealing. One sample was kept in the as-compacted condition for baseline comparisons. Figures 6(a)–6(d) show the Fe-Kα and Zn-Kα X-ray maps of sections of the as-explosively-compacted Fe-5 a/o Zn ( Figures 6(a) and 6(b)) and Fe-30 a/o Zn (Figures 6(b) and 6(c)) alloys. These maps clearly show iron-rich phases surrounded by zincrich phases. The hard iron-rich particles can be readily fragmented as can the zinc-rich region that was jetted around them during compaction. The attractive feature of explosive compaction is that, even though the local temperatures can be high, the compact cools extremely rapidly. In fact, most of the adiabatic work done on the sample appears as heat. In the as-compacted condition, there are two phases present. Figures 6(e) and 6(f)
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Figure 4. (a) 3D isometric view of irregularly shaped sintered compact, (b) cut parallel to y-z plane, (c) cuts parallel to y-z and x-z planes, and (d) cuts parallel to x-y, y-z, and x-z planes
Figure 6. Iron and zinc X-ray maps of sections of explosively compacted iron–zinc alloys: (a) to (d) before homogenization anneal; (e) and (f) after homogenization anneal
Figure 7. XRD pattern of Fe-30 a/o Zn alloy compact after homogenization anneal Figure 5. (a) Explosively compacted sample assembly and (b) optical micrograph of representative cross section of compact
show the Fe-Kα and Zn Kα X-ray maps in the Fe30 a/o Zn compact after the homogenization anneal, a process step for allowing complete mixing of the iron and zinc to achieve chemical unifor mity. The homogenization anneal was
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performed for 43 h at 873K and an additional 43 h at 973K. An XRD pattern of a Fe-30 a/o Zn sample that was annealed for 43 h at 873K and an additional 43 h at 973K after compaction is shown in Figure 7; only α-iron solid-solution peaks are present, indicating completion of chemical homogenization. Similar results were Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
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obtained for the other alloys. Figure 4(a) shows 3D isometric views of a piece of loose-powder-sintered Fe-30 a/o Zn. Figures 4(b)–4(d) show the 3D isometric views with sections cut parallel to (i) the y-z plane, (ii) the y-z and z-x planes, and (iii) the x-y, y-z, and x-z planes. The fragility of the as-loose-powder-sintered sample necessitated use of the irregularly shaped sample for XRT imaging. Figure 4 clearly shows the arrangement of the particles and the high level of porosity in the sample. As was shown in Figure 3(b), the particles are held together by the sintering necks. The melting and diffusion of zinc into the iron powder particles leaves voids in the iron–zinc alloy powder -particle network. Figures 8(a)–8(d) show 3D isometric views of a
Figure 8. (a) to (d) 3D isometric views of as-compacted Fe-30 a/o Zn alloy with images of sections cut parallel to x-y, y-z, and z-x planes. Dark regions are 3D voids/cracks in the compact. (e) 2D images of a series of selected sections along sample axis in a 3D XRT image of as-compacted Fe-30 a/o Zn alloy. Images show the gradual disappearance of the mach stem region as the step-size number decreases
Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
piece of the as-compacted Fe-30 a/o Zn alloy along with the images of sections cut parallel to the x-y, y-z, and z-x planes. Much of the copper tubing was machined off and the sample gently hand polished to a parallelepiped-shaped piece. For a resolution of 10 mm, each volumetric cell (voxel) has dimensions of 10 mm × 10 mm × 10 mm. For a 512 pixel × 512 pixel × 300 pixel image, the size of the sample will be ~5 mm × 5 mm × 3 mm and the memory needed would be 75 MB (300 MB for floating point data). For our data processing, a 300 MB file is convenient and the corresponding sample size is thus 5 mm × 5 mm × 3 mm. In Figure 8(e), the 2D images of selected sections along the sample axis are shown to illustrate the evolution of the structure in the direction of shock-wave propagation. Each number (lower left) represents the section number with successive sections 10 µm apart. Clearly, much of the compact is pore free except for the mach stem region near the center. It also illustrates the slow decrease in the size of the mach hole and the gradual elimination of this region. This is consistent with our observations in other compaction experiments, namely, that the mach-tube dimension decreases when moving along the sample axis in the direction of shock-wave propagation. This suggests a decrease in the compaction pressure as the shock wave propagates along the axis. Some fine cracks were observed that arise from the tensile component of the reflected shock waves. The same observations were made in all the iron–zinc alloy samples. These cracks and the mach tube for mation could be avoided by decreasing the E:M. Figures 9(a)–9(d) show 3D isometric views of a piece of the annealed Fe-30 a/o Zn alloy and images from sections cut parallel to the x-y, y-z, and z-x planes. In Figure 9(e), 2D images of selected sections along the sample axis are shown. It is seen that much of the annealed compact is pore free except for the mach stem region near the center. This piece was cut from a region of the compact close to the end that was near the blasting cap (Figure 2) during compaction and the mach-hole region persists through the entire length along the axis. Though the full 96 h treatment is not needed for complete homogenization, longer homogenization times strengthen the compact. XR T is an acceptable method for a detailed assessment of structural evolution during the var-
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ious stages of processing and optimization of the compaction process. Similar results were obtained for Fe-5 a/o Zn, Fe-10 a/o Zn, Fe-15 a/o Zn, and Fe-20 a/o Zn alloy compacts. The choice of explo-
Figure 9. (a) to (d) 3D isometric views of piece of annealed Fe-30 a/o Zn alloy compacts with images of sections cut parallel to x-y, y-z, and z-x planes. Dark regions are 3D voids/cracks in the compact. (e) 2D images of a series of selected sections along sample axis in a 3D XRT image of annealed Fe-30 a/o Zn alloy compact
sive compaction for consolidation was made based on the isostatic nature of the compaction process and the high attendant compaction pressures. Though the design of the explosive compaction assembly is complex, the authors have had extensive experience with such systems. It is likely that cold isostatic pressing is a possible alternative that could provide a more uniform product, though the maximum compaction pressure would be lower than that obtained during explosive compaction. Other techniques, such as swaging and hydrostatic extrusion before the extended homogenization anneal, can also lead to a uniform product but would have changed the preferred particle orientation obtained with the applied magnetic field during LPS. The iron–zinc alloy samples were also characterized using OIM to determine if the process used to induce texture in the samples was viable. Figure 10 shows an inverse pole figure of an Fe30 a/o Zn sample. It can be seen that there is definite texturing with a [100] preference along the sample axis; however, the texture is weak with 1.8 times random frequency. This weak level of texture is attributed to several factors. First, the alignment procedure itself will only line up the easy [100] direction of the largest grain in the particle. If there are other grains in the particle that are misoriented with respect to this grain, then these will be off axis and these grains will contribute to the low level of texturing. The second factor that leads to a weak texture is the LPS tem-
Figure 10. Inverse pole figure of texturing in Fe-30 a/o Zn alloy compact. [100] is along axis of compact
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shown in Table I. The saturation magnetostriction values show a large degree of scatter. It is likely that a less defective compact might have produced more consistent results for saturation magnetostriction. Nevertheless, the magnetostriction measurements do indicate that the enhancement in magnetostriction from the alloying of iron with zinc is, at best, small compared with the changes observed with gallium additions.
Figure 11. Saturation magnetization vs. zinc content in iron–zinc alloys after homogenization anneal
perature. With the increase in temperature as the powders sinter, the particles will begin to shift due to the thermal energy. This could rotate the particles slightly off axis from the direction of the magnetic field: the higher the temperature, the larger the misalignment. The final factor that affects texturing is the compaction process itself. As the sample is compacted, the particles are violently shifted. This will also probably contribute to the misalignment. Additionally, the severe plastic deformation that the sample undergoes during compaction builds up internal strain energy. When the sample is subjected to a homogenization anneal at 973K, some recrystallization is expected and texturing will be lost. The saturation magnetization data for iron–zinc alloys from VSM measurements are shown in Figure 11 as a function of zinc content. The specific magnetization varied continuously with concentration, decreasing from a value of 216.8 Am2/kg (emu/g) for pure iron to a value of 134.7 Am2/kg (emu/g) for the Fe-30 a/o zinc alloy. Magnetostriction data obtained for the compacted and homogenized iron–zinc alloys are TABLE I. SATURATION MAGNETOSTRICTION IN IRON–ZINC ALLOYS AS A FUNCTION OF ZINC CONTENT Zn (a/o)
Saturation Magnetostriction
0 5 10 15 20 25
22 -9 35 0 0 12
Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
CONCLUSIONS Textured bulk iron–zinc alloys with zinc contents ranging from 5 w/o to 30 w/o were successfully processed using a novel PM approach involving a combination of LPS, explosive compaction, and a homogenization anneal. The orientation of ferromagnetic iron powder by the application of a magnetic field during powder loading and subsequent low-temperature LPS promotes a [100] texture. XR T and optical microscopy show that explosive compaction results in compacts with negligible porosity, except in the central mach-tube region. XRT also illustrates the gradual elimination of the machhole region along the axis in the direction of shock-wave propagation, suggesting a decrease in compaction pressure along the axis. Reduction in this E:M to <6 is expected to avoid the formation of the central mach-stem region and microcracks. An alternative approach using cold isostatic pressing is expected to result in more uniform product but the maximum compaction pressure would be lower than that obtained during explosive compaction. Other techniques, such as swaging and hydrostatic extrusion (prior to the homogenization anneal), should also lead to a uniform product, but would change the preferred particle orientation obtained with the applied magnetic field during LPS. The magnetization measurements show a monotonic decrease in magnetization with zinc content, consistent with a reduction in the amount of iron. The magnetostriction measurement results show scatter but, nevertheless, suggest that there is no significant or dramatic change in the magnetostriction of iron due to the addition of zinc. ACKNOWLEDGEMENTS The authors acknowledge support of this work by the National Science Foundation under NSF DMR Award #0241603, and by the University of Utah.
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REFERENCES 1. S. Guruswamy, N. Srisukhumbowornchai, A.E. Clark, J.B. Restorff, and M. Wun-Fogle, “Strong, Ductile, and Low-Field Magnetostrictive Alloys Based on Fe-Ga,” Scripta Mater., 2000, vol. 43 no. 3, pp. 239–244. 2. N. Srisukhumbowor nchai, “Development of Highly Magnetostrictive Alloys Based on Fe-Ga and Fe-Ga-Al Alloys,” 2001, PhD Thesis, University of Utah, Salt Lake City, UT. 3. N. Srisukhumbowornchai and S. Guruswamy, “Large Magnetostriction in DS FeGa and FeGaAl Alloys”, J. Appl. Phys., 2001, vol. 90, pp. 5,680–5,688. 4. J. Emsley, The Elements, Third Edition, 1998, Clarendon Press, Oxford, UK. 5. B.P. Burton and P. Perrot, “Fe-Zn (Iron-Zinc)”, Phase Diagrams of Binary Iron Alloys, Monograph Series on Alloy Phase Diagrams No. 9, edited by H. Okamoto, ASM International, Materials Park, OH, 1993, pp. 459–466. 6. H.A. Wriedt and S. Arajs, “Ferromagnetic Curie Temperatures of Some Iron-Zinc Solid Solutions”, Phys. Stat. Sol., 1966, vol. 16, pp. 475–478. 7. F. Zhou, Y.T. Chou and E.J. Lavernia, “Formation of Supersaturated Single-Phase BCC Solid Solutions in FeZn Binary System by Mechanical Alloying”, Materials Transactions, 2001, vol. 24, no. 8, pp. 1,566–1,570. 8. G.R. Spiech, L. Zwell and H.A. Wriedt, “The Lattice Parameter and Alpha Phase Boundary of Ferritic IronZinc Alloys”, Met. Trans., 1964, vol. 230, pp. 939–940. 9. S. Guruswamy, M.K. McCarter and M.E. Wadsworth,
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10. 11. 12.
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16.
“Explosive Compaction of Metal-Matrix Composites and Deuterides,” Advances in Powder Metallurgy—1991, compiled by L.F. Pease III and R.J. Sansoucy, Metal Powder Industries Federation, Princeton, NJ, 1991, vol. 6, pp. 251–265. R.M. German, Powder Metallurgy Science, Second Edition, 1994, Metal Powder Industries Federation, Princeton, NJ. R.M. German, Liquid Phase Sintering, 1985, Plenum Press, New York, NY. L.E. Murr and K.P. Staudhammer, “Shock Wave Fundamentals: Effects on the Structure and Behavior of Engineering Materials”, Shock Waves for Industrial Applications, edited by L.E. Murr, Noyes Publications, Park Ridge, NJ, 1988, pp. 1–59. R. Prummer, “Powder Compaction”, Chapter 10, Explosive Welding, For ming and Compaction, edited by T.Z. Blazynski, Applied Science Publishers, London, UK, 1983, pp. 369–395. J. Pearson, “Energy Considerations in Explosive Metal Working”, Advanced High Energy Rate Forming Book 3, Proc. ASTME–Creative Manufacturing Seminar, American Society of Tool and Manufacturing Engineers, Detroit, Michigan, 1963, Paper No. SP63-26, pp. 1–29. Y. Zel’dovich and Y. Raizer, Physics of Shock Waves and High Temperature Hydrodynamic Phenomena, 1967, Academic Press, New York, NY. Explosives and Rock Blasting, 1987, Atlas Powder Company, Dallas, Texas, p.18. ijpm
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ENGINEERING & TECHNOLOGY
IMPROVEMENT IN FATIGUE PERFORMANCE OF POWDER-FORGED CONNECTING RODS BY SHOT PEENING Edmond Ilia, PMT*, Russell A. Chernenkoff**, and Kevin T. Tutton***
INTRODUCTION Shot peening is recognized as an important process for improving the fatigue performance of highly stressed components. This improvement is attributed to the formation of compressive residual stresses on the surface layer of a material by spherical media (shot) propelled by compressed air or centrifugal force. The shot strikes the surface of the part and causes plastic flow in the material, stressing it beyond its yield strength, resulting in a residual compressive stress on the surface of the part. The depth of the compressive layer is determined by the yield behavior of the material and the impact intensity of the shot. Impact intensity is specified by an Almen number.1 This number is determined by measuring the arc height of an Almen strip after the peening process. The influence of shot peening depends on the intensity, stability, and coverage of the residual stress induced in the part. The magnitude and depth of the residual stress can be assessed quantitatively by X-ray measurement techniques. Residual stresses are defined as "stresses present in a body that is free of external forces or thermal gradients."2 These stresses can be either beneficial or detrimental to the performance of a dynamically loaded component, depending on whether the stresses are compressive or tensile. It is widely recognized that fatigue cracks can only initiate or grow in a surface under tension,3 and that compressive residual surface stresses improve fatigue resistance. 1,4 This can best be described in terms of crack propagation. In general, fatigue cracks propagate when opened by a tensile load and do not propagate when closed by a compressive load. The influence of a residual stress on a component can shift the crack opening and closing point to a higher stress level or lower stress level during the loading cycle. When a load, within the elastic range of the material, is applied to a component with a compressive residual stress, the effective stress from the applied load is decreased by the amount of the compressive residual stress.5 In contrast, a tensile residual stress is added to the effective tensile
Shot peening is a surface treatment that is commonly used to improve the fatigue performance of structural components. Improvement is attributed to the formation of compressive residual stresses on the surface layer of the component by the impact of spherical shot. These stresses are known to reduce service tensile stress and therefore increase fatigue performance. To quantify the effect of shot peening, a series of powder-forged (PF) connecting rods were peened at several intensity levels, evaluated in fatigue, and the data correlated with compressive residual stress. Stress magnitude and depth were measured utilizing X-ray diffraction methods. Scanning electron microscopy (SEM) was also performed to determine the fracture initiation site(s).
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)
*Chief Metallurgist, ***Engineering Technician, Metaldyne Sintered Components, 1149 Rocky Road, Ridgway, Pennsylvania 15853, USA,
[email protected], **Senior Project Engineer, Metaldyne Sintered Components, 47603 Halyard Drive, Plymouth, Michigan 48170
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IMPROVEMENT IN FATIGUE PERFORMANCE OF POWDER-FORGED CONNECTING RODS BY SHOT PEENING
stress from the applied load. Since fatigue failures occur in areas of tensile stress, the crack initiation site of a component with a compressive residual stress at the surface is subsurface where tensile stresses that are offsetting the surface compressive stresses are higher. Even with these higher-than-normal tensile stresses in the subsurface region, the fatigue performance of the component is increased due to the additional load cycles needed to propagate the crack through the material under compressive stress to the surface of the component. Tensile residual stresses act in the opposite manner by increasing the tensile stress applied from a service load by the amount of the tensile residual stress. These residual tensile stresses can be detrimental to the fatigue performance of a part, especially if they are not accounted for in the design process. There are several ways of inducing a residual stress in a part; for example, machining, plating, and welding operations usually result in harmful tensile residual stresses. Processes such as surface rolling and shot peening induce beneficial compressive stresses on the surface of a part. A compressive residual stress that will best improve fatigue performance has an optimal balance of magnitude and depth. In this work, the beneficial effects of compressive residual stresses in PF connecting rods were evaluated and correlated with both higher fatigue strength and fatigue life. Bench fatigue tests were conducted on unpeened and shot-peened PF connecting rods to determine the effect of shot-peening intensity on fatigue performance. X-ray diffraction methods were utilized to determine the depth and magnitude of the induced residual stresses. Failure analysis on the fatigued connecting rods was also performed to determine the fracture initiation site(s). EXPERIMENTAL WORK Sample Preparation Three hundred consecutive PF connecting rods, manufactured under standard production conditions from PF-11C50: 2 w/o Cu-0.50 w/o C (asforged)-0.32 w/o MnS-balance Fe (MPIF Standard 35)6 were used in this study. These connecting rods were shot peened at five different intensities (50 connecting rods per intensity level), ranging from 10A to 24A, Table I. At least 100% coverage was obtained in all of the connecting rods during the shot peening process. Fifty connecting rods
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TABLE I: SHOT-PEENING TEST MATRIX Group 0A Group 10A Group 13A Group 17A Group 20A Group 24A
As-Forged Shot Peened at 10A Shot Peened at 13A Shot Peened at 17A Shot Peened at 20A Shot Peened at 24A
were evaluated without shot peening. Connecting rods from each group were subjected to metallographic analysis. All of the PF connecting rods evaluated met standard production requirements for density, microstructure, and hardness. The depth of the decarburized layer and its hardness were similar for the connecting rods analyzed from all of the six groups. A typical pearlitic–ferritic microstructure was observed in all of the connecting rods evaluated. Residual Stress Residual stress profiles for the PF connecting rods shot peened at the five different intensity levels, and for an unpeened connecting rod, are shown in Figure 1. Note that the depth of the compressive residual stress increases as the shotpeening intensity increases. The magnitude of the compressive stress was highest for specimens from Groups 17A and 20A, and the depths of the attendant compressive stresses were 0.462 mm and 0.488 mm, respectively. The depth of the compressive stress for specimens from Group 24A was significantly greater (0.693 mm) than at the other shot-peening intensity levels, but its magnitude at the surface was the least. The magnitude
Figure 1. Residual stress profiles (Groups 0A, 10A, 13A, 17A, 20A, and 24A)
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IMPROVEMENT IN FATIGUE PERFORMANCE OF POWDER-FORGED CONNECTING RODS BY SHOT PEENING
of the compressive stress was highest in the unpeened connecting rod due to work hardening during the forging process, but its depth was shallow. This combination is not conducive to the enhancement of fatigue resistance. A good balance between the magnitude of the compressive stress and its depth is required for optimal fatigue resistance. Fatigue Testing Fatigue strength was determined employing the staircase method (MPIF Standard 56).7 Axial constant-amplitude fatigue tests were conducted at room temperature using an MTS servo-hydraulic testing machine at a stress ratio R = -1. The criterion for runout was ≥107 cycles. Twenty connecting rods per group were tested. Fatigue life was determined employing the Weibull method, as specified by Lipson and Sheth,8 at a stress ratio R = -1. A sufficiently high stress level was applied in order to enhance failure as a result of the test. Eight connecting rods per group were tested. Fatigue Strength Results of staircase fatigue testing of specimens from Groups 0A through 13A and Group 24A are summarized in Figure 2. As shown, there is a significant improvement in fatigue strength as a result of shot peening; note the difference between specimens from Group 0A and Group 10A. Further improvements in fatigue strength are obtained by increasing the shot-peening intensity up to 13A. Note the difference between the fatigue results from Groups 10A and 13A. The difference in fatigue strength among the specimens from Groups 13A and 24A is not significant since the test results essentially overlap. For clarity, fatigue-test results for specimens from Groups 13A through 24A are summarized in Figure 3. As shown, the difference among these groups is minimal, in particular if shot-peening intensity levels 17A, 20A, and 24A are considered. The PF connecting rods from Group 13A sustained slightly lower stress levels during fatigue testing. The fatigue strength at 50% survival rate (SR), at 90% SR, and the respective scatter for each group are summarized in Table II. As shown in Table II, an improvement of approximately 31% in fatigue strength at 90% SR was obtained due to shot peening at an intensity Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
level 10A. An additional improvement of 11.5% was obtained by increasing the shot peening intensity level to 17A, resulting in a total improvement of approximately 46% in the fatigue strength at 90% SR. Further increases in fatigue strength were not observed, even though the shot-peening intensity level was increased up to 24A; it actually dropped slightly below the level reached at 17A. This is an indication that saturation conditions exist at this level of shot-peening intensity. The lowest scatter was observed for specimens in Group 17A, with Group 20A being similar, while the highest scatter was obtained for specimens from Group 24A. This result is attributed to peen-
Figure 2. Fatigue test results, R = -1 (Groups 0A, 10A, 13A, and 24A)
Figure 3. Fatigue-test results, R = -1 (Groups 13A through 24A)
TABLE II. FATIGUE STRENGTH (MPa)* Intensity
0A
10A
13A
17A
20A
24A
Fatigue Strength at 50% SR 220.9 285.9 310.5 315.0 314.9 314.3 Fatigue Strength at 90% SR 213.7 279.8 303.3 312.1 309.3 310.3 Scatter 5.6 4.8 5.6 3.7 4.3 6.4 *Based on 20 connecting rods per group
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TABLE III. FATIGUE-TEST RESULTS (WEIBULL TESTING AT R = -1) Cycles to Failure at ±317 MPa* Unpeened Group 10A Group 13A Group 17A Group 20A Group 24A 84,000 96,000 115,000 133,000 139,000 142,000 147,000 194,000
652,000 832,000 827,000 1,605,000 1,781,000 754,000 1,002,000 960,000 2,481,000 5,343,000 853,000 1,102,000 1,859,000 3,003,000 5,999,000 1,111,000 1,964,000 4,086,000 4,564,000 6,739,000 3,256,000 6,054,000 5,350,000 7,114,000 7,285,000 3,821,000 8,495,000 10,000,000 10,000,000 10,000,000 5,562,000 10,000,000 10,000,000 10,000,000 10,000,000 8,782,000 10,000,000 10,000,000 10,000,000 10,000,000
*Based on eight connecting rods per group ing surface extrusion faults (PSEF), as explained in detail by Chernenkoff, Mocarski, and Yeager.9
Figure 4. Effect of shot-peening intensity on the B10 life at 90% CL
Fatigue Life Weibull testing was conducted at a stress level of ±317 MPa. A summary of the test results is included in Table III. Note that a higher number of cycles to failure was obtained as the level of the shot-peening intensity increased. Weibull calculations were run on the groups to determine the slope (shape parameter β), the characteristic life (scale parameter η), the mean life (ML), and the B10 life at a 90% confidence level (CL). Since the Weibull distribution is generally not symmetric, the mean and the median values will not be the same, as is the case for the normal distribution.10,11 Therefore, the mean lives of each group were calculated from the formula: ML = γ + η Γ (1/β + 1)
(1)
where γ = parameter set at zero since experimental data points reflect a straight line fit, and Γ = gamma function (from tables).10 The rank regression (RR) method was used to evaluate the Weibull distribution parameters. All calculations were run using WeibullSMITH™, VisualSMITH™12 and proprietary software.13 The results of the calculations utilizing the RR method
are summarized in Table IV. As shown in Table IV, the slope of the Weibull lines was higher than unity in all of the groups, thus no “infant mortality” was observed during these tests. The regression coefficient was higher than 0.93 and all of the tests run by the software 12 confirmed the validity of the calculations. The slope was almost constant (1.16, 1.08, 1.09) for the first three shot-peening intensities (Groups 10A, 13A, and 17A) and it increased for Group 20A (1.57), reaching its highest value (2.05) for Group 24A. Both the B10 life at 90% CL and the characteristic life significantly increase with the level of the shot-peening intensity, as illustrated in Figure 4 in the case of B10 at 90% CL. This improvement in fatigue life is significantly more pronounced in the case of specimens from Groups 20A and 24A. FAILURE ANALYSIS To better understand the fatigue-test results, a closer look at the failure-initiation sites was taken. It was observed that in the case of specimens from Groups 0A and 10A all of the fatigue failures started at the surface, Figure 5(a). In specimens from the other groups, more cracks
TABLE IV. RESULTS OF PR CALCULATIONS (WEIBULL TESTING AT R = - 1)* °Intensity B10 at 90% CL Slope Regression Characteristic Life Mean Life
Unpeened
Group 10A
Group 13A
Group 17A
Group 20A
Group 24A
62,288 4.27 0.98 143,717 130,751
141,776 1.16 0.93 3,149,785 2,987,217
170,810 1.08 0.93 5,053,780 4,914,676
211,623 1.09 0.96 5,858,765 5,664,806
688,316 1.57 0.97 6,930,833 6,225,095
1,408,382 2.05 0.93 8,216,232 7,278,620
*Parameters obtained by statistical analysis of data in Table III
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IMPROVEMENT IN FATIGUE PERFORMANCE OF POWDER-FORGED CONNECTING RODS BY SHOT PEENING
started internally from subsurface sites, as shown in Figure 5(b). As a result of increasing the shot-peening
(a) Surface
(b) Subsurface
Figure 5. Surface and subsurface crack-initiation sites (scale marks omitted intentionally)
Figure 6. Effect of shot-peening intensity on crack-initiation site (surface vs. subsurface)
Figure 7. Depth of subsurface crack-initiation sites
intensity, initiation of the failure process shifts away from the surface (below the shot-peened layer), thus making surface imperfections less effective as fatigue-crack-initiation sites. Similar results were reported by Chernenkoff, Mocarski, and Yeager.9 Figure 6 summarizes the change in fatiguecrack-initiation sites from 100% surface in the case of the lowest shot-peening intensity (10A) up to 73% subsurface in the case of the highest shotpeening-intensity level (24A). Figures 7 and 8 show that the depth of subsurface crack-initiation sites increased with increasing shot-peening intensity. The average distance from the surface at which the crack initiated was 0 mm in the case of specimens from Groups 0A through 13A, reaching a maximum of 0.839 mm in the case of specimens from Group 24A. It is interesting to note that notwithstanding a slight increase in surface roughness with higher shot-peening intensities (Table V), the number of failures from surface crack-initiation sites was reduced. Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
Figure 8. Effect of shot-peening intensity on distance from surface of crackinitiation site
TABLE V. SURFACE ROUGHNESS* Intensity Ra (µm)
10A 2.30
13A 3.70
17A 3.90
20A 4.30
24A 4.80
*Based on a maximum of five connecting rods per group
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DISCUSSION The effects of increasing the level of intensity of shot peening above 17A result in: 1. No further improvements in fatigue strength 2. A higher number of cycles for crack initiation and propagation to failure 3. More subsurface crack-initiation sites 4. Deeper crack-initiation sites However, even though the level of the compressive residual stress stays essentially constant for specimens from Groups 17A, 20A, and 24A, the B10 life at 90% CL continues to increase significantly, as shown in Figure 9. In comparison, the depth of the compressive residual stress increases with the level of shotpeening intensity, as shown in Figure 10. It is primarily the depth of the shot-peened layer which affects the fatigue life of the PF connecting rods. Relaxation of compressive residual stresses takes place when, in a component, local stresses are higher than the elastic limit of the material and local plastic deformation occurs, resulting in lower fatigue per for mance. This relaxation depends on several factors includes load amplitude, number of loading cycles, and the state of the initial residual stress. It was proven that the relaxation is more severe in the case of shallow layers of compressive residual stresses than for deep shot-peened layers.14,15 This explains the higher B10 lives at 90% CL obtained in the case of specimens from Groups 20A and 24A. The fatigue strength at 90% SR as a function of the maximum compressive residual stress is shown in Figure 11. Ignoring the compressive stress in the as-forged connecting rods (Group 0A)*, there is a continuous increase in the maximum compressive residual stress up to a shotpeening-intensity level of 20A. However, at the highest shot-peening-intensity level (24A), the maximum compressive residual stress drops significantly due to excessive plastic deformation on the surface resulting in PSEF or cracks. As a result of this “over-peening,” optimal compressive stresses can not be formed at the surface of the material. Thus, it is primarily the depth of the compressive residual stress that controls the fatigue life of the connecting rods, rather than the maximum compressive residual stress. The latter likely controls the fatigue strength, as shown in Figure 11.
Figure 9. Dependence of fatigue life on maximum compressive residual stress
Figure 10. Dependence of fatigue life on depth of compressive residual stress
Figure 11. Dependence of fatigue strength on maximum compressive residual stress
*Attributed to work hardening during forging that produces a shallow layer at the surface in compression
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CONCLUSIONS 1. Increasing the level of shot-peening intensity resulted in a greater depth of compressive residual stress. 2. The fatigue strength of powder-forged connecting rods increased 31% when shot peened at an intensity level of 10A and up to 46% when shot peened at an intensity level of 17A. Increasing the shot-peening intensity (up to 24A) did not further improve fatigue strength. Maximum fatigue strength occurred for a shot-peening intensity level of 17A. 3. Weibull testing showed both the B10 life at 90% confidence level and the characteristic life increase significantly with the level of shot-peening intensity. 4. Subsurface fatigue-crack-initiation sites increased with an increase in the level of shot-peening intensity. The depth of the crack-initiation site was also greater as the level of shot-peening intensity increased, reaching an average depth of 0.839 mm in specimens from Group 24A. 5. The maximum compressive residual stress is essentially constant for shot-peening-intensity levels of 17A, 20A, and 24A; however, the B10 life at the 90% confidence level continues to increase significantly. 6. Residual stress relaxation was more significant in PF connecting rods shot peened at intensity levels of 10A, 13A, and 17A.
Volume 45, Issue 3, 2009 International Journal of Powder Metallurgy
REFERENCES 1. H.O. Fuchs and R.I. Stephens, "Self-Stresses and Notch Strain Analysis," Metal Fatigue in Engineering, John Wiley & Son, New York, 1980, pp. 125–147. 2. Metals Handbook—Desk Edition, Edited by H.E. Boyer and T.L. Gall, Am. Soc. Metals, Metals Park, OH, 1985, p. 1–31. 3. N.K. Burrell, "Controlled Shot Peening of Automotive Components," SAE Technical Paper No. 850365, SAE International, Warrendale, PA, 1985. 4. A. Niku-Lari, "Influence of Residual Stress Introduced by Shot Peening Upon the Fatigue Life of Materials," Experimental Techniques, 1983, vol. 7, no. 3, pp. 21–25. 5. J.O. Almen and P.H. Black, Residual Stresses and Fatigue in Metals, McGraw-Hill, New York, NY, 1963. 6. MPIF Standard 35, Materials Standards for P/F Steel Parts, Metal Powder Industries Federation, Princeton, NJ, 2007. 7. “Standard 56, Determination of Rotating Beam Fatigue Endurance Limit in Powder Metalurgy Materials”, Standard Test Methods for Metal Powders and Powder Metallurgy Products, Metal Powder Industries Federation, Princeton, NJ, 2006. 8. C. Lipson and N. Sheth, Statistical Design and Analysis of Engineering Experiments, McGraw-Hill, New York, NY, 1973, pp. 270–274. 9. R.A. Cher nenkof f, S. Mocarski and D.A. Yeager, “Increased Fatigue Strength of Powder-Forged Connecting Rods by Optimized Shot Peening”, SAE Technical Paper 950384, SAE International, Warrendale, PA, 1995. 10. D. Kececioglu, Reliability & Life Testing Handbook, PTR Prentice Hall, Englewood Cliffs, NJ, 1993. 11. R.B. Abernethy, The New Weibull Handbook, Third Edition, self-published, North Palm Beach, FL, 1998. 12. WeibullSMITH™, VisualSMITH™ software, Fulton Findings, San Pedro, CA. 13. E. Ilia, Weibull Software, proprietary, 1993. 14. K. Iida and K. Taniguchi, “Relaxation of Residual Stress Distribution Produced by Shot Peening under Fatigue Test”, Proceedings of the 6th International Conference on Shot Peening, edited by J. Champaigne, ISCSP, Mishawaka, IN, 1996, pp. 397–402. 15. J. Morrow, A.S. Ross and G.M. Sinclair, “Relaxation of Residual Stresses Due to Fatigue Loading,” SAE Transactions, 1960, vol. 68, pp. 40–48. ijpm
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MEETINGS AND CONFERENCES
2009 INTERNATIONAL METALLOGRAPHIC SOCIETY (IMS) ANNUAL CONVENTION July 26–30 Richmond, VA mm2009.microscopy.org/ BASIC PM SHORT COURSE July 27–29 State College, PA MPIF* ASM MATERIALS AND PROCESSES FOR MEDICAL DEVICES (MPMD) CONFERENCE & EXPOSITION August 10–12 Minneapolis, MN asmcommunity.asminternational.org/content/Events/ MPMD-09/ SHANGHAI INTERNATIONAL AUTOMOTIVE MANUFACTURING TECHNOLOGY AND MATERIAL SHOW August 18–21 Shanghai, China www.shanghaiamts.com THERMEC 2009: SIXTH INTERNATIONAL CONFERENCE ON ADVANCED MATERIALS AND PROCESSES August 25–29 Berlin, Germany SDMA 2009/ICSF VII—4TH INTERNATIONAL CONFERENCE ON SPRAY DEPOSITION AND MELT ATOMIZATION/7TH INTERNATIONAL CONFERENCE ON SPRAY FORMING September 7–9 Bremen, Germany www.sdma-conference.de/
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25TH ASM HEAT TREATING SOCIETY CONFERENCE AND EXPOSITION Co-located with AGMA Gear Expo 2009 September 14–17 Indianapolis, IN www.asminternational.org/ heattreat EURO PM2009 INTERNATIONAL POWDER METALLURGY CONGRESS & EXHIBITION October 12–14 Copenhagen, Denmark www.epma.com/pm2009 HIGHMATTECH 2009 INTERNATIONAL CONFERENCE October 19–23 Kiev, Ukraine www.hmt.kiev.ua CERAMITEC 2009 11TH INTERNATIONAL TRADE FAIR FOR MACHINERY, EQUIPMENT, PLAN, PROCESSES AND RAW MATERIALS FOR CERAMICS AND POWDER METALLURGY October 20–23 Munich, Germany www.ceramitec.de 96TH ASM ANNUAL MEETING AT MS&T 2009 October 25–29 Pittsburgh, PA www.matscitech.org 2009 CHINA (SHANGHAI) POWDER METALLURGY & ADVANCED CERMAICS EXHIBITION & CONGRESS November 9–10 Shangahi, China www.china-pmexpo.com/en
2010 POWDERMET2010: MPIF/APMI INTERNATIONAL CONFERENCE ON POWDER METALLURGY & PARTICULATE MATERIALS June 27–30 Hollywood (Ft. Lauderdale), FL MPIF* PRICM 7 7TH PACIFIC RIM INTERNATIONAL CONFERENCE ON ADVANCED MATERIALS AND PROCESSING August 1–5 Cairns, Australia www.materialsaustralia.com. au/scripts/cgiip.exe/ WService=MA/ccms.r?PageI D=19070 7TH INTERNATIONAL SYMPOSIUM ON ALLOY 718 & DERIVATIVES September 10–13 Pittsburgh, PA www.tms.org PM2010 WORLD CONGRESS October 10–14 Florence, Italy
*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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